Cover-Proceedings-create outline - Polymer Society of Thailand

-[ PC -5 ]-- nTTTT

International Polymer Conference of Thailand

PROGRAM BOOK


Annual Polymer Conference June 18-19, 2015

Pathumwan Princess Hotel, Bangkok Thailand

by


PCT-5: Another Step of Thai Polymer Society, A Strong

International Network

Suwabun Chirachanchai, Ph.D.

Professor

President of Polymer Society of Thailand

PCT-5 Chairman

The Polymer Society of Thailand (PST). PST was founded in 1999 by Thai academic and industrial

polymer scientists. The society plays its important role to be the catalyst among the scientists so that

the research, development and innovation of polymer science and engineering in the country can go

beyond. In order to achieve this goal, we realize the conference as an important stage to motivate and

accelerate the advancement of polymers along with updating information including the strengthen of

people network through the Plenary and Keynote lectures, general oral and poster presentations, and

the recreation programs. It is also the stage where we express our recognition to award the young

polymer scientists in the country, so-called 'Thai Polymer Society Rising Star'. Therefore, the Polymer

Conference of Thailand (PCT) has been held annually for the past 4 years. And for this year, the PST

is going to another step where we aim for the strong international networking and this results in the

International Polymer Conference of Thailand, PCT-5.

I would like to take this opportunity to acknowledge the guests from abroad who make the conference

be the real international one, the Plenary Speakers; Prof. Dr. Andreas Greiner (Universität Bayreuth,

Germany), Dr. Michel Wong Chi Man (Institut Charles Gerhardt Montpellier, France), Dr. Piyada

Charoensirisomboon (BASF Advanced Chemicals Co., Ltd. Shanghai, China), the Keynote Speakers;

Prof. Dr. Jun Li (National University of Singapore, Singapore), Prof. Dr. Rusli bin Daik (Universiti

Kebangsaan Malaysia), Asst. Prof. Dr. Michiya Matsusaki (Osaka University, Japan), Assoc. Prof. Dr.

Seiichi Kawahara (Nagaoka University of Technology, Japan), Prof. Dr. Masahiro Ohshima (Kyoto

University, Japan), Assoc. Prof. Dr. Moon Jeong Park (Pohang University of Science and Technology,

Korea), Prof. Dr. Yun Yan (Peking University, China), Prof. Dr. Seema Agarwal (Universität

Bayreuth, Germany), Prof. Dr. Masayuki Yamaguchi (Japan Advanced Institute of Science and

Technology, Japan) and Prof. Dr. Yuko Ikeda (Kyoto Institute of Technology, Japan).

The appreciations are also to the Thai Keynotes, the speakers from industries, the general participants,

who fill up the talk in two-day conference. I am also happy to see the instrument companies join the

talk by bringing in the updated information about instrumentation for the effective and efficient

research.

I, personally, look forward to the 'Rising Stars 2015' to be announced on the conference day and the

best poster awards which represent the research quality of the country.

On behalf of the Thai Polymer Society, I would like to express my deepest gratitude to the sponsors,

which are the Gold from the SCG Chemicals Company Limited, the BASF (Thai) Limited and the

PTT Global Chemical Public Company Limited, the Silver from the Thailand Convention and

Exhibition Bureau and the Bronze from the Bruker BioSpin AG.


Board of Polymer Society of Thailand

Advisory board

Asst. Prof. Dr. Krisda Suchiva

National Metal and Materials Technology Center (MTEC)

Prof. Dr. Pattarapan Prasassarakich

Department of Chemical Technology, Chulalongkorn University

Prof. Dr. Supawan Tantayanon

Department of Chemistry, Chulalongkorn University

Prof. Dr. Suda Kiatkamjornwong

Faculty of Science, Chulalongkorn University

President Prof. Dr. Suwabun Chirachanchai

The Petroleum and Petrochemical College, Chulalongkorn

University

Vice President Assoc. Prof. Dr. Pranee Phinyocheep

Department of Chemistry, Mahidol University

Vice President Dr. Veerapat Tantayakom

PTT Global Chemical Public Company Limited

Secretary Asst. Prof. Dr. Varawut Tangpasuthadol

Department of Chemistry, Chulalongkorn University

Treasurer Asst. Prof. Dr. Kanoktip Boonkerd

Department of Material Science, Chulalongkorn University

Committee Assoc. Prof. Dr. Ittipol Jangchud

Department of Chemistry, King Mongkut’s Institute of Technology

Ladkrabang

Committee Asst. Prof. Dr. Winita Punyodom

Department of Chemistry, Chiang Mai University

Committee Assoc. Prof. Dr. Pakorn Opaprakasit

Sirindhorn International Institute of Technology, Thammasat

University

Committee Dr. Asira Fuongfuchat


continued


Board of Polymer Society of Thailand (continued)

Committee Asst. Prof. Dr. Kannika Sahakaro

Faculty of Science and Technology, Prince of Songkla University,

Pattani Campus

Committee Dr. Prakaipetch Kitiyanan

BASF (Thai) Limited

Committee Dr. Pasaree Laokijcharoen


Committee Dr. Narin Kaabbuathong

PTT Public Company Limited

Committee Dr. Warayuth Sajomsang

National Nanotechnology Center (NANOTEC)

Committee Assoc. Prof. Dr. Vuthichai Ervithayasuporn

Department of Chemistry, Mahidol University

Committee Dr. Wonchalerm Rungswang

SCG Chemicals Company Limited


Committee Chairman Prof. Dr. Suwabun Chirachanchai The Petroleum and Petrochemical College,

Chulalongkorn University

International Advisory and Scientific Committee Prof. Dr. William H. Daly, Louisiana State University, USA Prof. Dr. Avraam I. Isayev, The University of Akron, USA Dr. Patrick Brant Exxon Mobil Chemical, USA Prof. Dr. Garry L. Rempel, University of Waterloo, Canada Prof. Dr. Yusuf Yagci Istanbul Technical University, Turkey Prof. Dr. Seema Agarwal Universität Bayreuth, Germany Prof. Dr. Andreas Greiner Universität Bayreuth, Germany Dr. Klaus Tauer Max Planck Institute of Colloids and Interfaces,

Germany Prof. Dr. Laurent Fontaine Université du Maine, France Prof. Dr. Philippe Daniel Université du Maine, France Dr. Michel Wong Chi Man Institut Charles Gerhardt Montpellier, France Prof. Dr. Masahiro Ohshima Kyoto University, Japan Prof. Dr. Masayuki Yamaguchi Japan Advanced Institute of Science and Technology,

Japan Prof. Dr. Yuko Ikeda Kyoto Institute of Technology, Japan Assoc. Prof. Dr. Seiichi Kawahara Nagaoka University of Technology, Japan Assoc. Prof. Dr. Yoshimasa Yamamoto Tokyo National College of Technology, Japan Asst. Prof. Dr. Michiya Matsusaki Osaka University, Japan Prof. Dr. Changwoon Nah Chonbuk National University, Korea Prof. Dr. In Joo Chin Inha University, Korea Prof. Dr. Hyoung Jin Choi Inha University, Korea Assoc. Prof. Dr. Moon Jeong Park Pohang University of Science and Technology, Korea Prof. Dr. Hongzhi Liu Shandong University, China Prof. Dr. Yun Yan Peking University, China Prof. Dr. Jun LI, National University of Singapore, Singapore Prof. Dr. Chi Wu Chinese University of Hong Kong, Hong Kong Prof. Dr. Braja Gopal Bag Vidyasagar University, India Assoc. Prof. Dr. Kinsuk Naskar Indian Institute of Technology Kharagpur, India Prof. Dr. Rusli bin Daik Universiti Kebangsaan Malaysia, Malaysia Assoc. Prof. Dr. Zulkifli Mohammad Ariff Universiti Sains Malaysia, Malaysia

Scientific Committee Mahidol University Assoc. Prof. Dr. Pranee Phinyocheep Chair of Scientific Committee Department of Chemistry Assoc. Prof. Dr. Vuthichai Ervithayasuporn Department of Chemistry Assoc. Prof. Dr. Taweechai Amornsakchai Department of Chemistry Assoc. Prof. Dr. Sombat Thanawan Department of Chemistry

Chulalongkorn University Prof. Suwabun Chirachanchai The Petroleum and Petrochemical College Prof. Dr. Pattarapan Prasassarakich Department of Chemical Technology Prof. Dr. Suda Kiatkamjornwong Faculty of Science Asst. Prof. Dr. Varawut Tangpasuthadol Department of Chemistry

Asst. Prof. Dr. Kanoktip Boonkerd Department of Material Science

http://chemistry.sc.mahidol.ac.th/en/people/faculty/vuthichai-ervithayasuporn/


King Mongkut’s Institute of Technology Ladkrabang Assoc. Prof. Dr. Ittipol Jangchud Department of Chemistry Asst. Prof. Dr. Chonlada Ritviruth Department of Chemistry Asst. Prof. Dr. Suparat Rukchonlatee Department of chemistry

Thammasat University Assoc. Prof. Dr. Pakorn Opaprakasit Sirindhorn International Institute of Technology Asst. Prof. Dr. Suwadee Kongparakul Department of Chemistry Dr. Nopparat Plucktaveesak Department of Chemistry Assoc. Prof. Dr. Cattaleeya Pattamaprom Department of Chemical Engineering Asst. Prof. Dr. Siwarutt Boonyarattanakalin Sirindhorn International Institute of Technology Asst. Prof. Dr. Wanwipa Siriwatwechakul Sirindhorn International Institute of Technology

Rajamangala University of Technology Thanyaburi Asst. Prof. Dr. Amorn Chaiyasat Department of Chemistry

Chiang Mai University Asst. Prof. Dr. Winita Punyodom Department of Chemistry Asst. Prof. Dr. Kanarat Nalampang Department of Chemistry Asst. Prof. Dr. Saengrawee Sriwichai Department of Chemistry Dr. Robert Molloy Department of Chemistry Dr. Kiattikhun Manokruang Department of Chemistry Dr. Paralee Waenkaew Department of Chemistry Dr. Patnarin Worajittiphon Department of Chemistry Dr. Runglawan Somsunan Department of Chemistry Asst. Prof. Dr. Jantrawan Pumchusak Department of Industrial Chemistry Dr. Datchanee Pattavarakorn Department of Industrial Chemistry

Naresuan University Assoc. Prof. Dr. Metha Rutnakornpituk Department of Chemistry,

Maejo University Asst. Prof. Dr. Tithinun Rattanaplome Faculty Engineering and Agro-Industry Dr. Worawan Pechurai Faculty Engineering and Agro-Industry Dr. Achara Kleawkla Department of Chemistry

University of Phayao Dr. Wijitra Meelua Department of Chemistry Dr. Boontharika Thapsukhon School of Science

Mae Fah Luang University Dr. Patchara Punyamoonwongsa School of Science

Rajamangala University of Technology Lanna Dr. Chinanat Witthayaprapakorn Faculty of Sciences and Agricultural Technology

Mahasarakham University Assoc. Prof. Dr. Yodthong Baimark Department of Chemistry

Suranaree University of Technology Assoc. Prof. Dr. Yupaporn Ruksakulpiwat School of Polymer Engineering Asst. Prof. Dr. Chantima Deeprasertkul School of Polymer Engineering Asst. Prof. Dr. Kasama Jarukumjorn School of Polymer Engineering Asst. Prof. Dr. Pranee Chumsamrong School of Polymer Engineering Asst. Prof. Dr. Wimonlak Sutapun School of Polymer Engineering Dr. Tatiya Trongsatitkul School of Polymer Engineering

Prince of Songkla University Assoc. Prof. Dr. Varaporn Tanrattanakul Department of Materials Science and Technology,

Hatyai Campus


Asst. Prof. Dr. Kannika Sahakaro Faculty of Science and Technology, Pattani Campus Asst. Prof. Dr. Anoma Titithammawong Faculty of Science and Technology, Pattani Campus Asst. Prof. Dr. Natinee Lopattananon Faculty of Science and Technology, Pattani Campus Dr. Sunisa Suchart Faculty of Science and Industrial Technology,

Suratthani Campus

Songkhla Rajabhat University Asst. Prof. Dr. Polphat Ruamcharoen Program of Rubber and Polymer Technology,

National Metal and Materials Technology Center Asst. Prof. Dr. Krisda Suchiva Dr. Asira Fuongfuchat Dr. Pasaree Laokijcharoen Dr. Chureerat Prahsarn Dr. Patcharee Larpsuriyakul

National Nanotechnology Center Dr. Warayuth Sajomsang

Private sectors Dr. Prakaipetch Kitiyanan BASF (Thai) Limited Dr. Veerapat Tantayakom PTT Global Chemical Public Company Limited Dr. Narin Kaabbuathong PTT Public Company Limited Dr. Wonchalerm Rungswang SCG Chemicals Company Limited

Organizing Committee Asst. Prof. Dr. Krisda Suchiva National Metal and Materials Technology Center Prof. Dr. Suda Kiatkamjornwong Faculty of Science, Chulalongkorn University Prof. Dr. Pattarapan Prasassarakich Department of Chemical Technology, Chulalongkorn

University Prof. Dr. Supawan Tantayanon Department of Chemistry, Chulalongkorn University Assoc. Prof. Dr. Pranee Phinyocheep Department of Chemistry, Mahidol University Assoc. Prof. Dr. Ittipol Jangchud Department of Chemistry, King Mongkut’s Institute

of Technology Ladkrabang Assoc. Prof. Dr. Pakorn Opaprakasit Sirindhorn International Institute of Technology,

Thammasat University Assoc. Prof. Dr. Vuthichai Ervithayasuporn Department of Chemistry, Mahidol University Asst. Prof. Dr. Kannika Sahakaro Faculty of Science and Technology, Prince of Songkla

University, Pattani Campus Asst. Prof. Dr. Winita Punyodom Department of Chemistry, Chiang Mai University Asst. Prof. Dr. Varawut Tangpasuthadol Department of Chemistry, Chulalongkorn University

Asst. Prof. Dr. Kanoktip Boonkerd Department of Material Science, Chulalongkorn University

Dr. Asira Fuongfuchat National Metal and Materials Technology Center Dr. Pasaree Laokijcharoen National Metal and Materials Technology Center Dr. Warayuth Sajomsang National Nanotechnology Center Dr. Prakaipetch Kitiyanan BASF (Thai) Limited Dr. Veerapat Tantayakom PTT Global Chemical Public Company Limited Dr. Narin Kaabbuathong PTT Public Company Limited Dr. Wonchalerm Rungswang SCG Chemicals Company Limited

Regional Organizing Committee Dr. Robert Molloy Department of Chemistry, Chiang Mai University Assoc. Prof. Dr. Yodthong Baimark Department of Chemistry, Mahasarakham University Assoc. Prof. Dr. Metha Rutnakornpituk Department of Chemistry, Naresuan University

http://chemistry.sc.mahidol.ac.th/en/people/faculty/vuthichai-ervithayasuporn/


Assoc. Prof. Dr. Jatuphorn Wootthikanokkhan Division of Materials Technology, King Mongkut’s University of Technology Thonburi

Asst. Prof. Dr. Chiraphon Chaibundit Department of Materials Science and Technology, Prince of Songkla University, Hat-Yai Campus

Asst. Prof. Dr. Chaiwat Ruksakulpiwat Department of Chemistry, Khon Kaen University Asst. Prof. Dr. Suttinun Phongtamrug Department of Industrial Chemistry, King Mongkut's

University of Technology North Bangkok Asst. Prof. Dr. Preeyaporn Chaiyasat Department of Chemistry, Rajamangala University of

Technology Thanyaburi Asst. Prof. Dr. Polphat Ruamcharoen Program of Rubber and Polymer Technology,

Songkhla Rajabhat University Asst. Prof. Dr. Rapee Gosalawit School of Chemistry, Suranaree University of

Technology Asst. Prof. Dr. Rukkiat Jitchati Department of Chemistry, Ubon Ratchathani

University Asst. Prof. Dr. Wanvimol Pasanphan Department of Materials Science, Kasetsart

University Dr. Sunisa Suchart Faculty of Science and Industrial Technology, Prince

of Songkla University, Suratthani Campus Dr. Achara Kleawkla Department of Chemistry, Maejo University Dr. Boontharika Thapsukhon School of Science, University of Phayao Dr. Patchara Punyamoonwongsa School of Science, Mae Fah Luang University Dr. Nattakan Soykeabkaew School of Science, Mae Fah Luang University


Contents

page

Main Event Program 1

Parallel Program 3

Plenary Lectures 8

PL-I Nanofibers by electrospinning – from a forgotten method to a major technique 9 Andreas Greiner

PL-II Bridged polysilsesquioxanes: synthesis and application fields 10 Michel Wong Chi Man

PL-III Creating chemistry for sustainable future in Asia 11 Piyada Charoensirisomboon

PST Rising Star Award Lecture 12

RS Amine-decorated polymeric colloidal particles :From syntheses to applications

13

Panya Sunintaboon

I. Session: Biomedical and Environmentally Friendly Polymers 14

KN-BIOEN-1 Supramolecular self-assembled polymers as novel biomaterials 15 Jun Li

KN-BIOEN-2 Control of cell surfaces by polymer/protein LbL films for fabrication of 3D-human tissue models

16

Michiya Matsusaki KN-BIOEN-3 Enzymatic degradation of oil palm empty fruit bunch biomass 17

Rusli Bin Daik KN-BIOEN-4 Chitosan dispersion as a pharmaceutical coating material 18

Satit Puttipipatkhachorn BIOENO Oral presentations 19

BIOENP Poster presentations 24

II. Session: Advances in Polymer Characterization 59

KN-CHAR-1

Preparation and properties of natural rubber with organic- inorganic nanomatrix structure

60

Seiichi Kawahara KN-CHAR-2 Chemically controlled self-assembly of gold nanoparticles by

site-selective protein immobilization: A model for antimalarial drug screening 61

Palangpon Kongsaeree

III. Session: Polymer Composites and Nanocomposites 62

KN-COMP-1 Interphase transfer of nanoparticles between immiscible polymer blends

63

Masayuki Yamaguchi KN-COMP-2 Hybrid porous polymers derived from octavinylsilsesquioxane 64

Hongzhi Liu

KN-COMP-3 Natural fiber reinforced rubber: recent advances toward high performance rubber matrix composites using pineapple leaf fiber

65

Taweechai Amornsakchai KN-COMP-4 Performance of aramid fiber in rubber compounds 66

Jutarat Phanmai

COMPO Oral presentations 67

COMPP Poster presentations 83


Contents

page

IV. Session: Advances in Polymer Processing 146 KN-PROC-1 Foam, (micro) foam, (nano)foam! - reality and dream 147

Masahiro Ohshima KN-PROC-2 Fiber design: A creation of fiber structure for feature and performance 148

Chureerat Prahsarn PROCO Oral presentations 149

PROCP Poster presentations 177

V. Session: Natural and Synthetic Rubbers 186

KN-RUBBER-1 New focus on rubber science and technology 187 Yuko Ikeda

RUBBERO Oral presentations 188

RUBBERP Poster presentations 206

VI. Session: Smart and Intelligent Polymers 229

KN-SMART-1 Coordination triggered division of vesicles 230 Yun Yan

KN-SMART-2 Self-assembled polymer electrolytes for future electrochemical devices 231 Moon Jeong Park

KN-SMART-3 Non-ionic thermoresponsive polymers of UCST-type in water: challenges and perspectives

232

Seema Agarwal KN-SMART-4 Polymer-based smart devices: Electronics on paper, plastic and textile 233

Teerakiat Kerdcharoen SMARTO Oral presentations 234

VII. Session: Polymer Research in Industry Sector 240

INDUS-I Development of bioplastics-based lamination application 241 Narin Kaabbuathong, PTT Public Co., Ltd.

INDUS-II From Bioscience to Polymer Science: Our Sustainable Prospects 242 Sukhgij Ysothonsreekul, PTT Global Chemical Public Company Limited

INDUS-III Optical characterization of thin films using a new Universal Measurement Accessory

243

Heng Soo Chin, Agilent Technologies, Inc.

INDUS-IV High performance composites for industry 244

Nopphawan Phonthammachai, SCG Chemicals Co., Ltd.

INDUS-V Synchrotron light for innovative polymers 245

Prae Chirawatkul & Wanwisa Limphirat, Synchrotron Light Research Institute Taweechai Amornsakchai

VIII. Session: Instruments 246

INS-I AXIS SUPRA, Cutting edge XPS technology for polymer 247 Bara Scientific Co. Ltd.

INS-II Simultaneous measurement of thermogravimetric differential thermal analysis and photoionisation mass spectrometry complexed through unique skimmer interface system, TG-DTA-PIMS

248

Crest Nanosolution (Thailand) Ltd. INS-VI Introduction of JAIMA and Investigation Project for Analytical Instruments

Related Industries

249

JAIMA

Size Exclusion Chromatography

SM Chemical

INS-VII Research laboratory for the production of high quality resorbable polymers 250

Polymer Research Laboratory, Department of Chemistry, Faculty of Science,

Chiang Mai University


1

Main Event PROGRAM

Thursday, June 18, 2015 (Day 1) 08:00-08:30 Registration (M Floor)

Ballroom A & B

08:30-08:40 Opening Remarks By Prof. Suwabun Chirachanchai President of The Polymer Society of Thailand (PST)

08:40-09:20 Plenary lecture I:

(Chair: Suwabun Chirachanchai) Prof. Andreas Greiner University of Bayreuth, Germany "Nanofibers by electrospinning – from a forgotten method to a major technique"

Refreshment Break

10:00-12:00 See Parallel Program

for details

Ballroom A Ballroom B Jamjuree 1 Jamjuree 2

Session: BIOEN1 Session: BIOEN2 Session: COMP1 Session: SMART1

Lunch@Citi Bistro (Ground Floor)

Ballroom A & B

13:00-13:10 PST Rising Star Awards Ceremony 1. Asst. Prof. Siripon Anantawaraskul Kasetsart University, Thailand

13:10-13:40 2. Asst.Prof. Panya Sunintaboon Mahidol University, Bangkok “Amine-decorated polymeric colloidal particles : from syntheses to applications”

Session: Instruments (Chair: Pasaree Laokijcharoen)

13:40-13:55 Bara Scientific Co. Ltd.

13:55-14:10 Crest Nanosolution (Thailand) Ltd.

14:10-14:25 Horiba (Thailand) Ltd.

14:25-14:40 JAIMA

14:40-14:55 LMS

14:55-15:10 SM Chemical

15:10-15:30 Research Laboratory for the Production of High Quality Resorbable Polymers Chiang Mai University

Refreshment Break

16:00-17:50 See Parallel Program

for details

Ballroom A Ballroom B Jamjuree 1 Jamjuree 2

Session: BIOEN1 Session: PROC Session: CHAR Session: SMART2


2

Main Event PROGRAM

Friday, June 19, 2015 (Day 2) 08:00-08:30 Registration (M Floor)

Ballroom A & B 08:30-09:10 Plenary lecture II:

(Chair: Suda Kiatkamjornwong) Prof. Michel Wong Chi Man Institute Charles Gerhardt Montpellier, France "Bridged polysilsesquioxanes: synthesis and application fields"

Refreshment Break

09:30-11:30 See Parallel Program for details

Ballroom A & B Jamjuree 1 Jamjuree 2

Session: COMP2 Session: RUBBER Session: PROC2

11:30-12:10 Polymer Society of Thailand

General Assembly

Lunch@Citi Bistro (Ground Floor)

Jamjuree 1 & 2

13:15-13:55 Plenary lecture III:

(Chair: Pattarapan Prasassarakich)

Dr. Piyada Charoensirisomboon Vice President, Innovation Campus Asia Pacific –Shanghai "Creating chemistry for sustainable future in Asia"

Session: Polymer Research in Industry Sector (Chair: Veerapat Tantayakom)

13:55-14:25 Dr. Narin Kaabbuathong PTT Research and Technology Institute, PTT Public Co., Ltd. ‘Development of bioplastics-based lamination application’

14:25-14:55 Dr. Sukhgij Ysothonsreekul PTT Global Chemical Public Company Limited “From Bioscience to Polymer Science: Our Sustainable Prospects”

14:55-15:25

Dr. Heng Soo Chin Agilent Technologies, Inc. ‘Optical characterization of thin films using a new Universal Measurement Accessory’

15:25-15:55 Dr. Nopphawan Phonthammachai SCG Chemicals Co., Ltd. ‘High Performance Composites for Industry’

15:55-16:20

Dr.Prae Chirawatkul, Dr. Wanwisa Limphirat, Assoc.Prof. Taweechai Amornsakchai Synchrotron Light Research Institute ‘Synchrotron light for innovative polymers’

16:20-16:35 Special guest

Dr. H. N. Chen Chair of International Activities Committee American Chemical Society “ACS International Activities and Collaboration”

Ballroom A & B

16:45-18:00 Poster Presentation

18:00-19:00 Poster Award Presentation

& Farewell Party


3

Parallel PROGRAM

BIOEN: Biomedical and Environmentally Friendly Polymers

Thursday, June 18, 2015 (Day 1)

Ballroom A BIOEN1 (Chair: Metha Rutnakornpituk) 10:00-10:20 BIOENO-01

A novel host-guest system and its supramolecular self-assembly and thermoresponsive micellization

Xia Song National University of Singapore, Singapore

10:20-10:40

BIOENO-02 Dual performances of benzoxazine dimers as metal ligand catalyst and as initiator for high efficient ring opening polymerization of lactide and branching poly(lactide)

Choltirosn Sutapin Chulalongkorn University, Thailand

10:40-11:10 KN-BIOEN-1 Supramolecular self-assembled polymers as novel biomaterials

Jun Li National University of Singapore, Singapore

11:10-11:30

BIOENO-03 Synthesis and characterization of medical grade poly(L- lactide-co-glycolide) for biomedical use as absorbable nerve guides

Pimwalan Techaikool Chiang Mai University, Thailand

11:30-11:50

BIOENO-04 Bioconjugation of anionic magnetite nanoparticle (MNP) with pyrrolidinyl peptide nucleic acid (PNA) for molecular biology technique

Sudarat Khadsai Naresuan University, Thailand

Ballroom B BIOEN2 (Chair: Warayuth Sajomsang) 10:00-10:20 BIOENO-5

Active ingredients with different water solubility loaded in fatty acid liposomes for sustained delivery

Han-Choi Yew University of Malaya, Malaysia

10:20-10:50 KN-BIOEN-2 Control of cell surfaces by polymer/protein LbL films for fabrication of 3D-human tissue models

Michiya Matsusaki Osaka University, Japan

10:50-11:10 BIOENO-06 Preparation and characterization of porous PEG/PEGDMA/GMA hydrogel scaffolds

Tharinee Theerathanagorn National Metal and Materials Technology Center, Thailand

11:10-11:30

BIOENO-07 Modulating the autofluorescence of silk to enhance analysis of cells and proteins by fluorescence imaging on silk-based biomaterials

Puay Yong Neo National University of Singapore, Singapore

11:30-11:50 BIOENO-08 Synthesis of positively charged poly(lactic acid) for preparation of electrospun fiber

Thanin Chalermbongkot Chulalongkorn University, Thailand

BIOEN continued


4


Ballroom A

BIOEN3 (Chair: Panya Sunintaboon)

16:00-16:30

KN-BIOEN-3 Enzymatic degradation of oil palm empty fruit bunch biomass

Rusli Bin Daik Universiti Kebangsaan Malaysia, Malaysia

16:30-16:50

BIOENO-09 Encapsulation of different log p anticancer drugs in 1,2- dioleoyl-sn-glycero-3-phosphoethanolamine-N-[methoxy- (polyethyleneglycol)-2000 (DOPE-PEG2000)-oleic acid liposome

Vicit Rizal Ehsuk University of Malaya, Malaysia

16:50-17:10

BIOENO-10 Study on covalent and ionic cross-linked in chitosan film by genipin and tripolyphosphate as potential material in medical applications

Siti Farhana Hisham Advanced Materials Research Centre (Amrec), Sirim Berhad, Malaysia

17:10-17:40 KN-BIOEN-4 Chitosan dispersion as a pharmaceutical coating material

Satit Puttipipatkhachorn Mahidol University, Thailand

Parallel PROGRAM

CHAR: Advances in Polymer Characterization


Jamjuree 1 (Chair: Taweechai Amornsakchai)

09:40-10:10

KN-CHAR-1 Preparation and properties of natural rubber with organic- inorganic nanomatrix structure

Seiichi Kawahara Nagaoka University of Technology, Japan

Jamjuree 1 (Chair: Kannika Sahakaro

16:00-16:20

CHARO-01 Long chain branching determination by triple-detector GPC

Thipphaya Pathaweeisariyakul SCG Chemicals, Thailand

16:20-16:40 CHARO-02 Mechanism of prevulcanization of isoprene rubber latex

Kewwarin Sae-heng Nagaoka University of Technology, Japan

16:40-17:10

KN-CHAR-2 Chemically controlled self-assembly of gold nanoparticles by site-selective protein immobilization: A model for antimalarial drug screening

Palangpon Kongsaeree Mahidol University ,Thailand

17:10-17:30 CHARO-03 The preparation and plausible structure of allylic bromination for phenyl-modified natural rubber

Nuorn Choothong Nagaoka University of Technology, Japan


5

Parallel PROGRAM

COMP: Polymer Composites and Nanocomposites


Jamjuree 1 COMP1 (Chair: Taweechai Amornsakchai)

10:10-10:40 KN-COMP-1 Interphase transfer of nanoparticles between immiscible polymer blends

Masayuki Yamaguchi Japan Advanced institute of Science and Technology, Japan

10:40-11:00 COMPO-01 Influence of pristine clay incorporation on strain-induced crystallization of natural rubber

Abdulhakim Masa Prince of Songkla University, Thailand

11:00-11:20 COMPO-02 Effects of organoclaytypes on morphological and mechanical properties of polyoxymethylene/polypropylene blends

Nipawan Yasumlee Silpakorn University, Thailand

11:20-11:50 KN-COMP-2 Hybrid porous polymers derived from octavinylsilsesquioxane

Hongzhi Liu Shandong University, China

Friday, June 19, 2015 (Day 2)

Ballroom A & B COMP2 (Chair: Chonlada Ritviruth)

09:30-09:50

COMPO-03 Study on model filler network in natural rubber matrix: Strain- induced crystallization behavior and dynamic mechanical Properties

Atitaya Tohsan Venture Laboratory, Kyoto Institute of Technology, Japan

09:50-10:10 COMPO-04 Preparation and characterization of TiO2/WO3/polythiophene composite

Nuttaporn Jaritkaun King Mongkut’s University of Technology Thonburi, Thailand

10:10-10:40

KN-COMP-3 Natural fiber reinforced rubber: recent advances toward high performance rubber matrix composites using pineapple leaf fiber

Taweechai Amornsakchai Mahidol University, Thailand

10:40-11:10

KN-COMP-4 Performance of aramid fiber in rubber compounds

Jutarat Phanmai Vice President - Marketing Trading Chemical Innovation Co., Ltd., Thailand

11:20-12:10 Polymer Society of Thailand- General Assembly

(All are welcome.)


6

Parallel PROGRAM PROC: Advances in Polymer Processing


Ballroom B PROC1 (Chair: Asira Fuongfuchat)

16:00-16:20

PROCO-01 Application of genetic algorithm in identifying ethylene/1- olefin copolymerization conditions from molecular weight distribution and chemical composition distribution

Uthane Nanthapoolsub Kasetsart University, Thailand

16:20-16:50 KN-PROC-1 Foam, (micro)foam, (nano)foam! - reality and dream

Masahiro Ohshima Kyoto University, Japan

16:50-17:10

PROCO-02 Determination of polymerization conditions for producing ethylene/1-olefin copolymers with tailor-made chain microstructures using artificial neural network

Thanutchoke Charoenpanich Kasetsart University, Thailand


Jamjuree 2 PROC2 (Chair: Kalyanee Sirisinha)

09:50-10:20 KN-PROC-2 Fiber design: A creation of fiber structure for feature and performance

Chureerat Prahsarn National Metal and Materials Technology Center, Thailand

10:20-10:40

PROCO-03 Simulation of morphological development during polymer crystallization: Effect of temperature gradient on the crystallization kinetics

Tharinee Teangtae Kasetsart University, Thailand

10:40-11:00 PROCO-04 Pressure slips casting: effect of pressure and time on green articles

Kittiya Jitklang King Mongkut's Universiti of Technology Thonburi, Thailand

11:00-11:20

PROCO-05 Comb-shaped polycarboxylate based copolymers with benzaldehyde derivative for molecular model of antimicrobial superplasticizer

Nalinthip Chanthaset Kasetsart University, Thailand

RUBBER: Natural and Synthetic Rubbers


Jamjuree 2 (Chair: Pranee Phinyocheep)

09:30-10:00 KN-RUBBER-1 New focus on rubber science and technology

Yuko Ikeda Kyoto Institute of Technology, Japan

10:00-10:20 RUBBERO-01 The use of modified palm oil as processing aids in tyre tread applications

Vorapot Thongplod Mahidol University

10:20-10:40

RUBBERO-02 Thermoplastic elastomers based on graft copolymers of natural rubber and poly(diacetone acrylamide)/polyamide-12

Gosalee Phersalaeh Prince of Songkla University Pattani campus, Thailand

10:40-11:00

RUBBERO-03 Thermoplastic vulcanizates based on natural rubber/propylene- ethylene copolymer blends; Influence of Viscosity and Ethylene content of the Copolymer on the properties

Toha Wohmang Prince of Songkla University Pattani campus, Thailand

11:00-11:20 RUBBERO-04 Morphology and properties of films prepared from different natural rubber clones

Treethip Phakkeeree Kyoto Institute of Technology, Japan


7

Parallel PROGRAM

SMART: Smart and Intelligent Polymers


Jamjuree 2 SMART1 (Chair: Robert Molloy)

09:40-10:10 KN-SMART-1 Coordination triggered division of vesicles

Yun Yan Peking University, China

10:10-10:40

SMARTO-01 Surface modification of polymer electrolyte membrane with heterocyclic brushes: a strategy to achieve effective proton transfer

Adisak Pokprasert Chulalongkorn University, Thailand

10:40-11:00 SMARTO-02 Preparation of microcapsules containing citronellal oil and galangal extract

Kankamon Sinpaksa Maejo University, Thailand

11:00-11:20

KN-SMART-2 Self-assembled polymer electrolytes for future electrochemical devices

Moon Jeong Park Pohang University of Science and Technology (POSTECH), Korea

11:20-11:40 SMARTO-03 Preparation of microcapsules containing citronellal oil and galangal extract

Benjawan Somchob Ubon Ratchatani University, Thailand

Jamjuree 2 SMART2 (Chair: Winita Punyodom)

16:00-16:30

KN-SMART-3 Non-ionic thermoresponsive polymers of UCST-type in water: challenges and perspectives

Seema Agarwal Universität Bayreuth, Germany

16:30-16:50 SMARTO-04 Layered-by-layered proton donor and acceptor polymers for effective and efficient proton transfer system

Chalanda Meemuk Chulalongkorn University, Thailand

16:50-17:10

SMARTO-05 Rapid reversible repeatable (RRR) mechanochromic-shape memory material: a unique combination of poly(ԑ-caprolactone) with melamine-benzoxazine network

Nattawat Yenpech Chulalongkorn University, Thailand

17:10-17:40 KN-SMART-4 Polymer-based smart devices: Electronics on paper, plastic and textile

Teerakiat Kerdcharoen Mahidol University, Thailand

Plenary Lectures

International Polymer Conference of Thailand 9

PL-I

Nanofibers by electrospinning – from a forgotten method to a major technique

Andreas Greiner*

Macromolecular Chemistry II of University of Bayreuth and Bayreuth Center for Colloids and Interfaces,

Universitätsstraße 30, 95440, Germany

Phone: +49 921 55-3399, Fax: +49 921 55-3393, E-mail: [email protected]

Abstract Electrospinning of polymer nanofibers was developed in the

early 19th

century but it took almost 80 years and until this unique

technique received within very few years enormous international

attention. Nowadays, electrospinning is the state of the state-of-the-art

technique for the preparation of polymer nanofiber nonwovens.

Numerous polymer systems were electrospun including water soluble

or dispersed polymers, polyelectrolytes, vinyl polymers,

polycondensates, biodegradable polymers, biopolymers, block

copolymers, blends, composites. As a result nanofibers of different

dimensions and shape were obtained including cylindrical, beaded,

barbed, core-shell and side-by-side fibers. Complex electrospinning

techniques led to e. g. macroscopically oriented nanofibers, three-

dimensional nonwovens, threads of nanofibers, core-shell nanofibers

or nanospring fibers. Further modifications of electrospun fibers can

be achieved by emulsion electrospinning, reactive electrospinning, or

chemistry on electrospun fibers. Near endless options are given by the

preparation of nanofiber composites including dyes and pigments,

catalysts, nano- and microparticles, carbon nanotubes and graphene,

pheromones, antibacterial compounds, drugs, enzymes, virus, living

bacteria, and cells. With this wealth of variation numerous

applications have been envisioned some of them, e. g. in air filtration,

are already realized. This contribution will discuss today´s challenges

of electrospinning of polymer nanofibers for real world applications

from the scientific point of view.

A selection of different shapes of electrospun nanofibers

For a review see: S. Agarwal, A. Greiner, J. H. Wendorff, Prog.

Polym. Sci. 2013, 38, 963.

Andreas Greiner

Date of Birth: 5. August 1959 Address: University of Bayreuth

Chair of Macromolecular Chemistry II Universitätsstraße 30, 95440 Bayreuth

Web: http://www.mcii.uni-

bayreuth.de/en/index.html

Education

1980 – 1986 Diploma in Chemistry,

Department of Chemistry, University of Marburg, Germany.

1986 – 1988 Ph. D. in Polymer Chemistry,

Department of Chemistry, University o of Marburg.

1989 – 1990 Postdoc at the University of

California, Santa Barbara, USA. 1990 – 1995 Habilitation for

Macromolecular Chemistry, Department of Chemistry, University of Marburg.

Professional experience

1995 – 1996 Adjunct professor, Department of Chemistry, University of Marburg.

1996 - 1999 University lecturer, Department of

Chemistry, University of Marburg. 1999 – 2000 Associate Professor for

Macromolecular and Organic Chemistry,

University of Mainz 2000-2012 Full professor for Macromolecular

Chemistry and Technology, University of Marburg

Since Oct. 2012 Full professor for

Macromolecular Chemistry, University of Bayreuth,

Since Nov. 2013 Head of Department “Future

Solutions” in New Materials Bayreuth

Research profile

General monomer and polymer synthesis,

reaction catalalysis, electrospinning of polymer nanofibers, polymer-functionalized

nanoparticles, artificial molecules, poly(p-

xylylene)s, functional polymer dispersions, polymers for coatings, filtration, textiles,

medicine, drug release, and agriculture, antibacterial, superhydrophobic polymers, light

weight foams, living composites, biobased

polymers.


PL-II

Bridged Silsesquioxanes: Synthesis and Application Fields

Michel Wong Chi Man

Laboratoire Architectures et Matériaux Nanostructurés, Institut Charles Gerhardt Montpellier, Ecole Nationale

Supérieure de Chimie de Montpellier, 34296 Montpellier (France)

Phone +33 467147219, Fax +33 467144322, *E-Mail: [email protected]

Abstract

Since the first reports on bridged silsesquioxanes (BS) in the

early nineties,1,2

growing interests are being paid towards this family

of silica-based hybrid materials which are obtained by the hydrolysis-

condensation (mainly by the sol-gel process) of organosilanes

precursors. The latter consists of at least two hydrolysable

trialkoxysilyl groups which are connected by bridging organic

functions. The combination of the organic component with the silica

network allows tuning their properties and to construct new smart

materials for targeted applications. BS thus prepared by a bottom-up

approach offer the possibility to obtain multifunctional materials for

challenging fields of interests provided the organosilylated precursors

are judiciously prepared.3

In this presentation, some examples of functional BS

conceived with sought-after properties for application in the following

fields of research will be discussed:

1 – Catalysis: the synthesis supported homogeneous catalysts

(organometallic and organic), their efficiency and recyclability will be

shown.4

2 – Structuring: the control of the structure of BS will be

demonstrated through self-assembly and molecular recognition.5-7

3 – Nanomedecine: mechanised BS can be triggered to deliver drugs

in cancer cells. 8,9

References:

1 - Shea, K.J.; Loy, D.A.: Webster, O.W., Chem. Mater, 1989, 1, 572 2 - Corriu, R.J.P.; Moreau, J.J.E.; Thépot, P.; Wong Chi Man, M., Chem. Mater., 1992,

4, 1217

3 - Bürglová, K.; Moitra, N.; Hodačová, J.; Cattoën, X.; Wong Chi Man, M., J. Org. Chem., 2011, 76, 7326

4 - Zamboulis, A.; Moitra, N.; Moreau J.J.E.; Cattoën, X.; Wong Chi Man, M., J.

Mater. Chem. 2010, 20, 9322 5 - Moreau, J.J.E.; Vellutini, L.; Wong Chi Man, M.; Bied, C., J. Amer. Chem. Soc.,

2001, 123, 1509

6 - Croissant, J.; Cattoën, X.; Wong Chi Man, M.; Dieudonné, P.; Charnay, C.; Raehm, L.; Durand, J-O., Adv. Mater., 2015, 27, 145

7 - Arrachart, G.; Creff, G.; Wadepohl, H.; Blanc, C.; Bonhomme, C.; Babonneau, F.;

Alonso, B.; Bantignies, J-L.; Carcel, C.; Moreau, J.J.E.; Dieudonné, P.; Sauvajol, J-l.; Massiot, D.; Wong Chi Man, M., Chem. Eur. J. 2009, 15, 5002

8 - Théron, C.; Gallud, A.; Carcel, C.; Gary-Bobo, M.; Maynadier, M.; Garcia, M.; Lu,

J.; Tamanoi, F.; Zink, J. I.; Wong Chi Man, M., Chem. Eur. J., 2014, 20, 9372 9 - Croissant, J.; Maynadier, M.; Mongin, O.; Hugues,V.; Blanchard-Desce, M.; Chaix,

A.; Cattoën, X.; Wong Chi Man, M.; Gallud, A.; Gary-Bobo, M.; Garcia, M.; Raehm,

L.; Durand, J-O. Small, 2015, 11, 295

Michel Wong Chi Man

Academic Qualifications

1987 - PhD in Chemistry (University

Montpellier 2- France)

2003 - Habilitation

Post-Doctoral Positions

Dec.1987-Dec.1988: (CNRS fellowship)

Laboratoire de Chimie Organométallique: University Montpellier 2 – France

March1989-Feb.1990: (Alexander von

Humboldt foundation fellowship) laboratory "Anorganisch-Chemisches Institut der

Universität" University of Heidelberg -

Germany

Employment History

March-Sept.1990: Associate Professor (University Montpellier 2 – France)

Oct.1990-Sept.2003: – Junior CNRS

scientist - Chargé de Recherches (University Montpellier 2 and Ecole Nationale Supérieure

de Chimie de Montpellier)

2003-2010: – Senior CNRS scientist - Directeur de Recherches 2ème classe (Institut

Charles Gerhardt Montpellier)

2010-present: – Senior CNRS scientist - Directeur de Recherches 1ère classe (Institut

Charles Gerhardt Montpellier)

Awards and Invited Positions

1989-1990: Alexander von Humboldt

fellowship July 2004: Invited Professor (University

Autonoma de Barcelona – Barcelona, Spain)

December 2006: Guest Scientist (Australian Nuclear Scientific & Technology

Organisation – Lucas Heights, Australia)

December 2011: Invited Scientist (University of Western Sydney – Sydney,

Australia)

September 2012: Guest Scientist (National Institute of Materials Science – Tsukuba, Japan)

14th July to 12th August 2014: Alexander

von Humboldt funding for 1 month stay in Germany (visits to Humboldt Univ zu Berlin,

Technische Univ of Berlin, Kiel Univ., Univ. of

Heidelberg, Univ des Saarlandes, Univ Kiel) 1st – 28th February 2015: FNRS Invited

sabbatical Scientist (University of Liège –

Liège, Belgium) 21st September – 21st December 2015:

Laureate of the "Programme de Chaires Franco-

Brésiliennes dans l’état de São Paulo" (São Paulo, Brazil)


PL-III

Creating Chemistry for Sustainable Future in Asia

Piyada Charoensirisomboon

BASF Advanced Chemicals Co., Ltd., GM/S 200137 Shanghai, China

Abstract

In 2050, more than nine billion people will live on our

planet. The world population and its demands will keep growing,

while the planet’s resources are finite. BASF looks ahead how we as a

company contribute to a sustainable future. BASF will continue to

develop, to innovate, to meet new challenges, to take advantage of

new opportunities and to succeed. Future-oriented innovation requires

market-driven research and development. BASF is building up

research platforms with a new focus and is creating a more global

research and development organization. This will enable BASF to

specifically address customer needs even better. The first phase of the

Innovation Campus Asia Pacific Shanghai has been inaugurated at the

end 2012 and the expansion of phase two is in good progress.

Innovation Campus Asia Pacific Shanghai is one of major steps on

BASF globalization of R&D. The Innovation Campus will help

intensifying the development of local scientific and technical talent

and to foster collaboration with universities and scientific institutes in

Asia Pacific. Scientists in close proximity to local markets work in

international and multi-disciplinary project teams. Example of

innovation will be given to restate the vision of creating innovation

from Asia for Asia and also for the world.

Piyada Charoensirisomboon

Vice President – Advanced Materials Innovation Campus Asia Pacific -Shanghai

• Strategic Business Management : St. Gallen

Business School

• Ph.D –Polymeric Materials, Tokyo Institute of

Technology

• Master- Materials Science & Engineering, Tokyo Institute of

Technology

• BSc –1st class honor, Industrial Chemistry KMITL

Professional Career at BASF

2013 Vice President – Innovation Campus Asia Pacific-Shanghai BASF Advanced

Chemicals Co, Ltd Shanghai, CHINA

2010 Global Business Manager – Internal

Start-Up, New Business Development-

Performance Chemicals BASF SE, GERMANY

2005.Head of Global R&D Styrenic Thermoplastics Group Global Polymer

Research BASF SE, GERMANY

2001 Regional Marketing Manager BASF

South East Asia Pte Ltd., SINGAPORE

2000 Research Fellow, BASF AG, GERMANY

Patents & Sci Publication: More than 25

patent & patent applications and 22 publications

PST Rising Star Award Lecture


RS

1. Asst.Prof.Dr. Siripon Anantawaraskul Department of Chemical Enginering

Faculty of Engineering, Kasetsart University

B.Eng. (2nd Honour, Chem. Eng.), Kasetsart University

M.Eng. (Chem. Eng.), McGill University

Ph.D. (Dean's Honour List, Chem. Eng.), McGill University

Visiting Scholar (Chem. Eng.), Univerity of Waterloo

Research areas:

Polyolefin Chain Microstructure and Characterization

Polyolefin Reaction Engineering

Modelling and Simulation in Polymer Science and Engineering

2. Asst.Prof.Dr. Panya Sunintaboon

Graduate Program Director (Polymer Science and Technology),

Department of Chemistry Faculty of Science, Mahidol University

B.Sc. (Chemistry) Chulalongkorn University (1997)

M.Sc. (Organic Chemistry) Chulalongkorn University(1999)

Ph.D. (Polymer Science) University of Akron, USA (2004)

Research areas:

Fabrication of amine-functionalized polymeric particles; Emulsifier-free emulsion polymerization.

Amine-decorated Polymeric Colloidal Particles: From Syntheses to Applications

Panya Sunintaboon

*

1 Polymer Science and Technology Program, Department of Chemistry, Faculty of Science, Mahidol University,

Salaya, Nakhon Pathom 73170, Thailand

Phone: +66-2441-9816, Fax : +66-2441-0511, *E-mail: [email protected]

Abstract

Polymeric colloidal particles bearing amine groups on their outer peripheries have attracted great deal

of attention because of their versatility for several applications, such as biomedicine (targeting drug/gene

delivery, diagnostics, and tissue engineering), chemo- or bio-sensor, catalysis, coatings, wastewater treatment ,

and so on . Thus, the fabrication of such particles with controllable characteristics (e.g. size and size

distribution) and well-defined amine accessibility is quite challenging and desirable. In this present work, the

syntheses of amine-functionalized colloidal particles via several pathways are illustrated. A wide variety of

particles’ physical and chemical properties (e.g. rigid, soft, pH-sensitive, thermo-sensitive, water swellable,

magnetic, or biocompatible) can possibly be tailored. In addition, some promising applications of the resulting

amine-functionalized colloidal particles prepared from these synthetic methods are presented.

Keywords: colloidal particle, amine-functionalized, scaffold, carrier, surfactant-free, microgel.

Biomedical and Environmentally Friendly Polymers

SESSION 1


15

KN-BIOEN-1

Supramolecular Self-assembled Polymers as Novel Biomaterials

Jun Li1,2*

1Department of Biomedical Engineering, National University of Singapore, Singapore 117574

2Institute of Materials Research and Engineering, A*STAR, Singapore 117602

Phone +65 6516 7273, Fax +65 6872 3069, E-Mail: [email protected]

Abstract The phenomena of molecular self-assembly have inspired

interesting development of novel functional materials for various

applications. Recently, we have successfully demonstrated the methods

for constructing self-assembled macromolecular systems based on

amphiphilic block copolymers and interlocked cyclodextrins (-CD, -

CD, and -CD), which can function as nano-carriers for potential drug

and gene delivery [1]. We developed a series of supramolecular

hydrogels formed by -CD and various triblock copolymers comprising

PEG and hydrophobic polyester blocks for controlled release of drugs,

as well as gene delivery. Amphiphilic star-block copolymers based on

polyester and PEG with adamantyl end-functionalization were

synthesized, which self-assembled into nanogel-like large compound

micelles, and transformed into vesicular nanostructures under the

direction of host-guest interaction between the adamantyl end and

dimethyl-β-CD [2]. The intracellular uptake of anticancer drug-loaded

nano-vesicles indicates that the nanovesicles could be potential drug

carriers for cancer therapy. We developed novel cationic

supramolecules self-assembled from cyclodextrins and block

copolymers as a new class of polymeric gene delivery vectors [3,4]. We

also developed redox-sensitive and targeted gene delivery systems for

cancer therapy, and multifunctional hybrid nano-carrier for

simultaneous dual therapeutics delivery and cellular imaging [5-6].

Other supramolecular polymer self-assembled nanostructures

developed in our lab include supramolecular nanocapsules based on

threading of CDs over polymer on gold nanoparticles [7],

supramolecular hydrogels formed by pyrene-terminated PEG star

polymers and -CD [8], and the star-star supramolecular architecture

and its thermosensitive hydrogel formation [9]. Most recently, we

demonstrated a supramolecular approach for building a multifunctional

gene carrier system with the functions of reduction-responsive

degradation and zwitterionic phosphorylcholine based extracellular

stabilization and favorable cellular uptake, and the supramolecular gene

carrier was applied to deliver the therapeutic p53 anti-cancer gene in

MCF-7 cells, showing great potential for cancer gene therapy

application. Compared to traditional covalent conjugation approach, the

supramolecular approaches are more convenient in building

complicated architectures with multiple functionalities integrated within

one system [10].

Keywords: Supramolecules, self-assembly, cyclodextrins, biomaterials

References [1] Li J and Loh XJ, Adv. Drug Deliv. Rev. 60, 1000 (2008).

[2] Zhu J, Liu KL, Zhang ZX, Zhang ZX, Li J, Chem. Commun. 47, 12849 (2011). [3] Li J, Yang C, Li HZ, et al., Adv. Mater. 18, 2969 (2006).

[4] Ping Y, Liu CD, Zhang ZX, Liu KL, Chen JH, and Li J, Biomaterials 32, 8328

(2011). [5] Ping Y, Hu Q, Tang G, and Li J, Biomaterials 34, 6482 (2013).

[6] Zhao F, Yin H, Li J, Biomaterials 35, 1050 (2014).

[7] Wu YL, Li J, Angew. Chem. Int. Ed. 48, 3842 (2009). [8] Chen B, Liu KL, Zhang ZX, Ni X, Goh SH, Li J, Chem.Commun. 48, 5638 (2012).

[9] Zhang ZX, Liu KL, Li J, Angew. Chem. Int. Ed. 52, 6180 (2013).

[10] Wen Y, Zhang Z, Li J, Adv. Funct. Mater., 24, 3874 (2014).

Jun Li

Department of Biomedical Engineering Faculty of Engineering

National University of Singapore

Singapore

Dr. Jun Li received his MSc in 1992

and PhD in 1995 in Macromolecular Science from Osaka University, Japan. From 1995 to

1998 he was a Special Postdoctoral

Researcher at RIKEN Institute in Japan. In 1998, he joined the Institute of Materials

Research and Engineering in Singapore as a

Research Scientist. From 2002, he became an Assistant Professor at Department of

Biomedical Engineering, National University

of Singapore, and was promoted to Associate Professor in 2007 and Professor in 2015. His

research interests include novel

supramolecular structures and block

copolymers as functional materials for

biomedical applications. He has developed

novel macromolecules with the ability to self-assemble into supramolecular structures

based on cyclodextrin/polymer complexes

and amphiphilic biodegradable block copolymers, as biomaterials (hydrogels,

nano-particles, micelles, nano-vesicles,

micro- and nano-encapsulation, surface coatings, etc.) for various applications such as

drug and gene delivery, and tissue

engineering. He has published 150 papers in SCI-indexed international journals, which

have received more than 8,600 citations with

an h-index of 49. He also holds 9 patents and a few book chapters.


16

KN-BIOEN-2

Control of Cell Surfaces by Polymer/Protein LbL Films for Fabrication of 3D-Human

Tissue Models

Michiya Matsusaki and Mitsuru Akashi

Graduate School of Engineering, Osaka University

2-1 Yamadaoka, Suita, Osaka 565-0871, Japan

E-mail: [email protected]

Abstract In vitro development of highly-organized three dimensional

(3D)-engineered tissues consist of multiple types of cells and ECM,

which possess a similar structure and function as natural tissues, is a

key challenge for tissue engineering and pharmaceutical assay.

Especially modulation of 3D-cell-cell interaction inside the 3D-

artificial tissues is one of the significant issues.

We have developed a simple and unique bottom-up approach,

“hierarchical cell manipulation”, using nanometer-sized Layer-by-

Layer (LbL) films consisting of fibronectin and gelatin (FN-G) as a

nano-extracellular matrix (nano-ECM) (Fig. 1) [1-5]. The FN-G

nanofilms were prepared directly on the cell surface, and we

discovered that at least 6 nm thick FN-G films acted as a stable

adhesive surface for adhesion of the second cell layer. We have also

developed a rapid bottom-up approach, “cell-accumulation

technique”, by a single cell coating using FN-G nanofilms, because

the fabrication of two-layers (2L) was limitation through the above

technique due to the time required for stable cell adhesion [6-9]. This

rapid approach easily provided more than twenty-layered (over 150

µm) 3D-tissues after only one day of incubation. Moreover, fully and

homogeneously vascularized tissues of 1 cm width and 100 µm height

were obtained by a sandwich culture of the endothelial cells. The

hierarchical cell manipulations will be promising to achieve one of the

dreams of biomedical field, in vitro automatic creation of artificial

3D-tissue models [10]. We are demonstrating in vitro reconstruction

of metastasis early processes, invasion, intravasation, mobilization

and extravasation of human invasive carcinomas using artificial 3D-

blood and lymphatic capillaries.

Michiya Matsusaki

Present Position:

Assistant Professor, Department of Applied

Chemistry, Graduate School of Engineering,

Osaka University

Education:

B.S. from Kagoshima University, March 1999 M.S. from Kagoshima University, March 2001

Ph.D. in Engineering, from Kagoshima

University, September 2003 (short period)

Academic Appointment:

April 2003 to March 2005: Japan Society for the Promotion of Science (JSPS)

Postdoctoral Research Fellow

January 2004 to March 2004: Visiting Scientist in Lund University (Prof. Carl A.K.

Borrebaeck Lab.) by Japan-Sweden Young

Researcher Exchange Program April 2005 to July 2006: Designated

Assistant Professor, Department of Applied

Chemistry, Graduate School of Engineering, Osaka University

August 2006 to present: Assistant

Professor, Department of Applied Chemistry, Graduate School of Engineering, Osaka

University

October 2008 to March 2011: PRESTO Researcher, JST (Concurrent position)

April 2012 to present: 18th Council member

of Japanese Society for Biomaterials Jan 2013 to present: Editorial Board

Member of PLoS ONE (PLOS Group)

April 2015 to present: Editorial Board Member of Scientific Reports (Nature

Publishing Group)

Academic Activities:

Members of The Society of Polymer Science,

Japan, The Chemical Society of Japan, Japanese Society for Biomaterials, American Chemical

Society, The Japanese Society of Artificial

Organs, The Japanese Society for Regenerative Medicine, The Kinki Chemical Society Japan,

and Research Group on Biomedical Polymers

Current Areas:

Polymer Science, Biomaterial, and Tissue

Engineering

Fig. 1. Schematic illustration of (a) hierarchical cell manipulation and

(b) cell accumulation technique.


17

KN-BIOEN-3

Enzymatic Degradation of Oil Palm Empty Fruit Bunch Biomass

Satriani Aga Pasma, Rusli Daik, Suria Ramli and Mohamad Yusof Maskat

School of Chemical Sciences and Food Technology

Faculty of Science and Technolgy

Universiti Kebangsaan Malaysia

43600 UKM Bangi, Selangor, Malaysia

[email protected]

Abstract

The objective of the study is to optimize the production of lignin

degradation products. Lignin from oil palm empty fruit bunch

(OPEFB) was extracted by using organosolv method and

directly isolated by three methods of isolation. Powder of lignin

was isolated from organsolv black liquor by using methanol,

acidified water and deionized water. Enzymatic hydrolysis was

carried on the lignin powder using laccase and cutinase. The

reaction was conducted in an incubator shaker for 24 hours with

phenol, water, and buffer as mediators. A total of 9 compounds

were found as Lignin OPEFB degradation products. They were

hydroxybenzoic acid, hydroxybenzaldehyde, vanillic acid,

vanillin, syringic acid, syingaldehyde, coumaric acid, ferulic

acid, and guaiacyl alcohol. Different mediators affected the

yield of degradation products. In water, ferulic acid was the

product found with the highest concentration (466 mg/L), and

this was followed by hydroxybenzoic acid (201 mg/L).

Whereas, vanillic acid was the product with the highest

concentration (126 g/L ) found in phenol. Guaiacyl alcohol was

detected in small amount when laccase was used in water and

phenol. For the cutinase, major compounds produced were

syringaldehyde (2493 mg/L) and syringic acid (4994 mg/L). For

the characterization of lignin and degradation products,

Thermogravimetric Analysis (TGA), Fourier transform infrared

(FTIR), and Field emission scanning electron microscope

(FESEM) were used. High Performance Liquid Chromatography

(HPLC) and Gel Permeation Chromatography (GPC) were used

to determine the quantity and molecular weight of degradation

products produced.

Professor Dr. Rusli Daik graduated from

Universiti Kebangsaan Malaysia (The

National University of Malaysia) majoring

in Chemistry. He obtained his PhD degree in

Polymer Synthesis from Durham University,

United Kingdom. A part from polymer

synthesis his research interest includes

polymer nanoparticle, polymer

nanocomposite, electroactive polymer,

biodegradable polymer from biomass-

derived monomer, colloidal polymer and

polymer nanofluid. Throughout his career he

published more than 200 research

manuscripts, and owned three patents. He

has also edited 5 books and written 15 book

chapters. He has received more than 20

awards / recognitions from national as well

as international organizations. He is an

Associate Editor of the Malaysian Polymer

Journal (published by Plastics and Rubber

Institute of Malaysia), and the Chief Editor

of the Journal of Polymer Science and

Technology (published by Polymer

Research Center, Universiti Kebangsaan

Malaysia). He is currently the Deputy Dean

(Research and Innovation) of the Faculty of

Science and Technology, Universiti

Kebangsaan Malaysia. He is also the

President (founding president) of the

Malaysia Polymer Society, and a Council

Member of the Federation of Asian Polymer

Societies representing Malaysia since 2012.


18

KN-BIOEN-4

Chitosan Dispersion as a Pharmaceutical Coating Material

Satit Puttipipatkhachorn1,2

1Department of Manufacturing Pharmacy, Faculty of Pharmacy, Mahidol University, Bangkok 10400

2Center of Excellence in Innovative Drug Delivery and Nanomedicine, Faculty of Pharmacy, Mahidol

University, Bangkok 10400

Phone +66 2644 8702, Fax +66 2644 8702, *E-Mail: [email protected]

Abstract

Chitosan is derived from chitin, the second most abundant

biopolymer in the nature. It is used as an excipient for pharmaceuticals;

for example, film former, binder, controlled releasing agent and a

carrier in drug delivery system. There is a drawback in using chitosan

as a pharmaceutical coating material as an acid is needed in preparation

of chitosan solution. This leads to a remaining of residual acid in

chitosan film which might activate a degradation of acid labile drugs

and cause an unpleasant smell as well as a damage of coating

equipment. These problems can be solved by converting conventional

chitosan solution to chitosan colloidal dispersion.

Aqueous chitosan colloidal dispersion composed of chitosan,

cetyl alcohol and polyvinyl alcohol were developed. Free films were

prepared using casting evaporation technique. Residual acetic acid

residue was analyzed by HPLC. The mechanical properties, water

vapor permeability, water uptake, weight loss and drug permeability

across free films were determined. In addition, the aqueous chitosan

colloidal dispersion was used to coat on acetaminophen tablets and drug

release studies in simulated gastric fluid and simulated intestinal fluid

were carried out. All experiments were compared with chitosan solution

in acetic acid. The free films with minimal remaining acid residue could

be obtained from aqueous chitosan colloidal dispersion. These free

films had higher tensile strength, higher percent elongation at break,

lower water vapor permeability, lower water uptake and weight loss,

and lower drug permeability than those prepared from chitosan solution

in acetic acid. The tablets coated with aqueous chitosan colloidal

dispersion had slower drug release rate in both simulated gastric fluid

and simulated intestinal fluid than those coated with the chitosan

solution in acetic acid.

In conclusion, the aqueous chitosan colloidal dispersion

provided free films with better characteristics such as improved

mechanical properties and reduced drug permeability, when compared

to the chitosan solution in acetic acid. The study demonstrated that

aqueous chitosan colloidal dispersion can be used as pharmaceutical

film coating material to produce modified release dosage form.

Keywords: Chitosan dispersion, Coating, Film, Drug permeability,

Drug release

Satit Puttipipatkhachorn

Dr. Satit Puttipipatkhachorn is currently

the Head of Department of Manufacturing Pharmacy and associate professor in

Pharmaceutics at Faculty of Pharmacy,

Mahidol University, Thailand. He obtained a B. Pharm. with the first class honors from

Chiang Mai University, Thailand (1979-

1984), M.Sc. in Industrial Pharmacy from Mahidol University, Thailand (1984-1987),

and Ph.D. in Pharmaceutical Sciences from

Chiba University, Japan (1987-1991). During Ph.D. study, he received the Japanese

Government Scholarship (Monbusho). After

graduation, he was trained at Sankyo Pharmaceutical Co.Ltd., and gained a short

experience in pharmaceutical R&D and

GMP. He started his academic career as a lecturer at Department of Manufacturing

Pharmacy, Faculty of Pharmacy, Mahidol

University in 1991 and was promoted to be Assistant Professor and Associate Professor

in 1993 and 1997, respectively. For

administrative work, he has been Head of Department of Manufacturing Pharmacy

(2001-2004, 2014-present) and Deputy Dean

on Graduate Studies (2004-2008). His research interest is in the area of

Soild Pharmaceutics, especially

physicochemical properties of drug substances and excipients, drug-polymer

interaction, relevance of physicochemical

properties and molecular interaction in the solid dosage form on drug product

performance including dissolution, stability and recently nanoparticle formation. Another

area of research is oral controlled-release

drug delivery system, nanoparticulate drug delivery system and new pharmaceutical

excipients from polysaccharides. To present,

he has published over 90 original articles in international journals. With recognition in his

research achievement, he received the

Ishidate Award in Pharmaceutical Research from the Federation of Asian Pharmaceutical

Associations (FAPA) in 2004, and also the

Research Award from the National Research Council of Thailand in 2005, 2006 and 2008.

Apart from academic work, he also gave

a contribution to pharmacy profession of Thailand as a secretary general of

Pharmaceutical Association of Thailand

(1992-2010).


19

BIOENO-08

Synthesis of Positively Charged Poly(Lactic Acid) for Preparation of Electrospun Fiber

Thanin Chalermbongkot, Worawan Bhanthumnavin and Varawut Tangpasuthadol

*

Organic Synthesis Research Unit, Department of Chemistry, Faculty of Science, Chulalongkorn University,

Bangkok 10330

Abstract

Poly(lactic acid), a biodegradable and biocompatible polyester, is used widely in many applications. The

high hydrophobic characteristic of PLA is, however, a drawback for some works that directly contact with water

such as drug delivery, or tissue engineering. Therefore, in order to enhance the hydrophilicity of PLA

electrospun fiber, PLAs having two positively charged end groups (PLAdi+) were synthesized by incorporating

glycidyl trimethylammonium chloride (GTMAC) into the polymer chain ends. Commercially available PLAs

doped with three types of PLAdi+, different in chain lengths, were electrospun to afford improved

hydrophilicity PLA fiber mats. The fiber diameter of PLA doped with 10 %wt PLAdi+ was found to decrease

with increasing amounts of the dopant, or with decreasing the molecular weight PLAdi+ used, as determined by

SEM of the fibers. Moreover, the hydrophilicity of the PLA doped with PLAdi+ was increased compared with

corresponding PLA fiber without doping, as measured by air-water contact angle measurement. The thin and

hydrophilic PLA fibers were successfully prepared and could potentially be used in applications related to

aqueous environment.

Keywords: Poly(lactic Acid), Electrospinning, Hydrophillicity, Quaternary ammonium salt.

1. Introduction

In recent years, poly(lactic acid) (PLA) has

received much attention due to its biodegradable and

biocompatible properties, which provide important

economic benefits. PLA is a biopolymer and renewable

polyester, which has been widely used in several

applications such as packaging materials, biomedical

materials, and fibers. However, PLA is highly

hydrophobic which provide less efficiency when used in

biomedical and biomaterial field that related to aqueous

media. Therefore, PLA having positive charges in its

structure should be one way to increase its

hydrophilicity, and thus incorporating other benefits such

as bactericidal properties.

Electrospinning is a key versatile method to

produce the non-woven nanofibers providing high

surface area. Using electrospinning technique, smaller

size and high surface area fibers were electrospun, which

are applied to various applications, for example, scaffold,

water filter, and wound dressing bandage.

Consequently, in this work, the PLA was

positively charged by modifying the polymer chain ends

via ring opening reaction between the chain ends of PLA

and GTMAC. Subsequently, the PLAdi+ was mixed into

commercial-grade PLA. The polymer mixture was then

electrospun into the nanofibers. Polymer concentration,

amounts of PLAdi+ dopant, and PLAdi+ species were

varied to assess the effect on electrospinning and

morphology of the fibers, which the no bead-like and the

smaller fibers will be obtained.

2. Experimental Methods

2.1. PLAdiCOOH Synthesis

The polymers were prepared as shown in Scheme

1 using a revised synthesis method as described in

previous work [1]. Firstly, 88 wt% lactic acid solution,

succinic acid, and half-portioned para-toluenesulfonic

acid (pTSA) as catalyst were weighed and put into three-

neck round bottom flask. Then, the reaction was firstly

carried out at 110OC with partial reduced pressure for 2

to 10 hours. This step was called dehydration or

oligomerization. The reaction was further proceeded to

140OC, 160

OC for 1 and 2 hours, respectively, with step-

wise reducing the pressure to remove water from the

reaction. The polymerization was further carried out at

180OC for 4 to 12 hours with reduced pressure until

reaching high vacuum while SnCl2.2H2O and another

half-portioned pTSA were added into the flask as co-

catalyst and catalyst. After the reaction was finished, the

polymer was cooled down and dissolved in


20

dichloromethane before precipitating in cold ethanol for

purification. The crude polymer was precipitated twice,

fully dried in vacuum and then characterized by nuclear

magnetic resonance (NMR) and gel permeable

chromatography (GPC) and finally calculated for their

molecular weight and yield.

2.2. PLAdi+ Synthesis

The synthesis scheme of PLAdi+ is shown in

Scheme 1. PLAdiCOOH was weighed into a three-neck

round bottom flask with 3 equivalents GTMAC, 1

equivalent triethylamine, and DMSO as solvent. The

flask was mounted with CaCl2 tube to trap the moisture

from the reaction. The reaction was performed at 70OC

for 24 hours. The resulting polymer was washed with DI

water and then centrifuged 3-5 times in order to purify

the polymer and remove the unreacted GTMAC before

freeze-dried and characterized by NMR.

2.3. Electrospinning

Commercial PLA, Mw = 103 kDa (Ingeo™ 4043D,

NatureWorks LLC) served as the major component of

the polymer blends due to its high molecular weight and

consequent high viscosity to afford the chain

entanglements and the ability to be electrospun

efficiently. The PLAdi+, were used as the dopants for the

electrospinning. All polymers were thoroughly dissolved

in chloroform/methanol mixture (3:1) before loading into

a 5 ml plastic syringe equipped tightly with disposable

26G blunt needle tip (0.45 mm diameter). The syringe

was immediately placed into the syringe pump and the

solution was electrospun horizontally at 1 ml/h flow rate,

20 kV voltage applied from a high voltage power supply,

and tip-to-collecter distance was 20 cm. The electrospun

fiber jet was collected consequently onto a 10cm x 10cm

aluminum foil as a ground collector. Electrospun mats

were carefully left to dry at ambient temperature

overnight and further dried via high vacuum for 4 hours

before investigation.

2.4. Characterization and Morphology Investigation

The synthesized PLAdiCOOH and PLAdi+ were

characterized by NMR and GPC. 1

H NMR integrations

of the methylene protons of succinic acid and PLA

methine protons were compared to determine the

repeating unit and molecular weight as well.

To assess the fiber size and diameter, fiber

morphology was characterized using a JEOL JSM-

6610LV field emission scanning electron microscope

(SEM). The reported fiber diameters of each sample

were averaged from 80 different measuring points in the

SEM micrographs.

2.5. Air-water Contact angle measurement

In order to determine the wettability and

hydrophilicity of prepared PLA electrospun mats, static

and dynamic contact angle were measured using Ramé-

hart™ standard goniometer. The water droplets were

carefully controlled to be equal in each sample.

Scheme 1. Synthesis pathway of PLAdiCOOH and PLAdi+


21 3. Results and Discussion

3.1. Synthesis

The dicarboxyl ends PLAs, PLAdiCOOHs, were

synthesized from lactic acid monomer and succinic acid

at 110OC to 180

OC using tin-catalyzed polycondensation

(Scheme 1). Dehydration time, mol% of succinic acid

added into reaction, and polymerization time were varied

to produce PLAdiCOOHs with high molecular weight as

confirmed by using 1H NMR and GPC analysis.

Comparison of peak integrations between the methylene

proton (-CH2-) of succinic acid and methine proton (-

CH-) of lactic unit in PLA backbone was used to

calculate the molecular weight as shown in Fig. 1. From

the results, it seems that the PLAdiCOOH molecular

weight increased as the mol% SA decreased, or as

polymerization time increased.

The PLAdiCOOH were further modified into

PLAdi+ with molecular weight of 4550, 6955, and

16,802 Da, as PLAdi+ dopants P1, P2, and P3,

respectively. The degree of substitution (DS%) of

GTMAC on PLAdiCOOH was calculated using the

integration of methyl proton signal of GTMAC as shown

in Eq. 1.

(1)

where a is methyl proton signal of quaternary ammonium

group, b is methine proton signal of usual PLA, c is

methylene proton signal of succinic acid.

Table 1. Molecular weight and Degree of substitution of

PLAdi+

Figure 1. 1H NMR spectra of PLAdiCOOH and PLAdi+

in CDCl3 show PLAdi+ was successfully synthesized. –

where a is methyl proton signal of quaternary ammonium

group, b is methine proton signal of usual PLA, c is

methylene proton signal of succinic acid, and d is

methine proton signal of lactic unit in PLA chain.

The DS of PLAdi+ were 98% (P1), 56% (P2), and

64% (P3) (see Table 1). Since P1 has the shortest

polymer chain, it therefore has the highest numbers of

carboxyl chain ends, creating more opportunity for the

carboxyl group to react with GTMAC.

3.2. Electrospinning

PLA/PLAdi+ solutions for electrospining were

prepared according to the amounts shown in Table 2.

Methanol was added into the polymer solutions in order

to raise their conductivity for providing the smaller

diameter fibers [2]. Specifically, we investigated how the

PLA concentration, dopant species, and dopant

concentration affected the fiber diameter and wettability

of the fiber mats.


22 Table 2. Electrospinning conditions of PLA and

PLA/PLAdi+

According to Figure 2, SEM images of

electrospun fibers reveal that the fiber diameter

decreased apparently as the polymer concentration was

reduced, due to the low polymer content and was

consistent with results obtained in previous studies [3].

As a results, PLA concentration at 5%w, entry 2, was

chosen for further test because it gave acceptably small

fiber diameter.

0.604 ± 0.121µm 0.688 ± 0.167µm 1.090 ± 0.146µm

Figure 2. SEM images of PLA fibers show effect of

concentration on the fiber diameters, the fiber diameter

increases with increasing the polymer concentration. (a)

entry 1, 3% PLA; (b) entry 2, 5% PLA; and (c) entry 3,

7% PLA.

Electrospun fibers from PLA doped with three

types of PLAdi+ are shown in Figure 3. The PLA/P1 mat

shows the smallest fiber diameter whereas PLA/P2 fiber

showed the largest fiber diameter. This could be

attributed to the fact that P1 was the positively-charged

PLA with the lowest MW prepared in this work. The

amount of positive charges in P1 was therefore the

highest comparing to the others. The high charge density

in PLA/P1 would improve the conductivity of the

polymer solution, resulting in smaller fiber diameter.

Lastly, in this work, it was hypothesized that

increasing amount of dopants would reduce the apparent

fiber diameter to provide the smaller size fibers and

would improve PLA hydrophilicity as well. The PLAdi+

P1 was selected and used as the dopant in this section,

because its high positive charge density. The P1 doping

Figure 3. SEM images of PLA/PLAdi+ electrospun

fiber- Tthe influence of PLAdi+ dopant species on fiber

diameters. (a) PLAdi+ P1 (4.5 kDa) (b) PLAdi+ P2 (7.0

kDa) (c) PLAdi+ P3 (16.8 kDa).

Figure 4. SEM images of PLA/PLAdi+ electrospun

fibers which %dopants were varied from 10%(a),

20%(b), 30%(c), 40%(d) to 50%(e). The figure shows

that the fiber diameter was reduced while the %dopant

increased.


23 content was varied from 10-50%. As illustrated by

Figure 4, the fiber diameter was reduced while the

percentage of dopant increased. These results indicate

that increasing amounts of PLAdi+ decrease the fiber

diameter because higher dopants contents induced the

higher conductivity of the polymer solution, therefore,

the solution jet was purged out at a faster rate to obtain

small fiber diameter.

3.3. Hydrophilicity of the fiber

As mentioned above, the hydrophilicity of the

fibers can be measured by water contact angle

measurement. PLA, PLA/P1, PLA/P2, and PLA/P3 fiber

mats (entry 1-4), were measured for contact angle using

the goniometer (see Table 3). According to the results,

the lowest to the highest contact angle was entry 2, 3, 4,

and 1, respectively. This result reveals that the presence

of PLAdi+ in the polymer blend helps increasing the

hydrophilicity of the fiber mats with extent depending on

the charge density of the PLAdi+. As mentioned before,

the P1 sample has low MW, it would therefore provide

the highest positive charge density when all dopant (P1

to P3) were added in equal amounts.

Table 3. Water-air contact angle of PLA and

PLA/PLAdi+

4. Conclusion

PLA with two carboxyl chain ends

(PLAdiCOOH) and low MW PLA with two positive

charge ends PLAdi+ were successfully synthesized. For

synthesis of PLAdiCOOH, mole percentage of succinic

acid and polymerization time played important role in

controlling the molecular weight and chain length.

PLAdi+ was further used as dopant in electrospinning

process of commercially available PLA to enhance the

polymer hydrophilicity and reduce fiber diameter.

Increasing amounts of PLAdi+ dopants and reducing

their molecular size played significant role in reducing

the fiber diameter, which provide higher total surface

area of the fiber, and also raising the hydrophilicity of

the fiber because of increasing positive charge density.

5. Acknowledgment

This research was kindly supported by Graduate

School Thesis Grant and the scholarship from Graduate

School, Chulalongkorn University to commemorate the

72nd

anniversary of His Majesty King Bhumibala

Aduladeja. The authors gratefully acknowledge

Natthaphon Suksamran for providing commercial-grade

PLLA from NatureWorks; Wilaiporn Kraisuwan for

electrospinning instruction; Nutjarin Pansombat for PLA

synthesis training. Finally, the authors would like to

express my gratitude for kindly support by all friends,

lab mates and my family.

References

[1] N. Pansombat, Synthesis of PLLA with two

positively charged chain ends, Program in

Petrochemistry and Polymer Science,

Chulalongkorn University, Bangkok, 2013.

[2] E. Luong-Van, L. Grøndahl, K.N. Chua, K.W.

Leong, V. Nurcombe, S.M. Cool, Controlled release

of heparin from poly(ε-caprolactone) electrospun

fibers, Biomaterials, 27 (2006) 2042-2050.

[3] Q.P. Pham, U. Sharma, A.G. Mikos, Electrospinning

of polymeric nanofibers for tissue engineering

applications: a review, Tissue Eng., 12 (2006) 1197-

1211.


24 BIOENP-04

Toughness Improvement of Poly(Lactic Acid) Using Modified Natural Rubber

Wasan Tessanan1 and Pranee Phinyocheep

1,2*

1Polymer Science and Technology Program, Department of Chemistry, Faculty of Science,

Mahidol University, Salaya Campus, Salaya, Nakhon Pathom 73170, Thailand 2Rubber Technology Research Centre, Faculty of Science, Mahidol University, Salaya Campus, Salaya,

Nakorn prathom 73170, Thailand

Abstract

The chemical structure of natural rubber (NR) was modified by hydrogenation and epoxidation,

respectively. The epoxidation was carried out using in-situ performic acid and the hydrogenation was done by

using diimide system. Epoxidized hydrogenated natural rubber containing 30 mol% epoxide content and 23 %

hydrogenation (EHNR-30) was prepared and explored for toughening poly(lactic acid) (PLA). The NR/PLA and

EHNR-30/PLA blends were prepared at various rubber concentrations from 1-10% by weight with a Haake

internal melt mixer. The impact property and morphology of the blends were investigated. The impact strength

of the NR/PLA and EHNR-30/PLA blends containing 10 wt% rubber contents were enhanced about 2 times and

8 times, respectively, compared with the neat PLA. The SEM micrograph of EHNR-30/PLA blend showed

smaller rubber particle size of the modified NR (EHNR-30), compared with the use of unmodified NR

indicating that EHNR-30 is partially compatible with PLA. The results revealed that the use of EHNR-30 is

better toughening of PLA than the virgin NR due to a good interfacial adhesion between EHNR-30 and PLA.

Keywords: toughness; poly(lactic acid); epoxidized hydrogenated natural rubber

1. Introduction

PLA is a thermoplastic aliphatic polyester

derived from renewable resources such as corn, sugar,

rice and wheat. The advantages of PLA are high

modulus, high tensile strength, excellent transparency,

biodegradability and biocompatibility. However, high

brittleness of PLA is a major disadvantage which limits

its wide application. Many strategies have been explored

for toughening of PLA including plasticization [1, 2],

copolymerization [3, 4] and blending with elastomers [5-

7]. The blending of PLA with elastomers was widely

investigated both in academic and industry because it is a

simple and the most effective and economical way for

improvement brittleness of PLA. Various elastomers

were used to blend with PLA such as natural rubber

(NR) [5], poly(ethylene-co-glycidyl methacrylate)

(PEGMA) [6] and polyamide elastomer (PAE) [7]. The

bio-based elastomers were usually attracted more than

petroleum-based elastomers.

NR is a green elastomer derived from Hevea

brasiliensis rubber trees. It has many advantages such as

high elasticity, high tensile strength, biocompatibility

and biodegradability [8]. NR has been widely interested

as a suitable candidate for toughening PLA [9].

However, the problem in the use of NR for blending with

PLA is phase separation between non-polar NR and

polar PLA which limits the efficiency in improvement

brittleness of PLA [8].

Modification of NR into a higher polarity by

epoxidation and improvement of thermal resistance of

NR by reducing its unsaturation using hydrogenation

reaction would result in a modified NR so-called.

epoxidized hydrogenated natural rubber (EHNR). This

would bring the EHNR having a potential to blend with a

polar polymer such as PLA.

In this work, the modification of NR into EHNR

was carried out and the modified structure was analyzed

using FT-IR and 1H-NMR. The melt blending of PLA

with NR and EHNR were prepared at various rubber

concentrations. Impact strength and morphology of the

blend were investigated.

2. Materials and Methods

2.1 Materials

PLA (2003D) was produced from NatureWorks,

USA, with melt flow index (MFI) of 6.0 g/10 min (190๐

C/2.16 kg) and a density 1.24 g/cm-3

. High ammonia

natural rubber (HANR) latex containing 60% dried

rubber was supplied by Thai rubber latex cooperation,

Thailand was used as the starting material for chemical


25 modification and blending with PLA. Poly(ethylene

oxide fatty alcohol) hexadecylether (Terric 16A-16) as

non-ionic surfactant was obtained from East Asiatic

Company, Thailand. Hydrogen peroxide (35%w/v) was

purchased from QRec, New Zealand. Hydrazine hydrate

was provided by Merck, Germany. Formic acid was

purchased from Carlo Erba Reagent, USA.

2.2 Preparation of modified NR

Modification of NR into EHNR was prepared in

the latex stage by 2 step of reaction. Hydrogenation

reaction as a first step was started with adding hydrazine

at 40 ๐C followed by adding hydrogen peroxide at 60 ๐C.

The reaction mixture was stirred throughout 3 h to

prepare hydrogenated natural rubber (HNR) latex.

Afterward, HNR latex was cooled down to room

temperature, then the latex was neutralized before

addition of hydrogen peroxide and formic acid at 60 ๐C.

The reaction mixture was stirred during 20 h before

coagulation in methanol. The EHNR obtained was then

washed with water before drying in oven at 40 ๐C. The

30 mol% epoxide content of EHNR was prepared which

was defined as EHNR-30.

Table 1 Blend Compositions

Sample code PLA

(wt%)

NR

(wt%)

EHNR-30

(wt%)

PLA

NR/PLA(1/99)

NR/PLA(3/97)

NR/PLA(5/95)

NR/PLA(7/93)

NR/PLA(10/90)

EHNR-30/PLA (1/99)

EHNR-30/PLA (3/97)

EHNR-30/PLA (5/95)

EHNR-30/PLA (7/93)

EHNR-30/PLA(10/90)

100

99

97

95

93

90

99

97

95

93

90

-

1

3

5

7

10

-

-

-

-

-

-

-

-

-

-

-

1

3

5

7

10

2.3 Preparation of PLA blends

The PLA were melt blended with NR and

EHNR-30 with various rubber concentrations at 1, 3, 5, 7

and 10wt% in Haake internal mixer. The blending was

carried out with a rotor speed of 50 rpm at 170 ๐C for 15

min. The composition of PLA blends were shown in

Table 1. The blend specimens were prepared with

compression molding machine at 170 ๐C for 2 minutes

under pressure of 15 kN for further testing.

2.4 Characterization

2.4.1Chemical structure characterization

Fourier transform infrared (FT-IR) spectrum was

carried out in the range of wavenumber 4000-400 cm-1

with number of scan 16 times and a resolution of 4 cm-1

.

1H-NMR spectrum was recorded using

tetramethylsilane as internal standard. The percentage of

epoxide and saturated content were calculated from

intensity ratio of peak area as following equations.

Saturated content (%) =

100

5.13A2.73A0.84A

0.84A

(1)

Epoxide content (%) =

100

5.1A3

0.84A

2.7A

2.7A

(2)

where A0.84, A2.7 and A5.1 represent the integrated peak

areas of signal of methyl protons of saturated unit,

methine protons of epoxide ring and unsaturated methine

protons of isoprene unit, respectively.

2.4.2 Impact property

All samples were performed according to ASTM

D256 using a mechanical impact tester. Six specimens of

each sample were investigated at room temperature with

V-notched izod mode.

2.4.3 Morphological study

The scanning electron microscope (SEM) at an

accelerating voltage of 15 kV was used to investigate the


26 morphology of the impact fractured surfaces and

cryogenic fracture of the blends.

3. Results and Discussion

3.1 Chemical structure of modified NR

Figure 1 FT-IR spectra of NR (a) and EHNR30 (b)

The chemical structure of NR and modified NR

(EHNR-30) were investigated with FT-IR as shown in

Figure 1. The FT-IR spectrum of NR exhibits main

absorption peak at 1664 and 835 cm-1

corresponding to

C=C stretching and C=C bending vibration, respectively.

As for FT-IR spectrum of modified NR into EHNR, it

displays new characteristic absorption band at 1253 and

873 cm-1

corresponding to C-O-C ring stretching and C-

O-C ring bending vibration, respectively.

The obtained EHNR was further confirmed with

1H-NMR as shown in Figure 2.

1H-NMR spectra of NR

and EHNR display main signal characteristic of

unsaturated methine protons of isoprene unit at 5.1 ppm.

Moreover, the synthesized EHNR-30 showed new

characteristic signal of methyl protons of saturated unit

and methine protons of epoxide ring at 0.84 and 2.7 ppm,

respectively. The percentage of epoxide and saturated

content were calculated from intensity ratio of peak area.

The calculated results were found 30 mol% epoxide

content and 23 mol% saturated content. These results

confirmed that EHNR-30 was successfully synthesized.

Figure 2 1H-NMR spectra of (a) NR and (b) EHNR-30

3.2 Impact property

The notched izod impact strength of the neat

PLA, NR/PLA and EHNR-30/PLA blends were

determined and the results are shown in Figure 3.

Figure 3 Impact strength of NR/PLA and EHNR-

30/PLA blends

It can be seen that the impact strength of EHNR-

30/PLA is higher than the NR/PLA blend in each ratio

investigated. The impact strength of NR/PLA and

EHNR-30/PLA blends with 10 wt% rubber content show

increasing about 2 times and 8 times, respectively, when

compared with neat PLA. This result indicated that

EHNR-30 is better toughening agent for PLA. Epoxide

rings on structure of EHNR-30 would improve the

polarity of rubber effecting on enhancement

compatibility with PLA.

(a)

(b)

1664 cm-1

1664 cm-1

873 cm-1

1258 cm

-1

835 cm-1


27 3.3 Morphology

The SEM micrograph of impact fractured surface of

the neat PLA and PLA blends containing 10 wt% rubber

content were investigated as shown in Figure 4. The

impact fractured surface of neat PLA shows smooth

surface which is typical brittle fracture surface. NR/PLA

blends exhibit many spherical voids which were formed

due to NR particles were easily debonded from PLA

matrix which indicated poor interfacial adhesion between

NR and PLA matrix. The impact fractured surface of

EHNR-30/PLA blend can be obviously seen more and

longer fibrils on the surface which is agreement with

impact strength results. This can be indicated that the

brittle fracture of PLA transformed to ductile fracture

with using EHNR-30.

The SEM micrograph of cryogenic-fractured

surface of the blend in Figure 5 shows homogeneous

phase of neat PLA. As for NR/PLA blend containing

with 10 wt% of NR, it exhibits phase separation between

spherical NR particles and PLA matrix. This indicates

that poor interfacial adhesion between the non-polar NR

and the polar PLA. In the other hands, the EHNR-

30/PLA blend containing with 10 wt% EHNR-30 shows

small rubber particle which were dispersed in PLA

matrix. The smaller rubber particle size of EHNR-30 in

the blend indicated higher compatibility than the virgin

NR. This result ascribes that the increase of epoxide ring

on structure of NR can enhance polarity of NR resulting

on more compatibility of NR with PLA.

Figure 4 SEM micrograph of impact fractured surface of (a) neat PLA, (b) NR/PLA(10/90) and (c) EHNR-30

/PLA(10/90) blends

Figure 5 SEM micrograph of cyrogenic fractured surface of (a) neat PLA, (b) NR/PLA(10/90) and (c) EHNR-30

/PLA(10/90) blends

(a) (b) (c)

(a) (b) (c)


28 4. Conclusion

The modification of NR into epoxidized

hydrogenated natural rubber containing 30 mol%

epoxide content (EHNR-30) was successfully carried

out. NR and EHNR-30 were melt blended with PLA by

various rubber concentrations from 1-10% by weight.

EHNR-30/PLA blends show higher impact strength than

NR/PLA in each blend ratio investigated. SEM

micrograph of the blends showed that the EHNR-30 has

a smaller rubber particles size which were dispersed in

PLA matrix than the virgin NR. This indicated that the

compatibility of modified NR with PLA was

significantly improved.

5. Acknowledgement

The authors would like to thank the financial

support from IRPC Public Company Limited. The

Science Achievement Scholarship of Thailand (SAST) to

W.Tessanan is also very much appreciated.

6. References

[1] Labrecque, L.V., Kumar, R.A., Dave, V., Gross,

R.A. and McCarthy, S.P., “Citrate esters as

plasticizers for poly(lactic acid)”, J. Appl. Polym.

Sci., 66(8): 1507-1513 (1997).

[2] Pillin,I., Montrelay, N. and Grohens, Y., “Thermo-

mechanical characterization of plasticized PLA: is

the miscibility the only significant factor?”,

Polymer, 47(13): 4676-4682 (2006).

[3] Hiljanen-Vainio, M., Karjalainen, T. and Seppälä, J.,

“Biodegradable lactone copolymers. I.

Characterization and mechanical behavior of ε-

caprolactone and lactide copolymers”, J. Appl.

Polym. Sci., 59(8): 1281-1288 (1996).

[4] Lan, P., Zhang, Y.P., Gao, Q.W., Shao, H.L. and

Hu, X.C., “Studies on the synthesis and thermal

properties of copoly(L- lactic acid/glycolic acid) by

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92(4): 2163-2168 (2004).

[5] Juntuek, P., Ruksakulpiwat, C., Chumsamrong, P.

and Ruksakulpiwat, Y., “Effect of glycidyl

methacrylate-grafted natural rubber on physical

properties of polylactic acid and natural rubber

blends”, J. Appl. Polym. Sci., 125: 745-754 (2012).

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copolymer”, Polymer, 50: 747-751 (2009).

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shape-memory effect of polylactide resulted from

toughening by polyamide elastomer”, Polymer, 50:

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Manchado, M.A., “Structure and properties of

polylactide/natural rubber blends”, Mater. Chem.

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International H.F. Mark-

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Thermoplastic Elastomers”, Vienna, 15-16

November, 2007.

[10] Zhang, C., Wang, W., Huang, Y., Pan, Y., Jiang,

L., Dan, Y., Luo, Y. and Peng Z., “thermal

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toughened by epoxidized natural rubber”, Mater.

Desi., 45: 198-205 (2013).


29 BIOENP-05

Synthesis and Characterization of Medical Grade Poly(L-lactide-co-Ɛ-caprolactone)

Jutamas Kongsuk, Pimwalan Techaikool, Kiattikhun Manokruang, Puttinan Meepowpan, Kanarat Nalampang,

Runglawan Somsunan, Patnarin Worajittiphon, Wathuka Booncharoen, Robert Molloy and Winita Punyodom*

Polymer Research Laboratory, Department of Chemistry, Faculty of Science, Chiang Mai University,

Chiang Mai, 50200, Thailand

Abstract

Poly(L-lactide-co-Ɛ-caprolactone) (PLLCL) is one the most attractive polymeric candidates for

fabricating devices for use in biomedical applications. PLLCL is both biocompatible and biodegradable and, by

varying its composition and microstructure, exhibits tunable mechanical properties. In this research, the main

aim has been to synthesize and characterize medical grade PLLCL with a copolymer composition of L-lactide

(LL): Ɛ-caprolactone (CL) = 70:30 mol %. This has been achieved via the bulk ring-opening polymerization

(ROP) of LL and CL at 120C for 96 hours using tin(II) n-butoxide, Sn(n-OBu)2, in liquid form as the initiator.

Following its purification, the various properties of the PLLCL which are required by the ASTM F1925-09 Test

Method (Standard Specification for Semi-Crystalline Polylactide Polymer and Copolymer Resins for Surgical

implants) for the copolymer to be used in biomedical devices have been determined. These properties include

molecular weight, composition, temperature transitions and the amounts of residual monomers and residual tin

from the initiator. The results have shown that by using only a very small amount of the Sn(n-OBu)2 initiator so

that the amount of residual tin in the copolymer, as confirmed by ICP-OES analysis, was within the limit

described in ASTM F1925-09, a high molecular weight copolymer with good physical properties suitable for

use in biomedical applications could be obtained. These results will be discussed in detail.

Keywords: Poly(L-lactide-co-Ɛ-caprolactone); medical grade; tin(II) n-butoxide; residual tin; biomedical

applications.

1. Introduction

The healthcare industry continues to evolve and

requires technological advances to meet today’s

healthcare providers in the field of biomedical devices

and to reduce the total cost of care for the end-patients.

During the past few decades significant advances

have been made in the development of biodegradable

materials for biomedical applications. A wide variety of

new synthetic polymers and biodegradable polymers

have been evaluated. The most common biodegradable

polymers such as polylactide (PL), polyglycolide (PG),

poly(Ɛ-caprolactone) (PCL) and their copolymers have

been widely used in drug delivery and tissue engineering

applications.

Interest in copolymers of L-lactide (LL) and Ɛ-

caprolactone (CL) has increased as their potential in a

wide range of biomedical have been adopted. The

biocompatible and biodegradation copolymers of LL and

CL are focused in this intensive study with regard to

organ and tissue regenerations. CL appears to be a

suitable comonomer for the preparation of a diversified

family of copolymers with mechanical properties ranging

from gummy and elastomeric to rigid solids. The

properties of poly(L-lactide-co-Ɛ-caprolactone) (PLLCL)

differ widely depending on the ratio of LL and CL. The

main advantages of PLLCL copolymer include relatively

fast degradation compared with both of its

homopolymers, and its well known highly elastic

properties [1-5]. PLLCL can be synthesized by ring-

opening polymerization (ROP). Tin(II) 2-ethylhexanoate

(Sn(Oct)2) is an initiator generally used for the ROP of

lactide and other cyclic ester monomers since it is a

highly efficient initiator allowing the complete

conversion of monomers to polymers [6-7]. However,

the usual amount of Sn(Oct)2 for ROP to get the desired

molecular weight is commonly over 150 ppm [8], which

is not acceptable by the ASTM Standard F1925-09 [9].

From literature reviews, Tin(II) n-butoxide (Sn(n-OBu)2)

also shows high efficiency for ROP to obtain high

molecular weight of polyester. Interestingly, the effective

amount of Sn(n-OBu)2 can be reduced to obtain the

residue tin lower than 150 ppm but it still gives high

efficiency in ROP [10-11]. To make the amount of

initiator acceptable by ASTM standard F1925-09, there

are two ways. The first is to reduce the amount of

initiator from the synthesis and control other factor to

obtain a high molecular weight copolymer. The second

method is to use an organic solvent for dissolution


30 followed by precipitation [11]. This research is aimed at

synthesizing a PLLCL copolymer LL: CL (70:30) with

an acceptably low amount of residual tin together with its

detailed characterization.

2. Experimental methods

Materials

Chloroform (RCI Labscan, 99.8%) and deuterated

chloroform (Aldrich, 99.8%) were used as received.

Ethyl acetate (Scharlau, 99.7%) was distilled before use.

Toluene was dried over sodium and stored under dry

nitrogen. LL monomer was synthesized from L-lactic

acid (Natureworks, 88%), purified by repeated

recrystallization from distilled ethyl acetate, dried under

vacuum at 55°C for 12 hours and stored under vacuum.

CL (Acros Oganics, 99.0%) monomer was purified by

fractional distillation under reduced pressure (boiling

point = 90°C / 7 mm Hg).

Initiator synthesis

The Sn(n-OBu)2 initiator was prepared as

described in the International Patent Application No.

PCT/TH2013 /000061(WO2014/0777785A1) [9]. A

stock solution of the initiator was prepared in toluene and

stored under dry nitrogen.

Synthesis of copolymers

For each copolymerization, 500 g of the

comonomers were accurately weighed into a round-

bottomed flask with a magnetic stirring bar. The stock

solution of initiator (0.010 mol %.) was added and the

copolymerization carried out at 120°C in a silicone oil

bath for 96 h. At the end of the copolymerization, the

flask was removed from the oil bath and quickly cooled

in an ice-bath to terminate any further polymerization.

The synthesis reaction is shown in Scheme 1.

Scheme 1 Preparation of PLLCL

Characterization of PLLCL

The infrared absorption spectra were collected at

25°C from 400-4000 cm–1

[12]. The spectra were

recorded on a Fourier-transform Infrared Spectrometer

(FT-IR), Bruker TENSOR 27.

1H-NMR (400 MHz) and

13C-NMR (100 MHz)

spectra were obtained from a Bruker DPX-300 NMR

Spectrometer. The samples were dissolved in deuterated

chloroform (CDCl3) at room temperature before analysis

[13].

Thermal analysis was performed by Differential

Scanning Calorimetry (Perkin Elmer DSC-7) in a 5-step

program: (heat, cool down, hold, cool down and heat)

from 0 to 200°C with a heating rate of 10°C/min under a

nitrogen atmosphere [14]. Thermogravimetric analysis

(TGA) was conducted on a Perkin Elmer TGA-7 under

N2 flow at a heating rate of 20°C/min from 50 to 600°C.

The intrinsic viscosity, [η], of the PLLCL was

determined from a single measurement of the relative

viscosity by using the Billmeyer relationship (1). The

experiment was conducted in chloroform as solvent at

30.0±0.1°C using a Schott-Geräte AVS300 Automatic

Viscosity Measuring System. The value of [ ] was

calculated from equation (1):

[ ] =

dl/g (1)

where [ ] = intrinsic viscosity, r = relative viscosity =

t/t0 (t0 and t are the flow-times of the chloroform solvent

and the copolymer solution respectively. C is the single

concentration of the copolymer solution (0.5 g/dl) [15].

The Sn and Pb contents in the PLLCL copolymer

samples were determined using an Inductively Coupled

Plasma Optical Emission Spectrometer (ICP-OES). The

sample was digested in a microwaved vessel, diluted

with nitric acid, and then measured following the

USEPA (2007) Method 3025.


A PLLCL copolymer with a composition of 70:30

mol % was successfully synthesized by ROP in the

presence of liquid Sn(OnBu)2 as an initiator. The


31 copolymer product was obtained as a white, elastomeric

solid obtained in high yield (92.7%).

The FT-IR spectrum confirmed its chemical

structure exhibiting C-H stretching around 2900 cm-1

,

C=O stretching around 1750 cm-1

, C-H bending in CH

and CH3 around 1450 and 1350 cm-1

respectively, C-O

stretching of acyl-oxygen around 1280 cm-1

, and C-O

stretching of alkyl-oxygen around 1090 cm-1

.

5001000150020002500300035004000

0.0

0.5

1.0

-CH3

bend

-C-O-C

stretch-C=O

stretch

-CH, CH3

stretch

-OH

stretch

Wavenumber (cm-1)

Tra

nsm

itta

nce (

%)

Figure 1 FT-IR spectrum of PLLCL copolymer

The final 92.7% conversion of the copolymer was

determined by weighing after rigorous purification and

drying to constant weight. The copolymer composition

(LL: CL mol %) was determined from the 1H-NMR

spectrum shown in Figure 2. The copolymer composition

could be calculated by taking the ratio of the peak areas

corresponding to the LL methine protons at δ = 5.0-5.3

ppm and the CL Ɛ-methylene protons at δ = 3.9-4.2 ppm

using equations (2) and (3). The calculated compositions

are given in Table 1 and show that the final copolymer

composition of LL: CL = 70:30 mol % (± 1%) was

identical with the initial comonomer feed.

Copolymer composition of PLLCL:

% mol of LL =

(2)

% mol of CL =

(3)

1.01.52.02.53.03.54.04.55.05.5

a

b

c d e f g

egc

b

a, d, f

O C C

CH3

H O

O CH2 CH2 CH2 CH2 CH2 C

Onm

Chemical shift (ppm)

Figure 2 1H-NMR (400 MHz) spectrum of PLLCL

Monomer sequencing in the PLLCL copolymer

was characterized by 13

C-NMR, specifically from the

expanded carbonyl carbon (C=O) region, as shown in

Figure 3. The various peaks were assigned to the C=O

carbons of the middle units of various triad sequences in

the copolymer chain. The appearance of the various

heterotriad peaks in between the homotriad CCC and

LLL peaks is a measure of the degree of randomness of

the monomer sequencing. Figure 3 suggests that the

monomer sequencing is only partly random and partly

blocky, due mainly to the differing monomer reactivity

ratios (LL > CL).

In the triad notations in Figure 3, L represents half

a lactide unit, -O-CH(CH3)-CO-, while C represents a

caprolactone unit, -O-(CH2)5-CO-. The carbon atom to

which each triad peak corresponds is that of the C=O

carbon of the middle unit.

Thermal analysis was performed by a

combination of DSC and TGA as shown in Figures 4 and

5. From DSC, it was found that the first run showed a Tm

melting peak at 163.0°C, while the second run (Figure 4)

showed Tg at 28.4°C, Tc at 110.0°C and Tm at 160.4°C.

The appearance of the Tm peaks confirm that the PLLCL

copolymer (70:30 mol %) was still semi-crystalline in

morphology despite its microstructural irregularity.


32

169170171172173174

(ppm)

LLL

CLL

LLCLCC/LCL

LCC

CCC

Figure 3 Expanded C=O carbon region of the 100 MHz

13C-NMR spectrum of purified PLLCL.

0 20 40 60 80 100 120 140 160 180 200

Tm= 160.4

Tc= 110.0

Tg=28.4

2nd

Run

N

orm

ali

zed

Hea

t F

low

En

do

Up

(W/g

)

Temperature(oC)

Figure 4 DSC thermogram (second run) of the PLLCL

copolymer

The copolymer’s TGA curve in Figure 5 shows

a single-step weight loss. After a small (<10%) initial

weight loss due to residual moisture and/or volatiles, the

main degradation onset temperature (Td) was observed at

275°C with degradation complete at 480°C. The

intermediate degradation temperature at 50% weight loss

was approximately 380°C. When the DSC and TGA

curves are considered together, it can be seen that the

copolymer has a wide melt processing range between Tm

(170°C) and Td (275°C) within which it can be safely

melt processed without accompanying thermal

degradation. This is very important for a polymer which

is intended for use in a biomedical application. The DSC

and TGA results are summarized in Table 1.

100 200 300 400 500 6000

20

40

60

80

100

Wie

gth

% (

%)

Temperature (oC)

Figure 5 TGA curve of PLLCL copolymer.

Table 1 Summary of PLLCL characterization results.

Abbreviation PLLCL

% Yield 92.7%

Physical appearance White solid

Intrinsic Viscosity

[] (dl/g)

2.10

Composition LL : CL 70:30 (mol %)

Monomer sequencing Partly random

Partly blocky

Thermal Properties

DSC ( 1st run)

( 2nd

run)

Tm = 163.0 oC

Tg = 28.4 oC

Tc = 110. oC

Tm = 160.4 oC

TGA Td (onset) = 275oC

Sn content (ppm) 136

Pb content (ppm) Not detected

The molecular weight of the PLLCL copolymer

was determined by a single measurement of solution

viscosity and the result shown in Table 1. The intrinsic

viscosity, [ ], of 2.10 dl/g indicates that the copolymer

has a high molecular weight, probably with Mn >

100,000.

Finally, the Sn content of the copolymer obtained

from ICP-OES is shown in Table 1. The amount of

residual tin from the Sn(n-OBu)2 initiator was lower than

the limit (150 ppm) described in ASTM F1925-09


33 specifications for a medical grade material. There was no

lead (Pb) content.

4. Conclusions

In this paper, the successful synthesis of a PLLCL

70: 30 mol % copolymer with a low amount of residual

tin has been described. The results have shown that the

ring-opening polymerization (ROP) of LL and CL

initiated by tin(II) n-butoxide (Sn(n-OBu)2) can be scaled

up to the 500 g scale of the comonomers with high %

conversion and molecular weight. The residual tin

content of 136 ppm, which is lower than the limit of 150

ppm set by the ASTM F1925-09 standard specifications,

shows that this synthesis and purification procedure is

suitable for the preparation of medical grade PLLCL for

use in biomedical applications.

5. Acknowledgments

This research was supported by the National

Research Council of Thailand (NRCT), National

Innovation Agency (NIA), Public Company Limited

(PTT), Chiang Mai University (CMU) and the Graduate

School, Chiang Mai University.

6. References

[1] Sodergard A, Stolt M. Properties of lactic acid based

polymers and their correlation with composition.

Prog Polym Sci 2002; 27: 1123-63.

[2] Bendix D. Chemical synthesis of polylactide and its

copolymers for medical applications. Polym Degrad

Stab 1998; 59: 129-35.

[3] Fernandez J, Etxeberria A, Sarasua JR. Effects of

repeat unit sequence distribution and residual

catalyst on thermal degradation of poly(L-lactide/Ɛ-

caprolactone) statistical copolymers. Polym Degrad

Stab 2013; 98: 1293-9.

[4] Woodruff MA, Hutmacher DW. The return of a

forgotten polymer-polycaprolactone in the 21st

century. Prog Polym Sci 2010; 35: 1217-56.

[5] Arbaoui A, Redshaw C. Metal catalysts for Ɛ-

caprolactone polymerisation. Polym Chem 2010; 1:

801-26.

[6] Sobczak M, Kolodziejski W. Polymerization of

cyclic esters initiated by carnitine and tin(II) octoate.

Molecules 2009; 14: 621-32.

[7] Nalampang K, Molloy R, Punyodom W. Synthesis

and characterization of poly(L-lactide-co-Ɛ-

caprolactone) copolymers: influence of sequential

monomer addition on chain microstructure. Polym

Adv Technol 2007; 18: 240-8.

[8] Thapsukhon B, Daranarong D, Meepowpan P, Suree

N, Molloy R, Inthanon K, Wongkham W,

Bunyodom W. Effect of topology of poly(L-

lactideco-ε-caprolactone)scaffolds onthe response

of cultured human

umbilical cord Wharton’s jellyderivedmesenchymal stem

cells andneuroblastoma cell lines. J Biomater Sci,

Polym Ed 2014; 25:1028-44.

[9] American Society for testing and Materials

(ASTM). 2009. Standard Specification for Semi-

Crystalline Poly(lactide) Polymer and Copolymer

Resins for Surgical Implants: F1925-09.

[10] Meepowpan, P.; Punydom, W.; Molloy, R.

International Patent Application No.

PCT/TH2013/000061 (WO 2014/0777785 A1)

2014.

[11] Dumklang M, Pattawong N, Punyodom W,

Meepowpan P, Molloy R, Hoffman M. Novel tin(II)

butoxides for use as initiators in the ring-opening

polymerisation of Ɛ-caprolactone. Chiang Mai J Sci

2009; 36: 136-48.

[12] Stjerndahl A, Wistrand AW, Albertsson AC.

Industrial utilization of tin-initiated resorbable

polymers: synthesis on a large scale with a low

amount of initiator residue. Biomacromolecles 2007;

8: 937-40.

[13] American Society for testing and Materials (ASTM).

2013. Standard Practice for General Techniques for

Obtaining Infrared Spectra for Qualitative

Analysis1: E1252-98.


2011. Standard Practice for Data Presentation

Relating to High-Resolution Nuclear Magnetic

Resonance (NMR) Spectroscopy1: E386-90.


34 [15] American Society for testing and Materials (ASTM).

2012. Standard Test Method for Transition

Temperatures and Enthalpies of Fusion and

Crystallization of Polymers by Differential Scanning

Calorimetry1: D3418 − 12

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(PET) by Glass Capillary Viscometer1: D4603-03.


35

BIOENP-06

Synthesis and Characterization of Polylactide-Poly(ethylene glycol)-Polylactide ABA-

Triblock Copolymers for Use in Medical Applications

Tichakorn Thornsri1,2

, Puttinan Meepowpan1,2

, Winita Punyodom1,2

and Kanarat Nalampang1,2*

1Department of Chemistry, Faculty of Science, Chiang Mai University, Chiang Mai 50200

2Materials Science Research Center, Faculty of Science, Chiang Mai University, Chiang Mai 50200

Abstract

Polylactide-poly(ethylene glycol)-polylactide (PLLA-PEG-PLLA) triblock copolymers, one of

biodegradable polymers, have been widely used in medical applications due to their biodegradability,

biocompatibility and tailor-made properties. In this work, PLLA-PEG-PLLA triblock copolymers were

synthesized by bulk ring-opening polymerization (ROP) of L-lactide and PEG at 120˚C for 24 hours using

stannous octoate (Sn(Oct)2) as a catalyst. The chemical structure and composition of obtained triblock

copolymers were characterized by various techniques such as Fourier transform infrared spectroscopy (FTIR)

and proton nuclear magnetic resonance spectroscopy (1H-NMR). Effect of molecular weight of PEG ( =

2000, 4000, 6000 and 8000) and effect of ratio between ethylene oxide (EO) and lactate (LA) (2:1 and 3:1) on

microstructure of copolymers were investigated. Moreover, thermal properties of copolymers were examined

by differential scanning calorimetry (DSC) technique. The results have shown that the copolymers were semi-

crystalline materials. Tm and ∆Hm were observed on dependence of block lengths of either PLLA or PEG in

copolymers.

Keywords: Biodegradable polymers; polylactide-poly(ethylene glycol)-polylactide; triblock copolymer

1. Introduction

Biodegradable polymers have been widely used in

medical applications. The advantages of biodegradable

polymers are not requiring surgical removal after they

serve their purposes. Poly(lactide) (PLA) and their

copolymers PLA are the most commonly used

biodegradable polymers. In general, the biodegradable

polyesters are strongly hydrophobic and this has caused

some limitations in certain applications. Hydrophilic,

poly(ethylene glycol) (PEG) is one of the most widely

employed, has been incorporated into the biodegradable

polyesters to modify the hydrophilicity. The advantages

of PEG are water-soluble polymer, non-toxic, non ionic,

biocompatible characteristic, and can be easily

eliminated from the body. Therefore, block copolymers

consisting of a hydrophobic polyester segment and a

hydrophilic PEG segment have been attracted large

attention due to their biodegradability, biocompatibility

and tailor-made properties [1-3].

Various types of block copolymer consisting of a

hydrophobic polyester segment and a hydrophilic PEG

segment have been developed such as AB diblock, ABA,

BAB- triblock, multi-block,

branched block, star-shaped

block, and graft block copolymers, A is a hydrophobic

block as biodegradable polyesters and B is a hydrophilic

PEG [4-11].

In this work, we focus on PLLA-PEG-PLLA

triblock copolymer. The obtained block copolymers were

synthesized by bulk ring-opening polymerization of L-

lactide and PEG using Sn(Oct)2 as a catalyst and

characterized by FTIR and 1H-NMR techniques. Effect

of molecular weight of PEG ( = 2000, 4000, 6000 and

8000) and various molar ratio of EO:LA (2:1 and 3:1)

were investigated. Moreover, thermal properties of

copolymer were examined by DSC technique. These

triblock copolymers will provide further benefits for use

as medical applications.

2. Experimental

2.1 Materials: Ethyl acetate (Scharlau) and

dichloromethane (J.T.Baker) was distilled before use.

PEG2000 (Aldrich), PEG4000 and PEG6000 (Ajax

Finechem), PEG8000 (Acros) and diethyl ether (ACI

Labscan) were used as received. Stannous octoate

(Aldrich) was distilled under reduced pressure to remove

octanoic acid before use. L-lactide was synthesized from

L-lactic acid (Aldrich) using Sn(Oct)2 as a catalyst [12].

The monomer was purified few times by recrystallization


36 in distilled ethyl acetate, dried under vacuum at 55

˚C for

12 hours and stored under vacuum ambient.

2.2 Polymerization: In this experiment, PLLA-PEG-

PLLA triblock copolymer were synthesized by using

PEG with different molecular weight ( = 2000, 4000,

6000 and 8000) and various molar ratio of EO:LA (2:1

and 3:1). An example of a polymerization (in case of

EO:LA = 3:1), 3.53 g of L-lactide and 6.47 g of different

molecular weight of PEG ( = 2000, 4000, 6000 and

8000) were introduced to a 30 ml round-bottomed flask

and 0.065 g of Sn(Oct)2 was added. Under vacuum, the

mixture was stirred at 120˚C in a preheated oil bath for

24 hours. Then the triblock copolymers were purified by

dissolved in dichloromethane and precipitated in diethyl

ether. Finally, triblock copolymers were filtered and

dried to constant weight in a vacuum oven.

2.3 Measurements: FT-IR spectra of homopolymers

and triblock copolymers were obtained using a Bruker

TENSOR 27 over the region 4,000 to 400 cm-1

. The

microstructures and copolymer compositions were also

confirmed by 400 MHz 1H-NMR Bruker DPX-400 using

deuterated chloroform as solvent, tetramethyl silane as

internal standard at room temperature, and sample

concentration of 3%(w/v). Thermal analysis was carried

out by DSC (Perkin-Elmer DSC7 Series). Prior to the

measurement, a stable baseline was established using

two empty reference pans. Block copolymers with a

typical mass of 5-10 mg was encapsulated in a sealed

aluminium pan. The sample was heated over a

temperature range of 0˚C to 250˚C under nitrogen

atmosphere with a heating rate of 10˚C/min.

3. Results and discussion

3.1 Synthesis and Characterization

Triblock copolymers were synthesized by ROP of L-

lactide and PEG using Sn(Oct)2 as a catalyst are shown

in Scheme 1. The obtained triblock copolymers were

interpreted to confirm the functional groups by FTIR

techniques as shown in Figure 1.

O

O

O

O

H3C

CH3

+ HOH2C

H2C O H

n

Sn(Oct)2

HOHC C O

H2C

CH3

O H2C O C

x n

OHC

CH3

O Hy

L-lactide PEG PLLA-PEG-PLLA

Scheme 1. ABA triblock copolymerization via ROP of

L-lactide and PEG using Sn(Oct)2 as a catalyst

Figure 1. FTIR spectra of PLLA (A), PEG2000 (B) and

the triblock copolymer of PEG2000 with lactide (C).

The FTIR spectra of the triblock copolymer exhibited

C=O stretching around 1750 cm-1

indicates the carbonyl

group as a part of ester in the lactate unit. The bands

around 3500 cm-1

showing O-H stretching and C-H

stretching at 2900 cm-1

are assigned to characteristic of

ethylene oxide segment [13]. Moreover, the chemical

structure and composition of obtained triblock

copolymers were characterized by 1H-NMR as shown in

Figure 2. The linkage of PEG with L-lactide in polymer chain can

be confirmed by 1H-NMR in which the chemical shift at

1.5–1.6 ppm is defined as CH3 (methyl group) of the

lactate units, at 3.5–3.8 ppm is CH2 (methylene group) of

the oxyethylene units, and at 5.0–5.3 ppm is CH

(methine group) of the lactate units. The PLLA/PEG

ratios of the copolymers can be calculated from the

integral ratios of the oxyethylene at 3.5–3.8 ppm and

lactate signals at 5.0–5.3 ppm. The number-average

molar mass, was calculated according to equation:

= x 44 + x 2 x 72 (1)

500 1000 1500 2000 2500 3000 3500 4000

Wavenumber (cm-1)

C=O stretching C-H stretching

O-H stretching

A

B

C


37 where = /44, = x (LA/EO)/2,

and 44 and 72 are the molar mass of EO and LA repeat

units [14].

HO CHC

O

CH3

OCHC

CH3

O

OCH2CH2 OCH2CH2 OCH2CH2 OCCH

O

CH3

OCCH

O

CH3

OH

m-1 n-1 l-1

Figure 2. 1H-NMR spectrum of the copolymer of PLLA-

PEG-PLLA with a molar ratio of EO:LA as 3:1.

Polymerization temperature was 120 ˚C and catalyst was

Sn(Oct)2.

3.2 Effect of molecular weight of PEG and ratio of

EO:LA

PLLA-PEG-PLLA triblock copolymer with

different molecular weight of PEG ( = 2000, 4000,

6000 and 8000) and various molar ratio of EO:LA (2:1

and 3:1) were obtained. The structure, composition,

molar mass ( ) and % yield of copolymers are shown

in Table 1. As the results, the copolymer compositions

(EO:LA) of triblock copolymers in polymer product

were higher than monomer feed for every ratios and

molecular weight of PEG. This finding might be implied

that conversion of lactide unit were not copolymerized

completely. Normally, unreacted lactide and probably

polylactide oligomer were eliminated by purification

step. Also, the difference of compositions between

polymer product and monomer feed increased with

increasing the ratio of EO:LA. In addition, a higher ratio

of EO:LA employed, less amount of lactide monomer

can react with the PEG hydroxyl end groups which

compare to a lower ratio of EO:LA under the same

synthesis condition of triblock copolymer. This is

unclear but probably because the viscosity of system

increases with increasing PEG content and diffusion

controlled might affect the copolymerization.

3.3 Thermal properties

Thermal properties of various triblock copolymers

were investigated by DSC (scanned immediately after

cooling at 10˚C/min.), as shown in Figure 3.

Figure 3 presents DSC thermograms of various

of PEG and theirs copolymers with lactide using

different compositions. It is clearly shown that PEG and

triblock copolymers are semi-crystalline. The

crystallization of triblock copolymers were reduced by

the presence of PLLA content. For example PEG6000

and theirs copolymer, PEG6000 homopolymer exhibited

Tm 66.7˚C and melting enthalpy (∆Hm) 182.9 J/g. In

contrast, Tm and ∆Hm of triblock copolymer decreased to

55.7˚C and 90.3 J/g for PLLA16PEG136PLLA16.

Moreover, the triblock copolymer containing longer

PLLA blocks (PLLA28/PEG136/ PLLA28), Tm and ∆Hm

also decreased to 54.0˚C and 58.0 J/g.

Figure 3. DSC thermograms (second scans, after cooling

at 10˚C/min) of various of PEG and their copolymers

with lactide using different compositions.

-2

3

8

13

18

23

28

33

38

43

0 20 40 60 80 100 120 140

en

do

Temperature (˚C)

PEG2000

PLLA5/PEG45/ PLLA5

PLLA8/PEG45/ PLLA8

PEG4000

PLLA9/PEG91/ PLLA9

PLLA16/PEG91/ PLLA16

PEG6000



PEG8000



a b d e

𝜹

e

b

a d

c

a b d e

c


38

The similar results can be obtained with other

molecular weight of PEG and theirs copolymers as

shown in Figure 3. Tm and ∆Hm were decreased with

decreasing the ratio of EO:LA at the same molecular

weight of PEG. These information can be confirmed that

the presence of PLLA blocks in the copolymer have a

profound influence on crystallization of PEG. The longer

PLLA block length could lower the degree of

crystallinity of PEG as reflected by decreasing of the

∆Hm values. Moreover, molecular weight of PEG ( =

2000, 4000, 6000 and 8000) really play an important role

on crystallinity of triblock copolymers since both Tm and

∆Hm increase when higher of PEG was used, as

shown in Table 2.

Conclusions

PLLA-PEG-PLLA triblock copolymers were

synthesized by ROP of L-lactide and PEG at 120˚C for

24 hours using Sn(Oct)2 as a catalyst. The functional

groups of the obtained triblock copolymers were

confirmed by FTIR and the chemical structure and

composition were characterized by 1H-NMR. The

copolymer compositions (EO:LA) of triblock

copolymers in polymer product were higher than

monomer feed for every ratios and molecular weight of

PEG and increased with an increasing of ratio of EO:LA.

In addition, DSC thermograms showed the triblock

copolymers were semi-crystalline. Tm and ∆Hm were

decreased with increasing PLLA block length in triblock

copolymers and this might because PLLA segment could

interfere crystallization of PEG.

Table 2. Thermal properties of PLLA-PEG-PLLA

triblock copolymer

Copolymer Tm (˚C) ∆Hm (J/g)

PLLA8/PEG45/ PLLA8 36.9 51.1

PLLA5/PEG45/ PLLA5 38.0 62.1

PLLA16/PEG91/ PLLA16 45.2 59.4

PLLA9/PEG91/ PLLA9 48.2 66.4

PLLA28/PEG136/ PLLA28 54.0 58.3

PLLA16/PEG136/ PLLA16 54.7 90.3

PLLA39/PEG182/ PLLA39 52.0 54.3

PLLA23/PEG182/ PLLA23 53.3 87.5

Acknowledgements

This research was supported from Department of

Chemistry, Faculty of Science and the Graduate School

Table 1. Effect of molecular weight of PEG and ratio of EO:LA

Structure PEG EO:LAa

c

d

e % Yield

PLLA8/PEG45/ PLLA8 PEG2000 2.8 (2.0) b 45 8 3132 80

PLLA5/PEG45/ PLLA5 PEG2000 4.6 (3.0) 45 5 2700 79

PLLA16/PEG91/ PLLA16 PEG4000 2.8 (2.0) 91 16 6308 86

PLLA9/PEG91/ PLLA9 PEG4000 4.9 (3.0) 91 9 5300 83

PLLA28/PEG136/ PLLA28 PEG6000 2.4 (2.0) 136 28 10016 89

PLLA16/PEG136/ PLLA16 PEG6000 4.4 (3.0) 136 16 8288 87

PLLA39/PEG182/ PLLA39 PEG8000 2.4 (2.0) 182 39 13624 93

PLLA23/PEG182/ PLLA23 PEG8000 3.9 (3.0) 182 23 11320 92

a Calculated from the integral ratios of the oxyethylene at 3.5–3.8 ppm and lactate signals at 5.0–5.3 ppm.

b Data in parentheses corresponding to EO:LA ratio in feed.

c = /44.

d = x (LA/EO)/2.

e = x 44 + x 2 x 72


39 of Chiang Mai University (CMU). National Research

Universities (NRU) Project under Thailand’s office of

the Higher Education Commission are also acknowledge.

References

[1] L. Chen, Z. Xie, J. Hu, X. Chen and X. Jing, Journal

of Nanoparticle Research, 9, 777–785 (2007).

[2] R. H. Kricheldorf and I. Kreiser, Macromolecular

Chemistry, 188, 1861- 1873 (1987).

[3] Ph. Dubois, C. Jacobs, R. JBrSme and Ph. TByssie,

Macromolecules, 24, 2266-2270 (1991).

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Journal of Pharmaceutics, 182, 187-197 (1999).

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Delivery Reviews, 54, 99-134 (2002).

[6] B. Jeong, Y. H. Bae and S. W. Kim, Journal of

Controlled Release, 63, 155-163 (2000).

[7] K.M. Huh and Y. H. Bae, Polymer, 40, 6147-6155

(1999).

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[9] J. S. Hrkach, M.T. Peracchia, A. Domb, N. Lotan

and L. Robert, Biomaterials, 18, 27-30 (1997).

[10] A. Breitenbach, Y. X. Li and T. Kissel, Journal of

Controlled Release, 64, 167-178 (2000).

[11] C. W. Lee, T. I. Manoshiro, Y. I. Hsu and Y.

Kimura, Macromolecular Chemistry and Physics,

213, 2174-2180 (2012).

[12] D.K. Yoo and D. Kim, Macromolecular Research,

14, 510-516 (2006).

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Macromolecules, 28, 2124-2132 (1995).

[14] S. Li and M. Vert, Macromolecules, 36, 8008-8014

(2003).


40

BIOENP-11

Studies on the Quaternization of Chitosan in Ionic Liquid

Maneerat Wangsiripaisarn1 and Varawut Tangpasuthadol

2*

1Program in Petrochemistry and Polymer Science, Faculty of Science, Chulalongkorn University,

Bangkok 10330, Thailand. 2Organic Synthesis Research Unit, Department of Chemistry, Faculty of Science, Chulalongkorn University,

Bangkok 10330, Thailand.

Abstract

Positively charged N,N,N-trimethyl chitosan (TMC) was synthesized via methylation at the amino

group of chitosan with iodomethane (CH3I) in the presence of base and 1-butyl-3-methylimidazolium chloride

(BMIMCl) as solvent. The degree of quaternization of TMC was determined by nuclear magnetic resonance

spectroscopy (NMR) and 2D 1H-

1H correlation (COSY). From these finding, it was found that potassium

carbonate (K2CO3) as organic base was more compatible with BMIMCl than sodium hydroxide. Moreover, the

degree of quaternization (%DQ) of TMC was affected by the type and amount of base. By using 9 equiv of

K2CO3 as base in BMIMCl as solvent, the highest DQ of 42% was obtained.

Keywords: Chitosan, Ionic Liquid, N,N,N-trimethyl chitosan (TMC)

1. Introduction

N,N,N-trimethyl chitosan (TMC) is one of the

most commonly studied chitosan derivatives. It was

developed to improve the properties of chitosan and to

overcome the main barrier in the use of chitosan in

pharmaceutical application, that is, its poor aqueous

solubility at physiological pH. TMC has a fixed positive

charge on the quaternary amino group and the derivative

is therefore highly soluble both in neutral and basic

environments [1]. Ionic liquids (ILs) have gained

attention in the recent years as novel chitosan solvents

[2,3]. Although it has been pointed out that ILs possess

some specific limitations, they bear huge potential as

reaction media for the chemical modification of chitosan

and have been applied for the synthesis of various

chitosan derivatives [4,5]. TMC is usually synthesized by

dispersing chitosan in 1-methyl-2-pyrolidole or NMP

[6,7]. In this work, however, an ionic liquid, 1-butyl-3-

methyl imidazolium chloride (BMIMCl), was studied as

an alternate solvent for the methylation of chitosan by

iodomethane (CH3I) (Scheme 1). In addition, type and

amount of base required in the reaction was also

investigated.

Scheme 1. Synthesis of N,N,N-trimethyl chitosan (TMC)

from chitosan

2. Experimental

2.1 Synthesis of TMC

NMP as a solvent

Chitosan (average Mw of 60 kDa, 92% degree of

deacetylation (DDA), was purchased from Seafresh

Chitosan (Lab.) Co., Ltd. Thailand) was dispersed in

NMP (2.5%w/v) at 55 °C and the mixture was stirred for

16 h Then 1M aqueous sodium iodide (2.68 ml, 4.5

equiv.) and 15% (w/v) aqueous sodium hydroxide (0.89

ml, 6 equiv.) were added and the solution stirred at 55 °C

for 15 min, followed by addition of methyl iodide (0.138

ml, 4 equiv.). The reaction was carried out in a closed

reaction vial at 55 °C. After 2 and 4 h, two more portions

of CH3I (0.138 ml, 4 equiv.) were added into the reaction

solution and the reaction was kept stirring for a total of

24 h. After methylation, the product was precipitated in

acetone. The isolation and purification of product was

perform as described in section 2.2.

BMIMCl as a solvent

Chitosan was completely dissolved in BMIMCl

(2.5%w/v) at 90 °C [4] and the mixture was stirred for 16

OOHO

OH

NH2n

CH3I, aq NaI

aq NaOH, NMP

CH3I

Base, BMIMCl

O

O

O

ON

N

HN

O

CH3

CH3

H3C

HO

OH CH3H3C

CH3

HO

OH

OH

HO O


41 h. Then potassium carbonate (K2CO3) (0.69 g, 9 equiv.)

was added and the mixture was stirred at 55 °C for 15

min, followed by the addition of methyl iodide (0.138

ml, 4 equiv.). The reaction was carried out in a closed

reaction vial at 55 °C. After 2 and 4 h, two more portions

of CH3I (0.138 ml, 4 equiv.) were added into the reaction

solution and the reaction was kept stirring for a total of

24 h. After methylation, the product was precipitated in

ethanol. The isolation and purification of product was

performed as described in section 2.2.

2.2 Isolation and purification of TMC

The solid product was dissolved in 15% (w/v)

NaCl solution in order to replace the iodide counter ion

with a chloride ion. The suspension was dialyzed with

deionized water for 2 days to remove inorganic materials

and then freeze-dried overnight, giving a white and fluffy

trimethylated chitosan

2.3 Measurements

The DQ of methyl group on chitosan was

calculated from the Eq. 1. For degree of dimethylation

(%di-) and monomethylation (%mono-), the peak

integration of protons from the N,N-dimethyl protons,

and N-methyl protons were used for calculation instead

of N,N,N-trimethyl amino, respectively (Eq. 2 and 3).

Equation 1 : Degree of quaternization (%DQ),

%DQ =

Equation 2 : Degree of dimethylation (%di-)

%di- =

Equation 3 : Degree of monomethylation (%mono-)

%mono- =


Dissolution of chitosan in BMIMCl

The dissolution of chitosan by BMIMCl was

accomplished by stirring the chitosan (2.5 %wt) at 90ºC

and 115ºC for 16 hr.[4] The dissolution of chitosan in

BMIMCl was caused by disruption of hydrogen-bonding

between chitosan chains by the BMIM+ and Cl

- ions (Fig.

1). This would cause the polymer chains to move further

apart and would allow easy access of chemical reagent to

react with the functional groups on the chitosan.

Figure 1. Proposed dissolution mechanism of chitosan in

BMIMCl

The solubility of chitosan in BMIMCl is better

than NMP. Due to the presence of anion and cation of

BMIMCl, it can separate inter- and intra-molecular

bonding among the macromolecular chains of chitosan.

But the dissolution of chitosan in NMP generally

requires a much longer time and mostly results in only a

suspension.

Figure 2 1H NMR spectra of (a) TMC (entry6, table1)

and (b) regenerated chitosan (D2O/CF3COOH)

Figure 2 displays 1H-NMR spectrum of TMC as

compared with chitosan. The signals of anomeric proton,

H1’ and H1, appeared at 5.40 and 5.05 ppm,

respectively. The proton signals at 3.65-4.35 ppm were

O

O

OO

O

ON

N HH

OO

H

N

N

HHH

H

H

N

N

Cl

Cl Cl

Cl

N

N

N

NN

N

O

N

N

O

O

OH

OH

HO

HONH2

HN

C O

H3C

O12

3

4

5

7

O O

O

O

O

6',6'

6,66,6 N

N

HN

O

CH3

CH3

H3C

HO

OH CH3

H3C

CH3

HO

OH

OH

HO

2

3'4

2'

34' 5' 5

8 8

810

11'

HOD

1

2

5 6

3 44’1’

5’ 6’

3

3’

2’

2

8 9

10

7

4 (a)

5 6,62’

7

(b)1


42 assigned to H2’,3,4,5,6,6 and the signals at 3.25, 3.0 and

2.8 ppm assigned to N,N,N-trimethyl protons, N,N-

dimethyl protons, and N-methyl protons of the GluN of

chitosan, respectively. Moreover, to confirm the proton

peak positions, COSY spectrum (2-D NMR) of TMC

was determined and shown in Fig. 3.

Figure 3. COSY spectrum of TMC (entry7, table3)

Synthesis of TMC

The study on the reaction condition of synthesis

of TMC was carried out in different types of solvent and

base as shown in Table 1. Comparison between the

reaction carried out in NMP (entry 1) and BMIMCl

(entry 2) at constant amounts of CH3I (12 eq.) and

NaOH/NaI (6/4.5eq.) revealed that in BMIMCl the

obtained %DQ was lower than that obtained from the

reaction occurred in NMP. This is somewhat unexpected

since chitosan was dissolved in the ionic liquid BMIMCl

more than in NMP. Our trials on pyridine (entry 3) and

imidazole (entry 4) as organic base in synthesis of TMC

appeared to be worse due to its weak base property. This

lower %DQ was most likely due to high viscosity effect

in the reaction mixture of BMIMCl, that retarded the rate

of molecular movement to achieve satisfied attacking

rate of the side chain amino groups onto CH3I. However,

Gericke reported that tosylation of cellulose in mixture

of ionic liquid and a co-solvent could be achieved in the

presence of pyridine and imidazole as base [8]. In

nucleophilic substitution of amine group by iodomethane

in organic solvent, K2CO3 was found as a suitable base

for deprotonation of amine group in acetonitrile [9].

From the result, %DQ was increased with increasing of

K2CO3 (entry 5-7). By using 9 equiv of K2CO3 as base

and BMIMCl as solvent, the highest DQ of 42% was

obtained.

Table 1 Different solvents and base conditions for N-

methylation of chitosan

Reaction condition a Product

Entry Solvent Base (equiv.) %DQ %di- %mono

1 NMP NaOH/NaI (6/4.5) 78 41 ND

2 BMIMCl NaOH/NaI (6/4.5) 32 78 ND

3 BMIMCl Pyridine (6) trace ND ND

4 BMIMCl Imidazole (6) trace ND ND

5 BMIMCl K2CO3 (2.5) 14 51 11

6 BMIMCl K2CO3 (9) 42 66 1

7 BMIMCl K2CO3 (12) 39 42 trace

a CH3I 12 equiv., reaction time 24 h and reaction temp.

55°C b ND is not detected

4. Conclusion

In this work, the method for TMC synthesis was

carried out by using BMIMCl, an ionic liquid, as a

solvent that could completely dissolve chitosan. Despite

this advantage, the %DQ on chitosan was about half of

that obtained the reaction in NMP solvent. This could be

attributed to the increase in viscosity of the reaction

mixture which BMIMCl was used as a solvent. Further

studies are currently performed in order to find an

optimized reaction condition in order to maximize the

%DQ on chitosan by reducing solution viscosity and the

recycle of BMIMCl

Acknowledgements

The financial support for this project was

provided by CU GRADUATE SCHOOL THESIS

GRANT.

Reference

[1] Benediktsdóttir, B. E., Baldursson, Ó., Másson, M.

"Challenges in evaluation of chitosan and

trimethylated chitosan (TMC) as mucosal

permeation enhancers: From synthesis to in vitro

12

3 4

5

6

1’2’

3’4’

5’ 6’

1’

2’

3’4’

5’6’

1

3

4

2

5

6

8

8

9

9

10

10

HOD

HOD


43 application", Journal of Controlled Release, 173,

18-31 (2014).

[2] Liu, L., Zhou, S., Wang, B., Xu, F., Sun, R.

"Homogeneous acetylation of chitosan in ionic

liquids", Journal of Applied Polymer Science, 129,

28-35 (2013).

[3] Wang, Z., Zheng, L., Li, C., Zhang, D., Xiao, Y.,

Guan, G., Zhu, W. "Modification of chitosan with

monomethyl fumaric acid in an ionic liquid

solution", Carbohydrate Polymers, 117, 973-979

(2015).

[4] Hua, D., Jiang, J., Kuang, L., Jiang, J., Zheng, W.,

Liang, H. "Smart Chitosan-Based Stimuli-

Responsive Nanocarriers for the Controlled Delivery

of Hydrophobic Pharmaceuticals", Macromolecules,

44, 1298-1302 (2011).

[5] Peng, P., Cao, X., Peng, F., Bian, J., Xu, F., Sun, R.

"Binding cellulose and chitosan via click chemistry:

Synthesis, characterization, and formation of some

hollow tubes", Journal of Polymer Science Part A:

Polymer Chemistry, 50, 5201-5210 (2012).

[6] Verheul, R. J., Amidi, M., van der Wal, S., van Riet,

E., Jiskoot, W., Hennink, W. E. "Synthesis,

characterization and in vitro biological properties of

O-methyl free N,N,N-trimethylated chitosan",

Biomaterials, 29, 3642-3649 (2008).

[7] Benediktsdóttir, B. E., Gaware, V. S., Rúnarsson, Ö.

V., Jónsdóttir, S., Jensen, K. J., Másson, M.

"Synthesis of N,N,N-trimethyl chitosan

homopolymer and highly substituted N-alkyl-N,N-

dimethyl chitosan derivatives with the aid of di-tert-

butyldimethylsilyl chitosan", Carbohydrate

Polymers, 86, 1451-1460 (2011).

[8] Gericke, M., Schaller, J., Liebert, T., Fardim, P.,

Meister, F., Heinze, T. "Studies on the tosylation of

cellulose in mixtures of ionic liquids and a co-

solvent", Carbohydrate Polymers, 89, 526-536

(2012).

[9] Vedejs, E., Kongkittingam, C. "Solution-Phase

Synthesis of a Hindered N-Methylated Tetrapeptide

Using Bts-Protected Amino Acid Chlorides:

Efficient Coupling and Methylation Steps Allow

Purification by Extraction", The Journal of Organic

Chemistry, 65, 2309-2318 (2000).


44 BIOENP-12

Synthesis and characterization of hyper-branched poly(L-lactide) by using polyglycidol

Nichakorn Pathumrangsan1, Atitsa Petchsuk

2 and Pakorn Opaprakasit

1*

1 School of Bio-Chemical Engineering and Technology, Sirindhorn International Institute of Technology (SIIT),

Thammasat University, Pathum Thani, 12121, Thailand 2 National Metals and Materials Technology Center (MTEC), Pathum Thani, 12120, Thailand

Abstract

Hyper-branched poly(L-lactide) (hbPLLA)s with various arm lengths are synthesized by ring-opening

polymerization of L-lactide (LLA) using polyglycidol (PG) as a macro-initiator. hbPLLAs are blended with linear

PLLA (l-PLLA) by varying the LLA branch content. Thermal and rheological properties, and optical transparency of

hbPLLAs and their blends with l-PLLA are investigated. All l-PLLA/hbPLLAs blends show slightly changes in of Tg

values, whereas the Tm is significantly unchanged. A single Tg is observed in all blends, indicating a completely

miscible system. All blends exhibit an increase in crystallinity, as the branch structure act as nucleating agent for

crystallization of l-PLLA. Viscosity of the blends decreases with the addition of l-PLLA. This provides easy processing

conditions. The blends also show high optical transparency, comparable to neat l-PLLA. Given these properties and

their biocompatibility, the blends can be used in biomedical applications.

Keywords: Polylactide, Branch structure, Rheology, Blend, Polyglycidol

1. Introduction

Nowadays, polylactide (PLA) has received much

attention, due to serious environmental problems on

plastic wastes. PLA is one of well-known degradable

polymers, which provides many good properties, such as

high mechanical strength, transparency, and

biocompatibility [1-3]. PLAs are widely used in many

applications, especially in biomedical field [4-6].

However, PLA-based materials possess certain

disadvantages which limit their use in some applications,

e.g., brittleness, and difficulty in controlling degradation

rates. Many approaches have been performed to

overcome these drawbacks, such as stereocomplexation,

introduction of branch-structured PLA, and blending

with other polymers. Among these, introducing of branch

structures into PLA matrix is a promising method to

solve these problems. Polymers with branch architectures

typically have lower glass transition temperature (Tg) and

melt viscosity than their linear counterparts of similar

molecular weight. Moreover, branch length is an

important parameter that affects the viscoelasticity of

fluidity range and crystallinity [7]. The use of various

hydrophilic cores have been reported in preparation of

branched PLA copolymers, such as, poly(ethylene glycol)

(PEG), poly(ethylene oxide) (PEO)[8-10], poly(amido amine)

(PAMAM) [11], and polyglycidol (PG) [7, 12-13]

In this study, hyper-branched PLLA (hbPLLA) is

developed for intended use in biomedical applications.

PG is chosen as a hydrophilic core because of its multi-

functionality, which can be used as a macro-initiator for

polymerization of PLA, and its biocompatibility [12].

Ring-opening polymerization in bulk is employed. The

resulting hbPLLAs is blended with linear PLLA (l-

PLLA) to optimize its physical and rheological

properties.

2. Experimental

2.1 Materials

L-Lactide (LLA) and tin octoate (Sn(Oct)2) were

purchased from Wako (Japan). Linear PLLA (l-PLLA)

( = 178 000 g/mol) was supplied by PURAC

(Netherland). PG macro-initiator was synthesized

according to a methodology reported earlier[14]. Ethyl

acetate, chloroform, ethanol and toluene solvents were

purchased from Lab Scan (Thailand).

2.2 Synthesis of hbPLLAs

The synthesis of hbPLLAs was performed by a

ring-opening polymerization in bulk, using Sn(Oct)2

catalyst and PG macro-initiator. Essentially, PG was

dried in a reactor under vacuum at 70oC for 1 h. After

drying, Sn(Oct)2 (1 %wt of macro-initiator) was added,

and the mixture was heated to 80oC for 1 h. After that,

LLA was added and reaction was further kept at 130oC

for 24 h. The feed ratios of LLA to PG were varied at

10/1, 20/1, 50/1, and 100/1. Finally, the reaction mixture


45 was dissolved in chloroform, and precipitate in a large

amount of ethanol to remove unreacted LLA and PG.

The powder precipitant of hbPLLAs was dried in a

vacuum oven at 50oC for 3-4 days.

2.3 Blending process

Blends of hbPLLAs with different structures and

l-PLLA were prepared by employing an internal mixer

(MX105-D40L50) using a rotor speed of 50 rpm and

blending time and temperature of 20 min and 170oC.

Blend ratios of hbPLLAs to l-PLLA of 10/90 was

employed. The blended samples were then presses into a

film form by a compression machine (PR2D-W300L300

HD-WCL).

2.4 Characterizations

Chemical structures of hbPLLA sampls were

characterized on an AVEN-CEIII 500 MHz digital

Nuclear Magnetic Resonance spectrometer (NMR) (AV-

500, Bruker Biospin), using CDCl3 solvent. For thermal

properties, hbPLLAs and their blended samples were

measured by differential scanning calorimetry (DSC)

(DSC822e Mettler Toledo) at a heating/cooling rate of

20 oC/min. All specimens were heated to 200

oC (first

scan) to erase their thermal history, and then cooled to -

20oC. The samples were then heated from -20 to 220

oC.

Rheological properties, in terms of complex

viscosity (*) of the blends were measured on a strain-

controlled rheometer (ARES, TA Inc., New Castle,

USA). Samples were prepared into a disc form with a

diameter of 25 mm and 1 mm thickness. The strain

amplitude was fixed at 0.5%. The samples were scanned

from 140 – 200 oC with a heating rate of 10

oC/min at a

frequency of 1 rad/s.


3.1 Chemical structures and properties of hbPLLAs

Hyper-branched PLLA is formed by a ring-

opening polymerization of LLA in a presence of

Sn(Oct)2 catalyst. PG is employed as a macro-initiator,

which subsequently serves as hyper-branched core

structure. The coordination-insertion mechanism is

proposed, as shown in Figure 1. Molecular exchange of

branched PG with the octoate ligands occurs, followed

by the coordination of LLA to the metal center. Insertion

of branched PG, followed by ring-opening generates a

linear monomer and starts propagation.

Figure 1 Coordination-insertion mechanism of Sn(Oct)2

catalyzed polymerization of LLA, where ROH represents

reactive sites of PG core.

Figure 2 Chemical structures and 1H-NMR spectra of

hbPLLAs synthesized from different feed ratios:

(A)hbPLLA101, (B)hbPLLA201, (C)hbPLLA501, and

(D)hbPLLA1001

The proposed structure and 1H-NMR spectra of

hbPLLAs, synthesized from various feed ratios are

shown in Figure 2. Signals due to PLLA structures are d,

e and f ( = 1.5, 5.1 and 4.3 ppm), which are assigned to

methyl, and methine protons in main chain and terminal

units, respectively. Signals a and b ( = 3.5 ppm) are

associated with methylene and methine protons of PG.

The ratio of the integral values of the signals e/f indicates

ROH (HPG)


46 an average arm length of LLA branches, while that of

e/(a+b) represents LLA/PG compositions in the chains.

The values for all hbPLLA samples are summarized in

Table 1.

Table 1. Results on chemical structures of hbPLLAs, in

terms of LLA⁄PG compositions and arm lengths of LLA

sequences.

Sample

composition Average arm

length of

LLA a

in feed in

chain

hbPLLA101 (A) 10/1 22/1 13

hbPLLA201 (B) 20/1 29/1 16

hbPLLA501 (C) 50/1 65/1 38

hbPLLA1001 (D) 100/1 187/1 71

a Calculated from integral ratios of e/f

3.2 Thermal properties of hbPLLAs and their blends

Thermal properties of hbPLLAs and their

blends are examined by DSC, in terms of glass transition

temperature (Tg), melting temperature (Tm), crystallization

temperature (Tc), heats of crystallization (Hc) and

melting (Hm). Results in Table 2 show that Tg of

hbPLLAs increases with an increase in their LLA arm

length, because chain mobility or segmental relaxation of

chains decreases. Melting temperature peak is not

observed in hbPLLAs with short LLA branches, i.e.,

hbPLLA101 and hbPLLA201. However, those with

longer LLA sequences possess Tm at 147 and 163 oC.

This indicates that a critical branch length of LLA

sequences is required to form crystal domains.

Effects of structures of hbPLLAs on thermal

properties of l-PLLA/hbPLLA blends are investigated by

keeping the blend composition at 90/10 wt%. The results

are summarized in Table 3. All blended samples show a

single Tg at a temperature comparable to or slightly

lower than that of l-PLLA matrix, indicating a complete

miscible blend system. The blends containing hbPLLAs

with short LLA sequences possess lower Tg values,

compared to those with the longer counterparts. This

indicates a lower synergetic effect of hbPLLAs and the l-

PLLA matrix, due to their relative higher PG contents in

the chains.

Table 2. Thermal properties of neat l-PLLA and

hbPLLAs with different structures.

Sample Tg

[oC]

Tm

[oC]

Hm

[J/g]

Tc

[oC]

Hc

[J/g]

l-PLLA 63 153 0.2

hbPLLA101 48 - - - -

hbPLLA201 47 - - - -

hbPLLA501 57 147 36.3 121 29.3

hbPLLA1001 61 163 41.0 121 34.9

Tm of all blends are significantly unchanged with

the addition of hbPLLAs. However, their crystallinity is

higher than neat l-PLLA. This reflects that hbPLLAs act

as a nucleating agent inducing crystallization of the l-

PLLA matrix. The degree of crystallinity ( ) is

calculated using the following equation:

=

(1)

Where: Hm is the heat of fusion of the samples

Ho

m is the heat of fusion of completely crystallized

PLA, i.e. 93 J g-1

[15].

Table 3. Thermal properties (2nd

heating scan) of l-

PLLA and l-PLLA/hbPLLAs (90/10) blends containing

different hbPLLAs

Blends Tg

[oC]

Tc

[oC]

Tm

[oC]

Hc

[J/g]

Hm

[J/g]

c

(%)

l-PLLA 63 - 153 - 0.2 0.2

l-PLLA/

hbPLLA101 61 131 156 15 15 16.1

l-PLLA/

hbPLLA201 61 125 153 24 25 26.9

l-PLLA/

hbPLLA501 64 132 154 12 11 11.8

l-PLLA/

hbPLLA1001 62 125 155 21 25 26.9

3.3 Rheological properties of l-PLLA/hbPLLA blends

Complex viscosity (*) of l-PLLA/hbPLLA

blends consisting of different hbPLLAs at the blend

composition of 90/10 wt. is measured, as a function of

temperature at a fixed strain of 0.5%. The results are


47 shown in Figure 3. At the temperatures below Tm of the

samples, blends containing short LLA branches exhibit

low * values, as this acts as a plasticizer in the l-PLLA

matrix. This is in accord with our previous results [14].

In contrast, blends consisting of long LLA branches,

especially hbPLLA501, show an increase in the values,

compared to neat l-PLLA, indicating strong interaction

between the blend components, likely due to higher

degree of chain entanglements. However, at the

temperatures range higher than Tm, where the samples

are in melt state, the samples is observed in an opposite

trend. This is likely due to the contribution of the

branched structure with longer arm lengths. The insight

into this property is essential in fabrication of PLLA

products for use in biomedical applications.

Figure 3. Temperature dependence of complex viscosity

(*) of l-PLLA/hbPLLA blends containing different

hbPLLAs.

3.4 Optical transparency of blended films

One of the unique characteristics of PLLA

products is their high optical transparency. This is

usually deteriorated upon blending with other

components, which leads to limitation in certain

applications. Effect of an addition of hbPLLAs on

transparency of the blends are examined using hot-

pressed films with a thickness ranging from 0.15 – 0.20

mm. The results, as shown in Figure 3, indicate that all

blends have comparable transparency to that of neat l-

PLLA.

Figure 4. Optical transparency of l-PLLA/hbPLLA

blends.

4. Conclusions

Hyper-branched PLLA is synthesized by bulk

polymerization of LLA using polyglycidol as a macro-

initiator. The copolymers are then blended with l-PLLA,

in which a complete miscible blend system is obtained.

Upon introducing of hbPLLAs, it is observed that

crystallinity of the blends increase, as the branch

structured act as a nucleate agent inducing crystallization

of l-PLLA matrix. The addition of hbPLLAs also play an

important role in rheological properties of the blends,

without affecting their optical transparency. Given these

properties and their biocompatibility, the blends can be

used in biomedical applications.

5. Acknowledgements

Financial support provided from the National Research

University (NRU) project of Thailand is gratefully

acknowledged. N.P. is thankful for a support from the

SIIT scholarship program.

References

[1] Ajioka, M., Enomoto, K., Suzuki, K., Yamaguchi,

A., “The basic properties of poly (lactic acid)

produced by the direct condensation polymerisation

of lactic acid”, J Environ Polym Degrad, 225-234

(1995).

[2] Tuominen, J., Kylma, J., Kapanen, A., Venelampi,

O., Itavaara, M., and Seppala, J., “Biodegradation of

lactic acid based polymers under controlled

composting conditions and evaluation of the

1000

10000

100000

140 150 160 170 180 190 200 210 220 230

*

[Pa.s

]

Temperature [oC]

Pure

Neat PLLA

PLLA/b-PLLA101_10

PLLA/b-PLLA201_10

PLLA/b-PLLA501_10

PLLA/b-PLLA1001_10

a


48 ecotoxicological impact”, Biomacromolecules,445-

455 (2002).

[3] Yujiang, F., Haruo, N.,Yoshihito, S., Yutaka, T.,

Takeshi, E., “Thermal degradation behaviour of

poly(lactic acid) stereocomplex”, Polym Degrad

and Stabil, 197-208 (2004).

[4] Uurto, I., Mikkonen, J., Parkkinen, J., Keski-Nisula,

L., Nevalainen, T., Kellomäki, M., Törmälä, P.,

Juha-Pekka, S., “Drug-eluting biodegradable poly-

D/L-lactic acid vascular stents: An Experimental

Pilot Study”, J Endovasc Ther, 371–379 (2005).

[5] John, A.O., and Patrick, W.S.S., “Bioabsorbable

Coronary Stents”, Circ Cardiovasc Interv, 255-260

(2009).

[6] Rahul, M.R., Amol, V.J., Douglas, E.H.,

“Poly(lactic acid) modifications”, Progress in

Polymer Science, 338–356 (2010).

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branched poly(lactide) using polyglycidol and

thermal, mechanical properties of its solution-cast

film”, Polymer, 429–434 (2006).

[8] Young, K. C., You, H. B., Sung, W. K., “Star-

shaped poly(ether-ester) block copolymers:

synthesis, characterization, and their physical

properties”, Macromolecules, 8766 – 8774 (1998) .

[9] Pistel, K.F., Bittner, B., Koll, H., Winter, G., Kissel,

T., “Biodegradable recombinant human

erythropoietin loaded microspheres prepared from

linear and star-branched block copolymers:

influence of encapsulation technique and polymer

composition on particle characteristics”, J

Controlled Release, 309 – 325 (1999).

[10] Salaam, L.E., Dean, D., Bray, T.L., “In vitro

degradation behavior of biodegradable 4-star

micelles”, Polymer, 310–318 (2006).

[11] Cai, Q., Zhao, Y., Bei, J., Xi, F., Wang, S. ,

“Synthesis and properties of star-shaped polylactide

attached to poly(amidoamine) dendrimer”,

Biomacromolecules, 828-34 (2003).

[12] Kainthan, R.K., Janzen, J., Levin, E., Devine, D.V.,

Brooks, D.E.., “Biocompatibility testing of branched

and linear polyglycidol”, Biomacromolecules, 703-

709 (2006).

[13] Jeffrey, L. A., Sergey, V., “Thermal properties and

degradation behavior of linear and branched poly(L-

lactide)s and poly(L-lactide-co-glycolide)s”,

Macromol Chem Phys, 924−929 (2012).

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Opaprakasit, M., Opaprakasit, P., “Preparation and

properties of multi-branched poly(D-lactide) derived

from polyglycidol and its stereocomplex blends”,

Express Polym Lett, 779-789 (2014).

[15] Phuphuak, Y., Chirachanchai, S., “Simple

preparation of multi-branched poly(L-lactic acid)

and its role as nucleating agent for poly(lactic

acid)”, Polymer, 572–582 (2013).


49

BIOENP-13

Synthesis and properties of hyper-branched polylactide employing polyethylene imine core

Narisara Jaikaew1, Atitsa Petchsuk

2, Pakorn Opaprakasit

1,*


Thammasat University, Pathum Thani, 12121, Thailand 2 National Metals and Materials Technology Center (MTEC), Pathum Thani, 12120, Thailand

Abstract Hyper-branched polylactide (hbPLA)s with various branched lengths are synthesized by ring-opening

polymerization of L-lactide (LLA) in bulk using poly(ethylene imine) (PEI) as a core macro-initiator. The polymer’s

branched lengths are varied by adjusting the PEI:LLA feed ratios. The synthesized hbPLLAs are used as additives for

properties enhancement of commercial linear PLLA (l-PLLA) by melt blending. Miscibility, thermal and rheological

properties of the resulting hbPLLA/l-PLLA blends consisting of different hbPLLAs are investigated. All blend samples

exhibit lower glass transition temperature (Tg), crystalline melting temperature (Tm), and complex viscosity than neat l-

PLLA, which can provide many advantages in processing of the materials. The blends properties can also be further

optimized for specific applications by varying the branched structured component.

Keywords: Degradable polymer, Hyper-branched polylactide, Poly(ethylene imine), Polymer blends

1. Introduction

Degradable polyesters, whose major advantages

being biocompatible and degradable, are a group of

materials of interest in biomedical, pharmaceutical, and

environmental applications for replacing traditional non-

degradable polymers. [1, 2]. These materials are

thermoplastic polymers consisting hydrolysable linkages

in their backbone [3]. Among these, polylactide or poly

(lactic acid) (PLA), which is classified as thermoplastic

aliphatic polyester, is one of the most attractive

materials, due to its degradability, good plasticity,

suitable processability, high mechanical strength,

relatively low cost of production, and renewability [4, 5].

PLA can be synthesized by either

polycondensation of lactic acid, or ring-opening

polymerization (ROP) of lactide, which is the most

effective process to obtain high molecular weight PLA.

However, the polymer exhibits poor thermal stability and

low melt strength. This limits its use in several ways [2,

6-8]. Recently, many attempts have been made to

improve PLA’s properties, such as copolymerization,

blending with other polymers or plasticizers, toughening

by rubber materials, introduction of branched structures,

stereocomplexation, and nano-composites reinforcement

[1, 9].

The introduction of branched structures into PLA

is proven as a promosing way to improve its properties

[2]. Branch-structured polymers, at the same molecular

weight, have lower solution and melt viscosity than its

linear-structured counterparts. An addition of these

materilas leads to improvements in physio-chemical

properties of the matrix [2, 10]. Dendrimic, and hyper-

branched polymers are macromolecules containing a

center core. This group of materials have received vast

attention, due to their unique properties of low viscosity

and a high degree of surface functionality, which can be

further designed and controlled. In addition, one of the

major advantages of hyper-branched polymers is that

they can be synthesized in a one-step process [11].

In this work, hyper-branched PLLAs

(hbPLLAs) are prepared by copolymerization of l-lactide

(LLA) using poly(ethylene imine) (PEI) as a macro-

initiator. The synthesized products are used to modify

properties of commercial l-PLLA by melt blending.

Effect of branching structures on properties of the blends

are studied by varying the PEI:LLA feed ratios. Thermal

and rheological properties of the blends as a function of

hbPLLA structures and compositions are investigated.

2. Experimental

2.1 Materials

L-Lactide (LLA) and l-PLLA (4043D) were

supplied by PURAC (Netherlands). Poly(ethylene imine)

(PEI) (Mn ~1,800 by GPC) was purchased from Aldrich


50 (USA). Ethyl acetate, chloroform, ethanol, and toluene

solvents were purchased from Lab Scan (Thailand). Tin

(II) octoate catalyst, Sn(Oct)2, was obtained from Wako

(Japan).

2.2 Synthesis and characterization

hbPLLAs were synthesized by ROP using PEI as

a macro-initiator and Sn(Oct)2 catalyst. The reaction was

performed in a round-bottom flask, equipped with a

condenser. The macro-initiator and monomer were first

dried for 2 hours under vacuum before polymerization.

The weight ratio of PEI to LLA was varied from 1:10,

1:20 to 1:100. PEI was stirred at 70˚C for 1 hour in the

reaction flask, before dissolving in dried THF. Sn(Oct)2

catalyst was then added, and the mixture was stirred at

80˚C for 2 hours, before adding LLA monomer. The

polymerization was kept at 110 ˚C for 5 hours with

continuous stirring. The mixture was finally dissolved in

chloroform, and then precipitated in a large amount of a

methanol/hexane mixture to obtain solid hbPLLA

products and remove unreacted LLA and PEI. The

synthesized products were blended with a commercial l-

PLLA by melt blending in an internal mixer at 170 ˚C for

20 minutes. The hbPLLA contents were varied from 5 to

20 %wt.

Chemical structures of hbPLLA products were

investigated by an AVEN CEIII 500 MHz digital nuclear

magnetic resonance spectrometer (AV-500, Bruker

Biospin), using CDCl3 (for hbPLLAs) and D2O (for PEI)

solvents. Thermal property of the products was evaluated

by differential scanning calorimetry (DSC822e Mettler

Toledo). The evaluation was performed by a heat-cool-

heat standard method from -20 to 200˚C, at a heating and

cooling rate of 20˚C/min. The first heating step is

employed to erase the sample’s thermal history.

Dynamic rheological measurements were carried out

using a strain-controlled rheometer (ARES, TA Inc.,

New Castle, USA) with a torque transducer capable of

measurement over the range of 2–200 g.cm. The strain

amplitude for dynamic measurements was fixed at 0.5%.

The samples were prepared by a compression molding

machine (Chareon Tut, Thailand) into a disc shape with a

diameter of 25 mm and 1 mm thickness. The shear

function mode for frequency dependence test was also

performed from low to high shear forces, at a constant

temperature of 180°C.


3.1 NMR spectroscopy

Chemical structures of hbPLLAs synthesized at 3

different PEI:LLA feed ratios are examined by 1

H NMR

spectroscopy. Figure 1 shows 1H NMR spectrum of the

materials obtained from a 1:10 ratio (denoted as

hbPLLA10). The spectrum shows characteristic

chemical shifts at 5.13 ppm (a, b), assigned to methine

protons in the main-chain repeat units, with the

integration of 1.00. The signal at 4.32 ppm(c) is assigned

to methine protons in terminal units of hbPLLA, whose

integration is 0.098. The chemical shifts at 1.55 ppm (d,

e) are assigned to the methyl group. A broad signal

covering the region of 2-3 ppm is associated with CH2

protons of PEI. This is compared with a 1H NMR

spectrum of neat PEI, as shown in Figure 2. The

integration ratios of these characteristic signals are

employed in the evaluation of chemical structures of the

products, in terms of LLA branch lengths ( ) and

PLLA/PEI molar ratios in the hyper-branched copolymer

chains ( ). The results are summarized in Table

1.

Figure 1. 1H NMR spectrum and chemical structure of

hbPLLA10.

(2)


51

Figure 2. 1H NMR spectrum and chemical structure of

PEI.

Table1. Results on average LLA lengths and PLLA/PEI

molar ratio in hbPLLA chains.

Sample PEI: LLA

feed

ratios

Average LLA

branch length

(units)

PEI :LLA

molar ratios

(in chain)

hbPLLA10 1:10 10 1:5

hbPLLA20 1:20 12 1:9

hbPLLA100 1:100 19 1:11

3.2 Thermal property

Thermal properties of the resulting hbPLLAs

are examined, in which their DSC thermograms are

shown in Figure 3. The results indicate that the

properties are strongly dependent on the polymer’s

chemical structures. hbPLLA with short LLA arm’s

length (i.e. hbPLLA10 and hbPLLA20) are complete

amorphous, as no crystalline characteristics are observed

in their DSC thermograms. In contrast, the copolymer

with longer arm’s length (i.e. hbPLLA100) processes a

semi-crystalline structure, where a melting characteristic

peak (Tm) is observed at 162°C. The glass transition

temperature (Tg) of the hyper-branched copolymer

increases with the length of LLA sequences from 36 to

45 and 47°C, as a result from a decrease in the chain

flexibility.

Thermal properties of l-PLLA/hbPLLA blends

containing various hbPLLAs, as a function of the blend

compositions, are investigated. All blends show lower Tg

and Tm than neat l-PLLA. Both Tg and Tm values

decrease with an increase in the hbPLLA content, and

also increase when the LLA arm’s length further

increases.

Figure 3. DSC thermogram (2nd

heating scan) of (a)

HbPLLA100, (b) HbPLLA20 and (c) HbPLLA10.

DSC thermograms of l-PLLA/hbLLA10 blends

as a function of the blend compositions are shown in

Figure 4. Single Tg is observed in all samples, indicating

complete miscible blend systems at all compositions. No

melting characteristic peak is observed in the

thermogram of hbPLLA10, due to its short LLA arm

length, while neat l-PLLA exhibits a Tm at 150°C. Up on

adding of hbPLLA10, a decrease in ΔHm of the blends is

observed. The degree of reduction in the values increase

with an increase in the hbPLLA10 content from 5 to

20 %wt, reflecting that this short-branched structure has

high synergic effect with the l-PLLA matrix, which

inhibits the crystal formation of the l-PLLA domains.

The corresponding thermograms of blends

consisting of hbPLLA100, which possess long LLA

sequences, are shown in Figure 5. Single Tg is also

observed in all samples, reflecting complete miscible

blend systems. hbPLLA100 exhibits a Tm at 162 °C,

which is higher than that of neat l-PLLA. This is likely

because its crystalline domains are derived from a close

packing of branches with optimum length. When the

copolymer is blended with l-PLLA, double Tm peaks

appear at the temperatures corresponding those of the 2

components. This indicates formation of 2 distinct

crystalline domains originated from hbPLLA100 and the

l-PLLA matrix.


52

Temperature (°C)

0 50 100 150 200

(a)

(b)

(c)

(d)

(e)

(f)

Figure 4. DSC thermograms (2nd


hbPLLA10 and l-PLLA/hbPLLA10 blends at various

(%wt) compositions: (b) 80/20, (c) 85/15, (d) 90/10, (e)

95/5 and (f) l-PLLA.

0 50 100 150 200

(f)

(e)

(d)

(c)

(b)

(a)

Temperature (°C)

Figure 5. DSC thermograms (2nd


hbPLLA100 and l-PLLA/hbPLLA100 blends at various

(%wt) compositions: (b) 80/20, (c) 85/15, (d) 90/10, (e)

95/5 and (f) l-PLLA.

3.3 Rheological property

The viscoelastic characteristics, in terms of

complex viscosity, of l-PLLA/hbPLLAs blends (at a

composition of 90/10) as a function of shear rate are

shown in Figure 6. The results indicate that all blends

have lower complex viscosity than that of neat l-PLLA,

as of result from the addition of the branch-structured

component. The lowest values are observed in the

blends containing hbPLLA10, because of it short LLA

sequences. When the arm length increases, the blend’s

complex viscosity increases, because the blend

component consists of branches that are long enough to

entangle with l-PLLA chains. Insight into the origin of

this property is essential, and can be applied in many

applications, especially in optimizations of flow

behaviors of PLLA during its processing and fabrication

processes. Therefore, hbPLLAs with different structures

can be employed as processing aids.

Figure 6. Complex viscosity of l-PLLA/hbPLLAs blends

containing different hbPLLAs, at a 90/10 composition.

4. Conclusions

hbPLLAs are successfully synthesized by bulk

polymerization of LLA with PEI core. The LLA branch

length and LLA contents of the copolymer chains

increase with an increase in the feed ratios from 1:10 to

1:100. The hbPLLA products are blended with l-PLLA.

The results indicate that all hyper-branched copolymers

are completely miscible with l-PLLA. The LLA arm

lengths and the blend compositions impose strong effect

on thermal properties and complex viscosity of the

blends. The hbPLLAs can be used as additive in

processing and fabrication of PLA products.

5. Acknowledgements

Financial support provided from the National Research

University (NRU) project of Thailand is gratefully

acknowledged. N.J. is thankful for a support from the

SIIT scholarship program.

References

[1] Zheng W, Li J, Zheng YF. "Preparation of poly(l-

lactide) and its application in bioelectrochemistry",

Journal of Electroanalytical Chemistry: 69-74

(2008).

0.01

0.1

1

10

100

1000

10000

100000

0.00001 0.001 0.1 10 1000

ƞ* [

Pa.s

]

ý [1/s]

l-PLLA

90/10 HbPLLA10

90/10 HbPLLA20

90/10 HbPLLA100


53 [2] Ouchi T, Ichimura S, Ohya Y. "Synthesis of

branched poly(lactide) using polyglycidol and

thermal, mechanical properties of its solution-cast

film", Polymer: 429-434 (2006).

[3] Nair LS, Laurencin CT. "Biodegradable polymers as

biomaterials", Progress in Polymer Science: 762-

798 (2007).

[4] Raquez J-M, Habibi Y, Murariu M, Dubois P.

"Polylactide (PLA)-based nanocomposites",

Progress in Polymer Science.

[5] Rhim J-W, Hong S-I, Ha C-S. "Tensile, water vapor

barrier and antimicrobial properties of

PLA/nanoclay composite films", LWT - Food

Science and Technology: 612-617 (2009).

[6] Imanaka M, Takeuchi Y, Nakamura Y, Nishimura

A, Iida T. "Fracture toughness of spherical silica-

filled epoxy adhesives", International Journal of

Adhesion and Adhesives: 389-396 (2001).

[7] Urbanczyk L, Ngoundjo F, Alexandre M, Jérôme C,

Detrembleur C, Calberg C. "Synthesis of

polylactide/clay nanocomposites by in situ

intercalative polymerization in supercritical carbon

dioxide", European Polymer Journal: 643-648

(2009).

[8] Najafi N, Heuzey MC, Carreau PJ, "Wood-Adams

PM. Control of thermal degradation of polylactide

(PLA)-clay nanocomposites using chain extenders",

Polymer Degradation and Stability: 554-565 (2012).

[9] Pongtanayut K, Thongpin C, Santawitee O. "The

Effect of Rubber on Morphology, Thermal

Properties and Mechanical Properties of PLA/NR

and PLA/ENR Blends", Energy Procedia: 888-897

(2014).

[10] Zhao R-X, Li L, Wang B, Yang W-W, Chen Y, He

X-H, et al. "Aliphatic tertiary amine mediated

synthesis of highly branched polylactide

copolymers", Polymer: 719-727 (2012).

[11] Zhang C-X, Wang B, Chen Y, Cheng F, Jiang S-C.

"Amphiphilic multiarm star polylactide with

hyperbranched polyethylenimine as core: A

systematic reinvestigation", Polymer: 3900-39009

(2012).


54 BIOENP-14

Glycolysis of poly (l-lactic acid) using microwave irradiation

T.Sripho 1, P. Opaprakasit

1*, and A. Petchsuk

2

1School of Bio-Chemical Engineering and Technology, Sirindhorn International Institute of Technology (SIIT),

Thammasat University, Pathumthani, 12121 Thailand

2National Metal and Materials Technology Center (MTEC), Pathumthani, 12120 Thailand

Abstract

Chemical-recycling process of polylactic acid (PLA) is developed by using a glycolysis reaction. Microwave

irradiation is employed to enhance the performance of the glycolysis reaction, compared to conventional heating.

Effects of glycolysis conditions on chemical structures and degree of polymerization (DP) of the glycolysed PLA

products (GlyPLA) are investigated. The results indicate that the reaction time can be dramatically decreased when

microwave irradiation is employed. The degree of depolymerization, i.e., chain length of GlyPLA, is affected by the

reaction temperature, reaction time, and EG: PLA feed ratios. GlyPLA products with shorter chain length is obtained,

when higher EG: PLA weight ratio is used. At low reaction temperature of 190°C, an increase in the reaction time leads

to a formation of smaller-sized products, whereas an opposite trend is observed when the reaction temperature is

increased to 200°C, as the oligomers undergoes transesterification. The insight into the reactions mechanisms can be

applied in the production of glycolized products for specific applications.

Keywords: Polylactic acid, Chemical recycling, Glycolysis, Microwave irradiation

1. Introduction

Polylactic acid (PLA) is well known as

degradable bio-based materials, in which its monomer

can be produced by fermentation of starch-rich plants,

such as corn, wheat and sugar beets. The polymer is

classified as an environmental-friendly material, as its

production and degradation after use produce lower

amount of carbon dioxide emission, compared to its

fossil-based counterparts(1, 2). The demand of PLA

consumption has been forecasted to rapidly increase in

the near future, because of its unlimited supplied sources

and its wide variety of applications (3, 4).

Although PLA is degradable polymer, an

effective process for recycling this material is still

required to increase efficiency of waste utilization and

decrease environmental impact. Recycling process of

PLA can be divided into two aspects; mechanical and

chemical recycling (3). Although mechanical recycling

process utilizes a simple process, this has some

limitations, especially deterioration of physical and

mechanical properties of the recycling products, i.e.,

decreases in stress and strain at break point (5).

Therefore, chemical recycling processes has attracted

vast attention in attempts to increase efficiency of PLA

recycling processes (6).

Solvolysis is one of the most important chemical

recycling methods for PLA, which includes alcoholysis,

hydrolysis, and glycolysis (7). Hydrolysis and

alcoholysis reactions have been widely employed in

recycling of PLA. . However, there are some

disadvantages from these recycling processes. For

instance, degradation rate and mechanism of hydrolysis

highly depends on acidity and alkaline conditions (8). If

not treated properly, these acidic or alkali solvents will

be released to the environment, and seriously affect the

eco-system. Moreover, intensive energy consumption is

another drawback of hydrolysis and alcoholysis

reactions, as both chemical recycling processes have to

operate under high pressure and temperature (9). Glycolysis recycling method is commonly employed in

recycling of aromatic polyester, i.e., poly(ethylene

terephthalate) (PET). Bis(hydroxyethyl) terephthalate

(BHET) and oligomers of PET are major products from

this recycling reaction(10). In our previous work (11-13)

the similar glycolysis reaction was employed in the

recycling of aliphatic PLA. The reaction was conducted

at high reaction temperatures at reaction times ranging

from 1-5 hours. The major products from the reaction

were oligomers of PLA.

Microwave heating involves the conversion of

electromagnetic energy into heat. The origin of the

microwave heating lies in the ability of the electric field

to polarize the charges in materials and the inability of

this polarization to follow extremely rapid reversals of


55 the changing electric field. The field interaction with the

molecular dipoles and charged ions causes rapid rotation

of these molecules or ions. Friction of this motion leads

to a heat evolution and increased temperature in the

materials(14, 15). There are many studies using

microwave irradiation in polymerized of PLA. However,

reports on effects of microwave irradiation in de-

polymerization of PLA are very rare. Hirao et al. studied

the utilization of microwave in de-polymerization of

PLA using hydrolysis and alcoholysis reactions. The

results showed that the microwave irradiation leaded to a

decrease in reaction time, compared to other

conventional heating processes (3, 16, 17).

The aim of this study is to enhance chemical

recycling efficiency of PLA by glycolysis reaction

employing microwave heat source. Effects of glycolysis

condition, which influence to chemical structure and

length of polymer chain, were studied.

2. Experimental

2.1. Materials and chemicals

PLA resin (Nature Work4043D), with weight-

average molecular weight (Mw) and number-average

molecular weight (Mn) of 300,000, and 84,000 g mol-1

,

respectively, was used as a starting material for

glycolysis reaction. All chemicals, including ethylene

glycol (EG), chloroform, and methanol, obtained from

Carlo Erba, were used without further purification.

2.2 Glycolysis reaction

A domestic microwave oven (Samsung, ME81Y

model, 850 watt) was used as a heating source in the

glycolysis reaction. The reaction flask was placed in the

oven, in which the top part of the compartment was

modified to connect with an external condenser. PLA

resin was treated with EG using different PLA: EG

weight ratios (1:2 and 1:0.5) via heterogeneous reaction.

Effects of temperatures and reaction times on structures

and properties of the glycolysed products (GlyPLA) were

studied. At a completion of reaction, GlyPLA was

recovered by dissolving the products in chloroform and

re-precipitating in methanol. The precipitant was further

dried to a constant weight in a vacuum oven at 60°C for

18 hours.


Chemical structures of the products were

characterized by a 500 MHz nuclear magnetic resonance

spectrometer (AV-500, Bruker Biospin) using CDCl3

solvent. Average molecular weight (Mn, Mw), and

degree of polymerization (DP) were characterized by

using a gel permeation chromatography (GPC), Waters

e2695 separations modules, using a combination of

differential viscometer (Viscotek model 270) and

refractive index (Viscotek model 3580) detectors. A

calibration curve was constructed using polystyrene

standards with average molecular weight between 4,490

and 1,112,000 g/mol. The samples were dissolved in

tetrahydrofuran (THF) (2 mg/ml) and filtered using a

nylon 66 membrane (pore size 0.45 μm). A mobile

phase flow rate of 1.0 ml/min was used.


3.1 Chemical structure of GlyPLA

Following Figure1, results from 1H-NMR spectra

of products from all glycolysis reaction conditions

illustrate a similar pattern, in which a selected spectrum

is shown in Figure 1. Four characteristic chemical shifts

are observed at 1.6, 3.8, 4.2 and 5.1 ppm. The signals

located at 1.6 and 5.1 ppm also appear in the original

PLA resin, which are assigned to resonances of methane

(-OCHCH3C=O) at 5.1 ppm and methyl (-

OCHCH3C=O) at 1.6 ppm. The signals at 3.8 ppm

(O=COCH2CH2OH) and 4.2 ppm (O=COCH2CH2 OH)

are only observed in the GlyPLA products, but not in the

starting material. These chemical shifts are associated

with methylene protons of EG. A 1H-NMR spectrum of

GlyPLA also shows another chemical shift at 4.3 ppm,

assigned to (HOCHCH3C=O) end groups. The proposed

mechanism of PLA glycolysis reaction via

transesterification and chemical structures of GlyPLA

products are shown in Figure 2, which agrees with our

previous reports (12, 13).


56

Figure1. 1H NMR spectrum and proposed chemical

structure of GlyPLA (PLA: EG 1:2, 5min, 600W)

Figure2. Mechanism of glycolysis reaction of PLA.

3.2 Effects of glycolysis conditions

The degree of polymerization (DP) and number-

average molecular weight (Mn) of the products obtained

at different reaction conditions are calculated by using

equations 1 and 2 (11). The results from this calculation

are summarized in Table 1.

(1)

(2)

when;

Ha is intensity of GlyPLA at chemical shift 4.3 ppm

Hb is intensity of GlyPLA at chemical shift 5.1 ppm

Hd is intensity of GlyPLA at chemical shift 3.8 ppm

Results on Mn of products show a decreasing

trend for all of GlyPLAs from different reaction

conditions (A1 to A6). A similar trend is also observed

from GPC result, as summarized in Table2. These reflect

that the PLA chains are degraded by successive

transesterification reactions of the hydroxyl group in EG

molecules.

Table1. Results on DP and Mn of GlyPLA products

obtained from glycolysis of PLA at various reaction

conditions.

No Temp.

(⁰C)

Time

(min)

PLA:EG

(%wt) DP Mn

A1 190 5 1:2 69 4,949

A2 190 7 1:2 44 3,183

A3 190 10 1:2 37 2,647

A4 200 7 1:2 52 3,776

A5 200 10 1:2 83 5,959

A6 190 10 1:0.5 72 5,213

The data on Mn of GlyPLA products prepared

from different glycolysis conditions clearly indicate that

the EG:PLA feed ratios, and glycolysis temperature and

time have strong influence on chemical structures and

Mn of GlyPLAs. . When higher EG: PLA feed ratios are

employed, GlyPLA products with shorter chain length

are obtained because the high diol content leads to higher

degree of trans-esterification.

When the reaction temperature is kept at 190°C,

depolymerization of PLLA via transesterification occurs

with formation of product chemical structure following

Figure 2. An increase in the reaction time at this

temperature produces the products with lower molecular

weight.

In contrast, an opposite trend is observed when

the reaction temperature is increased to 200°C. The

results from conditions A4 and A5 show that at this

temperature, products with higher Mn are obtained when

the reaction time increases. This is likely due to re-

combinations of short glycolized oligomers by

transesterification, which occurs at high temperature

conditions. When the same reaction time is applied in

conditions A2 and A4, products with lower Mn is

obtained at lower temperature (190°C). Similar results

+


57 are also observed in conditions A3 and A5. This

phenomenon can be explained by GPC results, as shown

in table 2

From GPC results, GlyPLA obtained at 200°C

shows bi-modal distribution of Mn at 10,112 and 1,160

g/mol, whereas that obtained at 190°C has only one

distribution peak. The lower Mn of A5 product, i.e.,

1,160 g/mol, is a result from glycolysis reaction, as

illustrated in figure 2. In contrast, the distribution mode

at higher Mn occurs from the reaction between end

groups of GlyPLA with another end group of GlyPLA,

as shown in figure 3. Proposed chemical structures of

the products from re-transesterification is also illustrated

in figure3. The similar chemical structure was reported in

our previous study(12).

Table2. Average molecular weight of GlyPLA (obtained

by GPC)

The results indicate that both glycolysis reactions

employing conventional heating or microwave

irradiation can be applied for recycling of PLA via

glycolysis reaction. The chemical structures of products

obtained from both heating sources are similar.However,

the remarkable difference of reaction time is observed.

At the same reaction temperature, GlyPLA of 5,000

g/mol can be achieved within 5 minutes by using

microwave irradiation. On the other hand, 60 minutes is

required in order to produce the products with the same

number average molecular weight(13).

4. Conclusions

Chemical-recycling process of polylactic acid

(PLA) is developed by using a glycolysis reaction.

Microwave irradiation is employed. The results show

that a dramatic decrease in the reaction time is achieved.

Therefore, this highly-effective technique is more

environmental friendly, as it requires lower energy

consumption, when compare with conventional heating

method. Results on effects of glycolysis conditions on

chemical structures of the glycolysed PLA products

indicate that the degree of depolymerization of PLA is

influenced by the reaction temperature, reaction time,

and EG: PLA feed ratios. The insight into the reactions

mechanisms can be applied in the production of

glycolized products for specific applications.

Figure3. Mechanism of GlyPLA reaction when reaction

temperature rising to 200°C

5. Acknowledgements

W.L. is thankful for a support from the TAIST-Tokyo

Tech scholarship program.

6. References

1. Carrasco F, Pagès P, Gámez-Pérez J, Santana OO,

Maspoch ML. Processing of poly(lactic acid):

Characterization of chemical structure, thermal

stability and mechanical properties. Polymer

Degradation and Stability. 2010;95(2):116-25.

2. Santosh Madival a RAa, *, Sher Paul Singh a,

Ramani Narayan. Assessment of the environmental

profile of PLA, PET and PS clamshell containers

using LCA methodology. Journal of Cleaner

Production. 2009:1183–94.

3. Hamad K, Kaseem M, Deri F. Recycling of waste

from polymer materials: An overview of the recent

works. Polymer Degradation and Stability.

2013;98(12):2801-12.

Reaction

condition

Peak

Number

Mn Mw Mw/Mn

A2 1 2,417 3,977 1.6

A5 1 10,112 17,135 1.7

2 1,160 1,247 1.1

+


58 4. Sin LT, Rahmat Abdul R, Rahman Wan AWA.

Overview of Poly(lactic Acid). 2013:11-54.

5. Le Duigou A, Pillin I, Bourmaud A, Davies P, Baley

C. Effect of recycling on mechanical behaviour of

biocompostable flax/poly(l-lactide) composites.

Composites Part A: Applied Science and

Manufacturing. 2008;39(9):1471-8.

6. Emig FSaG. Chemical Recycling of Polymer

Material:Review. Chem Eng Technol.

1998;21(10):778-81.

7. Navnath D. Pingale VSP, S. R. Shukla. Glycolysis

of Postconsumer Polyethylene

Terephthalate Waste. 2009:250-1.

8. Sinha V, Patel MR, Patel JV. Pet Waste

Management by Chemical Recycling: A Review.

Journal of Polymers and the Environment.

2008;18(1):8-25.

9. Song X, Zhang X, Wang H, Liu F, Yu S, Liu S.

Methanolysis of poly(lactic acid) (PLA) catalyzed

by ionic liquids. Polymer Degradation and Stability.

2013;98(12):2760-4.

10. Xi G, Lu M, Sun C. Study on depolymerization of

waste polyethylene terephthalate into monomer of

bis(2-hydroxyethyl terephthalate). Polymer

Degradation and Stability. 2005;87(1):117-20.

11. P. Sukpuang AP, P. Opaprakasit and M.

Opaprakasit. Synthesis and characterrization of

poly(lactic acid- co- ethylene terephthalate) from

glycolysed products. Pure and Applied Chemistry

International Conference2009.

12. Tounthai J, Petchsuk A, Opaprakasit P, Opaprakasit

M. Curable polyester precursors from polylactic acid

glycolyzed products. Polymer Bulletin.

2013;70(8):2223-38.

13. N. Nakruangsri PS, A. Petchsuk, P. Opaprakasit and

M. Opaprakasit. Glycolysis Reaction As a Chemical

Recycling Process for Poly(lactic acid).

International Conference on Green and Sustainable

Innovation; Chiang Rai2009.

14. Hynek Bene JS, Zuzana Walterová , David Rais a.

Recycling of waste poly(ethylene terephthalate) with

castor oil using

microwave heating. Polymer Degradation and Stability.

2013;98:2232-43.

15. Pelle Lidstrom JT, Bernard Wathey and Microwave

assisted organic synthesis—a review. Tetrahedob.

2001;57:9225-83.

16. Hirao K, Nakatsuchi Y, Ohara H. Alcoholysis of

Poly(l-lactic acid) under microwave irradiation.

Polymer Degradation and Stability. 2010;95(6):925-

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Hydrolysis of poly(l-lactic acid) using microwave

irradiation. Polymer Degradation and Stability.

2010;95(1):86-8.

SESSION 2

Advances in Polymer Characterization


60 KN-CHAR-1

Preparation and Properties of Natural Rubber with

Organic-Inorganic Nanomatrix Structure

Seiichi Kawahara

Department of Materials Science and Techonology, Faculty of Engineering,

Nagaoka University of Technology, 1603-1, Kamitomioka, Nagaoka,

Niigata 940-2188, Japan.

Abstract Natural rubber with filler nanomatrix structure was prepared

by forming chemical linkages between natural rubber particles and

filler nano-particles. The filler nanomatrix structure was formed by

graft-copolymerization of vinyltriethoxysilane (VTES) onto natural

rubber particles in the latex stage followed by casting of the latex to

prepare an as-cast film. The silica nano-particles were produced

during the graft-copolymerization through hydrolysis and

condensation, i.e. sol-gel reaction; as such, they linked to the natural

rubber particles. The nanomatrix structure was observed by

transmission electron microscopy, in which the natural rubber

nanomatrix. Tensile properties were significantly improved by

forming filler nanomatrix structure. The loss modulus and loss tangent

of the natural rubber with the filler nanomatrix structure were almost

independent of deformation frequency in the rubbery plateau region,

which was explained to be due to both the energetic elasticity and

entropic elasticity characteristics of the nanomatrix structure.

Seiichi Kawahara

Affiliation: Department of Materials

Science and Technology, Facluty of

Engineering, Nagaoka University of

Technology

Education:

1988 Bachelor, Faculty of Technology,

Tokyo University of Agriculture and

Technology

1992 Doctor, Graduate School of

Engineering, Tokyo University of

Agriculture and Technology

Job:

1992 Research Associate, Tokyo

University of Agriculture and Technology

1998 Associate Professor, Nagaoka

University of Technology

(1996-1997 Visiting Scientist, The

University of Akron)

Awards:

2000 Best Paper Award, The Society of

Rubber Industry, Japan

2004 The Award of Distinguished

Research Work, The Society of Rheology,

Japan

2006 Best Paper Award, The Society of

Rubber Industry, Japan

2010 Malaysian Rubber Board Service

Award

2010 Best Paper Award, Journal of

Rubber Research

2011 Wiley Award, The Society of

Polymer Science, Japan

2012 Best Presentation Award, The

Society of Rubber Science and

Technology, Japan

2014 Sparks-Thomas Award, Rubber

Division, American Chemical Society

2015 Science and Technology Award, The

Society of Rubber Science and

Technology, Japan

Keywords on Research work

Rubber, Elastomer, Rheology, NMR,

Characterization, Mechanical Properties,

Ion Conductivity, Nano-phase separated

structure


61

KN-CHAR-2

Chemically Controlled Self-assembly of Gold Nanoparticles by Site-selective Protein

Immobilization: A Model for Antimalarial Drug Screening

Palangpon Kongsaeree, Vasujin Numphud, Pattarapol Khongsuk

Department of Chemistry and Center for Excellence in Protein Structure and Function, Faculty of Science,

Mahidol University, Bangkok 10400, Thailand


Methods to efficiently control the activity of proteins in cells are

very useful to explore a variety of biological processes. Chemically

induced dimerization (CID) has proven used to be a powerful tool for

modulating protein-protein interactions by bringing two proteins of

interest into close proximity. Herein, a novel CID system with

synthetic ligands of methotrexate, a tight-binding ligand of

dihydrofolate reductase, was developed to investigate intermolecular

interactions on the surface of gold nanoparticles. Gold nanoparticles

can be used in various sensitive analytical purposes due to its high

molar absorption coefficient. In this work, dihydrofolate reductase-

functionalized gold nanoparticles (AuNP-DHFR) were prepared by

site-specific immobilization. Upon the addition of methotrexate-

based CID ligands, self-assembly of AuNP-DHFR was specifically

induced through the ligand-protein interactions, resulting in a clear

optical change of the colloidal color from red to violet-blue. This

surface plasmonic resonance (SPR) change was clearly observed by

naked eye and this phenomenon was concentration dependent of the

CID ligands. The aggregation of the AuNP-DHFR nanoparticles was

analyzed by transmission electron micrograph, zeta potential,

hydrodynamic radius, etc. Ultimately, a simple label-free

antimalarial assay based on this CID-SPR platform could be further

developed.

Palangpon Kongsaeree

Associate Professor Education

Ph.D., Cornell University , USA, 1998

Research Area

Chemical biology, X-ray crystallography,

Nanomaterials

Awards

The Royal Thai Government

scholarship under the Development and

Promotion of Science and Technology

Talent Project (DPST), 1985-1998;

Research assistance, Department of

Chemistry, Cornell University, 1992-

1997;

Young Scientist Award from

Foundation for Promotion of Science

and Technology under the Patronage of

H. M. the King, 2003

SESSION 3

Polymer Composites and Nanocomposites


63

KN-COMP-1

Interphase Transfer of Nanoparticles between Immiscible Polymer Blends

Masayuki Yamaguchi

School of Materials Science, Japan Advanced Institute of Science and Technology

Abstract

The nanofiller transfer technique was demonstrated, which can

be applicable to novel polymeric materials having nanoparticles on

surface and/or phase-separated polymer blends with uneven distribution

of fillers. It was found that nanoparticles immigrate from one polymer

to another owing to the difference in the interfacial tension. For

example, multi-walled carbon nanotubes (CNTs), which prefer

polycarbonate (PC) to polypropylene (PP), moved from PP to PC in the

molten state of the polymers. Since the transfer process requires

Brownian motion, more CNTs were transferred at high temperature and

for a long exposure time of annealing. The laminated sheets were

separated without any difficulty because of thin interfacial thickness

between the immiscible polymers. Consequently, a PC sheet having

CNTs only in the surface region was prepared. A similar phenomenon

was detected for nanoparticles of silica in the laminated sheets

composed of butadiene rubber (BR) and styrene-butadiene rubber

(SBR), in which silica particles moved from SBR to BR, and showed

nucleating ability for BR crystallization.

Masayuki Yamaguchi

Japan Advanced Institute of Science and

Technology

School of Materials Science

Professor, Doctor of Engineering

Director, Research Center for Highly

Environmental and Recyclable Polymers

Vice President of Japan Society of Polymer

Processing

Editor in chief, Journal of Society of

Polymer Processing

B.D. in Kyoto Univ. in 1987

M.S. in Kyoto Univ. in 1989 Tosoh Corporation from 1989 to 2005

Doctor degree in Kyoto Univ. in 1999 under

late Prof. Toshiro Masuda The Polymer Processing Institute in New

Jersey (USA) as a visiting scientist from 2000

to 2002. 2005, Associate professor in JAIST

2009, Professor in JAIST

Major research areas are polymer rheology and processing


64

KN-COMP-2

Hybrid Porous Polymers Derived from Octavinylsilsesquioxane

Hongzhi Liu

Department of Chemistry and Chemical Engineering, Shandong University, Jinan, China, 250100.

Abstract

Porous polymers have elicited considerable interest due to their extensive potential applications, such

as gas storage and separation, heterogeneous catalysis and sensors. The “bottom-up” topology combination has

proved to be an efficient strategy to construct porous materials.

Rational selection of suitable rigid building

blocks with different geometry will help to construct polymer networks with novel topology structures and

properties. Polyhedral ologomeric silsesquioxanes (POSS) are hybrid molecules with three–dimensional

nanometer–sized inorganic–organic hybrid structures and a general formula of (RSiO1.5)n (n = 6, 8, 10, 12).

Rigid cage make POSS become an ideal building block to construct novel inorganic-organic hybrid porous

materials. Some hybrid porous materials based on POSS have been synthesized by the reaction of other building

blocks with different geometries. Starting from cubic octavinylsilsesquioxane (OVS), we have successfully

prepared some hybrid porous materials by the combination with planar molecules, tetrahedral molecules or

polymer via the Friedel–Crafts reaction or Heck coupling reaction (Scheme).


65

KN-COMP-3

Natural Fiber Reinforced Rubber: Recent Advances toward High Performance Rubber

Matrix Composites using Pineapple Leaf Fiber

Taweechai Amornsakchai

Polymer Science and Technology Program, Department of Chemistry and

Center of Excellence for Innovation in Chemistry and

Center of Sustainable Energy and Green Materials,

Faculty of Science, Mahidol University,

Phuttamonthon 4 Road, Salaya, Phuttamonthon,

Nakhon Pathom 73170, Thailand

Abstract

Short fiber provides uniquely characteristic reinforcement to

rubber, i.e. a sharp rise in stress at relatively low strains, which is not

obtainable from particulate fillers. It can be incorporated directly

along with other additives into the rubber compound and further

processed using standard rubber-processing equipment. This makes it

more convenient and economical. Now the use of natural fiber has

attracted a great deal of attention due to environmental and

sustainability issues. In addition, some natural fibers have greater

modulus than fibers that are normally used in the rubber industry.

Thus by using natural fiber as the reinforcement, greener and higher

performance rubber composite should be obtained. Since rubber

matrix has extremely low stiffness and very high extensibility

compared to the reinforcing fiber, there can be various practical

problems in obtaining good rubber composites. In this presentation,

various approaches to development of high performance rubber

composite from short and fine pineapple leaf fiber (PALF) will be

presented. These include the use of microfiber, hybridization with

some particulate fillers and changing the vulcanization systems.

Taweechai Amornsakchai

Department of Chemistry, Faculty of Science,

Mahidol University, Phuttamonthon 4 Road,

Salaya, Nakhon Pathom 73170 Thailand

Educations 1989 B.Sc. (Industrial Chemistry) King

Mongut's Institute of Technology (First class

honor) Ladkrabang

1994 Ph.D. (Polymer Physics) University of Leeds

Working experiences 1994 Lecturer Dept. Chemistry, Faculty of

Science, Mahidol University.

1998 Assistant Professor Dept. Chemistry,

Faculty of Science, Mahidol University. 2003 Associate Professor Dept. Chemistry,

Faculty of Science, Mahidol University.

Research Experiences 1. Ph. D. training on the thesis entitled “A

study of the tensile properties and structure of polyethylene fibres” under the supervision of

Professor I. M. Ward.

2. JSPS Visiting scientist to Department of Chemistry, Gunma University, Japan for a

collaborated work with Associate Professor H.

Kubota on “Photo-grafting of some vinyl monomers on highly oriented polyethylene”.

3. EPSRC visiting fellowship to polymer physics group, JJ Thomson Physical

Laboratory, University of Reading, UK,

working on ‘structure and properties of new generation polyethylenes’ with Professors D. C.

Bassett (Reading) and I. M. Ward (Leeds)

4. JSPS Visiting scientist to Institute for Chemical Research, Kyoto University, Japan

for a collaborated work with Professor

Fumitaka Horii on “Solid state 13C NMR investigation of highly drawn polyethylenes”.


66

KN-COMP-4 Performance of Aramid Fibre in Rubber Compounds

Jutarat Phanmai

Chemical Innovation Co., Ltd., 18 Soi Ramkhamhaeng 30 (Ban Rao), Huamark, Bangkapi, Bangkok 10240

Phone +66 2375 5197, Fax +662374 6503, E-mail: [email protected]

Abstract

Aramid fibres are used in many applications such as automotive

components (brake pads, clutch, gasket and tires), antiballistic, and

fiber optic cable. The benefits of them are high strength, low

elongation, good heat resistance property, and light weight. The aim of

present work is to observe physical, mechanical, and rheological

properties of NR and NBR compounds containing aramid fibres. It was

found that modulus and tear resistance increased while tensile strength

dropped at high loading. Tan δ at 50 °C didn’t show any improvement

while that at 100 °C exhibits increment of the performance.

Keywords: Aramid fibres, High strength, Heat resistance, NR

compounds and NBR compounds.

Jutarat Phanmai

Education

1991-1994: Master’s Degree (Master of

Science in Polymer Science)

Mahidol University, Bangkok, Thailand Thesis: A Study of Toughed Epoxy

Adhesives with Epoxidised Liquid Natural

Rubber (ELNR) 1987-1990: Bachelor’s Degree (Bachelor of

Science in Chemistry)

Mahidol University, Bangkok, Thailand Senior Project: Polymerization of Stars

Polymer

Work Experiences

Jul. 2001-Present: Chemical Innovation

Co., Ltd. - Trading Company, Bangkok Vice President (Marketing Trading,

Jan. 2011-Present) Marketing Manager (Feb. 2006-Dec.

2010)

R&D Manager (Jul. 2004-Jan. 2006) Technical Manager (Jul. 2001-Feb.

2004)

Sept. 2000 -Jun. 2001: Siam Nippon Steel Pipe Manufacturer Steel pipe Manufacturer,

Bangkok

Section Chief (Marketing) Jan.1995 -Feb. 1999: PI Industry Ltd.-

Rubber Compounder, Bangkok

Research and Development Assistance Manager

Jul. 1994 -Dec. 1994 Chemical Innovation

Co., Ltd.-Trading Company, Bangkok Research and Development Assistance

Manager

Research and development areas

Cogents in P/O Cure

Metallic Cogent in Sulfur Cure

UV Curing Technology for Coating

Application

Rubbers & Chemicals Compounding Technology Coupling agents in Plastics Composite

Coupling agents in Biodegradable Polymer

Adhesion Promoters in Rubber

Compounds Aramid fibre in Rubber Compound


67

COMPO-01

Influence of Pristine Clay Incorporation on Strain-Induced Crystallization of Natural Rubber

Abdulhakim Masa,1 Hiromu Saito,

2 Tadamoto Sakai,

3 Azizon Kaesaman

1 and Natinee Lopattananon

1*

1 Department of Rubber Technology and Polymer Science, Faculty of Science and Technology,

Prince of Songkla University, Pattani 94000 Thailand. 2 Department of Organic & Polymer Materials Chemistry, Tokyo University of Agriculture and Technology,

Koganei-shi, Tokyo 184-8588, Japan. 3 Shizuoka University, 3-3-6 Shibaura, Minato, Tokyo, 108-0023, Japan.

Abstract

Strain-induced crystallization behavior of cured natural rubber (NR) nanocomposite filled with 5 phr clay was

investigated by means of tensile test and wide X-ray diffraction (WAXD) measurements. The dispersion of clay was

revealed by using X-ray diffraction (XRD) analysis. The uncured NR/clay nanocomposite and cured pure NR samples

were also prepared for comparison. It was found that the clay was intercalated by NR chains. Addition of clay

drastically affected the tensile behavior by changing the pattern of stress-strain curves. By comparing with the cured

neat NR, the stress upturn, corresponding to strain-induced crystallization, was developed at lower strain with the

presence of clay. WAXD results also showed that the onset strain for strain-induced crystallization of NR was

accelerated and the crystallinity was enhanced due to collaborative strain-induced crystallization by clay and

crosslinking point.

Keywords: natural rubber / clay / nanocomposite / strain-induced crystallization / WAXD

1. Introduction

Natural rubber (NR) is a renewable material, which

has been used to manufacture a wide range of industrial

products, i.e., automobile tires, vibration isolators,

medical gloves, etc. The interest in NR is basically

attributed to its ability to crystallize upon deformation

(strain-induced crystallization), which gives natural

rubber high tensile strength, resistance to cutting, tearing,

and abrasion [1], and stops crack propagation at ultimate

strain [2]. The strain-induced crystallization mechanism

of cured pure NR was explained on the basis of

inhomogeneous distribution of the network chain length,

i.e., highly crosslinked network region would favor the

molecular orientation of chains to induce crystallization

[3]. In NR-based nanocomposites, ability of strain-

induced crystallization was accelerated and increased in

the presence of clay [4, 5]. However, the change in

strain-induced crystallization behavior and the increase

of crystallization was not discussed. Therefore, in this

article the structure-property relationship of the cured

NR nanocomposite, uncured NR nanocomposite and neat

cured NR was examined in order to understand the

influence of clay on the development of strain-induced

crystallization process and the change in crystallization

behavior of the NR.

2. Experimental

Suspension of clay (Kunipia-F®, Kunimine

Industries Co., Ltd., Japan) in water was added into NR

latex (Yala Latex Co., Ltd., Thailand). The NR and clay

were mixed thoroughly under vigorous stirring at

ambient temperature for 30 min, and dried at 50C for 3

days to obtain uncured NR/clay nanocomposite. The

NR/clay nanocomposite was then compounded with

phenolic crosslinking agent (HRJ-10518) (Schenectady

International Inc., USA) and catalyst (SnCl22H2O)

(Carlo Erba Reagent, France) in a mixing chamber of a

miniature mixing machine (IMC-18D7, Imoto

Machinery Co., Ltd., Japan) at rotor speed of 140 rpm

and temperature of 100 °C for 20 min. The sample was

later melt pressed in a small hot-press machine (Imoto

Machinery Co., Ltd., Japan) at 180°C to obtain the cured

NR/clay nanocomposite film with a thickness of 1 mm.

In this study, the clay content was fixed at 5 phr. The

uncured and cured pure NR specimens were also

prepared by using the same procedure outlined above for

the preparation of the uncured and cured NR

nanocomposites, respectively. The samples were

characterized by using tensile test and wide angle X-ray

diffraction (WAXD) measurements.


68


Dispersion of clay in the uncured and cured NR/clay

nanocomposites is shown in Figure 1. It can be seen that

the clay layers were intercalated by the rubber chains,

indicated by the shift of reflection angle (2θ) at 7.04° for

the pure clay (Figure 1(A)) to lower ones for the NR/clay

nanocomposite (Figure 1(B)). Thus, the nanocomposites

of NR and clay with intercalated clay structure were

obtained.

Figure 2 shows stress-strain curves of the uncured

and cured pure NR, uncured and cured NR/clay

nanocomposites couple with the 2D WAXD images at

the selected strain. It can be seen from the stress-strain

curves that the strain, at which the stress was rapidly

upturned, indicating strain-induced crystallization, was

decreased to lower strain level by inclusion of clay in

both uncured and cured NR nanocomposite samples

when compared with their counterparts. For example,

the strain-induced crystallization for the NR

nanocomposites, determined from the change of

differential value in the plot of stress versus strain, was

found at strain of 270%, while those of their

corresponding pure NR was found at strain higher than

400%. This behavior could be explained by acceleration

effect of clay on crystallization behavior [6].

Considering the 2D WAXD images of

nanocomposite samples at 200% strain (Figure 2), it is

interesting to note that the highly oriented crystalline

reflection spots, corresponding to (200), (201) and (120)

planes, were found in both the uncured and cured NR

samples containing clay. However, these reflection spots

were not observed in the uncured and cured pure NR at

300% and 200% strain, respectively, suggesting that the

onset strain for strain induced crystallization of NR was

decreased by adding clay. These results confirmed that

the strain-induced crystallization process of the NR was

accelerated by the presence of clay.

Figure 1 XRD patterns for (A) pure clay and (B)

uncured and cured NR/clay nanocomposites

Based on the 2D WAXD image, the crystallinity

(Xc) of the stretched NR could be estimated by measuring

the WAXD intensity of the diffraction peaks,

corresponding to the (200) and (120) planes (Fig. 3).

The diffraction intensity in the equator direction was

normalized and azimuthally integrated in the range of the

azimuthal angles from 80º to 100º as shown in Fig. 3.

The area of the crystalline diffraction peaks assigned to

the (200) and (120) planes and the area of amorphous

halo were fitted by using Origin®9.1 software. The Xc

was calculated using the following equation (Eq. 1) [7];

%100

ac

cc

AA

AX

(1)

where Ac is the area of crystalline diffraction peaks,

corresponding to the (200) and (120) planes, and Aa is

the area of the amorphous halo. The data of crystallinity

(%) for various samples is shown in Figure 4.


69

Figure 2 Stress-strain curves of the uncured and cured

pure NR, uncured and cured NR/clay nanocomposites

coupled with their corresponding 2D WAXD images at

selected strain

It can be seen that the crystallinity of different

samples increased with increasing strain due to

molecular alignment of NR chain in the stretching

direction [8]. The crystallinity of the uncured and cured

NR/clay was initiated at lower strain, when the cured NR

without clay was compared. This indicated that the

crystallization induced from clay dispersion occurred

earlier than the one induced from crosslinking points.

This would be explained by the fact that the clays could

interacted with NR via adsorption of NR chains onto

their surfaces and/or intercalation of NR chains into their

galleries as discussed earlier, and the interaction between

clay and NR would served as additional crosslinks.

When the NR was deformed under tensile mode, a

number of NR chains adjacent to the clay surfaces were

probably over-strained, which led to crystallization at

around the clay particles [9].

Figure 3 Typical WAXD image (left) and WAXD

profile (right) as a function diffraction angle of stretched

cured NR/clay nanocomposite at 400%. Inset shows

integration limits from 80º to 100º.

From Figure 4, it is also found that the crystallinity

of the cured NR/clay nanocomposite was highest at any

given strains. This was because of collaborative

crystallization induced by clay and crosslinking points.

Figure 4 Evolution of crystallinity as a function of strain

for different samples

4. Conclusions

Influence of clay addition on strain-induced

crystallization of NR was studied. XRD observation

revealed that the clay was mainly intercalated by rubber

chains. Addition of clay into the NR drastically changed

the stress-strain behavior by accelerating the strain-

induced crystallization process, i.e., onset of


70 crystallization, of NR. It was also found that the

crystallinity of NR was greatest when the clay was

dispersed in the crosslinked network of NR due to

collaborative crystallization induced by clay and

crosslinking points.

Acknowledgements

Thailand Research Fund (TRF) through the Royal

Golden Jubilee Ph.D. Program (Grant No.

PHD/0052/2554), and Prince of Songkla Graduate

Studies Grant, Prince of Songkla University are

gratefully acknowledged.

References

[1] Ciullo, P. A. and Hewitt, N. The Rubber Formulary,

Noyes Publications, New York, (1999)

[2] Weng, G., Huang, G., Qu, L., Nie, Y., Wu, J., J.

Phys. Chem. B, 114, 7179-7188 (2010)

[3] Toki, S., Sics, I., Ran, S., Liu, L., Hsiao, B.S.,

Murakami, S., Senoo, K., Kohjiya, S.,

Macromolecules, 35, 6578-6584, (2002)

[4] Carretero-Gonzalez, J., Retsos, H., Verdejo, R.,

Toki, S., Hsiao, B.S., Giannelis, E.P., Lopez-

Manchado, M.A., Macromolecules, 41, 6763-6772,

(2008).

[5] Carretero-Gonzalez, J., Verdejo, R., Toki, S., Hsiao,

B.S., Giannelis, E.P., Lopez-Manchado, M. A.,

Macromolecules, 41, 2295-2298, (2008)

[6] Ray, S.S. Clay-Containing Polymer

Nanocomposites: From Fundamentals to Real

Applications, Elseview, Amersterdam. (2013)

[7] Hernandez, M., Lopez-Manchado, M.A., Sanz, A.,

Nogales, A., Ezquerra, T.A., Macromolecules, 44,

6574-6580, (2011)

[8] Hernandez, M., Sanz, A., Nogales, A., Ezquerra,

T.A., Lopez-Manchado, M.A., Macromolecules, 46,

3176-3182, (2013)

[9] Rault, J., Marchal, J., Judeinstein, P., Albouy, P.A.,

Macromolecules. 39, 8356-8368, (2006)

http://www.google.co.jp/search?tbo=p&tbm=bks&q=inauthor:%22Peter+A.+Ciullo%22&source=gbs_metadata_r&cad=3

http://www.google.co.jp/search?tbo=p&tbm=bks&q=inauthor:%22Norman+Hewitt%22&source=gbs_metadata_r&cad=3

http://www.researchgate.net/researcher/55235977_J_Carretero-Gonzalez

http://www.researchgate.net/researcher/79502787_H_Retsos

http://www.researchgate.net/researcher/8857640_R_Verdejo

http://www.researchgate.net/researcher/67688047_S_Toki

http://www.researchgate.net/researcher/9140645_B_S_Hsiao

http://www.researchgate.net/researcher/16174076_E_P_Giannelis

http://www.researchgate.net/researcher/71430493_M_A_Lopez-Manchado


http://www.researchgate.net/researcher/55235977_J_Carretero-Gonzalez

http://www.researchgate.net/researcher/8857640_R_Verdejo

http://www.researchgate.net/researcher/67688047_S_Toki

http://www.researchgate.net/researcher/9140645_B_S_Hsiao

http://www.researchgate.net/researcher/16174076_E_P_Giannelis



71

COMPO-02

Effects of Organoclay Types on Morphological and Mechanical Properties of

Polyoxymethylene/Polypropylene Blends

Nipawan Yasumlee, Sirirat Wacharawichanant*

Department of Chemical Engineering, Faculty of Engineering and Industrial Technology, Silpakorn University

NakhonPathom 73000, Thailand

Abstract

This work studied the effects of organoclay surface modified with three different types of dimethyl dialkyl (C14-

C18) amine (organoclay-DDA), trimethylstearyl ammonium (organoclay-TSA) and methyl dihydroxyethyl

hydrogenated tallow ammonium (organoclay-DHA) on mechanical and morphological properties of polyoxymethylene

(POM)/polypropylene (PP) blends. Scanning electron microscopy (SEM) results revealed the size of dispersed PP

phase decreased with increasing organoclay content. X-ray diffraction (XRD) results showed the formation of

exfoliated structure for POM/PP/organoclay (80/20/5phr) at all types of organoclay. Incorporation of organoclay

improved the Young’s modulus but dropped the percent strain at break of the blends. The POM/PP blends containing

organoclay-DHA revealed the highest tensile strength due to better exfoliated structure in the POM/PP/organoclay-

DHA

Keywords: Polyoxymethylene; Polypropylene; Organoclay; Morphology; Mechanical properties

1. Introduction

The polymer blends have been studied for many

years, develop new polymeric materials. However, most

polymer blends are immiscible and phase separation

leading to poor mechanical properties. Several articles

reported that the improvement miscible between polymer

blends by reduce interfacial tension of two phases, which

results in improved material properties [1].

Recently, several research groups have studied

about the application of organoclay in the polymer

blends.

By the incorporation of organoclay in blends, the

morphology can be affected by several factorssuch as the

composition of the dispersed phase, viscosity ratio and

the elasticity of the each phases [2]. For example, the

study of effect of organoclay in polypropylene

(PP)/polystyrene (PS) blends found that the addition

small amounts (2-5 wt%) of an organoclay to the blends

can reduce the dispersedsize of PS phase and improve

interfacial adhesion between two polymers.

From SEM study, it was found that the morphology of

the blends displays very fine particle size and good

interfacial adhesion between phases after of organoclay

addition. This indicates that the improvement elongation

at break of the PS/PP blends due to organoclay plays the

role of interfacial active agent.Sung et al. [3] studied the

effect of organoclay on the morphology of

poly(acrylonitrile-butadiene-styrene) (ABS) and PP.

They found that the dispersed PP phase changed from a

sphere to elongated structure, when the organoclay

content was increased due to the viscosity ratios of the

PP and ABS/clay decreased. Transmission electron

microscopy (TEM) and x-ray diffraction (XRD) results

found that the most of the organoclay existed in ABS

phase because of the good affinity between the ABS and

clay. Anup et al. [4] studied the morphology of PP/high-

density polyethylene (HDPE) blends by the addition of

organoclay. Form SEM results found that the addition of

the organoclay reduced the dispersed HDPE phase

because of the barrier effect of organoclay in a matrix,

including decrease in the viscosity ratio of HDPE and

PP. Sumana Mallick et al. [5] showed that the presence

organoclay improved the miscibility between Nylon6

and HDPE in Nylon6/HDPE blends. From XRD results

revealed that the organoclay were exfoliated in the

nylon6 and intercalated in HDPE matrix, indicating to

improve in tensile properties.

The incorporation of clay modification on the

mechanical properties of polyamide 6/PP showed that the

increase ofstiffness and storage modulus in the blends

[6]. Bitinis et al. [7] reported that the addition of the

organoclay in poly (lactic acid) (PLA)/ natural rubber

(NR) blends. They found that the elongation at break,


72 stiffness and modulus were improved when adding the

organoclay. The location of organoclay was found at

both the interface and PLA phases, it showed a similar

reinforcing effect.

Polyoxymethylene(POM) is the excellent

engineering thermoplastic. It is appreciated for their

good mechanical properties, elasticity fatigue and wear

resistances, good dimensional stability and low

coefficient of friction but lower elongation. On the other

hand, PP is a widely used commodity polymer because

its ease of processing, low cost, resistant to thermal and

toughness. The blends of POM and PP can combine the

easy processibility POM with good properties of PP.

These blends can combine the integral property between

POM and PP, provided the blend is miscible [8-9].

The objectives of this work studied the effects of

different organoclay on mechanical and morphological

properties of POM/PP blend at 80/20 (w/w) containing

organoclay 3 and 5 phr. The compatibility of POM/PP

blends in presence of organoclay was investigated by

SEM. The mechanical properties of polymer blends

without and with organoclay were investigated by tensile

test.


2.1 Materials

POM with melt flow rate of 8.9 g/10 min was

produced by Poly plastics Company. PP with melt flow

rate of 4 g/10 min was produced by HMC Polymers

Company. Three different types of surface modified with

dimethyl dialkyl (C14-C18) amine (organoclay-DDA),

trimethylstearyl ammonium (organoclay-TSA) and

methyl dihydroxyethyl hydrogenated tallow ammonium

(organoclay-DHA) were produced by Sigma Aldrich

Company.

2.2 Sample preparation

All types of polymers and organoclay were dried

before blending, POM was dried in an oven at 110oC for

3 h and organoclay was dried at 80oC for 24 h. POM/PP

and POM/PP/organoclay blends were prepared by melt

blending in an internal mixer at 200oC and a rotor speed

of 50 rpm for 20 min. The organoclay content was 3 and

5 phr.The standard dumbbell samples for tensile test

were prepared by a compression molding at 180oC for 15

min.

2.3 Sample characterization

Tensile test was carried out at a crosshead speed of

50.8 mm/min with a universal tensile testing machine

(LR 50KLloyd instruments, England) according to

ASTM D 638. Each value obtained represented the

average of five samples.

The morphology of the blends without and with

organoclay was studied by SEM. The samples were

carefully broken under a liquid nitrogen atmosphere. The

fracture surfaces of POM/PP blends and

POM/PP/organoclay composites were observed by SEM

(Maxim 2000S, Cam Scan Analytical, England). All

specimens were coated with gold before SEM study.

XRD was carried out by Bruker D8-Advance. The

diffractograms were obtained at the scattering angles

from 1ᵒ to 10ᵒ operated at 40 kV and 30 mA. The basal

spacing of organoclay was determined from the

peakposition (d001 reflection) in the XRD diffractograms

according to Bragg equation.


3.1 Mechanical properties

The Young’s modulus of POM/PP 80/20 (w/w)

without and with three types of organoclay is presented

in Figure 1. The results revealed that the Young’s

modulus of POM/PP/organoclay increased with

increasing the amount of organoclay (3 and 5 phr) at all

types of organoclay when respect POM/PP blends. The

improvement in Young’s modulus may be due to the

reinforcement effect of the organoclay and the

constraining effect of silicate layers on molecular motion

of polymer molecular chains [6]. At 5 phr of organoclay,

the organoclay-DDA showed higher the Young’s

modulus. This may be due to the longer alkyl chain of

organoclay-DDA.


73

Figure 1 Young’s modulus of POM/PP/organoclay

composites.

Figure 2 Tensile strength of POM/PP/organoclay

composites.

Figure 3 Stress at break of POM/PP/organoclay

composites.

Figure 2 represents the tensile strength of

POM/PP blends with and without organoclay, it was

shown that the tensile strength slightly decreased with

added organoclay. Sumana Mallick [5] reported that the

higher content of organoclay (at and above 3 phr), tensile

strength of the blends decreased and approached to that

of the blends when added 5 phr of organoclay may be

due to the agglomeration or poor dispersion of

organoclay in the matrix phases. However at 3 phr of

organoclay-DHA can be improved the tensile strength

and stress at break when compared to the blend, as

shown in Figure 2 and Figure 3, respectively.

The percent strain at break of the blends

decreased with increasing organoclay content as shown

in Figure 4. The incorporation of organoclay at 3 phr was

higher than the blends with organoclay 5 phr. This

suggests that the POM/PP blends became more brittle,

the reduction in ductility was attributed to the

constraining mobility of polymer chains in the addition

of the organoclay [5].

Figure 4 Percent strain at break of POM/PP/organoclay

composites.

3.2 Morphology

The SEM micrographs of the fractured surface of

pure POM, POM/PP blends with and without organoclay

at three different types of organoclay shown in Figure 5.

The morphology of pure POM (Figure 5a)

displayed the fractured surface of pure POM was quite

smooth which indicated the brittle properties of POM.

Figure 5(b) revealed the SEM image of POM/PP blends

(80/20) that showed the phase separation, which the

dispersed PP phase as spherical in POM matrix. It was

observed that the voids surrounding the dispersed PP

phases indicated weak interfacial adhesion between the

POM and PP phases, the spherical morphology is

expected because of the minimization of the interfacial

area [4]. The changes of phase morphology of the blends

as a function of organoclay content are investigated by

SEM. The morphology of POM/PP blends with the

addition of 3 and 5 phr of three organoclay types are


74 shown in Figure 5 (c-h). The dispersed PP phase size was

decreased with the increase of organoclay content. Form

SEM images, it can be seen the organoclay-DHA is more

finely dispersed in POM/PP blends than the organoclay-

TSA and organoclay-DDA. This could be attributed to

better exfoliated structure in POM/PP/organoclay-DHA

(3 phr) as shown in XRD result.

3.3 X-Ray Diffraction

The XRD patterns of the organoclay,

POM/organoclay (5phr) and PP/organoclay (5phr) shown

in Figure 6. The characteristic peak of organoclay-TSA,

organoclay-DDA and organoclay-DHA was observed at

2θ=3.67°, 3.49° and 4.80°, respectively, corresponding to

basal spacing of 2.40, 2.53 and 1.80 nm. The basal

spacing was calculated from XRD patterns and

summarized in Table 1. The shifting of the organoclay

peak position to lower 2θ in POM/organoclay (5 phr) of

organoclay indicated the intercalation of POM chains

inside the organoclay galleries, it is suggested that the

affinity between POM and organoclay. For the

PP/organoclay-TSA, PP/organoclay-DDA and

PP/organoclay-DHA (5 phr) the peak observed at about

2θ=3.82°, 3.80° and 4.89°, respectively, corresponding to

basal spacing of 2.31, 2.32 and 1.80 nm. The decrease of

the d-spacing in the PP/organoclay may be due to the

nonpolarity of the PP. It is indicated that the organoclay

in POM phase may be as a barrier for coalescence of the

dispersed PP phase [2]. The decrease in 2θ of

organoclay-TSA and organoclay-DHA in

POM/PP/organoclay (80/20/3phr) revealed an increase in

the d-spacing of organoclay-TSA and organoclay-DDA

in the presence of POM from 2.40 to 3.77 nm and 2.53 to

3.26 nm, respectively as shown in Figure 7. This

indicates that the polymer molecular chains have entered

the gallery of organoclay, and this blend with organoclay

has an intercalated structure. In contrast, the XRD

patterns of POM/PP/organoclay-DHA (80/20/3phr) did

not show any diffraction peaks, indicating that

organoclay-DHA platelets were completely exfoliated in

the blend with organoclay [10]. In POM/PP/organoclay

(80/20/5phr) at all types of organoclay, the characteristic

peak of organoclay disappears, indicating that the

structure is potentially highly exfoliated.

Figure 5 Scanning electron micrographsof pure POM and

POM/PP (80/20 w/w) blends with organoclay at 3 and 5

phr.

Figure 6 XRD patterns of organoclay, POM/organoclay

and PP/organoclay.

(a) (b)

(c) (d)

(e) (f)

(g) (h)

20 µm 20 µm

20 µm 20 µm

20 µm 20 µm

20 µm 20 µm

POM POM/PP 80/20

Organoclay-TSA 3phr Organoclay-TSA 5phr

Organoclay-DDA 3phr Organoclay-TSA 5phr

Organoclay-DHA 3phr Organoclay-DHA 5phr

(a) (b) POM POM/PP 80/20


75

Figure 7 XRD patterns of POM/PP (80/20 w/w) blends

with and without organoclay.

Table 1 Basal spacing (d001) of organoclay.

Sample 2θ (°) D001 (nm)

Organoclay-TSA

Organoclay-DDA

Organoclay-DHA

POM/Organoclay-TSA

POM/Organoclay-DDA

POM/Organoclay-DHA

PP/Organoclay-TSA

PP/Organoclay-DDA

PP/Organoclay-DHA

POM/PP/Organoclay-TSA

80/20/3

80/20/5

POM/PP/Organoclay-DDA

80/20/3

80/20/5

POM/PP/Organoclay-DHA

80/20/3

80/20/5

3.67

3.49

4.80

-

2.93

2.30

3.82

3.80

4.89

2.34

-

2.71

-

-

-

2.40

2.53

1.80

-

3.01

3.84

2.31

2.32

1.80

3.77

-

3.26

-

-

-

4. Conclusions

This work focused the effects of types of

organoclay surface modified on morphological and

mechanical properties in the POM/PP blends. The

morphology of the blend changed with added

organoclay, and the dispersed PP phase size was

reduced with increasing organoclay content. The addition

of organoclay-DHA displayed the mophology is more

finely dispersion of organoclay than organoclay-DDA

and organoclay-TSA.

The Young’s modulus of the blends can be

improves with organoclay at all types. The organoclay-

DDA in the blend showed the higher the Young’s

modulus with compared to organoclay-DHA and

organoclay-TSA. For high organoclay content, the

percent strain at break decreased may be due to the

decreasing of interfacial interaction between the filler

and the matrix. The XRD results revealed that the

intercalation of POM chains inside the organoclay

galleries, which can improve the dispersion of

organoclay in the POM/PP blends.

Acknowledgements

The authors would like to thank Silpakorn

University Research and Development Institute (SURDI)

and the Higher Education Research Promotion and

National Research University Project of Thailand, Office

of the Higher Education Commission for the financial

support of this project.

References

[1] Seahan, C.; Joung, S. H.; Seung, J. L.; Kyung, H. A.,

Jose, A. C.; Joao, M. M. Morphology and rheology

of polypropylene/polystyrene/clay nanocomposites

in batch and continuous melt mixing

processes.Macromolecular Materials and

Engineering 2011; 296: 341-348.

[2] Sung, Y. T.; Kim, Y. S. Lee, Y.S.; Kim, W. N.; Lee,

H. S.; Sung, J.Y.; Yoon, H.G. Effect of clay on the

morphology of poly(acrylonitrile-butadiene-styrene)

and polypropylene nanocomposites. Polymer

Engineering and Science 2007; 1673-1677.

[3] Suprakas, S. R.; Steve, P.; Mosto, B.; Leszek A., U.

Role of organoclay modified layerd silicate as an

active interfacial modifier in immiscible

polystyrene/polypropylene blends. Polymer 2004;

45: 8403-8413.

[4] Anup, K. D.; Jin, K. K.; Bhanu, B, K. Cocontinuous

phase morphology of asymmetric compositions of

polypropylene/high-density polyethylene blends by

the addition of clay. Applied Polymer Science 2011;

119:3080-3092.

[5] Sumana, M.; Khatua, B. B. Morphology and

properties of nylon6 and high density polyethylene


76 blends in absence and presence of nanoclay. Applied

Polymer Science 2011; 121: 359-368.

[6] Kusmono; MohdIshak, Z. A.; Chow, W. S.;

Takeichai, T.; Rochmadi. Effect of clay

modification on the morphological mechanical and

thermal properties of polyamide6/polypropylene/-

montmorillonite nanocomposites. Polymer

Composites 2010; 1156-1167.

[7] Bitinis, N.;Verdejo, R.; Maya, E. M.; Espuche,

E.;Cassagnau, P.; Lopez-Manchado, M. A.

Physicochemical properties of organoclay filled

polylactic acid/natural rubber blend

bionanocomposites. 2012: 305-313.

[8] Ajith, J. J; Alagar, M. Development and

characterization of organoclay-filled

polyoxymethylene nanocomposites for high

performance applications. Polymer Composites

2011; 1315-1324.

[9] Shashidhara, G. M. Effect of PP-g-MAH compa-

tibilizer content in polypropylene/nylon-6 blends.

2009; 63: 147-157.

[10] Yongjin, L. Co-continuous polyamide 6 (PA6)/

acrylonitrile-butadiene-styrene (ABS) nanocompo-

sites. 2005; 26: 710-715.


77

COMPO-04

Preparation and Characterization of TiO2/WO3/Polythiophene Composite

Nuttaporn Jaritkaun1, *

, Jatuphorn Wootthikanokkhan1, 2

, Pailin Ngaotrakanwiwat 2, 3

and Siriluk Chiarakorn1, 3

1School of Energy, Environment and Materials, King Mongkut's University of Technology Thonburi (KMUTT),

Bangkok 10140 2Nanotec-KMUTT Center of Excellence on Hybrid Nanomaterials for Alternative Energy (HyNAE), School of

Energy, Environment and Materials, King Mongkut’s University of Technology Thonburi, Bangkok 10140 3Department of Chemical Engineering, Burapha University, Chonburi 20131

Abstract

This research work has concerned a development of photocatalyst from hybrid metal oxides/polymer

composites. Titanium dioxide (TiO2) was blended with tungsten trioxide (WO3) at a molar ratio of 1/2 and

thiophene was in situ polymerized by using FeCl3 catalyst at 0.04 molar for 4 hours. The product was then

characterized by FTIR, TGA, XRD, SEM, EDX and UV/Visible spectroscopy techniques. Photocatalytic

activity of the composite catalysts under visible light illumination was then examined by following a

degradation of methylene blue as a function of time. It was found that photocatalytic activity of

TiO2/WO3/polythiophene composite was superior to that of the neat TiO2 and TiO2/WO3 (without polymeric

binder). The results were discussed in light of capability of the polymer in acting as a binder, facilitating charge

transport between metal oxides.

Keyword: Metal oxides, Conducting polymers, Polymerization, Binder, Photocatalysis

1. Introduction

Titanium dioxide (TiO2) is considered to be one

of the most widely used photocatalyst for pollution

abatement studies. This is because TiO2 is relatively

inexpensive, safe and chemically stable [1, 2, 3]. But its

band gap (Eg = 3.2 eV) is considerably high. This means

that it can be excited only by the ultraviolet light only.

However for indoor applications requiring the

degradation of organic compounds in visible light, some

modification of the metal oxide catalyst system is

necessary. In this regard, the use of TiO2 in combination

with the lower band gap transition metal oxides is of

interest and deserves consideration.

In this study, the use of WO3 coupled with TiO2 is of

interest. This is because bandgap energy value of TiO2

(3.2 eV) is higher than that of WO3 (2.6 eV) [4, 5], Upon

UV irradiation, photo-excited electron in the valence

band of TiO2 can transfer to the lower conduction band

of WO3. This means that WO3 can reduce a

recombination rate of the electron and hole.

In this regard, an intimate contact between TiO2

and WO3 (or the so-called active region) should be

created to ensure the charge transfer between the metal

oxide particles. To enhance the connectivity between

isolated metal oxide particles and the charge transfer

between phases, the use of a semiconducting polymer as

a polymeric binder for the hybrid metal oxides is

interesting and worth exploring.

Polythiophene (PTh) was used as a binder for

the TiO2/WO3 system in this study. This was due to the

fact that the energy band gap of PTh is lower than those

of many other conducting polymers [6].

In this study, the preparation, characterizations

and photocatalytic activity of TiO2/WO3/PTh composite

are of interest and focused on. An in situ oxidative

polymerization technique was used to prepare the hybrid

metal oxide/polymer composite. The aim of this work

was to investigate the effects of polythiophene on

photocatalytic activity of a TiO2/WO3/PTh composite

system under visible light. Comparisons of

photocatalytic activity of TiO2/WO3/PTh composite to

that of TiO2, and TiO2/WO3 (without polymeric binder),

are also of interest.

2. Experimental

2.1 Materials

Thiophene (99% reagent grade) was purchased

from Sigma-Aldrich. TiO2 nanoparticles (98.5% anatase


78 from Carlo Erba), WO3 (99.9% AR. grade from Fluka),

and methylene blue (Ajax Finechem) were used as

received without further purification. Chloroform (AR.

grade from Labscan) and anhydrous FeCl3 (Sigma-

Aldrich) were also used as received. Methanol and

distilled water were used to remove the residual FeCl3.

2.2 Preparation of hybrid metal oxides/polythiophene

composite

TiO2/WO3/PTh composite was prepared by in

situ oxidative polymerization. Given amounts of TiO2

(0.3445 g) and WO3 (2 g) were added to a chloroform

solution. The molar ratio of TiO2/WO3 was kept constant

at 1/2 throughout this work. The solution was sonicated

for 30 min. Then 0.0145 g (equivalent to 0.04 moles) of

thiophene was injected into the above suspension at 0 °C

and stirred for 10 min. After that, 50 ml of a saturated

anhydrous solution of FeCl3 in chloroform was added

drop-wise to the above cooled mixture. The reaction was

allowed to proceed at 0 °C for 4 h under a nitrogen gas

atmosphere. The content of the reaction flask was filtered

and washed with a large amount of methanol and

distilled water, respectively. Finally, the product was

dried at 60 °C for 24 h. For the purpose of comparison,

TiO2/WO3 composite was also prepared. In brief,

amounts of TiO2 (0.3445 g) and WO3 (2 g) were blended.

The molar ratio of TiO2 to WO3 was 1/2.

2.3 Catalytic activity

The catalytic activity of the various photoactive

materials were evaluated by monitoring the degradation

of methylene blue (MB) solution using a photo-reactor

with a distance between the lamp and the sample of

about 10 cm. The tests were carried out under visible

light illumination (with a 15-W fluorescent lamp). In a

typical experiment, 0.4 g of TiO2/WO3/PTh composite

was suspended in 400 mL of 10-5

M methylene blue

solution. Then the suspension was stirred in the dark at

room temperature for 12 h until an adsorption–desorption

equilibrium was reached. After irradiation, 9 mL of the

suspension was taken every 30 min during the irradiation

until the total illumination time was reached. All of the

sampling suspensions were centrifuged and the

concentration of the solution was determined by

measuring the absorbance of MB with a Genesys 10s

UV-vis spectrometer at a wavelength of 664 nm. The

degradation was evaluated in term of C/C0, where C0 and

C are the concentrations of MB before and after

irradiation, respectively.


ATR-FTIR spectroscopy (Thermo ScientificTM

Nicolet iS5 FT-IR Spectrometer) was used to detect

some functional groups of the synthesized

polythiophene. The weight percentage of polythiophene

in the composite was determined using thermal

gravimetric analysis (TGA) (Netzsch 409). The TGA

experiment was scanned over temperatures ranging

between 35 and 800 C under a nitrogen gas atmosphere

at a heating rate of 10 C/min. XRD patterns of metal

oxides and TiO2/WO3/PTh composite were investigated

on a Bruker Euler Cradle for D8 Advance X-ray

diffractometer.

The morphology of the synthesized products

was characterized using a scanning electron microscope

(SEM) (Nova NanoSEM 450, FEITM). Energy-

dispersive X-ray (EDX) measurements of various metal

oxides were also carried out using a JEOL JEM-2100

TEM. A dispersion of metal oxides and the polymer

composites in ethanol were deposited on a carbon-coated

copper grid.

The reflectance spectra of the samples were

measured using a UV-vis-NIR spectrophotometer (Cary

5000 Agilent Technologies). From the reflectance value,

the bandgap energy values of metal oxides and

semiconducting polymer were determined by using

Kubelka-Munk’s equation (1) and Planck’s equation (2):

F(R) = (1 – R)2/2R (1)

where F(R) is the Kubelka-Munk fraction, and R is

reflectance.

Eg = hC/λ (2)

where Eg is the band gap energy (eV), h is Planck’s

constant (J/s), C is the speed of light (m/s), and λ is the

cut-off wavelength (nm).


79 3. Results and discussion

3.1 Structural characterizations

Fig. 1 shows overlaid ATR-FTIR spectra of the

homopolymerized thiophene (PTh) and the product

obtained from in situ polymerization of thiophene with

metal oxides particles (TiO2 and WO3). It can be seen

that the several characteristic peaks representing the

chemical bonds in PTh molecules can be observed [7, 8].

These include C=C ring stretching vibration (1490 and

1433 cm−1

) and in-plane C-H bending (1324 cm-1

). Peak

at 1217 cm-1

can be ascribed to stretching vibration of in-

plane C-H whereas those appeared at 1087 and 1037

cm−1

can be related to the in-plane stretching vibration of

C-S. For the products obtained from in situ

polymerization of thiophene with hybrid metal oxides

particles (TiO2 and WO3), the FTIR peaks representing

the polymer were less obvious and the ATR-FTIR

signals were insufficiently strong. It was apparently that

the product yield obtained from the reaction was low.

However, when the amount of monomer was increased,

the FTIR peaks representing the polymer became greater

obvious. This suggested that in situ polymerization of

thiophene in the presence of the hybrid metal oxides

particles was also successful.

Fig. 1 FTIR spectra PTh and the TiO2/WO3/PTh

composite.

The weight percentage of PTh in the composite was

determined using TGA in Fig. 2. It can be seen that the

synthesized PTh homopolymer, the initial weight loss

that occurred over the temperature range between 80 and

160 °C could be attributed to elimination some moisture.

Next, a weight loss transition, occurred at the onset

temperature of 260 °C. This decomposition was

completed at 800 °C [8]. On the other hand, it can be

seen that the TiO2/WO3/PTh composite were not

completely decomposed. This was due to the fact that

both TiO2 and WO3 are thermally stable up to 800 °C.

For the TiO2/WO3/PTh composite, weight fractions of

polymer was 0.1%. These factors resulted to a formation

of the polymerized product with low molecular weight

and low percentage yield. In addition, some differences

between metal oxides and the polymer, in term of atomic

weight values, should also be taken into account. The

atomic weight of Ti (22) and W (74) metals are much

greater than that of the hydrocarbon, existed in the

repeating units of the polymer molecules. In this regard,

no further attempt was made to increase the monomer

content for the polymerization. It was believed that the

active surface area and photocatalytic activity of the

metal oxides might be reduced when the content of the

polymerized product was too high.

Fig. 2 TGA thermograms of PTh and the TiO2/WO3/PTh

composite (1/2/0.04 molar ratio).

3.2 Crystal structure and morphology

The XRD patterns of TiO2, WO3, and the

TiO2/WO3/PTh composite are illustrated in Fig. 3. For

the pattern of TiO2, the diffraction peaks correspond to

the anatase crystal of TiO2 [JCPDS ICDD 21-1272]. For

the pattern of WO3, the diffraction peaks correspond to

the monoclinic WO3 [JCPDS ICDD 83-0950]. The


80 similar XRD patterns were observed from the

TiO2/WO3/PTh composite. This indicates that the

deposition of polymer on the metal oxides particles did

not alter the crystal structure of the materials.

Fig. 3 XRD patterns of TiO2, WO3, and the

TiO2/WO3/PTh composites (1/2/0.04 molar ratio).

The morphologies of TiO2, PTh, TiO2/WO3, and

the TiO2/WO3/PTh composite were observed in Fig. 4.

Aggregated spherical shape TiO2 nanoparticles were

noted. The TiO2/WO3 particles seemed to be more

heterogeneous as compared to that of the neat TiO2

because it contains TiO2 and WO3 particles. SEM image

of the homopolymerized PTh was found that the

synthesized polymer was obtained in the form of

aggregated particles. For the TiO2/WO3/PTh composite,

the morphology of the metal oxides changed

significantly. The product became more agglomerate

particles as compared to that of the neat TiO2 and the

TiO2/WO3 particles.

Fig. 4 SEM images of (a) TiO2, (b) PTh, (c) TiO2/WO3

(1/2 molar ratio) and (d) TiO2/WO3/PTh (1/2/0.04 molar

ratio).

Results from TEM-EDX dot mapping technique

(Fig. 5) showed that distribution pattern of the W (L)

signal (representing the WO3 particles) were poorly

matched with that of the Ti (K) signal (represented the

TiO2 particles) for the neat TiO2/WO3 (without polymer).

This indicates that, in an absence of the polymer binder,

the two types of metal oxides were separated. For the

TiO2/WO3/PTh composite, the signals from K of S

atoms, representing polythiophene, can be noted. This

suggested that the polymerized polythiophene does exist

in the product. In addition, in the presence of PTh, the

distribution patterns of the Ti and W were in a good

agreement and matched. These refer to TiO2 and WO3

were attached together.

3.3 Opto-electrical properties

The energy band gap values of various materials

were summarized in Table 1. It was found that the band

gap energy values of TiO2 and WO3 are also close to

those reported in the literature [9, 10]. It can be seen that

the presence of WO3 and PTh on TiO2 nanoparticles

brought about a reduction of the band gap energy values

of TiO2. The above results suggest that the transition

metal oxide can excited by visible light. A similar effect

was observed in the literature [11, 12].

Fig. 5 TEM images and the corresponding EDX dot

maps of (a) TiO2/WO3 (1/2 molar ratio) and (b)

TiO2/WO3/PTh composites (1/2/0.04 molar ratio).


81 Table 1 Band gap energy values of the various

photoactive materials.

Samples Cut-off wavelength

(nm)

Eg

(eV)

TiO2 360 3.44

WO3 440 2.82

PTh 600 2.07

TiO2/WO3

(1/2 molar ratio)

407 3.05

TiO2/WO3/PTh composite

(1/2/0.04 molar ratio)

415 2.99

Fig. 6 Photodegradation of methylene blue (MB)

solution under visible light illumination by various

catalysts.

3.4 Catalytic activity

Figure 6 show photocatalytic activity under

visible light illumination of the various catalysts also

deserves consideration. It can be seen that the activity of

the neat WO3 was lower than that of the neat TiO2 for the

degradation of MB under visible light because the pure

WO3 has some limitations for use as a photocatalyst for

pollution abatement [13]. The reduction potential of

WO3 is relatively low, and thus the light energy

conversion efficiency of the material is small. Besides,

the photocatalytic activity of WO3 is also dependent on

other factors such as the morphology, particle size, and

crystalline structure [14, 15, 16]. By TiO2/WO3 (without

polymer), the photocatalytic activity of the metal oxides

increased as compared to that of the neat TiO2 and the

neat WO3. This can be ascribed to the WO3 is excited by

visible light, the holes will be formed in the valence band

of WO3. Then, the electrons in the valence band of TiO2

can move to that of WO3. Finally, the. holes generated in

the TiO2 can induce the photocatalytic oxidation

reactions [17].

After applying the PTh to the hybrid metal

oxides system, the catalytic activity of TiO2/WO3

increased. This reflects a better connection and charge

transfer between the isolated metal oxide particles.

4. Conclusions

The results are sufficient to confirm that the

TiO2/WO3/PTh composite have been successfully

prepared via in situ polymerization of thiophene. it was

apparently that TiO2 and WO3 the particles were adhered

together with the presence of PTh, acting as a kind of

polymeric binder. The catalytic activity of the metal

oxides based on TiO2/WO3 in visible light can be

enhanced by using an in situ polymerized polythiophene

as a binder. Performance and photo-catalytic activities of

the various catalysts can be related to their band gap

energy values and morphology of the materials.

5. Acknowledgements

This work has been supported by the

Nanotechnology Center (NANOTEC), NSTDA, Ministry

of Science and Technology, Thailand, through its

programme of the Center of Excellence Network.

6. References

[1] Hashimoto, K., Irie, H. and Fujishima, A., “Ti02

photocatalysis: a historical overview and future

prospects”, Japanese Journal of Applied Physics :

44, 8269-8285 (2005).

2] Legrini, O., Oliveros, E. and Braun, A.M.,

“Photochemical processes for water treatment”,

Chemical Reviews : 93, 671-698 (1993).

[3] Linsebigler, A.L., Lu, G.Q. and Yates, J.T.,

“Photocatalysis on Ti02 surfaces- principles,

mechanisms, and selected results”, Chemical

Reviews : 95, 735-758 (1995).

[4] Ngaotrakanwiwat, P., Saitoh, S., Ohko, Y., Tatsuma,

T. and Fujishima, A., “Charge-discharge behavior of

TiO2-WO3 photocatalysis systems with energy


82 storage ability”, Physical Chemistry Chemical

Physics : 5, 3234–3237 (2003).

[5] Tatsuma, T., Saitoh, S., Ngaotrakanwiwat, P., Ohko,

Y. and Fujishima, A., “Energy storage of TiO2-WO3

photocatalysis systems in the gas phase”, Langmuir :

18, 7777-7779 (2002).

[6] Dai, L., “Conducting Polymers”, In: 1 (ed.s)

Intelligent Macromolecules for Smart Devices:

From Materials Synthesis to Device Applications,

Engineering Materials and Processes, USA, acid-

free paper : 43 (2004).

[7] Gnanakan, S.R.P., Rajasekhar, M. and Subramania,

A., “Synthesis of polythiophene nanoparticles by

surfactant-assisted dilute polymerization method for

high performance redox super capacitors”,

International Journal of Electrochemical Science :

4, 1289–1301 (2009).

[8] Liu, R.C. and Liu, Z.P., “Polythiophene: synthesis in

aqueous medium and controllable morphology,

Chinese Science Bulletin : 54, 2028-2032 (2009).

[9] Geng, L., Huang, W., Zhao, Y., Li, P., Wang, S.,

Zhang, S. and Wu, S., “H2S sensitivity study of

polypyrrole/WO3 materials”, Solid-State

Electronics : 50, 723–726 (2006).

[10] Zhu, J., Wei, S., Zhang, L., Mao, Y., Ryu, J.,

Mavinakuli, P., Karki, A.B., Young, D.P. and Guo,

Z., “Conductive polypyrrole/tungsten oxide

metacomposites with negative permittivity”, Journal

of Physical Chemistry C : 114 16335–16342 (2010).

[11] Luo, Q., Li, X., Wang, D., Wang, Y. and An, J.,

“Photocatalytic activity of polypyrrole/TiO2

nanocomposites under visible and uv light”, Journal

of Materials Science : 46, 1646–1654 (2011).

[12] Dimitrijevic, N.M., Tepavcevic, S., Liu, Y., Rajh,

T., Silver, S.C. and Tiede, D.M. “Nanostructured

TiO2/polypyrrole for visible light photocatalysis”,

The Journal of Physical Chemistry C : 117, 15540–

15544 (2013).

[13] Asim, N., Badeiei,M., Ghoreishi, K.B., Ludin, N.A.,

Zonooz, M.R.F. and Sopian, K., “New

developments in photocatalysts modification: case

study of WO3”, Advances in Fluid Mechanics and

Heat & Mass Transfer : 110-116 (2012).

[14] Xin, G., Guo, W. and Ma, T., “Effect of annealing

temperature on the photocatalytic activity of WO3

for O2 evolution”, Applied Surface Science : 256,

165-169 (2009).

[15] Zhao, W., Wang, Z., Shen, X., Li, J., Xu, C. and

Gan, Z., “Hydrogen generation via

photoelectrocatalytic water splitting using a tungsten

trioxide catalyst under visible light irradiation”,

International Journal of Hydrogen Energy : 37, 908-

915 (2012).

[16] Purwanto, A., Widiyandari, H., Ogi, T. and

Okuyama, K., “Role of particle size for platinum-

loaded tungsten oxide nanoparticles during dye

photodegradation under solar-simulated

irradiation”,Catalysis Communication : 12, 525-529

(2011).

[17] Seung, Y.C., Yong, J.K. and Wan, I.L.,

"Photocatalytic WO3/TiO2 nanoparticles working

under visible light”, Journal of Electroceramics : 17,

909–912 (2006).


83

COMPP-02

Tensile Properties of Poly(Butylene Succinate) Reinforced with Rice Husk Silica

Apimook Laptippamon1, a

, and Pranut Potiyaraj2,b,*

1Interdisciplinary Program in Petrochemical & Polymer Science, Faculty of Science, Chulalongkorn University,

Bangkok 10330 Thailand 2Department of Materials Science, Faculty of Science, Chulalongkorn University, Bangkok 10330 Thailand

Abstract The objective of this study is to prepare rice husk silica reinforced poly(butylene succinate) composites

(PBS/silica). PBS composites were prepared by using a twin screw extruder. Glycidyl methacrylate grafted

poly(butylene succinate, PBS-g-GMA) were also used as a compatibilizer to improve the interfacial interaction

between PBS and silica. PBS composites were then processed into test specimens by injection molding. The

effect of type and ratio of silica on the tensile properties of the composites was investigated. It was found that

the tensile strength and elongation at break of composites without the compatibilizer was lower than neat PBS

due to agglomeration of silica. Nevertheless, tensile strength and Young’s modulus of PBS composites were

improved after the incorporation of PBS-g-GMA.

Keywords: poly(butylene succinate), silica, rice husk silica, composite, compatibilizer, mechanical property

Introduction

The development of biodegradable polymeric

materials with excellent material properties has received

much more attention worldwide. [1,2] Aliphatic

polyesters, such as poly(ε-caprolactone), poly(lactic

acid), poly(butylene succinate) (PBS), and poly(3-

hydroxybutyrate), are among the most promising

materials for the production of high-performance,

environmentally friendly, biodegradable materials. [3-5]

PBS, synthesized by the polycondensation of 1,4-

butanediol with succinic acid, has particularly attracted

increasing commercial interest because of its many

interesting properties, including biodegradability, melt

processability, and thermal and chemical resistance. [6]

However, its softness and low gas-barrier properties have

restricted further application of PBS. To improve PBS’s

properties, the conventional method is to use suitable

additives to improve its inferior attributes.

Rice husk, a waste product of the rice

production industry, is an environment threat causing

damage to the land and the surrounding area in which it

is dumped. However, rice husk contains about 20% of

ash which can be retrieved as amorphous, chemically

reactive silica. This rice husk ash (RHA) silica finds

broad applications such as filler, catalyst support,

adsorbent and a source for synthesizing high

performance silicon and its compounds. Various metal

ions and unburned carbon influence the purity and

color of the ash. Controlled burning of the husk after

removing these ions can produce white silica of high

purity. Husk contains 17%-20% silica in complex form

and RHA contains 85%-95% amorphous silica. [7]

The objective of this study is to prepare rice

husk silica reinforced poly(butylene succinate)

composites (PBS/silica). PBS-g-GMA was produced as a

compatibilizer via the reactive extrusion technique. The

compatibilizer was then used as a compatibilizer for

PBS/silica composites. The tensile properties of PBS

composites reinforced with rice husk silica and

commercial silica were then investigated and compared

with those of composites compatibilized by PBS-g-

GMA.

Experimental

Materials. PBS granules (AZ71TN) were of injection

grade purchased from Mitsubishi Chemical. PBS was

dried at 60°C for 24 h in a hopper dryer before use. Silica

(Siam silica Co.LTD) and ultrasil 9000 GR (Evonik

Industry AG) was dried at 60°C for 24 h and stored in

the desiccator before use. Dicumyl peroxide (DCP)

(Sigma Aldrich) was used as an initiator and glycidyl


84 methacrylate (GMA) (Sigma Aldrich) was used as a

reagent without further purification.

Synthesis of PBS-g-GMA. Firstly, PBS, GMA and DCP

were physically premixed. The reactive grafting process

was carried out in a twin-screw extruder model at 130°C

at 40 rpm. The obtained product was refluxed in

chloroform for 4 h, and the hot solution was filtered into

cold methanol. The precipitated polymer was washed

with methanol five times, in order to remove any

unreacted reagents, and followed by drying in an oven at

60°C for 24 h. The purified PBS-g-GMA was then

obtained.

Preparation of the PBS Composites. PBS composites

were prepared by the addition of 0, 0.5, 1, 2 and 3 phr of

silica with the presence of 0 and 5 phr of PBS-g-GMA

into PBS. PBS, silica and PBS-g-GMA were physically

premixed and melt-mixed in a twin screw extruder

model at 128°C with a screw speed of 40 rpm. After that,

the compounds were injected with an injection molding

machine BATTENFELD(Austria) model BA 250 CDC

40 Ton at 135°C into specimens for tensile testing.

FT-IR Characterization. For Fourier transform infrared

spectroscopy analysis was carried out on the PBS-g-

GMA at ambient temperature by using Thermo scientific

model Nicolet 6700 it was performed through the

scanning wavelength from 4000 to 400 cm-1

.

Mechanical Testing. Measurements of tensile properties

those are tensile strength, Young’s modulus and

elongation at break were performed using a Universal

testing machine LLOYD in accordance with ASTM

D638 with a gauge length of 115 mm and using the cross

head speed of 50 mm/min and a 10 kN load cell.

Results and Discussion

Characterization of PBS-g-GMA. The in situ melt

grafting of GMA onto the PBS matrix was performed in

a reactive twin-screw extruder using the concentration of

glycidyl methacrylate of 2 phr and dicumyl peroxide of

1.5 phr. After purification, the PBS-g-GMA was

analyzed by FTIR. The FTIR spectra are shown in Fig. 1.

The absorption at 1731 cm-1

is the characteristic band of

ester carbonyl group in PBS-g-GMA spectrum which

indicates the existence of GMA [9].

Fig. 1 FTIR spectra of neat PBS and grafted PBS-g-

GMA after purification

Mechanical Properties. Fig. 2 shows the tensile strength

of PBS/PBS-g-GMA/silica composites. The tensile

strength decreased with increasing concentration of RHA

silica. This decrease is an indication of poor adhesion

between the polymer matrix and inorganic nanoparticles

and the occurrence of large agglomeration [10]. A

similar trend for the composites with the commercial

silica was observed. At the same amount of filler, the

incorporation of PBS-g-GMA slightly improved the

tensile strength. The interactions between PBS and silica

were obviously promoted with the addition PBS-g-GMA

as demonstrated by the increase in the tensile strength.

Fig. 3 shows effect of fillers content on the tensile

modulus of composites [8].

There is significant increase in the Young’s

modulus of the composites as compared to the neat PBS.

The Young’s modulus increased as the silica contents

increased both with and without PBS-g-GMA and both

of silica because of the rigidity nature of silica. At the

same amount of filler, the incorporation of PBS-g-GMA

slightly improved the tensile modulus [8].


85

a.

b.

Fig. 2 The effect of silica content on tensile strength

a.Rh silica b. ultrasil 9000 GR

( without PBS-g-GMA, PBS-g-GMA 5 phr)

a. b.

Fig. 3 The effect of silica content on young’s modulus



a.

b.

Fig. 4 The effect of silica content on Elongation at break



In contrast, the elongation at break gradually

decreased when the filler content increased both with and

without PBS-g-GMA as shown in Fig. 4.

Conclusion

PBS/silica composites at various ratios of the

filler were prepared and their tensile properties were

investigated. It was found that the tensile properties of

PBS/silica decreased as the amount of silica increased

due to the agglomeration of silica, due to particle–

particle interactions of silica. Silica particles into the

PBS matrix were observed at low silica content, while

some small agglomerates formed at higher

concentrations and the incorporation of PBS-g-GMA

slightly improved the tensile strength. The interactions

between PBS and silica were obviously promoted with

the addition PBS-g-GMA as demonstrated by the

increase in the tensile strength. To improve matrix-filler

interaction, PBS-g-GMA was prepared through melt-

grafting reactive extrusion process to be used as a

compatibilizer. The presence of GMA on PBS was

confirmed by the FTIR analysis. It was found that the

37

38

39

40

41

0 0.5 1 2 3

Ten

sile

str

en

gth

(M

Pa

)

Rh silica content (phr)

37

38

39

40

41

0 0.5 1 2 3

Ten

sile

str

en

gth

(M

Pa

)

Ultrasil 9000 GR content (phr)

600

650

700

750

800

0 0.5 1 2 3 Yo

un

g's

mo

du

lus

(MP

a)

Rh silica contene (phr)

550

600

650

700

750

800

0 0.5 1 2 3 Yo

un

g's

mo

du

lus

(MP

a)

Silica 9000 GR content (phr)

0

5

10

15

0 0.5 1 2 3

% E

lon

ga

tio

n a

t b

rea

k

Rh silica content (phr)

0

5

10

15

0 0.5 1 2 3

% E

lon

ga

tio

n a

t b

rea

k

Silica 9000 GR content (phr)


86 tensile strength and tensile modulus slightly decreased

with the presence of PBS-g-GMA.

Acknowledgement

This research has been supported by the

Ratchadaphiseksomphot Endowment Fund 2013 of

Chulalongkorn University (CU-56-416-AM).

Reference

1. P.Pan, Y. Inoue, Prog. Polym. Sci. 2009, 34, 605

2. Gorrasi G., Vittoria V., Murariu M., Ferreira A.,

Alexandre M., Dubois P. Biomacromolecules

2008, 9, 984.

3. M. Avella*, M.E. Errico, P. Laurienzo, E.

Martuscelli, M. Raimo, R. Rimedio, Polymer 41

(2000) 3875–3881

4. P. Van De Witte, P. J. Dijkstra, J. W. A. Van

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M, (2003) Macromolecules 36:2355 49.

6. Song, J. B.; Song, C. L.; Wang, S. Y.; Zhang, H. F.;

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7. R. Ghosh, S. Bhattacherjee: J Chem Eng Process

Technol 2013, 4:4

8. A.A. Vassiliou, D. Bikiaris,K. El Mabrouk, M.

Kontopoulou J. App Polym Sci, Vol. 119, 2010–

2024 (2011)

9. C. S. Wu, H.T. Liao and J.J. Jhang: Polym. Bull. Vol.

70 (2013), p. 3443

10. A.A. Vassiliou, G.Z. Papageorgious, M.

Kontopoulou, A. Docoslis, D. Bikiaris. J. App

Polymer Sci, Vol. 54 (2013) 1018-1032.


87

COMPP-06

Effect of PGA:STY Ratio and Reaction Temperature on Coating PGA-STY on EPS

Beads

Jittranuch Jirapathomkul1 and Surachai Pornpakakul

1,2*

1Program in Petrochemistry and Polymer Science, Faculty of Science, Chulalongkorn University, Bangkok

10330 2Department of chemistry, Faculty of Science, Chulalongkorn University, Bangkok 10330

Abstract

In this research, effect of PGA:STY ratio (STY 2.5-7.5%w/v) and reaction temperature (60-80°C) on

coating rate and size of PGA-STY/EPS beads were investigated. We found that ratio of PGA:STY affected the

size of PGA-STY/EPS beads and coating rate. Increasing PGA:STY ratio in polymerization process resulted in

increasing coating rate. Using 7.5%w/v of styrene monomer, the size of PGA-STY/EPS beads obtained was

smallest (1.77 mm) and percentage of PGA-STY loading up to 91.5%. Reaction temperature affected the

percentage of PGA-STY loading onto EPS beads which the higher reaction temperature resulted in the higher

polymerization rate. Moreover, the morphology of the PGA-STY/EPS beads characterized by scanning electron

microscopy showed that the PGA-STY was successfully coated onto the EPS beads.

Keywords: Expandable polystyrene beads, Polyglutaraldehyde-styrene, Copolymer

1. Introduction

Glutaraldehyde is a linear 5-carbon dialdehyde.

It is an oily, colorless clear liquid with a pungent, fruity

odor. It can be polymerized to generated

polyglutaraldehyde through aldol condensation of

glutaraldehyde under basic condition. As a result of the

condensation, the polyglutaraldehyde contains a presence

of carbon-carbon double bond in its backbone and non-

conjugated aldehyde and conjugated aldehyde [1]. In

many research, the polyglutaraldehyde has been used as

new reagent for the immobilization of protein such as

antibodies or enzyme on solid substrates [2]. Several

studies on copolymerization of styrene and

polyglutaraldehyde such as Dizge et al. have successfully

prepared styrene-divinylbenzene-polyglutaraldehyde

(STY-DVB-PGA) copolymer by using High Internal

Phase Emulsion Polymerization (polyHIPE) technique

for enzyme immobilization [3]. Thanaporn has

successfully prepared styrene-polyglutaraldehyde

copolymer coated on styrene bead as support for enzyme

immobilization [4]. Since expandable polystyrene bead

(EPS) contains pentane as a blowing agent and can be

expended by heating at any higher temperature, the rate

of copolymer polymerization and of

copolymer coating should be measured by comparison

with expansion rate of EPS bead and any desired sizes of

particle can be achieved depending on temperature.

Thus, expandable polystyrene bead (EPS) was used as

support in this research and effect of

polyglutaraldehyde:styrene ratio and reaction

temperature on coating onto EPS beads were

investigated.


2.1 Preparation of polyglutaraldehyde (PGA) suspension

Polyglutaraldehyde suspension was obtained from

aldol condensation of glutaraldehyde solution (20%w/v)

catalyzed by 1 M NaOH solution at room temperature

for 5 min.

2.2 Preparation of polyglutaraldehyde-styrene coated

expandable polystyrene beads (PGA-STY/EPS beads)

Prior to use, the styrene monomer was treated

with 12% sodium hydroxide solution in a separatory

funnel three times in order to remove the anti-

polymerizer.

To mixture of styrene monomer, Tween 80® and

polyglutaraldehyde solution in a round bottom flask,

potassium persulfate as initiator was added. After rapid

stirring for 15 min, expandable polystyrene beads were


88 added and the mixture was continuously stirred at 60°C

for 2 h followed by heating at 80°C for the

polymerization time of 22 h. The flow chart of

experiment procedure was summarized as shown in

figure 1.

Figure 1 Flow chart for experimental procedure

2.3 Loading estimation

Percentage of polyglutaraldehyde-styrene coated

onto expandable polystyrene beads was calculated

as %loading using following equation:

2.4 Size of the polyglutaraldehyde-styrene coated

expandable polystyrene beads

The size of the polyglutaraldehyde-styrene coated

expandable polystyrene beads was measured by using

digital microscope (AD 4113TL-Dino-Lite Pro2) at 60

magnifications. The average diameter of the beads was

calculated by an image of one hundred support beads

from digital microscope photograph using dinocapture

2.0 software.

2.5 Attenuated total reflectance-Fourier transform

infrared spectroscopic (ATR-FTIR)

EPS beads and PGA-STY/EPS beads were

characterized by ATR-FTIR. The solid sample was

placed onto the Universal diamond ATR top-plate and

then applying pressure to a solid sample on the Universal

diamond ATR top-plate for characterization.


3.1 Characterization of PGA-STY/EPS beads

The ATR-FTIR spectra of EPS bead and PGA-

STY/EPS bead were compared as shown in Figure 2a,

2b. The spectrum of PGA-STY/EPS bead showed

absorption of aromatic and olefinic C-H stretching at

3079, 3056 and 3019 cm-1

, of aliphatic C-H stretching at

2924 and 2860 cm-1

, of C=O stretching at 1715 cm-1

for

non-conjugated aldehyde and at 1678 cm-1

for

conjugated aldehyde and C=C stretching at 1639 cm-1.

These results indicated that coating styrene-

polyglutaraldehyde onto EPS bead was successful

3.2 Effect of PGA:STY ratio and reaction temperature on

the size of PGA-STY/EPS beads

The effect of PGA:STY ratio and reaction

temperature were studied by using a fixed amount of

20%w/v polyglutaraldehyde, Tween 80®, K2S2O8 at 20

ml, 0.88 g and 2.59 mmol, respectively, and reaction

time of 24 h. The amount of styrene monomer was varied

from 2.5, 5 and 7.5%w/v and reaction temperature was

varied from 60, 70 and 80°C. The PGA-STY/EPS beads

prepared were shown in Figure 3

From Table 1 and Figure 3, it was found that

PGA:STY ratio affected the size of PGA-STY/EPS

beads. When the amount of styrene monomer was

increased at any reaction temperatures, the size of PGA-

STY/EPS beads was decreased from 2.53 to 1.77 mm.

These results suggested that the increasing styrene

monomer in polymerization process resulted in increase

of conversion of polymerization [5]. So the

polymerization rate of copolymer in condition 7.5%w/v

of STY monomer which coated on EPS beads was the

%loading =

amount of PGA-STY loaded

on polystyrene beads

amount of polystyrene beads

introduced

ˣ 100 (1)

Styrene 2.5%, 5%,

7.5%w/v

Tween 80® K2S2O8

GA

solution

20%w/v

Aldol condensation

of GA, at 25°C

for 5 min

EPS

beads

STY-PGA coating

PGA-STY/EPS

beads

PGA-STY/EPS

beads

PGA-STY/EPS

beads

T = 60°C

T = 70°C

T = 80°C


89 fastest resulted in the size of PGA-STY/EPS beads

obtained from the highest STY monomer was the

smallest. In table 1, comparison of the size of PGA-

STY/EPS beads (1.77 mm) obtained from 7.5%w/v of

STY monomer, 20%w/v of PGA at reaction temperature

80oC with the size of uncoated EPS beads (2.67 mm)

caused by heating at temperature 80oC (without PGA-

STY) also suggested that coating rate of PGA-STY onto

EPS beads should be faster than expansion rate of EPS

beads.

Figure 2 The ATR-FTIR spectra of (a) EPS bead (b) PGA-STY/EPS bead

Table 1 Effect of PGA:STY ratio and reaction

temperature on the size of PGA-STY/EPS beads

and %loading of copolymer

STY

(%w/v)

PGA

(%w/v)

Average diameter of

PGA-STY/EPSa (mm)

% loading of

copolymer

Temperature (°C)

60

70

80

60

70

80

2.5

20

1.46

2.08

2.53

47.8

52.6

62.4

5

20

1.96

1.99

2.02

66.8

74.8

81.9

7.5

20

1.77

1.81

1.77

69.1

76.8

91.5

Uncoated EPSb

1.35

1.48

2.67

-

-

-

a Average diameter of original EPS beads 1.2 mm

b Average diameter of uncoated EPS beads (without PGA-STY)

after heat treatment for 24 h measured by digital microscope.

Table 2 Average diameter of uncoated EPS beads after

heat treatment a

Average diameter of heated uncoated EPS beadsb (mm)

Temperature of heat treatment (°C)

60 70 80

2h 6h 24h 2h 6h 24h 2h 6h 24h

1.20 1.20 1.35 1.22 1.32 1.48 1.60 2.10 2.67

a Average diameter of heated uncoated EPS beads measured by

digital microscope.

b Average diameter of original EPS beads 1.2 mm

Using 5%w/v and 7.5%w/v of styrene monomer,

reaction temperature affected a little increase of the size

of PGA-STY/EPS beads while increase of reaction

temperature resulted in significant increase of %loading.

These should be caused by the increasing styrene

monomer in polymerization process resulted in fast

conversion rate of polymerization. Thus, the size of

PGA-STY/EPS beads obtained was rather constant at

reaction temperature 60°C, 70°C, 80°C. Also, the


90 reaction temperature affected the percentage of PGA-

STY loading onto EPS since the reaction temperature

increased leads to increase of decomposition rate of

initiator to give higher polymerization [5]. In case of

PGA-STY/EPS beads obtained from using 7.5%w/v of

styrene monomer, 20%w/v of PGA, 80oC of reaction

temperature, 91.5% loading and smaller size of the beads

than the beads obtained from 70oC of reaction

temperature were another evidence of the higher reaction

temperature, the higher polymerization rate and of faster

polymerization rate than expansion of EPS.

Figure 3 PGA-STY/EPS beads: (a) 2.5%W/V of STY at

(a1) 60°C, (a2) 70°C, (a3) 80°C (b) 5%w/v of STY at (b1)

60°C, (b2) 70°C, (b3) 80°C (c) 7.5%w/v of STY at (c1)

60°C, (c2) 70°C, (c3) 80°C

Moreover, Pictures and SEM micrograph of

PGA-STY/EPS bead (Figure 4) showed that the core of

PGA-STY/EPS bead was EPS bead and PGA-STY was

coated outside the EPS beads. This SEM micrograph also

supported how increasing styrene monomer in

polymerization process resulted in fast conversion rate of

polymerization.

Figure 4 Pictures of (a) heated uncoated EPS bead (b)

PGA-STY/EPS bead (cross section) and SEM

micrograph of (c) PGA-STY/EPS bead (d) PGA-

STY/EPS bead (cross section)

4. Conclusion

We found that PGA-STY was coated onto the

EPS bead. The PGA:STY ratio affected coating rate and

size of PGA-STY/EPS beads. Using 7.5%w/v of styrene

monomer in polymerization process resulted in the size

of PGA-STY/EPS beads obtained was smallest (1.77

mm) because coating rate of PGA-STY onto EPS beads

was faster than expansion rate of EPS beads and

percentage of PGA-STY loading up to 91.5%. Reaction

temperature affected the percentage of PGA-STY

loading onto EPS beads which the higher reaction

temperature, the higher polymerization rate.

References

[1] Rembaum, A., Margel S., “Design of Polymeric

Immunomicrospheres for Cell Labeling and Cell

Separation”, British Polymer Journal, 275-280

(1978).

[2] Rembaum, A., Margel S., Levy, J.,

“polyglutaraldehyde: A New Reagent for Coupling

Proteins to Microsphere and For Labeling Cell-

Surface Receptors”, Journal of Immunological

Methods, 239-250 (2009).

[3] Dizge, N., Keskinler, B., Tanriseven, A., “Biodiesel

Production from Canola Oil by Using Lipase

Immobilized onto Hydrophobic Microporous

(d)

(c) (a)

(b)

(a1) (a2) (a3)

(b1) (b2) (b3)

(c1) (c2) (c3)

(a) (c) (b) (d)


91 Styrene-Divinylbenzene Copolymer”, Biochemical

Engineering Journal, 220-225 (2009) .

[4] Jitrasing, T., Pornpakakul, S., “Polyglutaraldehyde-

Styrene Copolymer Modified Polystyrene Beads for

Improvement of Polymer Supports”, The 4th

Science

Research Conference, 284-287 (2012).

[5] Lin, H.R., “Solution Polymerization of Acrylamide

Using Potassium Persulfate as an Initiator: Kinetic

Studies, Temperature and pH Dependence”,

European Polymer Journal, 1507-1510 (2001).


92

COMPP-10

Effect of Sawdust Particle Sizes on Wheat Gluten-Based Biocomposites

Sawalak Jujai, Teerarat Kongprasert, Nattaya Tawichai, Uraiwan Inthata, Nattakan Soykeabkaew*

School of Science, Mae FahLuang University, Thasud, Chiang Rai 57100, Thailand

Abstract

The main objective of this research was to produce biocomposites as an alternative to conventional

materials based on petroleum. In this work, sawdust (SD) particles were used as reinforcement in the wheat

gluten (WG)-based composites. The effect of SD particle size on physical and mechanical properties of the

resultant composites was investigated. The as received SD was sieved into four ranges of particlesizes (i.e.,

<180 µm, 180–250 µm, 250–550 µm and 550–2500 µm), then hand-mixed with WG and glycerol

(plasticizer)and compression molded into the composite sheets. The bulk density, flexural properties and

fracture surface morphology of the composites were studied. The composite reinforced with SD particles size of

180-250 µm indicated the best mechanical performance. When the larger SD particle sizes were integrated,

mechanical properties of the composites tended to decline. From SEM images, an inferior interface between the

larger SD particles and WG-based matrix was revealed.

Keywords: biocomposites, sawdust, wheat gluten, particle size, mechanical properties

1. Introduction

Nowadays, the replacement of conventional (non-

degradable) plastics by biodegradable alternatives in a

short service life application has been increasingly

promoted [1]. New products derived from biodegradable

and renewable resources have been extensively

developed [2]. At present, bioplastics are being utilized

as packaging materials, horticultural products etc.

However, the use of these polymers is still limited due to

their high cost, inadequate performance, and/or strong

tendency to absorb moisture [1,3]. To improve

mechanical properties of bio-based polymers and extend

their applications, the addition of reinforcements such as

natural fibers to produce eco-friendly bio-composites is

one common approach [4,5] Natural fibers are mainly

consisted of cellulose, hemicellulose and lignin [5].

Instead, utilizing cellulose-based feed stocks from

agricultural wastes and byproducts obtained from

industrial processes, such as rice straw, corn cob,

coconut shell, pineapple stick, oil, bagasse palm shell

and sawdust will provide inexpensive, renewable, and

sustainable source [2,6]. Among the renewable natural

polymers, proteins from plants have been most studied

and develop for bioplastic manufacturing since they are

abundant, relatively low cost and rapidly biodegradable

[7]. In particular, wheat gluten (WG) has great potential

to substitute for conventional plastics because of its

suitable mechanical properties and interesting gas barrier

properties [8,9]. However, it has some drawbacks that

limit its widespread application, such as brittleness and

water absorption [10,11]. Incorporation of natural fibers,

e.g., hemp [4,5], jute [12], wheat straw [2], coconut

[10,11,13] and wood [3,5] fibers, into WG to improve

properties of the composite materials has shown to be

one of the effective ways. WG-based materials are

usually thermally processed between 80°C and 130°C

with plasticizer contents around 25-35% [4].

In this study, the WG-based biocomposites

reinforced with sawdust (SD) from the furniture

manufacturer in Lumpang Province, Thailand, were

prepared. It was well known that type, size as well as

size distribution of reinforcements had a major influence

on mechanical properties of the composite materials.

This present work aimed to investigate effect of SD

particle size on morphology, physical and mechanical

properties of the WG-based biocomposites.


2.1 Materials

Sawdust (SD )of Phyllocarpusseptentrionalis

Donn. SD was kindly supplied by the furniture

manufacturer in Lampang Province, Thailand. Vital

wheat gluten was purchased from Zhang Jia Gang

HengFeng Co. Ltd, China. Commercial grade glycerol


93 was used in this work. Magnesium nitrate salt used in the

controlled relative humidity (RH) chamber

of52.9%Mg(NO3)2 was supplied by Ajax Finechem Pty

Ltd., Australia.

2.2 Preparation of the biocomposites

The as received SD was firstly sieved into four

ranges of particlesizes, i.e., <180 µm, 180–250 µm, 250–

550 µm and 550–2500 µm. To prepare the

biocomposites, 62.5 g SD particles, 43.75 g wheat gluten

and 18.75 g glycerol (plasticizer) were hand-mixed until

homogeneous for 10 min. Then, the mixture was placed

into the metal mold with cavity size of 130 × 170 × 3.75

mm3. Each side of the mold was formerly covered with

the polyester sheets sprayed with mold release agent.

Next, the mold was deposited in the center of

compression molding machine (Labtech LP-S-80) heated

at 130°Cunder pressure of 26.6 MPa for 10 min, and

then, cooled for 5 min.


2.3.1Bulk density

Bulk density (D) of the SD particles in each range

of particle sizes were measured by weighting them (M)

using a 4-digit balance (Denver instrument, SI-234)in

container of know volume (V). For the biocomposites,

the samples were cut into the sizes of 20 mm × 100mm

and weighed (M). Then, each specimen dimension

(length, width, and thickness) was measured using a

digital caliper (TCM) for determining its volume (V).The

bulk density (D) was then calculated according to the

following equation:

v

mD (1)

where D = density (g/cm3), M = mass (g), V = volume

(cm3).

2.3.2 Scanning electron microscopy (SEM)

Surface morphologies of the SD particles as well

as the fracture surfaces of all biocomposites after failure

under mechanical test were examined by a scanning

electron microscope (SEM, LEO 1450 VP). All sample

surfaces were coated with gold sputtering prior to SEM

observation. The accelerating voltage of 10 kV was used.

2.3.3 Mechanical tests

The flexural test (4-point bending) of the

biocomposite specimens was performed using an Instron

5560universal testing machine with a load cell of 1 kN.

The rate of crosshead motion was2 mm/min according to

ASTM D6109-13.Before testing, the samples were

conditioned at the controlled relative humidity (RH)

chamber of 52.9% using a chamber of saturated Mg

(NO3)2 solution at 25°C (according to ASTM E104) for

40 h.The values of Flexural strength (S), maximum strain

(r) and modulus of elasticity (E) were reported as the

mean ±SD of five replicates for each biocomposite.

Flexural strength (S) for each specimen was

calculated according to the following equation:

2bd

PLS (2)

where S = stress in outer fiber throughout load span, psi

(MPa), P = total load on beam at any given point on the

load deflection curve, lb (N), L = supports pan (mm), b =

width of beam (mm), and d = depth of beam

(mm).Maximum strain for each specimen was calculated

by the following equation:

2

70.4

L

Ddr (3)

where r = maximum strain, D = midspan (mm)

deflection, d = depth of the beam, in. (mm) and L =

support span, in. (mm)

Modulus of elasticity (E) was determined from

the slope of the straight line that joins the originand a

selected point on the stress strain was calculated by the

following equation:

r

sE (4)

where E = Modulus of elasticity (MPa),r = maximum

strain and s = Strength (MPa)


The sawdust (SD) in four ranges of particle sizes,

i.e., <180 µm, 180–250 µm, 250–550 µm and 550–2500


94 µm are shown in Fig. 1.From SEM photos, the surfaces

of SD particleswere rough. In each range, SD particles

had various sizes and shapes. After seiveing, the portion

of SD particles in the range of 180–250 µm were

obtained in the highest quantity. The SD particles in this

range (Fig. 1b) also had the narrowest distribution of size

and shape when compared to the SD particles in the

other ranges.

Table 1: Bulk density of SD with different particle sizes

and their biocomposites

SD particles

sizes

SD particles

(g/cm3)

SD composites

(g/cm3)

<180 µm 0.317 ± 0.004 1.33 ± 0.02

180–250 µm 0.319 ± 0.003 1.34 ± 0.05

250–550 µm 0.305 ± 0.005 1.33 ± 0.01

550–2500 µm 0.281 ± 0.005 1.30 ± 0.01

From Table 1,the bulk density of SD particles in

the ranges of <180 µm, 180–250 µm, and 250–550 µm

were relatively close and obviously higher than that of

the SD particles of 550–2500 µm. However, the bulk

density of SD particles size of 180–250 µm was the

highest. Possibly, it was because the SD particles in this

range had the narrowest size and shape distribution

leading to the closest packing among these particles.

After compression molding, all WG-based

biocomposites reinforced with different SD particle sizes

showed the bulk density in the range of 1.30-1.34 g/cm3.

Voids between SD particles in these composites were

largely eliminated under high pressure during

compression [14]. Therefore, the different particle sizes

of the initial SD had almost no influence on the bulk

density of the resulting composites.

Fig. 1 Sawdust (SD) particles in the size range of (a)

<180 µm, (b) 180–250 µm, (c) 250–550 µm and (d) 550–

2500 µm.

400 µm

(a)

(b)

(c)

(d)


95

Fig. 2 Mechanical properties of WG-based

biocomposites reinforced with different SD particle

sizes.

Fig. 3 Stress-strain curves of WG-basedbiocomposites

reinforcedwith different SD particle sizes.

From the mechanical test, the composites

reinforced with SD particle size of <180 µm and 180–

250 µm showed both flexural strength and modulus

higher than the other two composites combined with the

larger SD particles (Fig. 2 and Fig. 3). SEM images also

revealed that their fracture surfaces had less pull-out of

SD particles from the WG-based matrix phase(Fig. 4),

suggesting a better interface in these composites. The

good fiber-matrix interfacial adhesion generally

improves mechanical properties of the composite [15].

Perhaps, it was due to the larger surface area of the

smaller SD particles and a narrow particle size

distribution which allowed a good interaction between

the SD particles and WG matrix in these

composites.Though, the composite reinforced with SD

particle size of 180–250 µm had both the highest flexural

strength and modulus.

On the other hand, from the stress-strain curves

(Fig. 3), it was obvious that the composite prepared with

the largest SD particles presented the lowestflexural

properties. Its fracture surface revealed more pull-out of

the large SD particles from the matrix phase and some

large and deep holes were also observed (Fig. 4d).This

indicated to a poor interface of this composite [10]. As

compared to the small SD particles, the large SD

particles generally had lower total surface area to interact

with the matrix phase. This reason possibly explained the

decline tendency in mechanical properties of the

composites as the size of the reinforcing SD particles

increased.

4. Conclusions

WG-based biocomposites reinforced with four

ranges of SDparticle sizes, i.e., <180 µm, 180–250 µm,

250–550 µm and 550–2500 µm, were prepared by

compression molding.It was found that the size of SD

particles had only little influence on the bulk density of

the composites. From the mechanical test results, the

composite reinforced with SD of particle size 180-250

µm exhibited the highest flexural strength and modulus.

As the size of SD particles increased, the mechanical

performance of the composites tended to decrease. On

the composites’ fracture surface observation, SEM

images revealed more SD particle pull-out from the

matrix in the composites prepared with the larger SD

particles, indicating to an inferior interfacial adhesion in

these composites. A reduction in total surface area of the

large SD particles to bond with the WG-based matrix

was thought to be one of the major causes.

Acknowledgements

This work was granted by the National Research

Council of Thailand (NRCT) and Mae Fah Luang

University, Chiang Rai, Thailand. The authors wish to

thank the Scientific & Technological Instruments Center

(STIC) staff for their assistances in mechanical tests and

SEM observation.

0.5

0.7

0.9

1.1

1.3

1.5

1.7

1.9

0

2

4

6

8

10

12

14

16

18

20

< 180 µm 180-250 µm 250-550 µm 550-2500 µm

Fle

xura

l m

od

ulu

s (G

Pa)

Fle

xura

l st

rength

(M

Pa)

Sawdust particle size (µm)

Strength

Modulus

0

2

4

6

8

10

12

14

16

18

20

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5

Str

ess

(MP

a)

Strain (%)

< 180 µm

180-250 µm

250-550 µm

550-2500 µm


96

Fig.4 Fracture surfaces of WG-based biocomposites

reinforced with SD particle size of (a) <180 µm, (b) 180–

250 µm, (c) 250–550 µm and (d) 550–2500 µm.

References

[1] Song Y, Zheng Q., “Structure properties of

methylcellulose microfiber reinforced wheat gluten

based green composites”, Industrial Crops and

Products, 446-454(2009).

[2] Beatriz M., Gabriela G., Emmanuelle G. and Patricia

T., “Biocomposites from wheat proteins and fibers:

Structure/mechanical properties relationships”,

Industrial Crops and Products, 545-555(2012).

[3] Beg M., Pickering K. and Weal S., “Corn gluten

meal as a biodegradable matrix material in wood

fibre reinforced composites”, Materials Science and

Engineering A, 7–11(2005).

[4] Kunanopparat T., Menut P., Morel M. and Guilbert

S., “Plasticized wheat gluten reinforcement with

natural fibers: Effect of thermal treatment on the

fiber/matrix adhesion”, Composites: Part A, 1787–

1792(2008).

[5] Kunanopparat T., Menut P., Morel M. and Guilbert

S., “Reinforcement of plasticized wheat gluten with

natural fibers: From mechanical improvement to

deplasticizing effect”, Composites: Part A, 777–

785(2008).

[6] Zhang X., Wu X., Haryono H. and Xia K., “Natural

polymer biocomposites produced from processing

raw wood flour by severe shear deformation”,

Carbohydrate Polymers, 46–52(2014).

[7] Rafieian F., Shahedi M., Keramat J. and Simonsen

J., “Mechanical, thermal and barrier properties of

nano-biocomposite based on gluten and

carboxylated cellulose nanocrystals”, Industrial

Crops and Products, 282–288(2013).

[8] Martinez D., Partal P., Martinez I. and Gallegos C.,

“Gluten-based bioplastics with modified controlled-

release and hydrophilic properties”, Industrial Crops

and Products, 704-710(2012).

[9] Hemsri S, Alexandru D., Grieco K. and Richard S.,

“Biopolymer composites of wheat gluten with silica

and alumina”, Composites: Part A, 1764–

1773(2011).

(d)

(c)

(b)

(a)


97 [10] Hemsri S, Alexandru D., Grieco K. and Richard S.,”

Wheat gluten composites reinforced with coconut

fiber”, Composites: Part A, 1160–1168(2012).

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Parnas, “Toughened wheat gluten and treated

coconut fiber composite”, Composites: Part A, 90–

97(2013).

[12] Reddy N. and Yang Y., “Novel green composites

using zein as matrix and jute fibers as

reinforcement”, bio mass and bio energy, 3496-

3503(2011).

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Siriwattanayotin S., “Effect of lignin removal on the

properties of coconut coir fiber/wheat gluten

biocomposite”, Composites: Part A, 173-179(2010).

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“Optimized material composition to improve the

physical and mechanical properties of extruded

wood–plastic composites (WPCs)”, Construction

and Building Materials, 120–127(2011).

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“Effects of interfacial modification and fiber content

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98

COMPP-11

Rice Husk Reinforced Wheat Gluten-Based Composites: Effect of Particle Size

Supannee Meema, Renuka Pholkaset, Uraiwan Intatha, Nattaya Tawichai, Nattakan Soykeabkaew*

School of Science, Mae Fah Luang University, Chiang Rai 57100, Thailand

Abstract

Rice husk (RH) is a waste from rice milling processes which is abundant in Thailand, low cost,

biodegradable and environmental friendly. Wheat gluten (WG) is also obtained as a byproduct from wet-milling

of wheat flour. In this work, the bio-based composites of RH (50 wt%) reinforced in WG-based matrix were

prepared by thermo-molding at 130°C and 26.6 MPa for 10 min. The effect of RH particle size (i.e., 100-550

µm, 550-1200 µm and the as received RH) on structure and mechanical properties of the resulting composites

was studied. Morphologies of the RH particles and their composites were examined by optical and scanning

electron microscopies. The bulk density of the composites integrated with different RH particle sizes was in the

range of 1.30-1.35 g/cm3. The composite reinforced with the RH particle size of 550-1200 µm presented the

highest flexural strength (14 MPa), modulus of elasticity (1.5 GPa) and strain at break (2.3%). On the other

hand, the composite combined with the as received RH showed the lowest mechanical performance. From its

fracture surface, SEM image revealed more RH particle pull-out, indicating a poor interfacial interaction

between the as received RH particles and WG-based matrix in this composite.

Keywords: bio-based composites, rice husk, wheat gluten, scanning electron microscopy, mechanical properties

1. Introduction

Plastics play an important role in our daily life

and are used in wide-ranging applications. With today’s

technology, we can produce the plastics as required at

low price with a number of good properties, e.g.,

adequate strength, light-weight, heat resistance,

waterproof and available in many colors. However,

synthetic plastics are difficult to decompose (non-

degradable) and recycle, presenting a serious threat to the

environment [1]. Bio-based polymers are now widely

accepted as an alternative to conventional plastics in

some common applications [2]. Researchers as well as

industrials have now focusing on polymers from agro-

resources because utilization of agricultural wastes,

byproducts and co-products as feedstock to develop

biodegradable and renewable materials provides a great

benefit to economics and environment at the same time

[3, 4].

Rice husk (RH) is a waste from rice milling

processes. Therefore, it is highly available in Thailand.

The main components of RH are 25-35% cellulose, 18-

21% hemicellulose, 26-31% lignin, 15-17% amorphous

silica and waxes [5]. RH has low price, low density,

good specific mechanical properties, high toughness and

are fully biodegradable. From these advantages, RH can

be considered as a potential reinforcement in bio-based

composite materials [6, 7].

Wheat gluten (WG) is a complex protein obtained as

a byproduct from wet-milling of wheat crop or wheat

flour in industrial processes and converted to powders

[8,9]. It consists of 75% to 80% of protein, which are

mostly the two protein types, i.e., gliadins and glutenins. WG is readily available in large quantities, relatively low

price, water insoluble, highly viscoelastic, tough,

renewable and can be rapidly degraded [10]. Domenek et

al. (2004) reported that all gluten materials were fully

degraded within 50 days in farmland soil [11]. Thus, WG

has great potential to substitute for conventional plastics

[12].

The purpose of this research was to prepare the bio-

based composites of RH reinforced in WG-based matrix.

The mechanical properties of composite materials are

known to depend on several factors. Type of

reinforcement and its morphological characteristics are

also considered major influences. Therefore, the effect of

size of reinforcing RH particles on structure and

properties of the bio-based composites was investigated

in this study.


99 2. Materials and methods

2.1 Materials

Rice husk (RH) was kindly supplied from the rice

milling in Chiang Rai province, Thailand. Three ranges

of RH particles were used as reinforcements, i.e., as

received RHs, 550-1200 µm and 100-550 µm. The last

two ranges were prepared by reducing the size of the as

received RHs using a kitchen blender (Hongking, QF-

0515SG) and then sieving them into the size ranges.

Next, the RH particles were dried to remove moisture in

a hot-air oven at 80°C for 10 min and kept in an air-tight

container. Vital wheat gluten powder was purchased

from Zhang Jia Gang Heng Feng Co. Ltd, China.

Glycerol of commercial grade was supplied by Union

Science Co., Ltd., Thailand.

2.2 Preparation of bio-based composites

RH particles, wheat gluten and glycerol were

hand-mixed until homogeneous for 5 min. The ratio of

RH particles to wheat gluten and glycerol was 5:5 in

weight, while the ratio of wheat gluten to glycerol was

7:3 in weight. After mixing, the resultant materials were

compression molded (Labtech LP-S-80) at 130°C and

26.6MPa for 10 min. Before testing, the samples were

conditioned at relative humidity of 50±5% in a chamber

of saturated Mg (NO3)2 solution at 25°C (according to

ASTM E104) for 40 h.

2.3 Density

Bulk density of the RH particles of different

ranges of particle size was measured by weighting them

(M) in a known volume container (V). Then, the bulk

density (D) was calculated according to D = M/V in unit

of g/cm3. For bulk density of the bio-based composites,

the cut specimens were weighted using a 4-digit balance

(Denver Instrument, SI-234) and measured for their

dimensions (length, width, and thickness) using a digital

caliper (TCM).The bulk density (D) of the composites

was then calculated according to their weights and

volumes.


The flexural test (4-point bending) of the

composites was performed on a universal testing

machine (Instron 5560) equipped with a load cell of 1 kN

according to the standard method ASTM D6109-13. The

rate of crosshead motion was 2 mm/min. Given flexural

strength, modulus of elasticity, and strain at break (%)

are the means of five replicates for each bio-based

composites. Standard deviation was also calculated and

the results were expressed in mean ±SD.

2.5 Scanning electron microscopy (SEM)

The surface morphology of the RH particles as

well as the fracture surfaces of composites after failure

under flexural test were examined by scanning electron

microscope, LEO 1450 VP, with 10 kV accelerating

voltage. All sample surfaces were sputter coated with

gold before the observation.

3. Results and discussions

3.1 Morphological observation of the RH particles

From optical microscopy, the RH as received was

in yellow-brown color with the average length of 9.7 mm

and width of 3.9 mm (Fig. 1a). From SEM photo (Fig.

2a), it was clear that the outer surface RH particles were

rough with tiny ridges and hairs. The germs were also

observed at the tip of each particle. After the size

reduction, the RH particles in range of 500-1200 µm

mostly appeared to be as the thin flakes with irregular

shapes (Fig. 1b). The inner surface of RH particles was

smooth, but the outer surface was rough. However, they

seemed to be flatter than the as received RH particles and

hair was no longer seen on the outer surface of the

particles (Fig. 2b). On the other hand, the majority of the

RH particles in the range of 100-550 µm were in

rectangular shapes mixed with a few of large flakes (Fig.

1c and Fig. 2c). The size distribution of the RH particles

in this range seemed to be the broadest.


100

Fig. 1 Rice husk (RH); (a) as received, (b) particle

sizes of 550-1200 μm, and (c) 100-550 μm

observed by optical microscopy.

3.2 Bulk density

From Table 1, as the size of RH particles

decreased, their bulk density increased. Typically, the

smaller RH particles could arrange themselves in a closer

fashion.

Fig. 2 Rice husk (RH); (a) as received, (b) particle

sizes of 550-1200 μm, and (c) 100-550 μm

observed by scanning electron microscopy.

The absence of hairs on their outer surfaces of the

RH should be also another reason that allowed a closer

packing among them. This caused the air gaps between

the smaller particles decreased. For the bio-based

composites, after the size reduction step, it was found

that the bulk density of the composites tended to

increase. Possibly, the smaller RH particles could also

packed themselves better in the composites leading to a

reduction of voids in this composite material structures

[13].

(a)

(b)

(c)

(a)

(b)

(c)


101 Table 1: Bulk density of the rice husks and their bio-

based composites

Particle sizes

(µm)

Bulk Density (g/cm3)

Rice husk particles Composites

As received 0.14±0.01 1.30±0.01

550-1200 0.23±0.01 1.31±0.03

100-550 0.38±0.02 1.35±0.04

3.3 Mechanical properties and fracture surfaces

From the flexural tests, the composite prepared

with the RH of particle size 550-1200 μm showed the

highest flexural properties when compared to the

composites prepared with the other two RH sizes; the as

received RH and the RH of particle size 100-550 μm

(Fig. 3).

Fig. 3 Mechanical properties the bio-based composites

prepared with RH particles of different sizes.

Fig. 4 Stress-strain curves the bio-based composites

prepared with RH particles of different size.

Fig. 5 Fracture surfaces of the bio-based

composites prepared with (a) as received RH, (b)

RHs of particle size 550-1200 μm, and (c) 100-550

μm.

The composite prepared with the as received rice

husks had the lowest performance and, from its stress-

strain curve, the earliest material failure was also found

(Fig. 4). Perhaps, the roughest surface with hairs of the

as received RHs inhibited their good bonding with the

matrix phase [14], resulting in a poor interfacial

interaction between the two components in the composite

structure [15]. In line with the SEM results, the most

pull-out of RH particles with deep holes were observed

on the fracture surface of this composite (Fig. 5a) as

0

2

4

6

8

10

12

14

16

0 1 2 3 4

Str

ess

(MP

a)

Strain (%)

As received

500-1200 m

100-550 m

(a)

(b)

(c)

µ

µ


102 compared to those of the other composites (Fig. 5b and

c).

Lastly, the composite prepared with the RH of

particle size 100-550 μm presented a slightly lower

flexural strength and modulus of elasticity than the one

prepared with particle size 550-1200 μm. This was

possibly due to the broad size distribution of the RH

particles in the range of 100-550 μm as previously seen

in Fig. 1c and Fig. 2c [16].

4. Conclusions

The bio-based composites reinforced with RH

particles in different sizes were prepared by thermo-

molding technique. The as received RHs were in yellow-

brown color and their outer surfaces were rough with

tiny ridges and hairs, whereas the inner surfaces of RH

particles were smooth. After the size reduction step, the

outer surface of RH particles became less rough with no

hair presented. The bulk density of the bio-based

composites prepared with different sizes of RH particles

was similar in the range of 1.30-1.35 g/cm3. From

mechanical test results, the composite reinforced with the

RH of particle size 550-1200 µm exhibited the highest

flexural properties. Meanwhile, the composite prepared

with the RH of particle size 100-550 µm presented a

slightly lower flexural strength and modulus possibly due

to the broad size distribution of the RH particles in this

range. On the other hand, the composite combined with

the as received RHs showed the lowest mechanical

performance. Perhaps, the roughest outer surface with

the presence hairs of the as received RHs inhibited their

good bonding with the WG-based matrix. SEM images

also revealed more pull-out of RH particles on the

fracture surface of this composite, indicating an inferior

interface when compared to the other two composites

prepared with the smaller RH particles.

Acknowledgements

This work was granted by the National Research

Council of Thailand (NRCT) and Mae Fah Luang

University, Chiang Rai, Thailand. The authors wish to

thank STIC staff for their assistances in mechanical tests

and SEM characterization.

References

[1] Mohanty A.K., Misra M., Drzal L.T. Sustainable

biocomposite from renewable resources:

Opportunities and challenges in the green materials

world, Journal of polymer and the environment, 10:

19-26 (2002).

[2] Duval A., Molina-Boisseau S,. Chirat C.

Comparison of Kraft lignin and lignosulfonates

addition to wheat gluten-based materials:

Mechanical and thermal properties, Industrial Crops

and Products, 49: 66-74 (2013).

[3] Reddy N., Yang Y. Novel green composites using

zein as matrix and jute fibers as reinforcement, Bio

mass and bio energy, 35: 3496- 3503 (2011).

[4] Zhang X., Wu X., Haryono H., Xia K. Natural

polymer biocomposites produced from processing

raw wood flour by severe shear deformation,

Carbohydrate Polymers, 113: 46-52 (2014).

[5] Johnson A.C., Yunus N.B. Particleboards from Rice

Husk: A Brief Introduction to Renewable Materials

of Construction. P. 12-15 (2009).

[6] Hemsri S., Grieco K., Asandei A.D., Parnas R.S.

Wheat gluten composites reinforced with coconut

fiber, Composites: Part A, 43: 1160-1168 (2012).

[7] Monta˜no-Leyva B., Silva G.G.D., Gastaldi E.,

Torres-Chávez P., Gontard N., Angellier-Coussy H.

Biocomposites from wheat proteins and fibers:

Structure/mechanical properties relationships,

Industrial Crops and Products, 43: 545-555 (2013).

[8] Yuan Q., Lu W., Pan Y. Structure and properties of

biodegradable wheat gluten/attapulgite

nanocomposite sheets, Polymer Degradation and

Stability, 95: 1581-1587 (2010).

[9] Zárate-Ramírez L.S., Romero A., Bengoechea C.,

Partal P., Guerrero A. Thermo-mechanical and

hydrophilic properties of polysaccharide/gluten-

based bioplastics, Carbohydrate Polymers, 112: 24-

31 (2014).


103 [10] Day L. Wheat gluten: production, properties and

application, Food and Nutritional Sciences, 10: 267-

288 (2006).

[11] Domenek S., Feuilloley P., Gratraud J., Morel M.,

Guilbert S. Biodegradability of wheat gluten based

bioplastics, Chemosphere, 54: 551-559 (2004).

[12] Diao C., Dowding T., Hemsri S., Parnas R.S.

Toughened wheat gluten and treated coconut fiber

composite, Composites: Part A, 58: 90-97 (2014).

[13] Tiago J.C.D. Effect of Rice Husk Ash Particle Size

in Lime Based Mortars, Instituto Superior Técnico,

P: 1-11 (2011).

[14] Sumaila M., Amber I., Bawa M. Effect of fiber

length on the physical properties and mechanical

properties of random oriented, nonwoven short

banana (musa blabisiana) fibre/ epoxy composite,

Mechanical Engineering Department, 2: 39-46

(2013).

[15] Syafri R., Ahmad I., Abdullah I. Effect of Rice Husk

Surface Modification by LENR the on Mechanical

Properties of NR/HDPE Reinforced Rice Husk

Composite, Sains Malaysiana, 44: 749-756 (2011).

[16] Zhang Y., Ghaly A.E., Li B. Physical properties of

Rice residues as affected by variety and climatic and

cultivation ondition in three conditions, American

Journal of Applied Sciences, 9: 1757-1768 (2012).


104

COMPP-12

Mechanical Properties of Poly(butylene succinate) Films Reinforced with Silica

Nanthaporn Sangviroon1 and Pranut Potiyaraj

2*

1Interdisciplinary Program in Petrochemical & Polymer Science, Faculty of Science, Chulalongkorn University,

Bangkok 10330 2Department of Materials Science, Faculty of Science, Chulalongkorn University, Bangkok 10330

Abstract

Poly(butylene succinate) or PBS is a biodegradable polymer which has good processability, chemical

resistance and environmental friendly However, the application was limited by its mechanical properties. The

aim of this research is to improve mechanical properties of PBS bioplastic films by reinforcing with silica. The

composite films were prepared with the presence of glycidyl methacrylate grafted poly(butylene succinate)

(PBS-g-GMA) as a compatibilizer at a fixed ratio of 5 wt%. PBS and silica were mixed in a twin screw extruder

at the amount of silica between 0-3 wt%. The obtained compounds were then processed into films by a chill roll

cast extruder. The effects of compatibilizer and silica loading on mechanical properties of the prepared

composite films were investigated. It was found that the mechanical properties of PBS/silica composite films

were improved when 1%wt of silica was added. However, the mechanical properties decreased with increasing

silica loading due to the agglomeration of silica particles. The results also indicated that the PBS/silica

composite films with the presence of PBS-g-GMA possessed improved mechanical properties over the films

without the compatibilizer.

Keywords: Poly(butylene succinate), Silica, Composite, PBS-g-GMA, Mechanical properties

Introduction

Nowadays, biodegradable polymeric materials

attract much more attention due to the fact that plastic

waste has caused the environment pollution.

Poly(butylene succinate) or PBS is an aliphatic

biodegradable polyester which is now commercially

available. PBS is synthesized by polycondensation of

succinic acid with 1,4-butanediol. Despite its good

processablity and chemical resistance, its softness and

poor barrier gas thus limit the applications of PBS. [1]

Some researchers have thus attempted to improve

properties of PBS by mixing with reinforcing filler i.e.

silica. However, it is difficult to disperse the silica

particles in polymer matrix homogeneously because silica

particles have strong tendency to agglomerate, which is

caused by the poor interaction between silica particles

and polymer matrix, resulting in reduced mechanical

properties. [2]

One way to overcome these drawbacks and to

prepare materials with enhanced properties is to

incorporate a compatibilizer in order to enhance the

dispersibility of silica particles and improve interfacial

adhesion between silica particles and polymer matrix. [3]

Glycidyl methacrylate grafted polymers have been

used as compatibilizer in some previous research because

of its epoxide functional groups, which is highly

electrophilic and capable of reacting with a variety of

functional groups such as carboxylic acids, amides, and

alcohols. [4]

In this study, PBS-g-GMA was prepared by the

reactive melt-grafting extrusion method. Then composites

were processed into films by a chill roll cast extruder.

The effects of PBS-g-GMA and silica loading on

mechanical properties of PBS/silica composite films were

investigated.

Experimental

Materials

PBS granules (GS PlaTM

FZ91PD) were of film

grade purchased from Mitsubishi Chemical. Silica

(ULTRASIL® 9000 GR) was kindly supported by Evonik

with a specific surface area of 235 m2/g and an average

particle size of 13.3 Dicumyl peroxide (DCP) (Sigma

Aldrich) was used as an initiator and glycidyl

methacrylate (GMA) (Sigma Aldrich) was used as a

reagent without further purification.


105 Preparation of PBS-g-GMA

Firstly, PBS, GMA and DCP were physically

premixed. The reactive grafting process was carried out

in a twin-screw extruder (Prism TSE 16 TC, Thermo,

UK) at 140°C at 30 rpm. The amount of GMA and DCP

used were fixed at 10 and 2 phr respectively. The degree

of grafting of GMA onto PBS is 2.95 % as determined

through the titration method.

Preparation of Composite Films

PBS, PBS-g-GMA pellets and silica were initially

dried in the oven at 60ºC for 24 h prior to further

processing. Polymers were physically premixed at the

ratios shown in Table 1 and melt-mixed in a twin-screw

extruder at 160°C with a screw speed of 30 rpm. Then,

the pelletized compounds were dried in the oven at 60ºC

for 24 h before casting into film by chill-roll cast extruder

(LCR-300HDCO-EX, Labtech Engineering, Thailand).

Characterization

FT-IR Analysis

The Fourier transform infrared spectroscopy

analysis was carried out on the PBS-g-GMA at ambient

temperature by using a Perkin-Elmer One FT-IR

Spectrometer (USA). It was performed through the

scanning wavenumber from 4000 to 400 cm-1

with

resolution of 64 bit.

Mechanical Testing

Tensile properties that are tensile strength, tensile

modulus and elongation at break were measured

according to ASTM D882 using a Universal Testing

Machine (LLOYD LR100K, West Sussex, UK) with a

gauge length of 125 mm and using the cross head speed

of 12.5 mm/min and a 100 N load cell. The tear strength

was measured according to ASTM D1938 using a

Universal Testing Machine (LLOYD LR100K, West

Sussex, UK) with the cross head speed of 250 mm/min.

and a 100 N load cell.

Morphological

To reveal the dispersion of silica particles in the

PBS matrix, the fractured samples from tensile testing of

the neat PBS, PBS/silica and PBS/PBS-g-GMA/silica

composite films were sputter-coated with a thin layer of

gold and then examined for morphological structure by

Scanning Electron Microscope (SEM, QUANTA 250,

FEI, USA) operated at 20 kV with resolution of 3000×.

Table 1. Composition of PBS composite compounds.

Sample (%wt) PBS PBS-g-GMA Silica

PBS 100 - -

PBS/G5 95 5 -

PBS/S1 99 - 1

PBS/S2 98 - 2

PBS/S3 97 - 3

PBS/G5/S1 94 5 1

PBS/G5/S2 93 5 2

PBS/G5/S3 92 5 3


106 Results and Discussion

Fig. 1. FT-IR spectra of neat PBS and PBS-g-GMA

after purification.

FT-IR spectra

FTIR spectra for neat PBS and PBS-g-GMA are

shown in Fig. 1. Characteristic peaks for PBS at 3,200–

3,700, 1,700–1,750, and 500–1,600 cm-1

appeared in both

polymers. Two extra shoulders characteristic of ester

carboxyl groups were observed at 1,731 cm-1

in the

modified PBS-g-GMA spectrum. The shoulders represent

free acid in the modified polymer PBS-g-GMA and thus

indicated the successful grafting of GMA onto PBS. [5]

Mechanical Testing

The mechanical properties including tensile

strength, Young’s modulus, elongation at break and tear

strength of PBS and its composite films were displayed in

Fig. 2. and 3. The composite films around 0.12 mm thick

were tested along in the machine direction (MD) and in

the transverse direction (TD).

For the composite films without PBS-g-GMA,

the tensile strength and tensile modulus increased with

the addition of 1%wt of silica. However, as then amount

of silica increased, tensile strength and tensile modulus

were reduced. The elongation of composite films was

lower than neat PBS at every ratio. Due to the

aggregation of silica particles, interfacial adhesion

between silica particles and PBS was weakened thus

contribute to crack propagation and hence potential

composite failure. [6]

Scheme 1. Grafting reaction of GMA onto PBS.

When the PBS-g-GMA was incorporated, tensile

strength, tensile modulus as well as elongation at break of

composite films were higher than those of neat PBS film.

The tear strength of composite films was lower than that

of neat PBS film except for the films with 1%wt silica

which possessed a slightly higher tear strength comparing

with neat PBS.

The addition of PBS-g-GMA can improve

interfacial adhesion by forming covalent bonds between

silica particles, hence PBS phase can be transfers stress to

silica particles resulting increased mechanical properties.

[7] The grafting reaction of GMA onto PBS is illustrated

in Scheme 1.

The influence of orientation on the mechanical

properties was investigated. Tensile strength and

elongation at break in the MD were higher than in the TD

for all samples tested due to preferential molecular

orientation in the MD. In the MD of stretching covalent

bonding is predominant, while the TD van der Waal will

be dominant bonding type [8], thus the MD of stretching

required a higher force to propagate a crack.

On the other hand, tear strength in the TD was

higher than in the MD for all samples tested because

preferential molecular orientation in the MD that vertical

to the test, leading hard to tear and yields higher tear

strength.

O

CH2

CH2

CH2

CH2

O

C O

CH2

CH2

C O

n

O

CH2

CH2

CH

CH2

O

C O

CH2

CH2

C O

n

CH2

C CH3

C O

O

CH2

CH

CH2

O

O

CH2

CH2

CH

CH2

O

C O

CH2

CH2

C O

n

CH2 CH

CH3

C O

O

CH2

CH

CH2

O

C

CH3

CH3

O O C

CH3

CH3

C

CH3

CH3

O

+

DCP

PBS PBS-g-GMAGMA


107

Fig 2. Mechanical properties of PBS and their composite films. (A) Tensile strength, (B) Young’s modulus

and (C) Elongation at break.

Fig 3. Tear strength of PBS and their composite films.


108

Fig. 4. SEM micrographs of the fracture surface of the PBS/silica composite films at 2 wt% (A) without PBS-g-GMA

(B) with PBS-g-GMA and (C) neat PBS

Morphological analysis

Fig 4. presents the SEM micrographs which

were obtained on tensile fracture surfaces of the

composite films. Generally, hydrophilic silica particles

easily aggregated because of particle-particle interaction,

and the aggregated silica particles were also found in the

PBS/S2 composite films as shown in Fig. 4A While Fig.

4B shown the PBS-g-GMA was introduced to promote

better silica dispersion resulting mechanical properties

were increased. The possible reason may be that the

reaction occurred between the epoxy group of PBS-g-

GMA and silanol groups on silica particles, which is also

effective in improving the compatibility of silica

particles and PBS phase. [9]

Conclusion

PBS/silica composite films were prepared at

various ratios of the filler and their mechanical properties

were investigated. It was found that the mechanical

properties of PBS/silica composite films were improved

when 1%wt of silica was added. However, the

mechanical properties decreased with increasing silica

loading due to the agglomeration of silica particles. The

results also show that improvements in the mechanical

properties were obtained when PBS-g-GMA was used as

a compatibilizer as the filler dispersion and filler-matrix

interfacial interactions were enhanced as observed by

SEM technique.

Acknowledgement

This research has been supported by the

Ratchadaphiseksomphot Endowment Fund 2013 of

Chulalongkorn University (CU-56-416-AM).

References

[1] Bian, J., Han, L., Wang, X., Wen, X., Han, C.,

Wang, S., and Dong, L., “Nonisothermal

crystallization behavior and mechanical properties

of poly (butylene succinate)/silica nanocomposites”

Journal of Applied Polymer Science, 116(2): 902-

912 (2010).

[2] Lim, J. S., Hong, S. M., Kim, D. K., and Im, S. S.,

“Effect of isocyanate - modified fumed silica on the

properties of poly (butylene succinate)

nanocomposites” Journal of Applied Polymer

Science, 107(6): 3598-3608 (2008).

[3] Vassiliou, A. A., Chrissafis, K., and Bikiaris, D. N., “In situ prepared PBSu/SiO2 nanocomposites. Study

of thermal degradation mechanism” Thermochimica

Acta, 495(1): 120-128 (2009).

[4] Papke, N., and Karger-Kocsis, J., “Determination

methods of the grafting yield in glycidyl

methacrylate-grafted ethylene/propylene/diene

rubber (EPDM-g-GMA): Correlation between FTIR

and 1H-NMR analysis”

Journal of applied polymer science, 74(11): 2616-2624

(1999).

[5] Wu, C. S., Liao, H. T., and Jhang, J. J., “Palm fibre-

reinforced hybrid composites of poly(butylene

succinate): characterisation and assessment of

silica agglomerate

(A) (B) (C)


109 mechanical and thermal properties” Polymer

Bulletin, 70(12): 3443-3462 (2013).

[6] Han, S. I., Lim, J. S., Kim, D. K., Kim, M. N., and

Im, S. S. “In situ polymerized poly (butylene

succinate)/silica nanocomposites: physical

properties and biodegradation” Polymer

Degradation and Stability, 93(5): 889-895 (2008).

[7] Karger-Kocsis, J., and Fakirov, S., “Nano- and

Micro-mechanics of polymer blends and

composites” Hanser: 107-113 (2009).

[8] Klein, R. “Material properties of plastics” Laser

Welding of Plastics: Materials, Processes and

Industrial Applications: 3-69 (2011).

[9] Xiuju, Z., Juncai, S., Huajun, Y., Zhidan, L., and

Shaozao, T. “Mechanical properties, morphology,

thermal performance, crystallization behavior, and

kinetics of PP/microcrystal cellulose composites

compatibilized by two different compatibilizers”

Journal of Thermoplastic Composite Materials,

24(6): 735-754 (2011).


110

COMPP-13

Tobacco Cellulose Nanofibril As Filler For Poly(Lactic Acid)/Poly(Vinyl Acetate)

Blends

Pitchayasinee Komontreea, Krisana Siralertmukul

b, Kawee Srikulkit

c

aMajor of Petrochemistry and Polymer Science, Faculty of Science, Chulalongkorn University, Bangkok,

Thailand bMetallurgy and Materials Science Research Institute (MMRI), Chulalongkorn University, Bangkok, Thailand

cDepartment of Materials Science, Faculty of Science, Chulalongkorn University, Bangkok, Thailand

Abstract This work presented the preparation of tobacco cellulose nanofibril and its dispersibility of cellulose

nanofibrils (CNB) in poly(lactic acid)/poly(vinyl acetate) blend or PLA/PVAc blend. In this experiment,

cellulose nanofibrils was prepared by acid hydrolysis of tobacco cellulose in the presence of chitosan as an

anticoagulant. The addition of CNB into PLA/PVAc blend was carried out by solution casting technique.

Compatibility study between CNB and PLA/PVAc was investigated by scanning electron microscopy, Fourier

transform infrared spectroscopy (FTIR), and Thermogravimetric analysis (TGA).

Keywords: Cellulose nanofibrils, Poly(lactic acid)/poly(vinyl acetate) blend

1. Introduction

At the presence, the world is facing the serious

problem concerning an environmental issue. Of course,

petroleum plastics are the major culprit due to the fact

that they are extremely inert and non-degradable. The

combustion produces toxic gas including dioxin and

carbon monoxide. Plastic buried deep in landfills can

leach harmful chemicals that contaminate groundwater.

Thus, nowadays researchers are interested in

biodegradable plastics particularly poly(lactic acid)

(PLA) and polybutylene succinate (PBS). PLA is

produced from lactic acid obtained by microbial

fermentation of agricultural products mainly the

carbohydrate such as starch from potato, corn[1]. Lactic

acid is the most widely occurring hydroxyl carboxylic

acid due to its versatile application in food,

pharmaceutical, textile.[2]

Figure1 Chemical structure of lacticacid[3]

Lactic acid containing bifunctional groups

( hydroxyl group and a carboxyl group) (Figure1)is able

to undergo self-condensation reaction in the presence of a

catalyst, resulting in aliphalic polyester.[4]

The most common way to obtain high-molecular-

weight poly(lactic acid) is through ring-opening

polymerization of lactide monomer, a reactive cyclic

intermediate.[5]

Figure 2 Polymerization of PLA[6]

Poly(lactic acid) is a clear, colorless

thermoplastic.[8]The weakest characteristic of PLA

includes brittleness and poor thermal resistance.

Therefore, mechanical properties of processed PLA are

far more interior to petroleum plastics like PE and PP

products.

To improve the properties and reduce the cost of

production, polymer blending with the addition of filler is


111 one of research approaches.[9] In this work, the research

interest is focused on the utilization of cellulose

nanofibril as a reinforcing material to improve properties

of PLA. However, cellulose nanofibrils (Figure 3) by

nature tend to strongly adhere each other through

intermolecular hydrogen bonding, causing hard aggregate

which is known as a bad reinforcing filler.

Figure3 Cellulose molecule [10]

In this research, cellulose nanofiber in wet gel

form was applied directly through solution casting. The

dispersion of cellulose nanofibril in PLA/PVAc blend

was investigated in details.

2. Experimental

2.1Chemicals and Materials; Chemicals and materials

included tobacco stalk, poly(lactic acid), Poly(vinyl

acetate), ethanol (20 %w/w) sodium

hydroxide(10%w/w), sodium hypochlorite,(NaClO

6 %w/w), hydrogen peroxide(H2O2) 30 %w/w, methanol,

sulfuric acid (H2SO4), deionized water, and

dichloromethane(CH2Cl2).

2.2Characterization;

2.2.1 Fourier transform infrared spectroscopy

(FTIR) - Identification of functional groups of the

polymers were confirmed using Nicolet 6700 FTIR.

2.2.2Scanning electron microscope(SEM)-The

surface morphology of sample was characterized using

scanning electron microscope (SEM) on a Joel (JSM

6400)

2.2.3 Thermogravimetric analysis (TGA)-The

TGA was conducted from 30 oC to 600

oC using a

METTLER TOLED TGA/SDT851e with a heating rate of

10oC/min under 20 mL/min nitrogen and oxygen gas

flow.

2.3 Extraction of cellulose from tobacco

The tobacco stalk 20 kg.was cut to 1.5 inch in

length and dried by sunlight for 15 days and then treated

in 20 L. of 20% (w/w) ethanol 20 L for 15 days. The soft

tobacco stalk was boiled in 10% w/w of NaOH3L for 2 h

at 108 oC and followed washing by water until pH value

became neutral. Next, cellulose fiber was treatedin

10L.NaClO 6 % (w/w) at room temperatureovernight and

washed by water until pH value becameneutral.Further

treatment of NaClO bleached cellulose pulp was carried

using 10 L of 30% (w/w) H2O2 overnight. Purified

cellulose pulp was obtained. To avoid aggregation,

solvent exchange with methanol was conducted twice

before drying at ambient temperature.

2.4 Preparation of cellulose nanofibrils.

Cellulose nanofibril was prepared by acid

hydrolysis 75% (w/w) H2SO4 was prepared and kept at

zero temperature overnight. Then, cellulose pulp 20 g

was added into a beaker containing 400 mL of 75%

(w/w) H2SO4. The beaker was put into an ice bath to

control the temperature of acid hydrolysis reaction and to

prevent the degradation reaction of cellulose. After being

left standing for 12 h, cellulose nanofibril was obtained

and washed thoroughly until pH value became neutral.

Cellulose nanofibril in gel form was achieved.

2.5 PLA/PVAc Blend in the Presence of CNB

PLA/PVAc blends having various CNB loadings

were prepared according to the composition shown in

Table 1. Firstly, PVAc was well-mixed with CNB to

obtain good dispersion. Then, CH2Cl2 solubilized PLA

was added and mechanically mixed. The mixture was cast

into petridisk and solvent was allowed to evaporate.

After complete drying, PLA foam was obtained.


112 Table1Ratio of PLA/PVAc/Cellulose

3.Result and discussion

3.1 Tobacco Cellulose Nanofibril

Figure4 The picture of (A) Tobacco stalk (B) Cellulose

pulp (C) Cellulose gel (cellulose nanofibril) (D) SEM of

cellulose nanofibril

By sequential treatments according to sections 2.3

and 2.4, cellulose nanofibril present in gel form (Figure

4C) is achievable. By using SEM analysis, densely

packed nanofibrils are observed (Figure 4D). Percent

yield was measured to be about 17.301%.

3.2 Characterizations of PLA/PVAc blends in the

presence of CNB

3.2.1Fourier transform infrared spectroscopy (FTIR)

FTIR spectra of PLA/PVAc blends with CNB

loadings are shown in Figure 4. Characteristic peaks for

PLA at about 1756 and 1180 cm-1

which are belong to

the C=O and C-O-C stretching of PLA. The characteristic

absorption band of cellulose appears at 1735 cm-1

,

representing of C=O The IR bands of PVAc at about

1749 cm-1

and 1370 cm-1

correspond to C=O stretching,

and CH3, respectively. These absorption bands decrease

gradually with a decrease in the fraction ratio of PVAc.

Figure 5FTIR of (A)PLA2:PVAc1:CNB1 (B)PLA1.5:PVAc1:CNB1 (C)

PLA1:PVAc1:CNB1 (D) PLA2:PVAc0.75:CNB1 (E) PLA2:PVAc0.5:CNB1

(F)PLA2:PVAc0.25:CNB1

3.2.2 Scanning electron microscope(SEM)

SEM images of PLA/PVAc blends are presented in

Figure 6. As seen, cellulose nanofibril is well mixed with

PLA/PVAc blends in all case, judged by smooth surface

of SEM images. Further TGA evidence is provided to

support this claim. It is thought that PVAc acts as

compatibilizer or coupling agent between hydrophobic

PLA and hydrophilic CNB since PVAc by nature is

compatible with both PLA and cellulose nanofibril.

Sample PLA

(g)

PVAc(g) Cellulose(g)

PLA2:PVac1:CNB1 200 100 100

PLA1.5:PVac1:CNB1 150 100 100

PLA1:PVac1:CNB1 100 100 100

PLA2 :PVac0.75:CNB1 200 75 100

PLA2:PVac0.50:CNB1 200 50 100

PLA2:PVac0.25:CNB1 200 25 100


113

Figure 6 SEM images of (A)PLA1:PVAc1 (B)PLA2:PVAc1:CNB1 (C)PLA1.5:PVAc1:CNB1 (D) PLA1:PVAc1:CNB1 (E) PLA2:PVAc0.75:CNB1 (F)

PLA2:PVAc0.5:CNB1 (G)PLA2:PVAc0.25:CNB1


114 3.2.3 Thermogravimetric analysis (TGA)

Figure 7FTIR of (A)PLA2:PVAc1:CNB1 (B)PLA1.5:PVAc1:CNB1 (C) PLA1:PVAc1:CNB1 (D) PLA2:PVAc0.75:CNB1 (E) PLA2:PVAc0.5:CNB1

(F)PLA2:PVAc0.25:CNB1

Table2 the Tonset and Td of PLA/PVac blend in the presence of CNB

Sample

Tonset (oC)

Td(oC)

PLA2:PVAc1:CNB1 310.99 354.78

PLA1.5:PVAc1:CNB1 325.78 357.35

PLA1:PVAc1:CNB1 320.42 353.17

PLA2 :PVAc0.75:CNB1 328.55 356.61

PLA2:PVAc0.50:CNB1 322.93 350.33

PLA2:PVAc0.25:CNB1 339.58 363.61


115 Figure 3 shows TGA thermograms of PLA/PVAc

blends at various CNB loadings. Thermograms indicate

that percent weight loss curve exhibits a single step of

decomposition, indicating the homogenous blending.

Therefore, TGA results confirm that PLA, PVAc and

CNB could be homogenously mixed together.

Conclusion

From the experimental, the results showed that

tobacco cellulose nanofibril was achieved. Thus prepared

CNB was well-dispersed in PLA/PVAc blend. It was

believed that PVAc acted as a good compatibilizer.

Acknowledgment

The researchers are thankful to Department of

material science, Chulalongkorn university.

References

[1] John. R.P., Nampoothiri, A., 2006, Solid-state

fermentation for L-lactic acid production from agro

wastes using Lactobacillus using Lactobacilli. Appl.

Biochem (41) 759 – 763

[2] Vickroy, T. B., 1985. Lactic acid. In: Moo-Young,

Comprehensive Biotechnol. Pub: DicToronto:

Pergamon Press, pp. 761 – 776

[3] Nampoothiri, K. M., Nair, A. R., John, R. P., 2010,

An overview of the recent developments in

polylactide(PLA) research. BioresourceTehnology

(101) 8493 – 8501

[4] Sarasua, J. R., Prud’homme, R. E., Wisniewski, M.,

LeBorgne, A., Spassky, N., 1998. Crystallization

and melting behavior of polylactides.

Macromolecules (31) 3895 – 3905

[5] Gupta, A. P., Kumar, V., 2007. New emerging

trends in synthetic biodegradable polymer –

polylactide: a critique, Eur. Polym. J. (10) 4053 –

4074

[6] Averous, L., 2008, Polylacitc acid: synthesis,

properties and applications.(21) 433 – 435

[7] Schneider and Wilmington, 1955. Properties of high

melting lactide. US Patent (2) 771 – 777

[8] Auras, R., Harte, B., Selke, S., 2004, An overview

of polylactides as packaging materials.

Macromol.Biosci (4) 835 – 864

[9] Mohanty, A. K., Misra, M., Hinrichsen, G. 2000

Biofibers. Biodegradable polymers and

biocomposite: An overeview. Macromolecular

Materials and Engineering (276/277) 1 – 24

[10] Siqueira, G., Bras, J., Dufresne, A., 2010 cellulosic

Bionanocomposites: A Review of Preparation,

Properties and Applications. Polymers (2) 728 – 765

[11] Berglund, L., 2005 Natural fibres, Biopolymers and

Biocomposite. (pp. 808) CRC. Press

[12] Chabba. S., Mathews, G. F., Netravali, A. N., 2005

Green composites using cross-linked soy flour and

flax yarns. Green Chemistry, (2) 576 – 581

[13] Dufresne, A., Cavaille, J., Y., Vignon, M., R., 1997

Mechanical behavior of sheets prepared from sugar

beet cellulose microfibrils. Journal of Applied

Polymer Science (64) 1185 - 1194


116

COMPP-18

Effect of Fiber Treatment of oil Palm Mesocarp Fiber on the Properties of Wood

Composites

Polphat Ruamcharoen1*

,Chor Wayakorn Phetphaisit2 and Jareerat Ruamcharoen

3

1Rubber and Polymer Technology Program, Faculty of Science and Technology,

Songkhla Rajabhat University, Songkhla, 90000, Thailand 2Department of Chemistry, Faculty of Science, Naresuan University, Pitsnulok,65000, Thailand

3Department of Science, Faculty of Science and Technology, Prince of Songkla University,

Pattani campus, Pattani, 94000, Thailand

Abstract

This study focused on the effect of fiber treatment of oil palm mesocarp fiber (OPMF) on the properties

of OPMF-unsaturated polyester composites. OPMF has been treated with 5% sodium hydroxide solution for 24,

48 and 72 h. The fibers were first mixed with unsaturated polyester resin and shaped by compression moulding

technique. FTIR analysis for OPMF indicated that lignin and residue palm oil had been removed and SEM

micrographs showed the roughness of fiber surface. The change of physical properties with treatment time

related to the change in surface morphology. It was also found that the wood composites with 72 h treated

OPMF gave the best properties.

Keywords: oil palm kernel fiber, composite, treatment.

1. Introduction

Natural fibers are considered to be reinforcing

fillers in polymers. The possibility of using natural

fibers as reinforcement in composites has yielded to

studies indicating many advantages such as good

mechanical performances, low density and

biodegradability. In this study, the reinforcing filler used

was oil palm mesocarp fiber (OPMF). OPMF consists of

about 60% of cellulose and 11% of lignin [1]. Generally

OPMF is a waste material after oil extraction. This waste

material creates a significant environmental problem.

Currently, OPMF is used as a mulching medium, a boiler

fuel source, and as a fiber source for composite used in

furniture. Thus finding new useful utilization of the

OPMF will surely alleviate environmental problems

related to the disposal of oil palm wastes.

Unsaturated polyesters are extremely versatile in

properties and applications and have been a popular

thermoset used as the polymer matrix in composites.

They are widely produced industrially as they show

many advantages compared to other thermosetting resins

including low temperature cure capability, good

mechanical properties and transparency, and adaptability

to be transformed into large composite structures due to

no by-product is formed during the curing reaction.

The reinforcement of polyesters with cellulosic

fibers has been widely reported, for instance, polyester-

jute [2,3],polyester-sisal [4], polester-coir [5]. In general,

utilization of biomass in lignocellulosic composites has

been attributed to several advantages such as low

density, biodegradability and low cost. However, in

producing a good lignocellulosic composite, the main

obstacle to be

resolved is the compatibility between the fiber and

matrix. In addition, disadvantage of this material is oil

palm residue on the surface and smoothness of surface.

Chemical modifications lead to major changes in

the fibrillar structure of fibers and remove the amorphous

components causing changes in the deformation

behavior. In this paper, we present the influence of fiber

surface treatment by sodium hydroxide solution of

OPMF on the fiber structure and related physical

properties of OPMF composites.

2. Experimental Methods

2.1 Materials


117 The OPMFs were obtained from United Palm Oil

Industry Public Company Limited. Sodium hydroxide

was supplied by BDH Laboratory. Unsaturated polyester

(ER 2428AP-12) with cobalt naphthenate as an

accelerator was obtained from Eternal Resin Co. Ltd.

2.2 Fiber preparation and wood composite preparation

OPMFs were washed, dried and ground into small

Particles. The sieve (Retsch model AS200) was used to

separate the particles into different sizes. The filler sizes

were between 106 and 500 m. The OPMF were

immersed in 5% sodium hydroxide solution for 24, 48

and 72 h. Fibers were then washed with distillated water

until the pH of washed water was approximately 7 and

then were dried in an oven at 105oC for approximately

24 h. The dried fiber and binder were mixed together

with a desired ratio in a 40:60 by weight. Each of

formulated samples was hot-pressed in a compression

moulding machine (Gotech model GT-7014-A10C) at

75oC for 30 min with a pressure of 50 kg.cm

-2. The

samples were then post-cured in an oven at 100oC for 24

h.

2.3 Characterization and Mechanical Testing

Fourier transform infrared (FTIR) spectroscopy

(Shimazu 8900 series) with KBr method was used to

characterize the fibers. The scanning electron

microscopy of OPMFs was performed with Jeol (JSM-

5800LV). The dried samples were sputter-coated with

gold prior to SEM examination. For mechanical

testing, the prepared composites were cut into tensile and

flexural test samples. Tensile tests and flexural test were

conducted according to ASTM D3039 and ASTM D790

respectively using Universal Testing machine (Instron

5567) at a cross-head speed of 5 mm.min-1

. For flexural

properties, the speed was 2.8 mm.min-1

.


3.1 Chemical structure and morphology of OPMF

The scanning electron micrographs of the

untreated and treated fibers surfaces with various times

revealed the changes in the porous structure of fibers as

shown in Fig.1. It corresponded to the previous

reports[1,6]. This could be explained that the treatment

results in the leaching out of the amorphous waxy cuticle

layer [1,6,7] including with lignin [8].

(a)

(b)

Fig.1 Scanning electron micrographs of OPMF (x150),

(a) untreated (b) 72h.

Fig. 2 shows the FTIR spectra for OPMF

(Fig.2a) and treated OPMF (Fig.2b).

Fig. 2 FTIR spectra of OPMF(a) and treated EFB(b).

In Fig.2(a), the bands at 3450 cm-1

was due to the

OH stretching, 1728 cm-1

for C=O stretching of oil

coated on the fiber surface1, 1436 cm

-1 and 1254 cm

-1 for

bending of CH2 in lignin [9,10] and for stretching of C-O

of acetyl in lignin were obtained. After fiber treatment

(Fig.2(b)), the peak at 1728 cm-1

including with 1436 cm-

1728 cm-1

1436 cm-1

1254 cm

-1 3450 cm

-1

3422 cm-1

1045 cm-1


118 1 and 1254 cm

-1 was disappear. This result confirmed that

the treatment of OPMF was resulted in lignin and

remaining oil elimination.

3.2 The Effect of Fiber treatment on the properties of

Wood composite

It can be seen in Table 1 that density slightly

increased while the water absorption and thickness

swelling decreased with treatment time. This could be

due to the greater adhesion between polyester resin and

cellulose. This corresponded to the previous reports

[11,12]. Jacob and coworkers reported that the treatment

of sisal fibers resulted in the decrease in water absorption

[13].

Table 1 Physical properties of OPMF-unsaturated

polyester composites.

Properties Untreated treatment time (h)

24 48 72

Density (kgm-3) 980 992 997 998

Water absorption (%) 20.3+1.3 18.0+1.0 16.7+1.1 7.7+1.1

Thickness swelling (%) 10.0+0.7 7.9+1.1 7.7+1.1 7.3+1.1

Tensile strength (MPa) 6.0+1.1 9.3+0.9 14.4+0.4 17.2+0.2

Young’s modulus (GPa) 1.40+0.02 1.47+0.07 1.53+0.03 1.6+0.09

Flexural strength (MPa) 30.30+4.0

8

50.3+4.69 60.9+6.5 67.0+6.92

Flexural modulus(GPa) 1.69+0.70 2.9+0.4 4.2+0.3 5.0+0.7

The longer time of treatment led to the increase

of flexural strength and modulus. This may be related to

the fact that the treatment by sodium hydroxide was the

lignin elimination and also leaching out the amorphous

waxy cuticle layer as seen in Fig.1. The scanning

electron micrographs of untreated and treated fiber

surfaces resulted in the changes in the porous structure of

the fiber. This led to the mechanical interlocking with

polyester resin.

4. Conclusion

The results showed that the treatment by sodium

hydroxide solution played a significant role in the change

of the properties. The properties increased when the time

of treatment increase. This relates to the changes in the

porous structure of the fibers resulting from surface

treatment.

Acknowledgement

The authors would like to thank Eternal Resin Co. Ltd.

for chemicals support.

References

[1] Sreekala, M.S., Kumaran, M.G. Thomas, S.(1997)

Journal of Applied Polymer Science, 66, 821-35.

[2] Roe, P., Ansell, M.(1985) Journal of Material

Science, 20, 4015-20.

[3] De Alburquerque A., Joseph K., Hecker de

Carvalho L., Morais d’Almeida, J.(1999) Journal of

Composite Science and Technology, 60, 833-44.

[4] Pal S., Mukhopadhayay D. Sanyal S. Mukherjea R.

(1988) Journal of Applied Polymer Science, 35,

973-85.

[5] Owolabi, O. Czvikovszky, T. Kovacs, I.(1985)

Journal of Applied Polymer Science, 30, 1827-36.

[6] Sreekala, M.S.,Thomas, S.(2003) Composites

Science and Technology, 63, 861-869.

[7] Shinoj, S., Visvanathan, R., Panigrahi, S.,

Kochubabu, M. (2011) Industrial Crops and

Products, 33, 7-22. [8] Mohanty, A.K., Misra,

M., Dzral, L.T. (2001) Composite Interfaces, 8,

313-343.

[9] Ray, D. Sarkar, B. K. (2001) Journal of Applied

Polymer Science, 80,1013-20.

[10] Mwaikambo, L.Y., Ansell, M.P. (2002) Journal of

Applied Polymer Science, 84, 2222-34.

[11] Aziz, S. H., Ansell, M. P., Clarke, S. J., Panteny, S.

R. (2005), Composite Science and Technology, 65,

525-535.

[12] Zadorecki, P. and Flodin, P. (1985) Journal of

Applied Polymer Science, 30, 3971-83.

[13] Jacob, M., Varughese, K. T., Thomas, S. (2005)

Biomacromolecules, 6, 2669-79.


119

COMPP-19

Physical Properties and Morphology of Banana Flour and Cassava Starch

Reinforced with Bentonite Clay

Jareerat Ruamcharoen1*

, Kanittha Totawee1, Tipaporn Saengpan

1 and Polphat Ruamcharoen

2

1Department of Science, Faculty of Science and Technology, Prince of Songkla University, Pattani 94000 2Rubber and Polymer Technology, Faculty of Science and Technology, Songkhla Rajabhat University,

Songkhla, 90000

Abstract

The biocomposites of banana flour and cassava starch were prepared by incorporation of varying

amount of bentonite clay via solution blending. The physical properties and morphology of biocomposite films

were investigated. The moisture content and water adsorption of the composite films decreased with the

addition of bentonite to the starch. The biocomposite film of cassava starch presented an increase in the

modulus and tensile strength, respectively. The improvement in the properties of starch-based composites is due

to the strong interfacial interaction between matrix and clay which has a high modulus and change the

morphology of the matrix of starch.

Keywords: banana flour, cassava starch, bentonite, biocomposites

1. Introduction

Nowadays, the environmental pollution from

consumed polymers has become serious, particularly

from packaging materials and single-use plastic bags and

cups. Hence, there is a considerable interest in replacing

some of the synthetic plastics by biodegradable materials

in many applications. It is one promising way to solve

environment pollution problems caused by polymer

wastes.[1-3] Starch is a biodegradable polymer produced

in abundance from many renewable resources, easily

available and very cheap. Starch is a semicrystalline

polymer which is composed of repeating 1,4--D

glucopyranosyl units: amylose and amylopectin. The

amylose is linear chain, in which the repeating units are

linked by (1-4) linkages; the amylopectin has an (1-

4)-linked backbone and ca. 5% of (1-6)-linked

branches.2 The relative amounts of amylose and

amylopectin depend upon the plant source. Several

studies are considered on the development of starch-

based materials.[2-6] Unfortunately, the starch has some

drawbacks, such as the strong hydrophilic behaviour

(poor moisture barrier) and less mechanical properties

than the conventional plastic films.3

Recently the application of the nanocomposite

concept has proven to be a promising option in order to

improve mechanical and barrier properties. In this work

biocomposite films were carried out by homogeneously

dispersing layered silicates (clay minerals) in

thermoplastic starch via solution technique. These films

were made by using two types of starch matrices i.e.,

banana four and cassava starch. The physical properties

and morphology of the biocomposite films were

investigated.

2. Materials and methods

2.1 Materials

The materials used for the preparation of composite

film are banana flour and cassava starch. The bentonite

clay with a cationic exchange capacity (CEC) of 93

mequiv/100 g was obtained from Aumarin Clay Factory,

(Thailand).

2.2 Preparation of starch-based biocomposite films

Starch-based biocomposite films of different clay

contents (1, 2, 3, 4 and 5% of starch weight) were

prepared by solution blending. Granular starch was

dispersed in water (3% w/w) and gelatinized by stirring

and heating to 80oC for 30 min. A clear, viscous solution

was obtained. To this solution, glycerol (as plasticizer for

the starch), 20% w/w relative to starch on a dry basis,

was added. This was then poured on polystyrene dishes

and dried at 50oC over 3 days.


120 2.3 Physical characterization

Moisture measurement

The moisture contents of the starch and the

starch-based composite films were determined by a

gravimetric method, where by samples were dried at

105oC in an oven until constant weight was achieved.

Water swelling test

Swelling behavior was reported in terms of

changes in weight after immersion in distilled water at

ambient temperature. The samples were removed at

specified time intervals. The weight of each swollen

sample was recorded.

Tensile property test

Tensile test, to investigate the ultimate properties

(strength, elongation), along with the modulus was

performed at room temperature according to ASTM D

882 by a Hounsfield universal testing machine at a

crosshead speed of 10 mm/min.

Composite characterization and morphology study

To obtain information of the interaction

between bentonite clay and starch, Fourier transform

infrared spectroscopy was conducted in a Bruker

equipment. Thin film of starch and starch reinforced with

bentonite were taken at 16 scans in the range of 4000-

400 cm-1

with 4 cm-1

resolution of spectra. For

morphology study, transmission optical microscopy was

performed on the neat starch and starch-based

biocomposite film surface employing a Olympus BX51

with magnification of 10X.


3.1 Moisture content

The water content of all the biocomposite films

was lower than that of the neat starch films. The

interesting aspect of this observation is the decrease of

moisture content of biocomposite film as the bentonite

content increased as shown in Fig. 1. These results

indicated that the addition of clay can improve the water

resistance of starch. It could be explained that the starch

is able to form hydrogen bonds with the hydroxyls of the

silicate layers and this strong structure could reduce the

diffusion of water molecules in the composite film.[2]

Figure 1. Moisture content of neat starch and composite

films from cassava starch (CSB) and banana flour (BFB).

Figure 2. The percent of swelling of banana flour film

(BFB0) and composite film prepared from banana flour

reined with bentonite 4%wt (BFB4).

3.2 Swelling behavior

It was observed in Fig. 2 that the incorporating

bentonite clay into the banana flour cause a decrease in

the equilibrium swelling in water. This proves that the

bentonite helped increase the water resistance of the

overall system due to the presence of silicate layers.

3.3 Tensile properties

The tensile properties of banana flour (BFB)

and cassava starch (CSB) reinforced with bentonite clay

as varying amount of 1, 2, 3, 4 and 5 wt% of starch are

shown in Fig. 3. It was found that a obviously increase of

100% in Young's modulus when 2 wt% of bentonite was

added to the cassava starch. This behavior was expected

and was attributed to the resistance exerted by the clay

itself and to the orientation and aspect ratio of the

intercalated silicate layers. In addition, the stretching

0

2

4

6

8

10

12

14

% M

ois

ture c

on

ten

t

Bentonite (%w/w)

CSB BFB

0

10

20

30

40

50

60

70

80

0 200 400 600 800 1000

% S

well

ing

Time min

BFB0 BFB4


121 resistance of the oriented backbone of the amylose and

amylopectin chain in the gallery bonded by hydrogen

interaction also contributed to enhancing the modulus

and the stress. On the other hand, a slight decrease in

Young's modulus and tensile strength value was

observed in the banana flour film added with bentonite

when compared with the controlled film without clay.

This is probably due to incomplete intercalation in the

banana flour; thus, aggregates or sites for nucleation

could be present, producing an amorphous structure and

ductile material when bentonite clay was added into the

polymeric matrix.4

Figure 3. Young's modulus (a) and tensile strength (b) of

composite films from banana flour (BFB) and cassava

starch (CSB) reinforced with bentonite clay.

3.4 FTIR analysis

FTIR spectra for the bentonite, neat starch and

composite films were shown in Fig. 4 and 5. The shift of

band at 3627 cm-1

, from free OH groups of silicate layer

surface, to a lower frequency, 3288 cm-1

, indicated the

interaction between starch and bentonite clay. The

similar results have been reported.[5,6] The double

peak of O-C stretching band at 1021 cm-1

to 998 cm-1

results from bending both ‘O’ of C–O–H and ‘O’ of

anhydrous glucose ring in starch molecules. The IR

spectrum of bentonite clay revealed mainly two bands

corresponding to the Si-O stretching vibration at the

1001 cm-1

, and the Si-O bending vibration at 448 cm-1

.

The peak found at 997-1001 cm-1

for banana flour/clay

(Fig. 4) and cassava starch/clay (Fig. 5) composites

corresponding to Si-O stretching indicated the

incorporation of clay into starch.

Figure 4. FTIR spectra of the bentonite clay, banana

flour and the biocomposite film of 4 wt% clay content.

Figure 5. FTIR spectra of the bentonite clay, cassava

starch and the biocomposite film of 2 wt% clay content.

3.5 Surface morphology of biocomposite film

The surfaces of the biocomposite films were

observed by means of transmission optical microscopy.

The micrographs (Fig. 6 and 7) showed a change in the

morphology of the specimen surfaces with the bentonite

content. From micrographs shown in Fig. 6b-f and Fig.

7b-f the cassava starch and banana flour added with

bentonite displayed coarse phase morphology.

0

50

100

150

200

250

300

350

400

Yo

un

g's

Mo

du

lus

(MP

a)

Bentonite (%w/w)

BFB CSB (a)

0

1

2

3

4

5

6

7

8

Ten

sile

str

eng

th (

MP

a)

Bentonite (% w/w)

BFB CSB (b)

Bentonite clay

Ab

s

BFB0

BFB4

Wavenumber (cm-1)4000 3600 3200 2800 2400 2000 1600 1200 800 400

1001 cm-1

Si-O stretching

1021 cm-1

C-O stretching

1001 cm-1

Si-O stretching

3627cm-1

4000 3600 3200 2800 2400 2000 1600 1200 800 400

Wavenumber (cm-1)

Bentonite clay

CSB0

CSB2

Ab

s

3302cm-1

3289 cm-1

1001 cm-1

Si-O stretching

1010 cm-1

C-O stretching

999 cm-1

Si-O stretching


122

Figure 6. Optical micrographs of composite films from

cassava starch (a), CSB1 (b), CSB2 (c), CSB3 (d),

CSB4 (e) and CSB5 (f).

Figure 7. Optical micrographs of composite films from

banana flour (a), BFB1 (b), BFB2 (c), BFB3 (d), BFB4

(e) and BFB5 (f).

This agglomeration of particles may reduce the

tensile strength of composite films. The OM image of

the cassava starch reinforced with 2 wt% of bentonite

(Fig. 6c) shows that the neat starch granules were

destroyed and formed as a continuous phase because the

clay layers were uniformly dispersed in starch matrix.

This leaded to the improved mechanical

properties explained by the interactions at the phase

boundaries upon incorporating silicate layer [2,6].

However, some agglomeration of bentonite clay particles

can also be seen in Fig. 6f and 7f. This agglomeration of

particles may have reduced the tensile properties of

composite film.

Conclusion

Bentonite clay can be used to improve the

properties of starch-based composite films. The water

absorbed by the composites measured was reduced by

the addition of bentonite to the starch. The biocomposite

film of cassava starch presented an increase in the

modulus and tensile strength, respectively. From the

FTIR spectra, formed hydrogen bonds and the interaction

among starch/bentonite were evidenced by peaks

associated with -OH stretching located at 3288 cm-1

and

1001 cm-1

. The main reason for the properties

improvement in starch based composite film is the strong

interfacial interaction between matrix and clay which has

a high modulus and change the morphology of the matrix

of starch.

References

[1] Avella, M., De Vlieger, J. J., Emanuela Errico, M.,

Fischer, S., Vacca, P. and Grazia Volpe, M. Food

Chemistry. 93: 467- 474 (2005).

[2] Cyras, V. P., Manfredi, L. B., Ton-That, M-T. and

Vazquez, A. Carbohydrate Polymers. 73: 55-63

(2008).

[3] Lu, P., Zhang,M., Qian, P and Zhu, Q. Polymer

Composites. 33: 889-896 (2012).

[4] Rodríguez-Marín, M. L., Bello-Pérez, L. A., Yee-

Madeira, H., Zhong, Q. and González-Soto, R. A.

Materials Science and Engineering C. 33: 3903-

3908 (2013).

[5] Liu, H., Chaudhary, D., Yusa, S-I. and Tadea M. O.

Carbohydrate Polymers. 83: 1591-1597 (2011).

[6] Majdzadeh-Ardakani, K., Navarchian, A. H. and

Sadeghi, F. Carbohydrate Polymers. 79: 547-554

(2010).

ba

c d

e f

a b

c d

e f


123

COMPP-21

Zeolite/Silver Nanoparticle Composites: Preparation, Antimicrobial Activity and

Feasible Application in Active Packaging Film

Puttinun Kiatjiranon

and Rangrong Yoksan

*

Department of Packaging and Materials Technology, Faculty of Agro-Industry, Kasetsart University, Bangkok

10330, Thailand

Abstract

Recently, zeolite/silver nanoparticle (zeolite/Ag0) composites have been received much attention in

various fields such as catalyst, water treatment, medical materials as well as detergent due to the specific

properties of zeolites (catalyst and ion exchange) and silver nanoparticles (antimicrobial activity). However, the

reduction of silver ions was mostly carried out using sodium borohydride and hydrazine, which are harmful and

non-ecofriendly reducing agents. Therefore, the objective of this research is to prepare composites between

zeolite and silver nanoparticles via a two-step method of ion exchange and reduction using vitamin C as a

reducing agent. The successful preparation of zeolite/Ag0 composites was confirmed by UV-Visible

Spectroscopy, Scanning Electron Microscopy and Energy Dispersive X-ray Spectroscopy. The antimicrobial

activity of zeolite/Ag0 composites against Escherichia coli was also investigated. Zeolite/Ag

0 composites

showed Surface Plasmon Resonance (SPR) band of silver nanoparticles at 430 nm. The silver nanoparticles

attached on the external zeolite surface possess spherical shape with an average diameter in the range of 19-21

nm. The amount of silver nanoparticles increased with increasing initial silver concentration. Zeolite/Ag0

composites exhibited excellent antimicrobial activity against Escherichia coli with MICs of 0.313 mg/mL when

initial concentration of silver used for preparation was 0.02 M. The as-prepared zeolite/Ag0 0.02 composites

showed the feasibility to be applied as an additive for improving tensile strength and stiffness of starch-based

film.

Keywords: Zeolites, Silver nanoparticles, Antimicrobial activity

1. Introduction

Zeolite is a nanoporous aluminosilicate

compound, in which its structure composes of Si and Al

atoms connected by O atom to form porous three-

dimensional structure. Owing to various outstanding

properties, such as ion exchange and adsorption [1],

zeolites have been used in many applications as

adsorbent and molecular sieve [2,3], catalyst [4] as well

as plastic additives [5]. The combination of zeolites with

silver nanoparticles is reported as an alternative to

enhance the antimicrobial activity of zeolites.

Previously, many research groups have demonstrated

that zeolite/Ag0

composites showed excellent

antimicrobial activity against Gram-negative and Gram

positive bacteria [6-8]. The combination of zeolites with

silver nanoparticles can expanded zeolite applications,

including coating, drinking water and biomedical devices

[8]. Therefore, the objective of this research is to provide

on antimicrobial capability to zeolites by incorporating

silver nanoparticles. The obtained zeolite/Ag0

composites were expected to be used as antimicrobial

ion exchange

material or adsorbent and as reinforcing additive for

polymeric films.


Zeolite/Ag0 composites were prepared using a

two-step method of silver ion exchange and reduction.

Dried zeolite powder (Zeolite A with a dimension of

45 m, a pore size of 0.5 nm and a moisture content of

0.998 %wt, 60 mg) was magnetically stirred in 0.02 M or

0.05 M of AgNO3 solution (3 mL) at ambient

temperature for an hour. After that the mixture was

centrifuged and the precipitate was collected followed by

drying at 60 °C overnight. The obtained powder of

zeolite/silver ion (zeolite/Ag+) composites was stirred

regularly in vitamin C solution (0.5 mol equivalent to

mol of Ag) at ambient temperature for an hour. After that

the mixture was centrifuged to collect the precipitate,

which was then dried at 60 °C overnight to obtain


124 zeolite/Ag

0 composites. The as-prepared composites

were characterized by UV-Vis spectroscopy, Scanning

Electron Microscopy (SEM) and Energy Dispersive X-

ray Spectroscopy (EDS). The antimicrobial activity of

the composites was examined by a tube dilution method

against Escherichia coli (E. coli, TISTR 361) using

microbial concentration of 105 CFU/mL in order to

determine the Minimum Inhibitory Concentration (MIC).

The obtained zeolite/Ag0 composites were then

incorporated into starch film by a solution casting

method. Zeolite/Ag0 composites with various

concentrations, i.e. 1%, 2.5% or 5% w/w of starch were

dispersed in distilled water (100 mL) for 15 min. After

that cassava starch (5%) was added to the dispersion.

The mixture was mixed together at 80 °C for 1 h to

completely gelatinize starch; then glycerol (30% w/w of

starch) was added as a plasticizer and the mixture was

stirred extensively for 30 min. The obtained mixture was

poured onto acrylic plate and dried in a hot air oven at

60 °C overnight. The film was peeled from the plate and

then cut into a square shape with a dimension of 1.5 cm

9 cm before testing tensile properties.


UV-Vis spectra of zeolite and zeolite/Ag0

composites were shown in Figure 1. Zeolite does not

exhibit UV absorption at the wavelength range of 300-

800 nm. In contrast, zeolite/Ag0 composites showed

Surface Plasmon Resonance (SPR) band, which is the

characteristic of silver nanoparticles at 430 nm, implying

the existing of silver nanoparticles in the samples. The

band intensity of zeolite/Ag0 0.05 composites was higher

than that of zeolite/Ag0 0.02 composites, indicating that

silver nanoparticle content in the composites increased

with increasing the initial silver concentration. This

result was in agreement with the one reported by Shameli

et al. (2011) [8].

Figure 1. UV-Vis spectra of (a) zeolite, (b) zeolite/Ag0

0.02 composites and (c) zeolite/Ag0 0.05 composites.

The morphology observed by a scanning electron

microscope indicated that the external surface of zeolite

was slightly rough (Figure 2a), while those of both

zeolite/Ag0 0.02 composites and zeolite/Ag

0 0.05

composites possessed higher roughness due to the

appearance of the spherical silver nanoparticles (Figure

2b-c). The diameter of silver nanoparticles was in the

range of 20-30 nm and increased with increasing the

initial silver concentration. Shameli et al. (2011) also

found similar result [8].

Figure 2. Scanning electron micrographs at 30 kV

(100000) of (a) zeolite, (b) zeolite/Ag0 0.02 composites

and (c) zeolite/Ag0 0.05 NCs.

(a)

(b)

(c)

400 500 600 700 800

Wavelength (nm)

(b)

0.20

0.15

0.10

0.05

0

(a)

(c)

(a)


125

(a)

Figure 3. EDS mapping images of (a) zeolite/Ag0 0.02

composites and (b) zeolite/Ag0 0.05 composites

After Ag+ was exchanged with Na

+ and H

+ of

zeolites [9], followed by chemical reduction, Ag0 could

be formed on the surface of zeolite as seen in EDS

mapping images (Figure 3). Silver nanoparticles

exhibited good distribution (red spots) on the external

surface of zeolite/Ag0 0.02 composites (Figure 3a) and

zeolite/Ag0 0.05 composites (Figure 3b). The elemental

compositions of zeolite and zeolite/Ag0 composites

analyzed by EDS were presented in Table 1. Silver

content in the composites increased with increasing

initial silver concentration.

Antimicrobial activity of the as-prepared

composites was determined against E. coli using a tube

dilution method. Minimum inhibitory concentration

(MIC) of the samples was evaluated by observing the

turbidity and transparency of the culture media after

incubation at 37 °C for 24 h, in conjunction with

considering the absorbance at 600 nm of those media.

After incubation, the transparent media will be obtained

if the microorganism growth is inhibited. In contrast, the

culture media will be turbid if the microorganism is alive

[10].

Table 1. Elemental compositions of zeolite, zeolite/Ag0

0.02 composites and zeolite/Ag0 0.05 composites.

Elements

Weight (%)

Zeolite Zeolite/Ag0 0.02

composites

Zeolite/Ag0 0.05

composites

Si 24.39 24.00 23.58

Al 12.00 11.63 11.50

O 57.74 56.13 55.78

Na 1.94 2.28 0.90

Ca 3.92 4.33 2.97

Ag - 1.62 5.27

Figure 4. Appearances of culture media containing

E. coli (105 CFU/mL) and samples: (A) zeolite, (B)

zeolite/Ag0 0.02 composites and (C) zeolite/Ag

0 0.05

composites with various sample concentrations: (a) 5, (b)

2.5, (c) 1.25, (d) 0.625, (e) 0.313, (f) 0.156, (g) 0.078 and

(h) 0.039 mg/mL after incubation at 37 °C for 24 h.

Figure 4A shows that the culture mixtures

containing E. coli and zeolite are turbid for all zeolite

concentrations, implying that zeolite was not able to

inhibit the growth of E. coli, whereas the culture

mixtures containing E. coli and zeolite/Ag0 composites

became transparent when the concentration of

zeolite/Ag0 composites increased up to of 0.313 mg/mL

and 0.625 mg/mL when the initial silver concentrations

of 0.02 M and 0.05 M were used. The lowest

concentration of sample that provides transparent culture

media is defined as MIC.

(C)

(a) (b) (c) (d) (e) (f) (g) (h)

(B)

(A)

(b)

(a)


126

Figure 5. Absorbance at 600 nm of culture media after

incubation at 37 °C for 24 h of (a) zeolite/Ag0 0.02

composites and (b) zeolite/Ag0 0.05 composites.

Figure 4 and 5 confirms that MIC of zeolite/Ag0

0.02 composites is 0.313 mg/mL, whereas that of

zeolite/Ag0 0.05 composites is 0.625 mg/mL. The result

indicated that the as-prepared composites could inhibit

the growth of E. coli. The lower the MIC corresponds to

the higher the antimicrobial activity. Even though the

silver content of zeolite/Ag0 0.02 composites was lower

than that of zeolite/Ag0 0.05 composites (Table 1), it

showed greater antimicrobial activity. This might be

possible that the smaller size of silver nanoparticles in

zeolite/Ag0 0.02 composites (Figure 2) possessed higher

surface area resulting in higher antimicrobial action [11].

Figure 6. (A) Tensile strength, (B) modulus and (C)

elongation at break of starch film containing different

concentrations of zeolite/Ag0 0.02 composites. The data

is reported as mean ± SD, n = 3–5. The different small

letters indicate significant difference at p < 0.05

(Duncan's new multiple range test).

The zeolite/Ag0 0.02 composites were then

incorporated into starch film; the effect of the composites

on tensile properties of the film was also investigated.

Tensile strength, modulus and elongation at break of the

films are shown in Figure 6. Incorporating zeolite/Ag0

0.02 composites with the concentrations of 1%, 2.5% and

a

a

(B) a

b

120

100

80

60

40

20

0 0 1 2.5 5

Zeolite/Ag0 0.02 composites content (%)

(a)

a

b

0 1 2.5 5


140

120

100

80

60

40

20

0

(C)

b

b

(a)

b

2.5

2.0

1.5

1.0

0.5

0

0 1 2.5 5


a

a a

b

(A)

(b)

MIC

5 2.5 1.25 0.625 0.313 0.156 0.078 0.039

Concentration of zeolite/Ag0 0.05 composites

(mg/mL)

1.2

1.0

0.8

0.6

0.4

0.2

0

MIC (a)

1.2

1.0

0.8

0.6

0.4

0.2

0

MIC

5 2.5 1.25 0.625 0.313 0.156 0.078 0.039

Concentration of zeolite/Ag0 0.02 composites

(mg/mL)

(a)

(b)


127 5% resulted in increased tensile strength and modulus,

while decreased elongation at break. The concentration

of zeolite/Ag0 0.02 composites hardly affected tensile

properties of the film. The results implied that

zeolite/Ag0 0.02 composites could improve tensile

strength and stiffness. This might be explained by the

surface interaction through hydrogen bonds formed

between zeolites and polysaccharide chains, resulting in

strong interfacial adhesion in composites and reduced

starch chain mobility; as a result macroscopic rigidity of

the film was enhanced [11,12]. The reduction of

extensibility of starch film by incorporating zeolite/Ag0

0.02 composites was due to the interruption of polymer

chain mobility by zeolite [13].

Conclusions

Zeolite/Ag0 composites were successfully

prepared via a two-step method of silver ion exchange

and reduction, using vitamin C as a reducing agent. The

characteristic SPR band of silver nanoparticles in the

composites appeared at 430 nm. The average diameter of

the silver nanoparticles was in the range of 20-30 nm and

those silver nanoparticles showed good distribution on

the external zeolite surface. The diameter and content of

silver nanoparticles on the zeolite surface increased with

increasing the initial silver concentration. The

composites exhibited antimicrobial activity against the

growth of E. coli with the MICs of 0.313 mg/mL for

zeolite/Ag0 0.02 composites and 0.625 mg/mL for

zeolite/Ag0 0.05 composites. The incorporation of

zeolite/Ag0 0.02 composites into starch films led to

improvement of tensile strength and stiffness.

Antimicrobial activity of starch film incorporating

zeolite/Ag0 composites will be further studied.

References

[1] Lesthaeghe, D., Delcour, G., Speybroeck, V. V.,

Marin, G. B. and Waroquier, M., “Theoretical study

on the alteration of fundamental zeolite properties

by methylene functionalization”, Microporous and

Mesoporous Materials 96: 350-356 (2006).

[2] Tso, C. Y., Chan, K. C., Chao, C. Y. H. and Wu, C.

L., “Experimental performance analysis on an

adsorption cooling system using zeolite 13X/CaCl2

adsorbent with various operation sequences”,

International Journal of Heat and Mass Transfer 85:

343-355 (2015).

[3] Ahmed, M. J. and Theydan, S. K., “Modeling of

propane separation from light hydrocarbons by

adsorption on 4A molecular sieve zeolite”, Journal

of Natural Gas Science and Engineering 18: 1-6

(2014).

[4] Han, W., Zhang, P., Tang, Z. and Lu, G., “Low

temperature CO oxidation over Pd–Ce catalysts

supported on ZSM-5 zeolites”, Process Safety and

Environmental Protection 92: 822-827 (2014).

[5] Thipmanee, R. and Sane, A., “Effect of zeolite 5A

on compatibility and properties of linear low-density

polyethylene

[6] Ferreira, L., Fonseca, A. M., Botelho, G., Aguiar, C.

A. and Neves, I. C., “Antimicrobial activity of

faujasite zeolites doped with silver”, Microporous

and Mesoporous Materials 160: 126-132 (2012).

[7] Flores-López, N. S., Castro-Rosas, J., Ramírez-Bon,

R., Mendoza-Córdova, A., Larios-Rodríguez, E. and

Flores-Acosta, M., “Synthesis and properties of

crystalline silver nanoparticles supported in natural

zeolite chabazite”, Journal of Molecular Structure

1028: 110-115 (2012).

[8] Shameli, K., Ahmad, M. B., Zargar, M., Yunus W.

M. Z. W. and Ibrahim, N. A., “Fabrication of silver

nanoparticles doped in the zeolite framework and

antibacterial activity”, International Journal of

Nanomedicine 6: 331-341 (2011).

[9] Akgul, M., Karabakan, A., Acar, O. and Yurum, Y.,

“Removal of silver (I) from aqueous solutions with

clinoptilolite”, Microporous and Mesoporous

Materials 94: 99-104 (2006).

[10] Yoksan, R. and Chirachanchai, S., “Silver

nanoparticle-loaded chitosan–starch based films:

Fabrication and evaluation of tensile, barrier and

antimicrobial properties”, Materials Science and

Engineering C 30: 891-897 (2010).


128 [11] Ghosh, A., Ma, L. and Gao, C., “Zeolite molecular

sieve 5A acts as a reinforcing filler, altering the

morphological, mechanical, and thermal properties

of chitosan”, Journal of Materials Science 48:3926-

3935 (2013).

[12] Belibi, P. C., Daou, T. J., Ndjaka, J. M. B.,

Michelin, L., Brendlé, J., Nsom, B. and Bernard, D.,

“Tensile and water barrier properties of cassava

starch composite films reinforced by synthetic

zeolite and beidellite”, Journal of Food Engineering

115: 339-346 (2013).

[13] Chang, B. P., Akila, H. M. and Nasir, R. B.,

“Mechanical and tribological properties of zeolite-

reinforced UHMWPE composite for implant

application”, Procedia Engineering 68: 88-94

(2013).


129

COMPP-22

Extraction and characterization of cellulose nano-whiskers from rice bran waste

Witita Laosamathikul1, Kittiwut Kasemwong

2, and Pakorn Opaprakasit

1


Thammasat University, Pathum Thani, 12121, Thailand 2National Nanotechnology Center, National Science and Technology Development Agency,

Pathum Thani, 12120, Thailand

Abstract

Rice bran waste (RBW), a residue from rice bran oil production process, is used as a raw material to

produce cellulose nanowhiskers (CNWs). Rice bran cellulose (RBC) was isolated from the raw material by

using alkali treatment and bleaching treatment to remove hemicellulose and lignin. The diameter of CNWs is

smaller than RBC and RBW with the highest aspect ratio, since amorphous domains are eliminated. This leads

to the higher crystallinity index of CNWs. TGA results suggest that CNWs has higher thermal property than its

starting materials. Given these properties and its biocompatibility, the materials can be used as reinforcing

agents for biodegradable polymers for use in biomedical or packaging applications.

1. INTRODUCTION

Rice bran oil is oil product extracted from hard outer

brown layers of rice after chaff (rice husk). This is very

popular cooking oil in many countries, since it provides

many health benefits such as cholesterol busting power,

cancer protection, immunity boost or antioxidant. 1.

There are some residues from rice bran oil process that

contain high amount of cellulose, which can be extracted

and used in many applications, especially cellulose and

its derivatives. Cellulose is one of the most interesting

materials for use in reinforced bio-composites, because

of its outstanding mechanical properties, low density,

and renewability.

Nano-sized cellulose is considered an attractive

reinforcing agent because of its good mechanical

properties, such as, high tensile strength, high flexural

strength, and high stiffness with low density 2,3

. Cellulose

nano-whiskers (CNWs) are generated by acid hydrolysis

to dissolve amorphous fractions off, leading to an

increase in the percentage of crystalline domains 4. There

are many reports on extractions of cellulose and

productions of CNWs from various types of plants and

even wastes, such as biomass and agricultural residuals 5-

7.

The aim of this research is to develop a process to

produce cellulose nano-whiskers from rice bran waste.

The material’s structures and properties are then

characterized for use as a reinforcing agent for bio-

materials.

Table 1 Rice bran waste composition

Composition Percentage (%)

Moisture 9.06

Protein 12.47

Fat 1.35

Ash 7.80

Total Carbohydrate 69.32

Dietary Fiber 27.14

2. EXPERIMENTAL

2.1 Materials

Rice bran waste (RBW) was supplied by a rice bran

oil production company in Nonthaburi, Thailand.

Sodium hydroxide pellets (AR grade, Quality Reagent

Chemical), Sodium Chlorite 80% (technical grade,

Sigma-Aldrich), Acetic acid glacial with more than 80%

acid by mass (Mallinckrodt Chemicals), Sulphuric acid

(96%, Carlo Erba Reagents) were used without further

purification.

2.2 Preparation of rice bran cellulose (RBC)

2.2.1 Alkali treatment

The alkali treatment is employed by using a modified

method reported by Deepa et al. the rice bran waste raw

material was sieved at 180 mesh, and then purified by

removing lignin and hemicellulose. 30 g of rice bran was


130 treated with 750 ml of 4 wt% NaOH in an Erlenmeyer

flask. The mixture was autoclaved with high pressure

steam sterilizer (TOMY SX-700, Japan) at 120 °C for 1

h. Finally, the solid samples were filtered with electric

aspirator (JEIO TECH, Korea), and neutralized by

washing several times with distilled water 8.

2.2.2 Bleaching process

A bleaching process is modified from LaCourse et

a1. Essentially, the alkali-treated sample was soaked in a

1 wt% aqueous chlorite solution. Acetic acid was then

added to adjust the pH to 5±0.2 (SUNTEX SP-2200,

Microprocessor pH meter, Taiwan). The mixture was

heated to 80°C with constant agitation by a magnetic

stirrer (IKA RCT basic, Malaysia) overnight, and the pH

was checked for a few times until the reaction was

complete. After the material’s color turned to be white, it

was allowed to cool to room temperature, and then

filtered by a suction filtration. The solid product was

mixed with ethanol in a ratio of 1:3 (sample over liquor)

for a few hours to remove the remaining water, before

being filtered and washed again with acetone. The

product was finally oven dried in a hot-air oven

(BINDER, Germany) at 50 °C for overnight. Snowy

white powder product was obtained 9.

2.3 Preparation of cellulose nano-whiskers (CNWs)

The preparation of CNWs was modified from that of

Lu and Hsieh. Bleached rice bran was treated by a

concentrated a sulfuric acid solution (64 wt% sulfuric

acid in water) at 45 °C for 45 min under constant stirring

using a shaking incubator (Labnet, USA). The ratio of

rice bran to acid solution was 1–20 g/ml. The suspension

was then diluted 10 times with deionized water, and kept

in a refrigerator to prevent further reaction. Residual

sulfuric acid in the suspension was removed by repeated

centrifugation (TOMY MX-305 High speed refrigerated

micro centrifuge) at 10,000 rpm for 25 min until the

supernatant was turbid. The supernatant (CNW

suspension) was collected and dialyzed with molecular

porous membrane tubing at molecular weight cutoff of

12,000–14,000 (Spectrum Labs, USA) against deionized

water for 4–5 days, to remove the remaining acid until it

reaches a constant pH. Finally, CNWs were obtained in

deionized water solution 10

.


2.3.1 X-ray diffraction

The X-ray diffraction patterns of RBW, RBC, and

CNWs were obtained with 1.54 ∝ radiations at 40

kV and 30 mA. The scanning region of the two-theta

angle (2 ) was from 5 to 40° with a scan rate of 1°/min.

(Bruker D8 Advance). The Crystallinity Index (Cr.I.)

was determined by

Cr.I. (%) = (Sc / St) • 100

where: Sc is area of the crystalline domain

St is area of the total domain

11

2.3.2 Thermogravimetric analysis (TGA)

TGA analyses were performed to observe thermal

stability of materials (TGA, METTLER TOLEDO,

USA). Samples were weight around 5-10 mg and placed

in a clean ceramic plate and the data were recorded by

heating the samples at 10 ◦C/min from 30°C to 600 °C

in N2 with a purging rate of 50 mL/min 12

.

2.3.3 Fourier transform infrared (FTIR) spectroscopy

Chemical structures of RBW, RBC, and CNWs

products were examined by fourier transform infrared

(FTIR) spectroscopy (NICOLET 6700, Thermo

Scientific, USA). The sample in powder form was

grinded and mixed with KBr at a concentration of ca.

1 % wt. The mixture was then compressed to be in a disc

form by using a hydraulic machine at a pressure of 9.5

tons for 10 sec. All spectra were recorded from 4000-400

cm-1

, with 32 scans at a resolution of 4 cm-1

.

2.3.4 Scanning electron microscopy

Morphology of RBW, RBC, and CNWs samples

were examined on a HITACHIS-3400N scanning

electron microscope, Illinois, USA with voltage of 30

kV.


131 2.3.5 Transmission electron microscopy

The dimensions of CNWs were investigated on a

JEM-2100 transmission electron microscope, JEOL,

USA with an accelerating voltage of 80 kV13

. CNWs

samples were prepared by dropping into a copper grid

and straining with 2% uranyl acetate, and dried at an

ambient temperature before analysis.

3. RESULTS AND DISCUSSION

3.1 Crystal structure analysis

XRD spectra of RBW, RBC, and CNWs samples are

shown in Figure 1. All materials can be classified as

cellulose I or native cellulose, as reflected by their

diffractogram’s profiles with peaks at 2θ angle of 15°,

17°, 21°, 23°, and 34° 14,15

. The results indicate that the

crystallinity index (Cr.I.) of RBW, RBC, and CNWs

samples are 28.2, 46.5, and 47.8%, respectively. RBC

has higher Cr.I., compared to that of the original RBW,

as a result from the removal of amorphous non-cellulosic

components induced by the alkali and bleaching

treatments performed in the purification process. This is

reflected by its narrower and shaper XRD peak patterns

13. The crystallinity of CNWs also increases, compared to

RBC, as the acid hydrolysis treatment further removes

some of amorphous components in the samples, leading

to much sharper and more intense crystalline peaks.

3.2 Thermal stability analysis

TGA thermograms of the samples are compared in

Figure 2. All samples show a decomposition step

covering 30-170 °C, which is attributed to vaporization

of moisture 16

. RBW has 3 steps of degradation, in which

the first step is due to hemicellulose and cellulose

decomposition. Since hemicelluloses are the most

thermally unstable components of biomass, because of

random amorphous structures and reactive acetyl groups,

this decomposes at an onset temperature of 161-249 °C,

with a weight loss of 4.89%. Is second decomposition

step is mainly due to cellulose and some lignin at 248-

358 °C, with 38.31% mass loss. Finally, the final mass

loss stage at 359 to 487 °C represents lignin

decomposition with 11.77% weight 17-19

.

Fig.1. X-ray diffraction patterns of RBW, RBC, and CNWs.

Fig.2. TGA (A) and DTG (B) curves of RBW, RBC, and

CNWs

The corresponding thermogram of RBC shows only

one step of sample degradation, due to linear polymer

chains of glucose in cellulose structures. This occurs at

RBW

RBC

CNWs

A

B


132 higher decomposition temperature, because lignin and

hemicellulose are already removed during the alkaline

treatment and bleaching process. Also, because some

amorphous domains are removed from the sample, RBC

with higher crystallinity, exhibits higher thermal

stability, and decompose at a temperature range of 175 to

494°C 19

.

In contrast, CNWs shows different TGA patterns,

compared to its starting RBC material. This difference is

likely due to the presence of sulfate groups in outer

surface structure of the nano-crystalline particles20

.

CNWs has lower degradation temperature than RBW and

RBC, in which the first degradation stage at 111-190 °C

might originate from its highly-sulfated amorphous

domains that influenced easier degradation. Another

decomposition at 190 to 234 °C is due to the unsulfated

crystalline domains 21

. It is also observed that CNWs has

the highest amount of residue (59.1%), compared to

RBW and RBC with the residues of 35.54% and 23.74%,

respectively. This is likely because the associated sulfate

groups act as a flame retardant agent or protective barrier

on the burning surface 3,21

.

3.3 Chemical structures analysis

Chemical structures of the products are examined by

FTIR spectroscopy, as shown in Figure 3. RBW show

characteristic bands at 1700 cm-1

, assigned to acetyl and

ester groups in hemicellulose or carboxylic acid groups

in ferulic and p-coumaric components of lignin. The

peak at 1546 cm-1

is attributed to C=C vibration, due to

the presence of lignin15

. The band disappears in spectra

of the other spectra, as the alkaline and bleaching

treatments can remove noncellulosic materials13,22

. The

spectra of RBC and CNWs show bands at 1060 and 898

cm−1

, corresponding to C–O stretching and C–H rock

vibrations of cellulose. The increase in intensity of these

peaks reflects an increase in the percentage of cellulosic

components. This implies that after chemical treatment,

the sample has higher cellulose content or almost pure

cellulose. The high intensity of this mode is also an

indicative of high crystallinity of the cellulose sample

3,15,23.

Fig. 3 FT-IR spectra of RBW, RBC, and CNWs.

The CNWs spectrum, after acid hydrolysis, also

shows a band due to sulfate groups at 1250 cm−1

( asymmetric S−O vibration) and 833 cm−1

(symmetric

C−O−S vibration) 24

. This agrees with those observed in

TGA thermograms.

3.4 Morphological structures

Morphological structures of the samples are

examined by SEM, as shown in Figure 4. RBW has

agglomerated irregular shape fibrils and a rough surface

morphology. After alkaline treatment and bleaching

treatment, the morphology changes to rod-like structures.

Evidences of partial removal of impurities hemicellulose

and lignin after chemical treatment of RBC are observed,

i.e., cementing components around the fiber-bundles

disappear, and only cellulose remains in the sample5.

In alkaline treatment, it is expected that mainly

hemicellulose is removed while the bleaching treatment

is conducted to remove rice bran fiber-bundles and

separate them into individual fibers. After chemical

treatments, the diameter of the fibers decreases from 40-

580 µm to 40-300 µm. The size reduction is mainly

attributed to the separation of the fibers’ primary cell

wall due to the removal of hemicellulose and lignin.

RBC also has higher fiber aspect ratio (length/diameter)

around 2.40-9.17 compare to 1.0-2.4 in RBW 13

.


133

Fig.4. Scanning electron micrograph of RBW (A) and RBC (B)

Figure 5 shows TEM image of CNWs. This

confirms that the size is reduced to nano-scale with the

shape like whiskers. The amorphous regions of cellulosic

microfibrils are removed as a result from the acid

hydrolysis, leaving behind only straight-shaped

crystalline domains13

. The results strongly reflect the

success of nano-whiskers formation.

Fig.5. Transmission electron micrograph of CNWs

4. CONCLUSIONS

Rice bran cellulose is extracted from rice bran waste

by alkali and bleaching treatments to remove non-

cellulosic constituents and break down microfibrils to

individual fibers. By using acid hydrolysis, amorphous

domains are removed, the remaining crystalline cellulose

presents in the form of cellulose nano-whiskers. The

materials can be used as reinforcing agents for

biomaterials, because of their high thermal stability,

crystallinity, hydrophilicity, and high aspect ratio with

small diameter. The materials can be used to composite

with biodegradable PLA to improve its physico-

mechanical properties for use in biomedical or packaging

applications.

5. ACKNOWLEDGEMENTS

W.L. is thankful for a support from the TAIST-Tokyo

Tech scholarship program.

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135

COMPP-24

The Preparation of Vanadium Oxides Films via A Polymer Assisted Deposition

Onruthai Srirodpai1*

, Jatuphorn Woothikanokkhan1,3

and Saiwan Nawalertpanya 2,3

1School of Energy, Environment and Materials, King Mongkut's University of Technology Thonburi

(KMUTT), Bangkok 10140 2Department of Chemical Engineering, Faculty of Engineering, King Mongkut's University of Technology

Thonburi (KMUTT), Bangkok 10140 3Nanotec-KMUTT Center of Excellence on Hybrid Nanomaterials for Alternative Energy, Bangkok, 10140

Abstract

Vanadium dioxide films, to be used as a thermochromic material for smart glazing, were prepared and

fabricated on glass substrate via a polymer assisted deposition (PAD). Poly(vinyl pyrrolidone) (PVP) and

poly(vinyl alcohol) (PVOH) were used as the film former to control the viscosity of precursor solution and to

induce an interaction with vanadium ions. The film-forming mechanism, phase determination and surface

morphology were studied by using FT-IR, X-ray diffraction and scanning electron microscopy techniques,

respectively. The results showed that PVOH had greater interaction with VO2+

than PVP. A variety of other

vanadium oxides were found, including VO2 . The temperature had significant influence on morphology of VOx

films. In this study, the best condition that might be used to prepare VO2 (M) film is that by using the PVOH as

a precursor solution and annealing at 450oC for 3 h.

Keyword: vanadium dioxide film, polymer assisted deposition, poly (vinyl alcohol), poly (vinyl pyrrolidone)

1. Introduction

Nowadays, a smart or intelligent window

technology has been evolved and developed to decrease

the energy consumption and increase energy efficiency

in building [1,2]. Chromogenic glazing, is a kind of

smart glass capable of changing color upon an external

stimuli and controlling solar transmittance. These include

thermochromic, photochromic and electrochromic

glazing in which the color was changed with

temperature, light intensity and electricity, respectively

[3-5]. It has been reported that transmission of infrared

radiation through the thermochromic glazing is lower

than that through the photochromic glazing. The

development of thermochromic glazing is also less

expensive and the fabrication process is uncomplicated,

compared with those of the electrochromic glazing [3,6].

Several thermochromic materials have been

used to develop the smart glazing. These include V2O3,

V2O5, V5O9, V4O7, V6O13 Ti2O3 and VO2 [7, 23].

Particularly, VO2 is the most interesting because of its

lower transition temperature (68 oC) and capability of

exhibiting fully reversible phase transition between M

phase and R phase. Below the phase transition

temperature, it has a monoclinic lattice with the P21/c

space group (M phase). Above the phase transition

temperature, the monoclinic phase of VO2 transforms to

tetragonal lattice with the P42/mnm rutile space group (R

phase). This transformation change the electrical

conductivity, optical transmittance and reflectance in

near infrared region of the VO2. These properties can be

applied for smart window, temperature sensing device,

modulator, etc. Moreover phase transition temperature of

VO2 can be further reduced by doping it with some

metals such as tungsten (W) [7-11].

Techniques used to fabricate VO2 film can be

divided into two main systems. These are a vacuum

system and a solution based (or sol-gel) system. VO2

film obtained from vacuum system is highly uniform but

the process is complex and expensive. For the solution

based system, the fabrication is less complicated and

inexpensive. However, the obtained film is usually less

uniform and prone to cracking. Recently, a new solution

based technique, namely a polymer assisted deposition

(PAD), was developed to overcome the above

limitations. PAD is considered uncomplicated and

inexpensive, compared with vacuum system. The main

concept of PAD is the use of the polymer to bind metal

ions and to control the solution viscosity, leading to a


136 homogeneous distribution and uniformity of metal films.

[12,13].

In this work, the preparation of VO2 films on

glass substrate via PAD was studied. The aim of this

work was to investigate the effects of fabrication

parameters, which are polymer types and annealing

temperature, on oxidation number and microstructure of

the vanadium oxides product.

2. Experimental

2.1 Chemicals

Vanadiam pentoxide (V2O5, 99.5 % pure) was

obtained from Aldrich Co.Ltd. Hydrazine monochloride

(N2H4•HCl,analytically > 98 % pure) and Poly(vinyl

pyrrolidone) (PVP K90, average molecular weight:

1,300,000) were obtained from Acros Co.Ltd. Poly(vinyl

alcohol) (PVOH, average molecular weight: 205,000 was

obtained from Aldrich Co.Ltd. Hydrochloric acid (HCl,

analytical pure) was obtained from Merck Co.Ltd. All of

chemicals were used without further purification.

2.2 Preparation of precursor solution

The precursor solution of vanadyl oxydichloride

(VOCl2) was prepared by adding 1 g of N2H4•HCl in 6

ml of HCl into suspension V2O5 3.5 g containing 50 mL

of deionized water. After stirring until blue solution was

formed and that was then filtered. A clear VOCl2

solution (pH~1) was obtained. The concentration of

VOCl2 solution was adjusted to 0.1 mol/L, then PVP and

PVOH were added to the solution (6 wt%). These were

used as the precursor solution for VOx dipping.

2.3 Preparation of VOx Films

Films were coated on the glass substrate by

dipping process. The dipped film was then dried at 60 °C

for 30 min to get rid of some excess solvent. A smooth

thin film precursor was formed. The dried films were

annealed under nitrogen gas flow for 3 h in various

temperatures.


Functional group of the precursor solutions

were determined by FT-IR (Thermo scientific, Nicolet

IS5) in wavenumber ranged between 500 and 4000 cm-1

.

X-ray diffraction patterns were recorded by an X-ray

diffractometer (XRD, Bruker AXS D8-Discover) in the

2θ range of 10–80° using Cu-Kα radiation (λ=1.54178

Å). The accelerating voltage and the current used were

40 kV and 40 mA, respectively. A field emission gun

scanning electron microscopy (FEG-SEM) technique

was used to determine the microstructure of VOx films.


3.1 Stability of precursor solutions and the film-forming

mechanism

Fig. 1 shows FTIR spectra of precursor solution

prepared by using PVP and PVOH. It can be seen that

the stretching vibration of -C=O bond in PVP molecules

was shifted from 1,644 cm-1

to 1,639 cm-1

when PVP was

added into precursor solution (Fig 1a). This was because

of a loosening of the -C=O bond by the coordination

between carbonyl group and vanadium cations (VO2+

)

[14,15]. The vibration peak at 1,289 cm-1

representing

-C-N of PVP was unchanged after adding into the

precursor solution. This indicates that there was no

interaction between vanadium cations and amine

functional groups in PVP. In the case of PVOH, it was

found that both stretching vibration of -O-H and -C=O in

PVOH was shifted from 3,300 cm-1

to 3,271 cm-1

and

from 1,724 cm-1

to 1,637 cm-1

, respectively. This was

owing to the strong electrostatic interactions between

VO2+

and some functional groups (-O-H and -C=O) in

the PVOH molecules [16].


137

Fig. 1 FTIR spectra of (a) PVP film, (b) PVP with precursor solution, (c) PVOH film and (d) PVOH with precursor

solution.

Fig. 2 Schematic illustration of the film-forming mechanism of (a) PAD of VOx with PVP (b) PAD of VOx with PVOH.

A schematic diagram illustrating of the film-

forming mechanism between polymers and VO2+

was

proposed and showed in Fig. 2 [16,17]. The schematic

diagram suggests that the orbit conjugation in PVP

molecules was occurred by the partial donation of lone-

pair electrons in nitrogen atoms to adjacent carbon atoms

[18]. This led to negatively charged carbonyl groups and

positively charged amine groups [18,19], resulting to

formed electrostatic interactions between carbonyl

groups in PVP and VO2+

in the precursor solution. Before

the addition of PVOH into the precursor solution, the O-

H groups are related in intramolecular and intermolecular

hydrogen bonding with -C=O groups on the PVOH

backbone [20]. The addition of PVOH in precursor

solution leading to both O-H and -C=O groups are

strongly interacting with VO2+

in coordination complex

shown in Fig. 2b.

Moreover, it has been reported that the strong

interaction between metal ions and polymer contributed

to a greater viscosity and film formability of the solution.

This lead to the better homogeneity of the film products

[16,17].

3.2 Phase determination

The XRD spectra of VOx films prepared by

using PVP precursor solution and annealed under various

temperature are illustrated in Fig. 3a. The broad

background peaks from 15° to 35° in XRD spectra are

attributed to glass substrates. There was no crystal

monoclinic VO2(M) detected from the XRD patterns of

products. However, the metastable VO2 (B) phase

[JCPDS No. 81392] can be detected at the annealing

temperature of 450, 500 and 550 oC. In addition, many

other structures of vanadium oxides with different

valences such as V5O9 were also noted [21].

XRD patterns of VOx films prepared from

PVOH precursor solution were illustrated in Fig. 3b.

Crystal monoclinic VO2(M) phase [JCPDS No. 72-0514]

was detected at the annealing temperature of 450 and 550

oC. The metastable VO2 (B) was obtained at all the

annealing temperature. Beside, the XRD peaks

representing VO2 (FCC), V5O9, V3O5 and V2O5 were also


138 observed. This is similar to the results found in Fig. 3a

and that can be explained in a similar fashion.

Fig. 3 XRD patterns of VOx films prepared by (a) PVP

precursor solution and (b) PVOH precursor solution at

the annealing time 3 h.

Noteworthy, our result are different from that

was reported by Li et al. in a study on PAD of the VO2

with PVP [17]. In our opinion, the discrepancy could be

attributed to the fact that the different annealing

temperature were used. In addition, the type of furnace

also play role. It could be possible that by using the

different furnace, volumn of nitrogen gas flow used were

different and that affected the oxidation of vanadium

precursor. We belived that other vanadium oxides were

found during preparation of VO2(M) because in the V-O

phase diagram, there are more than 15 stable vanadium

oxide phases, such as VO, V6O13, and V7O13. The

formation of VO2 accounted for a small fraction due to its

narrow processing window in term of oxygen partial

pressure [22]. Only the M/R phase of VO2 shows

thermochromic properties while more than 10

polymorphs of VO2 lattice were found [23]. From the

above discussion, the amount of VOx crystal obtained by

using PVOH precursor was greater than that form in PVP

precursor solution. This was because PVOH has stronger

interaction with VO2+

than PVP.

Fig. 4 SEM images of VOx films prepared via PVOH

precursor solution (a) dried film, film annealed at (b) 450

oC, (c) 500

oC and 600

oC.

3.3 Morphologies of precursor gels and VOx films

The morphologies of VOx films prepared via

PVOH precursor solution were observed in Fig. 4. After

VOx films were dried at 60 oC, a smooth and uniform

films were observed without any microcracks or

precipitates. This was because the formation of a uniform

polymer mixed with VOCl2. A lot of pores among the

grain resulting from the degradation of polymer and

crystallization of VOx were observed after annealing

process [17]. The morphology of VOx films have been

changed due to the different annealing temperatures. For

the films annealed at 450 oC, the rod and irregular prisms

were appeared. When the annealing temperature was


139 increased the microstructures were changed as see in Fig.

4c. Moreover, the shrinkage of microstructures was

observed at the high annealing temperature.

4. Conclusion

Vanadium oxides films were prepared by

polymer assisted deposition (PAD), using PVP and

PVOH. It was found that the use of PVOH is better than

PVP due to a greater interaction with VO2+

. A variety of

vanadium oxides including VO2(M), VO2(B), VO2(FCC),

V3O5 and V5O9 were observed from XRD patterns of the

products. Temperature also had significant influence on

morphology of VOx films. The crystalline shape and

shrinkage changed when the temperature was increased.

In this study, the best condition that might be used to

prepare VO2 (M) film is that by using the PVOH as a

precursor solution annealed at 450 oC for 3 h.

5. Acknowledgements

This work has been supported by the

Nanotechnology Center (NANOTEC), NSTDA, Ministry

of Science and Technology, Thailand through its program

of the Center of Excellence Network.

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141

COMPP-25

Mechanical Properties and Thermal Resistance of Poly (Butylene Succinate) Reinforced

with Halloysite Nanotubes

Chanakarn Chucheapchuenkamol1 and Kalyanee Sirisinha

1,*

1Department of Chemistry, Faculty of Science, Mahidol University,

Phutthamonthon 4 Road, Salaya, Nakhon Pathom 73170.

Abstract

Poly (butylene succinate) (PBS)/halloysite nanotubes (HNTs) composites were fabricated via melt-

compounding in a twin-screw extruder. The effects of silanization of HNTs and stabilizer addition on the

mechanical and thermal properties of PBS were investigated using tensile and impact testings, heat distortion

temperature testing, differential scanning calorimetry, and thermogravimetric analysis. The results showed that

the reinforcement of PBS with HNTs increased both stiffness and toughness of PBS, without loss of tensile

strength. HNTs could serve as a nucleating agent for the PBS. The combined use of HNTs and mixed stabilizers

could enhance the thermal resistance of PBS significantly. The heat distortion temperature increased more than

10 C in the stabilized composite, compared to the neat PBS.

Keywords: Poly (butylene succinate), Composites, Halloysite nanotube, Properties

1. Introduction

In light of depleting landfill space and increasing

demand for disposal packaging usage, there is a need for

polymers that are biodegradable. Poly (butylene

succinate) or PBS is an aliphatic polyester which can be

degraded by hydrolysis pathway. Its properties are

similar to those of polyethylene. As with many

thermoplastics, PBS has a decreasing mechanical

strength with increasing temperature. Therefore, it would

be desirable to have biodegradable PBS-based products

which have greater resistance to deformation at higher

temperatures. This would be useful during storage and

transportation in summer time periods, and when

contacting with hot drink and hot food.

Halloysite nanotubes or HNTs are naturally

occurring inorganic nanotubes. They belong to kaolin

group of clay minerals and chemical formulation is

Al2(OH)4Si2O5.2H2O [1]. HNTs have high tensile

strength and modulus. They have been considered as

excellent reinforcing material for plastics. HNTs also

have lower density than other fillers such as talc and

calcite, making them convenient for preparing light-

weight polymer composites. Compared to carbon

nanotubes, the price of HNTs is much lower. Erdogan et

al. reported that the incorporation of 5% HNTs improved

the mechanical properties of polyamide-6 significantly

[2]. There is also a report on the effect of HNTs on

polypropylene where HNTs showed the direct stabilizing

effect on PP during thermal degradation [3,4].

This work aims to improve the mechanical

properties and thermal stability of PBS by using HNTs as

reinforcing filler. PBS/HNTs nanocomposites were

prepared via melt-compounding. The effects of HNTs

with and without surface modification on the properties

of the composites were compared. Also, the properties

changes as influenced by stabilizer addition were

included in this work. The composite properties were

characterized using universal tensile testing, impact

testing, thermogravimetric analysis, heat distortion

temperature testing, and differential scanning

calorimetry.

2. Experimental

2.1 Materials

Poly (butylene succinate), GS Pla FZ91PD, was

the product of Mitsubishi Chemical Corporation.

Halloysite nanotubes, DRAGONITE-XRTM

, were from

Applied Minerals Inc. Vinyl trimethoxysilane (VTMS)

was from Sigma-Aldrich. Two stabilizers used were

Irganox 1010 (Merit Solution Co., Ltd.) and Irgafos 168

(Ciba Specialty Chemicals.)

2.2 Surface modification of HNTs

HNTs were surface-treated with VTMS. The ratio

of VTMS to HNTs was 1:10. The silanization of HNTs


142 was carried out in a rounded-bottom flask containing a

solution of 500 ml of ethanol, 50 ml of distilled water

and 2 ml of acetic acid. The reaction was performed at

60 C for 2 hrs. After the reaction, the treated HNTs

were washed with distilled water and dried in an oven.

The dried HNTs were then grounded using a porcelain

mortar and pestle.

2.3 Preparation of PBS/HNTs composites

PBS nanocomposites containing 5% weight of

HNTs were prepared in a twin-screw extruder, using a

screw speed of 60 rpm. The temperatures from feed zone

to die were 120, 130, 140, 150, and 150 C, respectively.

The composites were cooled in a water bath before

pelletizing.

2.4 Preparation of test specimens

The specimens for tensile, impact, and HDT tests

were prepared by compression moulding at 165 C for 5

mins.

2.5 Characterizations

Tensile testing of PBS and PBS nanocomposites

was performed on a Universal tensile testing (INSTRON

5566, USA), according to ASTM D638, with a 1kN load

cell and a crosshead speed of 50 mm/mins. Tensile

strength, elongation at break and Young’s modulus were

recorded. The impact testing was conducted by Zwick

5102, USA. The notched samples were prepared

according to ASTM D256. The results are reported as

notched izod impact strength.

Differential scanning calorimetry measurements

were conducted. 7-9 mg of the samples was heated from

40 to 150 C with a heating rate of 20 C/mins and the

temperature was held at 150 C for 5 mins. Then, the

melted samples were cooled to 40 C with the same rate.

Heat distortion temperature testing was performed in the

edgewise position of a rectangular bar, according to

ASTM D648. 0.455 MPa of load was applied and raising

temperature was 2±0.2 C/min. The temperature at which

the specimen is distorted for 0.25 mm was recorded.

Thermal stability of the PBS composites was analyzed

on a thermogravimetric analyzer. Sample was heated

from 40 to 800 C at a heating rate of 10 C/mins.

Figure 1 (a) structure of HNTs, (b) surface modification

of HNTs


3.1 Tensile and impact properties of PBS/HNTs

composites

Figures 2a - 2d show tensile strength, modulus,

elongation at break and impact strength of pure PBS and

its composites with 5% wt HNTs. The effects of mixed

stabilizers (Irganox and Irgafos mixture) on those

properties are also included. Pure PBS shows a tensile

strength of 37.27 MPa, modulus of 542.42 MPa and high

elongation of nearly 200%. The incorporation of only 5%

wt HNTs could increase the modulus of PBS by 25%.

This confirms a reinforcing effect of HNTs in the

system. Tensile strength of PBS changes very slightly

after HNTs have been incorporated. However, a large

drop in extensibility of the composites is clearly seen. An

increasing in composite stiffness together with a drop in

elongation is commonly found in filled thermoplastic

systems. From the work of Leong et al. the incorporation

of talc resulted in significant drops in elongation at break

and increases of modulus of polypropylene. This

indicated that the incorporation of a filler restricts the

polymer chain mobility and deformability of the matrix

polymer [5].

Interesting results can be found for the effect of

HNTs on the impact strength of PBS. A significant

improvement in impact strength of PBS is achieved by

HNTs addition. The improvement in impact strength of

(b)

=

(a)


143 polymer by rigid filler could be due to a number of

reasons. In mineral-reinforced semicrystalline polymer,

the deformation processes that have been identified as

energy dissipating mechanisms include crazing,

cavitation or debonding of minerals followed by

microvoid formation and fibrillation [6]. In this study,

the addition of nanoscale HNTs is believed to alter the

micromechanism of deformation of PBS in the similar

ways. From the work of Yuan and co-workers, the

addition of nanoscale calcium carbonate to high density

polyethylene (HDPE) altered the deformation of HDPE

from crazing to fibrillation in the nanocomposite. This

led to an improvement in impact strength of HDPE by

approximately 15% for the system with 5% nanocalcium

carbonate [7].

In this work, the PBS has a notched Izod impact

strength of 4.50 KJ/m2. After 5% HNTs were added, the

impact strength of 7.34 KJ/m2 is resulted which is

approximately 65% greater than the unfilled one. These

results point out that the biodegradable PBS composites

with enhancement of both stiffness and toughness are

obtained by HNTs addition approach.

The results of Figure 2 also show that silane

treatment of filler surface has only minor effects on

tensile and impact properties of PBS. Slight

improvement of tensile strength and elongation at break

is found after silane modification. In the case of

composites with high filler loading, filler agglomeration

is usually occurred and this would result in a severe

dispersion problem. As a consequence, the composites of

poor properties are obtained. By treatment of filler

surface with appropriate chemical agents, better

dispersion of filler is promoted. From the previous work,

the mechanical properties of polyamide-6 were

significantly improved after surface-treated HNTs were

used [2]. In this study, the surface treatment of HNTs by

VTMS does not show a profound effect on the composite

properties due to two main reasons. The first reason is

that the amount of silane treated on the filler surface is

rather low. The second one is due to the fact that the

loading of HNTs in the PBS composites was only 5% wt.

Therefore, the problem of filler agglomeration is not so

intense in this study.

Figure 2 (a)tensile strength, (b)Young’s modulus,

(c)elongation at break, and (d)impact strength

(a)

PBS

PBS with 5% unmodified HNTS

PBS with 5% modified HNTS

PBS with 5% modified HNTS and stabilizer

(b)

(c)

(d)


144 Influence of stabilizers on the composite

properties is demonstrated in Figure 2. The mixture of

two stabilizers was used here with the aim to inhibit the

degradation of PBS during processing at high

temperature and promote long-term thermal stabilization.

Our results show that the nanocomposites containing

stabilizers have similar mechanical properties to those

without stabilizers.

3.2 Thermal properties of PBS/HNTs composites

Table 1 summarizes the DSC, HDT, and TGA

results. From the DSC analysis, the effects of HNTs on

crystallization temperature (Tc), onset temperature of

crystallization (Tc onset), melting temperature (Tm), and

total crystallinity (%Xc) of the PBS nanocomposites

were determined. HNTs of only 5%wt show strong

effects on crystallization behaviors of the PBS. Tc and Tc

onset increase significantly after HNTs were added to the

composites. This points out that HNTs act as nucleating

agent for crystallization of the PBS. The silane-treated

HNTs seem to have stronger effect on those properties

than the untreated one. This may be partly attributed to a

reduction in HNTs agglomeration in the modified

system. In the other words, more nucleating sites exist in

the composites of treated HNTs.

Compared to the pure PBS, the filled composites

exhibit lower values of crystallinity. This is mainly

owing to the dilution effect. However, when

recalculating the percentage of crystallinity based solely

on the PBS portion in the systems, one can see that all

composites exhibit similar amount of crystallinity (%

normalized Xc).

The effects of HNTs on thermal resistance of PBS

were investigated using HDT and TGA analysis. The

results are reported in the terms of the heat distortion

temperature (HDT), the onset (Td onset) and the

temperature of decomposition (Td). The results of Table

1 clearly reveal that the HDT of PBS obviously increases

from 85 to 95 C when 5% wt of HNTs was added. No

difference in the HDT values is found between the use of

surface-treated and untreated HNTs. Further

improvement of thermal resistance of the composites

could be done by adding stabilizers to the systems. The

combined use of HNTs and stabilizer results in a

composite with an HDT of 97 C, which is twelve

degrees higher than the neat PBS. Similar trend of

thermal stability enhancement can be seen from the TGA

results where the Td onset of the stabilized nanocomposites

shifts to a higher temperature of 346.5 C, compared to

that of pure PBS (325.0 C).

Table 1 Crystallization temperature (Tc), onset temperature of crystallization (Tc onset), melting temperature (Tm), total

crystallinity (%Xc), percentage of normalized Xc , heat distortion temperature (HDT), onset (Td onset) and temperature

of decomposition (Td)

Samples Tc

( )

Tc onset

( )

Tm

( )

Xc

(%)

Normalised Xc

(%)

HDT

( )

Td onset

( )

Td

( )

PBS 66.6 72.6 115.1 55.1 55.1 85 325.0 424.0

PBS with 5%

unmodified

HNT

81.6 86.6 115.1 52.8 55.6 95 339.0 411.0

PBS with 5%

modified HNT 83.2 88.2 114.9 53.0 55.8 95 339.0 404.9

PBS with 5%

modified HNT

and stabilizer

N/A N/A N/A N/A N/A 97 346.5 405.7


145 Unfortunately, opposite effect is found in the case

of the Td values. The Td values of PBS/HNTs composites

are lower than that of the neat PBS. A decrease of Td in

filled polymer composites has been reported earlier,

since filler can exhibit either a barrier or a catalytic effect

on polymer degradation [8,9].

4. Conclusions

Compounding of PBS nanocomposites containing

5%wt of HNTs was carried out in a twin-screw extruder.

The effects of HNTs addition, surface treatment of

HNTs, and presence of stabilizers on the mechanical

properties, thermal resistance, and crystallization

behaviors of PBS were investigated. Significant

enhancement of composite modulus and impact strength

was achieved by adding only 5% HNTs to the PBS.

HNTs could act as a nucleating agent for the

crystallization of PBS. Tc and Tc onset increase

significantly after HNTs were added. Silane treatment of

HNTs is believed to promote a better filler dispersion to

some extent. Not only the mechanical properties, but the

thermal resistance of PBS also improved by the

combined use of HNTs filler and mixed stabilizers. HDT

and Td onset of the stabilized composites are much

higher than those of the neat PBS.

Acknowledgements

The authors would like to thank Dr. Supakij

Suttiruengwong of Silpakorn University for providing

the PBS materials. The technical support from the

Research and Development Centre for Thai Rubber

Industry (RDCTRI) Mahidol University is greatly

appreciated.

References

[1] Rawtani D, Agrawal YK. Multifarious applications

of halloysite nanotubes: a review. Review Advanced

Material Science 2012; (30): 282-295.

[2] Erdogan AR, Kaygusuz I, Kaynak C. Influences of

aminosilanization of halloysite nanotubes on the

mechanical properties of polyamide-6

nanocomposites. Polymer Composites 2014; 35(7):

1350–1361.

[3] Prashantha K, Lacrampe MF, Krawczak P .

Processing and characterization of halloysite

nanotubes filled polypropylene nanocomposites

based on a masterbatch route: effect of halloysites

treatment on structural and mechanical properties.

eXPRESS Polymer Letters 2011; 5(4): 295-307.

[4] Wang B, Huang HX. Effects of halloysite nanotube

orientation on crystallization and thermal stability of

polypropylene nanocomposites. Polymer Degra-

dation and Stability 2013; 98(9): 1601–1608.

[5] Leong YW, et al. Comparison of the mechanical

properties and interfacial interactions between talc,

kaolin, and calcium carbonate filled polypropylene

composites. Journal of Applied Polymer Science

2004; 91(5): 3315–3326.

[6] Tanniru M, Yuan Q, Misra RDK. On significant

retention of impact strength in clay–reinforced high-

density polyethylene (HDPE) nanocomposites.

Polymer 2006; 47(6): 2133–2146.

[7] Yuan Q, et al. On Processing and Impact

Deformation Behavior of High Density Polyethylene

(HDPE)–Calcium Carbonate Nanocomposites.

Macromolecular Materials and Engineering 2009;

294(2): 141–151.

[8] Henrist C, et al. Toward the understanding of the

thermal degradation of commercially available fire-

resistant cable. Materials Letters 2000; 46(2-3):

160–168. [9] Zhao C, et al. Mechanical, thermal and

flammability properties of polyethylene/clay

nanocomposites. Polymer Degradation and Stability

2005; 87(1): 183–189.

SESSION 4Advances in Polymer Processing

Advances in Polymer Processing


147

KN-PROC-1

Foam, (Micro)Foam, (Nano)Foam! - Reality and Dream

Masahiro Ohshima

Department of Chemical Engineering, Kyoto University

Email: [email protected]

Abstract

The foaming technique has been contributing significantly to

material saving with lower emission of greenhouse gases.

Nowadays, microcellular foaming technique using CO2 and/or N2

as physical blowing agent attracts great attention from both

environmental and value-added material viewpoints. Microcellular

foams refer to polymer foams with cell size of the order of 10 m

and cell density higher than 108 cell/cm

3. Several physical foaming

techniques, such as solid state foaming, foam extrusion and foam

injection molding, have been developed to produce fine cellular

foams from various kinds of amorphous and semi-crystalline

polymers. To make this technology stable and productive in

industry, not only the foam processing techniques but also the

modifications of polymer properties and machine development

have been synergistically conducted. In this talk, the state of the art

in the world of fine cellular foaming filed is discussed by taking

several examples of Japanese industrial foam products and our

laboratory experimental results. Figure 1 illustrated some of the

recent developments of foams in Japanese Industries and our

laboratory.

The reality of the foam science and engineering is that the foaming

technology is an old technology but it is an advancing technology,

by which the foams with 5-100 m cell size can be produced and

the weight reduction and heat insulation performances are achieved

using environmentally benign foaming agents, N2 or CO2, instead

of HFC or butane. The high reflection of optical light and the high

sound energy absorption are new performances that the foam is

expected to have. Many attempts have been also conducted to

reduce the cell size further down to nanoscale levels. The cell size

could be reduced to nanoscale level, however, the expansion ratio

could not be kept high. The transparent plastic nanocellular

foam is still a dream.

Keywords: Microcellular Foam, Nanocellular Foam, Crystals

Nucleating Agent, Bubble Nucleating Agent, Fibrous Network,

Cell Structure.

Professor Ohshima started his

academic career as an Instructor of

Chemical Engineering at Kyoto

University in 1986, just after

graduating the Ph.D course of the same

University. Then year 1994, he

became an Associate Professor of

Computer Science and Systems

Engineering at Miyazaki University,

which is located in southern part of

Japan. Two years later, he returned to

Kyoto University and was promoted to

the full Professor in 2001. Since then,

he has been serving as a Professor of

Chemical Engineering at Kyoto

University and the leader of Material

Process Engineering laboratory. From

the beginning of his academic career,

he has devoted himself to researches in

both the process control and polymer

processing, especially polymer

foaming. He received several best

paper awards in both areas. . In 2011,

he obtained the technical award (Aoki

Katashi award) from Japan Society of

Polymer Processing (JSPP). He is a

Fellow of Society of Plastic Engineers

(SPE) and now served as the president

of JSPP.

.

a) Foamed Engine Cover and b), Foamed Drink bottle and c) PP foam with nanopores on the wall Its Cell morphology, its Cell morphology prepared by coreback injection mold

Fig. 1. Current microcellular Foams product

30％Weight reduction

Courtesy of Mazda & Daikyonishikawa

200μm

Ou

ter surfa

ce

Inn

er surfa

ce

Courtesy of Toyo Seikan

High light refection


148

KN-PROC-2

Fiber design: A creation of fiber structure for feature and performance

Chureerat Prahsarn

National Metal and Materials Technology Center,

National Science and Technology Development Agency, Pathumthani, 12120

Phone +66 2564 6500, Fax +66 2546 6446, Email: [email protected]

Abstract

Functional fibers are recognized for their performances

suitable for the desired applications. Their features and

performances are achieved via structural design and fabrication.

In this talk, examples of structural design concept in natural and

man-made fibers will be given. Some of our research work on

fibers’ structural design and their resulting properties, such as

soft-to-touch fabrics, self-crimped fibers, and nonwoven of

microfibers, will also be discussed, based on selection of

polymers, additives, and fiber processing techniques.

Keywords: functional fibers, bicomponent fibers, shaped

fibers, nonwovens

Chureerat Prahsarn

Researcher (Polymer, Textiles)

National Metal and Materials Technology Center

(MTEC)

National Science and Technology Development

Agency

Ministry of Science, Technology and Environment,

Thailand.

Education

2001 Ph.D. in Fiber and polymer science,

College of Textiles, North Carolina State

University, USA. 1995 M.S. in Macromolecular science,

Case Western Reserve University, USA. 1992 B.Sc. (Chemistry), Faculty of

Science, Khon Kaen University, Thailand

Research and Professional Experience

2002-Present Researcher at National

Metal and Materials Center (Thailand)

Nov.-Dec. 2009 Guest researcher at The

National Institute of Advanced Industrial

Science and Technology, AIST (Japan)

2005-2007 Guest researcher at

Mitsui Chemicals, Inc. (Japan)

1998-2001 Research Assistant (College

of Textiles, North Carolina State

University)

Specialize: Fiber processing, Bicomponent

fibers, Textiles (moisture transport in

fabrics), Nonwovens

Interest: Functional fibers and nonwovens,

Structural and property design in fibers,

Nonwovens for technical textiles, Comfort-

related moisture transport in fabrics


149

PROCO-01

Application of Genetic Algorithm in Identifying Ethylene/1-Olefin Copolymerization

Conditions from Molecular Weight Distribution and Chemical Composition

Distribution

Uthane Nanthapoolsub and Siripon Anantawaraskul*

Department of Chemical Engineering, Kasetsart University, 50 Phaholyothin Rd, Jatujak, Bangkok, Thailand

10900.

Abstract

For a known polymerization system, chain microstructures of polymers produced at a certain

polymerization conditions can be estimated using the polymerization kinetic model. However, an inverse

problem of identifying adequate polymerization conditions for producing polymers with desired chain

microstructures is rather complicated as mathematically it cannot be solved directly. In this work, genetic

algorithm (GA) was proposed as an optimization tool for identifying polymerization conditions. The considered

chain microstructures include molecular weight distribution (MWD) and chemical composition distributions

(CCD). The approach was validated with the model two-site-type catalytic system of ethylene/1-butene

copolymerization. The results showed that GA can help adequately identify the copolymerization conditions

from given MWD and CCD information.

Keywords: Chemical composition distribution (CCD), Ethylene/1-olefin copolymerization, Genetic algorithm

(GA), Molecular weight distribution (MWD), Polyethylene.

1. Introduction

The physical properties of polyethylene depend on

chain microstructures, such as average molecular weight,

molecular weight distribution (MWD, average

comonomer content, and chemical composition

distribution (CCD) [1-3]. Catalytic system and

polymerization conditions can strongly affect these chain

microstructures. Therefore, relationships between

polymerization conditions and chain microstructures are

very important for tailor-making polymer structures and

properties.

Several approaches based on polymerization kinetic

model can be used to describe the effect of

polymerization conditions on chain microstructures [4-

7]. However, the inverse problem on identifying

polymerization conditions for producing polymers with

desired chain microstructures is a rather complex

problem that cannot be solved directly.

Although average chain microstructural information

(e.g., average molecular weight, average comonomer

content, polydispersity index) were generally used in

industry, this information is insufficient for determining

the polymerization conditions. This is because there are

several forms of MWDs and CCDs (from different

polymerization conditions) that produce polymers with

the same average microstructural information. This can

easily lead to the multiple solution problems.

MWD and CCD, which provided extensive detail

on chain microstructures, could be more appropriate

information for determining polymerization conditions.

MWD can typically be obtained from gel permeation

chromatography (GPC) [8] and CCD can be obtained

from temperature rising elusion fractionation (TREF) [9-

10] or crystallization analysis fractionation (Crystaf)

[10]. To identify the polymerization conditions for

producing polymers with specified MWD and CCD, one

can perform a large scale optimization with the help of

polymerization kinetic model. The objective function to

be minimized is the sum of square error between the

simulated and experimental MWD and CCD.

Genetic algorithm (GA) is a global optimization tool

based on the natural evolution process (i.e., inherit,

crossover, and mutation) that can efficiently solve a large

scale problem by performing a global search without

relying on initial guesses. GA has been applied to several

problems in polymer science and engineering [11-15].

In this work, a genetic algorithm was used to

perform global search to determine appropriate


specified MWD and CCD. A series of the model data of


150 known ethylene/1-butene copolymerization with two-

site-type catalytic system were used to validate the

proposed method.

2. Mathematical modeling

Chain microstructures of copolymers produced on

each active site type were assumed to follow

Stockmayer’s bivariate distribution [17] (see Appendix

A and B). Flory’s distribution, which describes a weight

distribution function of kinetic chain lengths (r), can be

expressed as follows:

2( ) expj j jw r r r (1)

Where r is the chain length, j is number active site type,

j is related to the Mnj of site type j and calculated from

the polymerization reaction (see Appendix A and B).

Similarly, the CCD of polymer produced on each

site type can be described by using CCD component of

Stockmayer’s distribution as follows:

1 5/22

1 1

3( )

( - )4 2 1

2

j

j

j j

j j

w FF F

(2)

1 1 1 1 1 2(1- ) 1- 4 (1- )(1- )j j jj j j jF F F F r r (3)

F1 is the mole fraction of monomer, is the

average mole fraction of monomer made on site type j

and r1j and r2j are the reactivity ratios for

copolymerization.

Polymer chain microstructures produced on

multiple-site-type catalytic systems can be considered as

a mixture of polymer chain microstructures made on

each active site type. Therefore, the overall of MWD and

CCD were calculated by using the expressions:

1

( )n

j j

j

W r m w r

(4)

1

1

( )n

j j j

j

W F m w F

(5)

where n is the number of active site types and mi is the

mass fraction of polymers produced on site type j.

Genetic algorithm (GA) was developed to

minimize the objective function. The initial population

was randomly generated by uniform distribution in

search space. The number of population size of each

generation was 300 individuals for increasing

opportunity to search the best individual. Each individual

was a set of strings of estimated parameters (mi, [Co-cat],

[M1], [M2], [H2], Temp.). The search space was adjusted

to be in the range from 0 to 1 for all parameters. The

total parameter in each individual depends on the number

of catalytic size type as 5 + (j-1) because total

summation of mass fraction is equal to 1.

The MWD and CCD chain microstructures were

calculated from the individual by using rate reaction

equations of coordination polymerization model (see

Appendix A and B) and the individuals were evaluated

by the objective function as fitness function from chain

microstructures (MWD and CCD).

The objective function (i.e., the sum of the

squares of differences between model profiles and

simulation results) to be minimized is

1

2

model simulation

,

1Objective function

r F

W WN

(6)

where N is the total number of data for MWD or CCD. If

the criteria are not acceptable, this population was set as

a parent generation to generate the next generation using

three mechanisms: selection, crossover, and mutation.

The tournament selection was used to select the

parent for the crossover step. The individual as a player

was competed in tournament by using the objective

function. 4 randomly chosen individuals together with

individuals having the best fitness value (or lowest

objective function) were selected as parents in the

crossover step and maintained at 300.

In crossover step, 50% of parent is randomly

chosen for the crossover process. Heuristic crossover

was applied to create population of next generation. The

child was generated on the direction of the search

between two parents. It can specify distant of child from

the better parent by ratio (R=1.2). If Parent1 is better the

fitness value than Parent2. The child are created

according to


151 Child1 = Parent1

Child2 = Parent2 +R*(Parent1-Parent2)

where child1 and child2 are the individual in next

generation.

The mutation process helps increasing

opportunity to search space and prevents stagnation in

the local minimum [18]. The adaptive feasible mutation

was applied to create the next generation. The randomly

number was added to generate the child population in the

feasible bound. The directions are adaptive with respect

to the last successful or unsuccessful generation. The

new generation is then reevaluated and the process is

repeated until the process reaches 50,000 generations or

the value of objective function is less than 1x10-15

.


Genetic algorithm (GA) was applied to identify


specific information of chain microstructures (MWD and

CCD) in two-site-type system. The polymerization

conditions of model dataset are given in Table 1. It was

found that the conditions identified from GA (the

simulation results) are very close to the model dataset.

Figure 1 showed comparison of MWD and CCD results

obtained from model and simulation (GA estimated)

results.

Table 1 Polymerization condition parameters between

model and simulation results of ethylene/1-butene

x Model Simulation results

(GA estimation)

mi 0.4 0.400

[Co-cat] 1.0x10-3

0.988x10-3

[M1] 7.00 7.001

[M2] 3.00 2.997

[H2] 1.00 0.999

Temp. (C) 80 79.9

Objective function 1.4870x10-24

Figure 1. MWD and CCD of model and simulation

results.

3. Conclusions

Genetic algorithm (GA) was successfully

applied to determine ethylene/1-butene copolymerization

conditions for producing polymers with desired

molecular weight distribution (MWD) and chemical

composition distribution (CCD) in two-site-type system.

This new approach was validated with model

polymerization dataset. It was found that GA can

estimate polymerization conditions that are very close to

the model parameters.

Appendix A: Polymerization Mechanism

The terminal model was used to describe the

polymerization mechanism in a CSTR. The type of

active site on a multiple-site catalyst is indicated by (j) in

these equations.

Site activation:

)(0

)(

)( j

jf

j Pk

AC Polymer chain initiation:

)(1

)(1,

1)(0 j

ji

jP

kMP

)(2

)(2,

2)(0 j

ji

jP

kMP

)(1

)(1,

1)( j

ji

jH Pk

MP

3 3.5 4 4.5 5 5.5 60

0.5

1

1.5

log(MW)

Wei

gh

t

MWD

Model

Simulation result

0 1 2 3 4 5 60

0.5

1

1.5

% mole of comonomer

Wei

gh

t

CCD

Model

Simulation result


152

)(2

)(2,

2)( j

ji

jH Pk

MP

Propagation:

)(1

)(11,

1)(1 j

jp

jP

kMP

)(2

)(12,

2)(1 j

jp

jP

kMP

)(1

)(21,

1)(2 j

jp

jP

kMP

)(2

)(22,

2)(2 j

jp

jP

kMP

Chain transfer to monomer:

DPk

MPj

jM

j

)(1

)(11,

1)(1

DPk

MPj

jM

j

)(2

)(12,

2)(1

DPk

MPj

jM

j

)(1

)(21,

1)(2

DPk

MPj

jM

j

)(2

)(22,

2)(2

Chain transfer to hydrogen:

DPk

HPjH

jH

j

)(

)(1,

2)(1

DPk

HPjH

jH

j

)(

)(2,

2)(2

Chain transfer to co-catalyst:

DPk

APj

jA

j

)(0

)(1,

)(1

DPk

APj

jA

j

)(0

)(2,

)(2

β-Hydride elimination:

DPk

PjH

j

j

)(

)(1,

)(1

DPk

PjH

j

j

)(

)(2,

)(2

Catalyst deactivation:

DCk

Pjd

jd

j

)(

)(

)(0

DCk

Pjd

jd

jH )(

)(

)(

DCk

Pjd

jd

j

)(

)(

)(1

DCk

Pjd

jd

j

)(

)(

)(2

Table A1. Kinetic constants of ethylene/1-butene

copolymerization at T = 360K

Mechanism Kinetic

constant Site 1 Site 2

Site activation

(L/mol.s)

kf 2.6E-01

2.6E-01

Initiation ki,1 2.8E+02 2.8E+02

(L/mol.s) ki,2 2.6E-01 2.6E-01

Propagation kp,11 6.8E+00 4.2E+00

(L/mol.s) kp,12 4.2E-01 2.1E-01

kp,21 3.4E+00 6.8E+00

kp,22 4.2E-01 4.2E-01

Transfer to kM,11 1.7E-03 1.3E-04

monomer kM,12 8.4E-03 5.1E-04

(L/mol.s) kM,21 8.4E-04 1.0E-04

kM,22 8.4E-03 5.1E-04

Transfer to H2 kH,1 8.2E-03 4.2E-02

(L/mol.s) kH,2 2.6E-03 1.0E-02

Transfer to kA,1 8.9E+00 2.8E+00

cocatalyst

(L/mol.s)

kA,2 2.2E+00

5.5E-01

β-Hydride kβ,1 1.4E-11 1.8E-11

Elimination

(1/s)

kβ,2 3.5E-12

3.5E-12

Deactivation

(1/s)

kd

5.5E-16

5.5E-16

Appendix B: Model Development

Pseudo-kinetic polymerization

The simple rate of polymerization per site type is given

by the equation:

( ) ( ) ( )[ ]p j p j jR k M

Therefore, the polymerization yield for each site type in

a CSTR with residence time equal to tr is simply given

by:

( ) ( ) ( )[ ]j p j j rY k M t

Finally, the mass fraction of polymer produced by each

catalyst site type is given by the expression:

( )

( )

( )

1

j

j n

i

i

Ym

Y


153 where n is the total number of active site types in the

reactor.

Molecular weight distributions (MWD)

The number average chain length for the polymer made

on site type j is related to the polymerization kinetic

constants by the expression:

( ) ( ) ( ) ( ) 2

( ) ( )

[ ] [ ] [ ]1

[ ]

j A j M j H j

n j p j

k k A k M k H

r k M

( )

1j

j

n j n j

mw

M r

Chemical Composition Distributions (CCD)

The chemical composition component of Stockmayer’s

bivariate distribution for each site type is related to the

polymerization kinetic constants by the expression:

)(21,

)(22,

)(2

)(12,

)(11,

)(1 ,

jp

jp

j

jp

jp

jk

kr

k

kr

and F1 is the molar fraction of comonomer type 1 in the

copolymer, and )(1 jF is the average molar fraction

monomer type 1 in the copolymer made on site type j,

given by the Mayo-Lewis equation:

)(21)(2

2

1)(2)(1

1

2

1)(1)(1

)1(2)2(

)1(

jjjj

jj

rfrfrr

ffrF

References

[1] Ferdinand Rodrigues, Claude Cohen, Christopher K.

Ober and Lynden A. Archer., PRINSIPLES OF

POLYMER SYSTEMS. 5th edition. United States

of America: Taylor & Francis Books, Inc., (2003).

[2] Andrew J. Peacock. HANDBOOK OF

POLYETHYLENE Structures, Properties and

Applications. United States of America :Marcel

Dekker, Inc., (2000).

[3] Shirayama, Kenzo, Shin-Ichiro Kita and Hiroshi

Watabe. Effect of Branching on some Properties of

Ethylene/alpha-Olefin Copolymers. Die

Makromolekulare Chemie, 151, pages 97-120

(1971).

[4] H. Hatzantonis, H. Yiannoulakis, A. Yiagopoulos,

and C. Kiparissides, "Recent developments in

modeling gas-phase catalyzed olefin polymerization

fuidized-bed reactors: The efect of bubble size

variation on the reactor's performance," Chemical

Engineering Science, vol. 55, pp. 3237-3259,

(2000).

[5] Bo Kou, Kim B. McAuley, C. C. Hsu, David W.

Bacon, and Zhen K. Yao, "Mathematical Model and

Parameter Estimation for Gas-Phase Ethylene

Homopolymerization with Supported Metallocene

Catalyst," Industrial & Engineering Chemistry

Research, vol. 44, pp. 2428-2442, (2005).

[6] Antnio G. Mattos Neto, Marcelo F. Freitas, Mrcio

Nele, and Jos Carlos Pinto, "Modeling Ethylene/1-

Butene Copolymerizations in Industrial Slurry

Reactors," Industrial & Engineering Chemistry

Research, vol. 44, pp. 2697-2715, (2005).

[7] Tuyu Xie, Kim B. McAuley, James C.C. Hsu, and

David W. Bacon, "Modeling Molecular Weight

Development of Gas-Phase Alpha-Olefin

Copolymerization," AIChE Journal, vol. 41, pp.

1251-1265, (1995).

[8] Y. V. Kissin, "Molecular Weight Distributions of

linear Polymers: Detailed Analysis from GPC

Data ," Journal of Polymer Science part A: Polymer

Chemistry, vol. 33, pp. 227-237, (1995).

[9] J. B. P. Soares and A. E. Hamielec, "Temperature

rising elution fractionation of linear polyolefins,"

Polymer, vol. 36, no. 8, pp. 1639-1654, (1995).

[10] Siripon Anantawaraskul, João B. P. Soares, and

Paula M. Wood-Adams, "Fractionation of

Semicrystalline Polymers by Crystallization

Analysis Fractionation and Temperature Rising

Elution Fractionation," Advances in Polymer

Science, vol. 182, pp. 1-54, (2005).

[11] Gujarathi, A. M., & Babu, B. V. Multiobjective

Optimization of Industrial Processes Using Elitist

Multiobjective Differential Evolution (Elitist-

MODE). Materials and Manufacturing Processes,

26, pp 455-463, (2011).


154 [12] Mu, Y., Zhao, G., Wu, X., & Zhang, C. An

optimization strategy for die design in the low-

density polyethylene annular extrusion process

based on FES/BPNN/NSGA-II. Journal of

Advanced Manufacturing Technology, 50, pp 517-

532, (2010).

[13] Nanthapoolsub, U., Anantawaraskul, S., &

Saengkhamkhom, K. Simultaneous Deconvolution

of MWD and CCD of Ethylene/1-Olefin

Copolymers Using Genetic Algorithm.

Macromolecular Symposium, 330, pp142-149,

(2013).

[14] Padhiyar, N., Bhartiya, S., & Gudi, R. D. Optimal

Grade Transition in Polymerization Reactors: A

Comparative Case Study. Industrial & Engineering

Chemistry Research, 45, pp 3583-3592, (2013).

[15] Yousefi, F., & Karimi, H. Application of equation

of state and artificial neural network to prediction of

volumetric properties of polymer melts. Journal of

Industrial and Engineering Chemistry, 19,pp 498-

507, (2013).

[16] Soares, J. B. Mathematical modelling of the

microstructure of polyolefins made by coordination

polymerization: a review. Chemical Engneering

Science, 56, pp 4131-4153, (2001).

[17] Stockmayer, W. H. Distribution of Chain Lengths

and Compositions in Copolymers. The Journal of

Chemical Physics, 13, pp 199-207, (1945).

[18] Singh, Gurmeet, Sukhdeep Kaur, Dhananjay G.

Naik, Virendra K. Gupta. Evolutionary Computing

Approach for Evaluating Flory Distribution Curves

in Gel Permeation Chromatography:Study of the

Poly(1-octene) System. Journal of AppliedPolymer

Science, Vol. 117, pp 3379-3385, (2010).


155 PROCO-02

Determination of Polymerization Conditions for Producing Ethylene/1-olefin

Copolymers with Tailor-made Chain Microstructures using Artificial Neural Network

Thanutchoke Charoenpanich1, Siripon Anantawaraskul

1,2* and João B. P. Soares

3

1Department of Chemical Engineering, Faculty of Engineering, Kasetsart University, Bangkok, Thailand 10900

2Center for Advanced Studies in Nanotechnology and Its Applications in Chemical, Food and Agricultural

Industries, Kasetsart University, Bangkok, Thailand 10900 3Department of Chemical and Materials Engineering, University of Alberta, Edmonton, Alberta, Canada T6G

2V4

Abstract

Two artificial neural network (ANN) models were developed for describing ethylene/1-butene

copolymerization with two-site-type catalytic system by using the datasets calculated from copolymerization

kinetic model. The forward model was first applied to mimic the kinetic model which predicts chain

microstructures (i.e., molecular weight distribution and chemical composition distribution) from defined

polymerization conditions (i.e., ethylene concentration, 1-butene concentration, cocatalyst concentration,

hydrogen concentration, and polymerization temperature). The results are in a good agreement with the

theoretical results. The inverse model was applied to predict the polymerization conditions from the desired

chain microstructures. Although the results showed large deviations due to the multiple patterns of solution, the

unique solution can be obtained by defining one of three key variables (ethylene concentration, 1-butene

concentration, and polymerization temperature) as a constant.

Keywords: artificial neural network (ANN), chemical composition distribution (CCD), copolymerization

kinetic model, ethylene/1-butene copolymerization, molecular weight distribution (MWD)

1. Introduction

Linear low density polyethylene (LLDPE),

synthesized by copolymerization of ethylene with α-

olefin, is one of the most widely used commodity

polymers. Its properties can be related to chain

microstructures that depend on copolymerization

conditions and catalytic system. Polymers with high

molecular weights and narrow molecular weight

distributions can be synthesized by using metallocene

catalysts, which have single-site-type nature. A

combination of two metallocene catalysts can be used for

controlling structure of polyolefin with high versatility.

Although the kinetic model of ethylene/1-olefin

copolymerization (typically as a system of ordinary

differential equations or algebraic equations) can be used

to estimate the chain microstructures from a specific set

of polymerization conditions in a given catalytic

system[1]

, it cannot be used in an inverse function to

determine appropriated conditions to yield desired chain

microstructures.

An artificial neural network (ANN) is a model

which mimics a pattern recognition process of the

nervous system. This model can be used to solve highly

non-linear problems by learning from examples, like a

human being. ANN needs only input and output datasets

and then can find optimum relationships between them

by training and testing process. Recently, the ANN

models have been applied to solve various problems in

polymer applications.[2, 3]

This non-phenomenological

model can predict the results with a high accuracy.

In this work, ethylene/1-butene copolymerization

with two single-site-type catalytic system was studied.

ANN was applied as both “forward model” and “inverse

model”. The forward model was used to predict the chain

microstructures obtained from a specific set of

polymerization conditions, mimicking the kinetic model.

On the other hand, the inverse model was used to

determine the polymerization conditions to produce

copolymers with desired chain microstructures. This

model can be used to tailor-make the final properties of

polymers by controlling polymerization conditions. The

steady state kinetic model of ethylene/1-butene

copolymerization was used to provide the input and

output datasets for training and testing process of ANN.


156 2. Model development

2.1 Copolymerization mechanisms

The general mechanisms of ethylene/1-butene

copolymerization are shown in Equation (1) to (23).[1]

Site activation:

( ) 0( )

( )f

j j

jkC A P

(1)

Polymer chain initiation:

0 ( ) 1 1( )

,1( )

j j

i jkP M P

(2)

0 ( ) 2 2 ( )

,2( )

j j

i jkP M P

(3)

( ) 1 1( )

,1( )

H j j

i jkP M P

(4)

( ) 2 2 ( )

,2( )

H j j

i jkP M P

(5)

Propagation:

1( ) 1 1( )

,11( )

j j

p jkP M P

(6)

1( ) 2 2 ( )

,12( )

j j

p jkP M P

(7)

2 ( ) 1 1( )

,21( )

j j

p jkP M P

(8)

2 ( ) 2 2 ( )

,22( )

j j

p jkP M P

(9)

Chain transfer to monomer:

1( ) 1 1( )

,11( )

j j

M jkP M P D

(10)

1( ) 2 2 ( )

,12( )

j j

M jkP M P D

(11)

2 ( ) 1 1( )

,21( )

j j

M jkP M P D

(12)

2 ( ) 2 2 ( )

,22( )

j j

M jkP M P D

(13)

Chain transfer to hydrogen:

1( ) 2 ( )

,1( )

j H j

H jkP H P D

(14)

2 ( ) 2 ( )

,2( )

j H j

H jkP H P D

(15)

Chain transfer to cocatalyst:

1( ) 0 ( )

,1( )

j j

A jkP A P D

(16)

2 ( ) 0 ( )

,2( )

j j

A jkP A P D

(17)

β-Hydride elimination:

1( ) ( )

,1( )

j H j

jkP P D

(18)

2 ( ) ( )

,2( )

j H j

jkP P D

(19)

Catalyst deactivation:

0 ( ) ( )

( )

j d j

d jkP C D

(20)

( ) ( )

( )

H j d j

d jkP C D

(21)

1( ) ( )

( )

j d j

d jkP C D

(22)

2 ( ) ( )

( )

j d j

d jkP C D

(23)

Where j represents the active site of each catalyst

(j=1, 2), M1 and M2 represent ethylene monomer and 1-

butene comonomer, respectively.

2.2 Molecular weight distribution (MWD)

According to Flory's most probable distribution,

the MWD with log scale can be given by:

2

( )( ) ( )

(log ) 2.3026( ) exp(- ) jn j n j

MW MWw MW

M M

(24)

where: mMW rW

(25)

( ) ( ) ( )n j n j jM r mw

(26)

( ) 1 1 2 2( ) ( )

( ) ( )j j jmw F mw F mw (27)

Note that r is the chain length; Wm is an average

weight of repeating unit; mw1 and mw2 are the molecular

weight of ethylene and 1-butene, respectively; rn(j) is the

number average chain length:

2( ) ( ) ( ) ( )

( ) ( )

[ ] [ ] [ ]1

[ ]

j A j m j H j

n j p j

k k A k M k H

r k M

(28)

For a mixture of Ns site types, the MWD for

copolymer can be given by:

( ) ( )

1

(log ) (log ) sN

j j

j

W MW m w MW

(29)

where m(j) is the mass fraction of polymer product by site

type j

2.3 Chemical Composition Distribution (CCD)

CCD of each active site can be calculated by

using the chemical composition component of

Stockmayer’s bivariate distribution:

1 ( ) 2

1 1 ( )( ) ( ) 5/2

( ) ( )

3( )

( - )4 2 [1 ]

2

j

n jj j

n j j

w FF F r

r

(30)


157 where:

( ) 1 1 1 1 1( ) 2( )( ) ( ) ( ) ( )(1- ) 1- 4 (1- )(1- )j j jj j j j

F F F F r r

(31)

2

1( ) 1 1

1 2( ) ( ) 2

1( ) 2( ) 1 2( ) 1 2( )

( -1)1

( 2) 2(1 )

j

j j

j j j j

r f fF F

r r f r f r

(32)

11( )

1( )

12( )

p j

j

p j

kr

k

, 22( )

2( )

21( )

p j

j

p j

kr

k (33)

For a mixture of Ns site types, similar expression

is used to obtain the CCD shown below:

1 ( ) 1 ( )

1

( ) ( ) sN

j j

j

W F m w F

(34)

2.4 Artificial neural network (ANN)

An artificial neural network (ANN) is one of

black-box models that can solve highly non-linear

problem. ANN is a model that learns from examples (i.e.,

various datasets) likes a human brain; it can be used to

find empirical relationships between input and output

variables. Furthermore, this model can identify and

respond to patterns, which are similar but have never

been trained before.

In this study, the feed forward neural network

with back propagation training algorithm was applied for

two cases as a “forward model” and an “inverse model”.

The forward model was used to mimic the

copolymerization kinetic model by predicting chain

microstructures from a specific set of polymerization

conditions. The scheme of this model was shown in

Figure 1. The considered polymerization conditions,

which include the ethylene concentration, 1-butene

concentration, cocatalyst concentration, hydrogen

concentration, and polymerization temperature, were

considered in an input layer. The chain microstructures

including weight fractions of MWD and CCD calculated

from Equations (29) and (34) were considered in an

output layer.

The inverse model was used to determine the

polymerization conditions from the desired chain

microstructures, as shown in Figure 2. Table 1 showed

the kinetic parameters used in this study which obtained

from literatures[4]

. The range of polymerization

conditions was summarized as shown in Table 2.

Figure 1. Forward ANN model scheme

Figure 2. Inverse ANN model scheme

The datasets were randomly separated into two

groups, which were 85% for training and 15% for

testing. All weights and biases were first defined

randomly in every interconnecting of neurons. The

training datasets were now trained in the model to get the

adequate weighting system. Subsequently, the testing

datasets were applied as unknown datasets to check

accuracy in the adequate system. These can prevent for

overtraining of the network.


3.1 Forward model

The forward model was developed to mimic the

kinetic model of ethylene/1-butene copolymerization,

predicting MWD and CCD from a given set of

polymerization conditions. The optimum topology of the

forward model is 5-50-50-30. The ANN prediction

results for MWD and CCD were shown in Figure 3 and

Figure 4, respectively.


158 Table 1. Kinetic constants of ethylene/1-butene

copolymerization at T = 360K[4]

Mechanism Kinetic

constant Site 1 Site 2

Site activation kf 1 1

Initiation ki,1 1 1

ki,2 0.14 0.14

Propagation kp,11 8.5 8.5

kp,12 2 15

kp,21 64 64

kp,22 1.5 2.26

Transfer to kM,11 0.0021 0.0021

monomer kM,12 0.006 0.11

kM,21 0.0021 0.0021

kM,22 0.006 0.11

Transfer to H2 kH,1 0.088 0.37

kH,2 0.088 0.37

Transfer to kA,1 0.024 0.12

cocatalyst kA,2 0.048 0.24

β-Hydride kβ,1 0.0001 0.0001

elimination kβ,2 0.0001 0.0001

Deactivation kd 0.0001 0.0001

Table 2. Range of polymerization conditions used to

train ANN model

Polymerization

condition

Unit Minimum Maximum

[Ethylene] mol L-1

0.2 3.74

[1-Butene] mol L-1

0.03 1.05

[Hydrogen] mol L-1

0.0001 0.01

[Cocatalyst] mol L-1

0.0003 0.0143

Temperature °C 70 90

Figure 3 compares MWD obtained from kinetic

model and MWD obtained from ANN model at sampling

data points. The results were expressed in form of

normalized data within the range between -1 to 1.

Similar to Figure 3, Figure 4 compares CCD obtained

from kinetic model and CCD obtained from ANN model

at sampling data points. In both cases, both of training

and testing datasets were closed to diagonal line, which

indicates an excellent agreement between kinetic model

and ANN model.

Figure 3. Comparison between MWD obtained from

kinetic model and ANN model, sampling data point at

W(log(MW))1, 9, 15

Figure 4. Comparison between CCD obtained from

kinetic model and ANN model, sampling data point at

W(F1)1, 9, 15

3.2 Inverse model

The inverse model was applied to determine the

conditions to produce the desired MWD and CCD.

Figure 5 shows the comparison between expected

polymerization conditions and estimated polymerization

conditions from ANN, indicating unexpected spreading

from diagonal lines. Although the topology was changed,

the errors cannot be decreased. It was found that the


159 source of these errors is multiple solution problems;

some of MWD and CCD can be obtained from various

polymerization conditions. Therefore, the ANN cannot

identify the unique solution.

To solve multiple solution problems, some of key

conditions must be fixed as a constant. It was found the

errors can be minimized when ethylene concentration, 1-

butene concentration, or the polymerization temperature

was fixed. Examples of these results were shown in

Figure 6 which fixing ethylene concentration at 3.74 mol

L-1

. The results were now much closer to the diagonal

line.

Figure 5. Comparison between polymerization

conditions obtained from kinetic model and ANN model

Figure 6. Comparison between polymerization

conditions obtained from kinetic model and ANN model

when ethylene concentration is kept constant at 3.74 mol

L-1

The investigation was found that those three

variables had more influence for synthesizing MWD and

CCD than the hydrogen and cocatalyst concentrations.

Hence, the multiple patterns can be caused by those three

variables and can be reduced by fixing one variable as a

constant.

3.3 Case study

The polymerization conditions used in the case

study were defined in Table 3. This case reduced the

multiple patterns by fixing polymerization temperature at

80°C. The forward model was first applied to estimate

MWD and CCD, as shown in Figure 7. This illustrated

the good agreement of ANN prediction comparing with

kinetic model. Then the inverse model was applied to


160 predict the polymerization conditions from these

distributions. The results were compared with the first

defined conditions as shown in Table 3. All variables

were closed to the defined conditions which validate the

concept of ANN model. The model can be further

improved in the future by training with more refined

dataset.

Table 3. Polymerization conditions defined in the case

study, compared with results from inverse model which

fixing polymerization temperature at 80°C

Condition [Ethylene]

(mol L-1)

[1-Butene]

(mol L-1)

[Hydrogen]

(mol L-1)

[Cocat]

(mol L-1)

Defined 1.720 0.322 0.0080 0.0020

Inverse model 1.724 0.324 0.0079 0.0019

Figure 7. Case study of forward model comparing with

kinetic model using the defined conditions in Table 3, a)

MWD, b) CCD

4. Conclusion

An ANN model is proposed as an alternative

approach for describing ethylene/1-butene

copolymerization system. The forward ANN model was

used to predict MWD and CCD from the polymerization

conditions, which agree well with the kinetic model. The

inverse ANN model was applied to determine conditions

to produce a specific MWDs and CCDs. The results

exhibited a large deviation due to the multiple solutions.

However, the unique solution can be obtained by fixing

one of three parameters as a constant. These key

parameters include ethylene concentration, 1-butene

concentration, and polymerization temperature.

Acknowledgements: Mr. Thanutchoke Charoenpanich

thanks The Graduation School, Kasetsart University and

The International Affairs Division, Kasetsart University

for their kind financial supports.

References

[1] Soares J. B. P., “Mathematical modelling of the

microstructure of polyolefins made by coordination

polymerization: a review”, Chem. Eng. Sci. (2001),

56, 4131.

[2] Anantawaraskul S., Toungsetwut M. and Pinyapong

R., “Determination of Operating Conditions of

Ethylene/1-Octene Copolymerization Using

Artificial Neural Network (ANN)”, Macromol.

Symp. (2008), 164, 157.

[3] Anantawaraskul S. and Chokputtanawuttilerd N.,

“Estimation of Average Comonomer Content of

Ethylene/1-Olefin Copolymers Using Crystallization

Analysis Fractionation (Crystaf) and Artificial

Neural Network (ANN)”, Macromol. Symp. (2009),

282, 150.

[4] Shawa B. M., McAuley K. B. and Bacon D. W.,

“Simulating Joint Chain Length and Composition

Fractions from Semi-Batch Ethylene

Copolymerization Experiments”, Polym. React. Eng.

(1998), 6(2), 113.


161

PROCO-03

Simulation of Morphological Development during Polymer Crystallization: Effect of

Temperature Gradient on the Crystallization Kinetics

Tharinee Teangtae, Siripon Anantawaraskul, Thitiporn Sooksod, and Chattong Pornpiriyayotha

Department of Chemical Engineering, Faculty of Engineering, Kasetsart University

Bangkok, Thailand, 10900

Abstract

Industrial processes of polymer solidification to form articles with desired shape often involve

crystallization under temperature gradient. During the crystallization, polymer morphology is developed. As

crystallization temperature affects both nucleation and crystal growth process, the temperature gradient could

significantly influence the crystallization kinetics and polymer morphology. This work investigated effect of

temperature gradient on crystallization kinetics and developed the theoretical parameter to evaluate the

crystallization rate. The results showed that temperature gradient significantly affects the crystallization kinetics,

average crystallite size and crystallite size distribution. The average Avrami rate constant corresponding to

relative crystallinity; therefore, it can be used to quantitatively evaluate the crystallization rate, especially at the

early crystallization.

Keywords: Avrami rate constant; Crystallize morphology; Crystallization Kinetics; Monte Carlo simulation

1. Introduction

The crystallization of semi-crystalline polymers is

a major field in polymer physics because the properties

of polymer products strongly depend on the morphology

formed and the extent of crystallization occurring during

the processing. [1]

Generally, the polymer crystallization

process consists of two steps: nucleation and crystal

growth. In the first step, nuclei are formed with a critical

size when the thermodynamic conditions are satisfied. [2]

Subsequently, the nuclei are growth as crystallites on the

crystalline region by adding other chains segment to the

nuclei center. Both of nucleation and crystal growth rate

are a function of crystallization temperature; therefore,

processing conditions can strongly influence

crystallization kinetics and final morphology.

In the industrial process, the crystallization

temperature not only depends on time for complete

crystallization but also depends on the position within a

bulk. Hence, the crystallization of polymer often relates

to temperature gradient, leading to complicate

crystallization kinetics. [3]

In this work, the effect of

temperature gradient on crystallization kinetics and

morphology was investigated in 2D simulation using

Monte Carlo model. The theoretical parameter to

evaluate the rate of crystallization was also developed.

2. Theoretical background

Avrami model is proposed to describe the kinetics

of polymer crystallization under an isothermal condition.

[4-5] The relative crystallinity, θ(t), is a function of

crystallization time, t, as follows:

1- exp - 0,1an

at k t

(1)

where ka is Avrami rate constant and na is Avrami

exponent (related to crystal geometry). The parameter ka

is related to the total concentration of the predetermined

nuclei (Ntot) and the crystal growth rate (G) according to

the following equation:

2

a totk N G (2)

To investigate the kinetics for semi-crystalline

polymer, the spherulite growth rate is assumed to depend

on only temperature and follows the Laurizen-Hoffman

theory:

*

0exp exp

g

c c

KUG G

R T T T T f

(3)

where G0 is a temperature-independent pre-exponential

factor, U*

is the activation energy for the transportation

of segments of molecules across the melt/solid surface

boundary (usually given a value of 1500 cal/mol), T

signifies the cessation of long-range molecular motion

(i.e.,g

T =T -30

, where Tg is the glass-transition


162 temperature), R is the universal gas constant, ΔT is the

degree of undercooling (i.e., 0

m cΔT=T -T , where

0

mT is

the equilibrium melting temperature and Tc is the

crystallization temperature), f is a correction factor for

the temperature dependence of the enthalpy of fusion

(i.e., 0

c c mf=2T / T +T ) and

gK is the nucleation

exponent. Practically, the temperature dependent of a

k

can be described similar shown as eq. (4):

0exp exp

G

c

c c

KT

R T T T T f

(4)

where c

Ψ T and 0

Ψ are the overall crystallization rate

parameter (a

k ) and the pre-exponential parameter (a0

k ),

respectively, is a parameter relates to the activation

energy characterizing the molecular transport across the

melt/solids interface and G

K is a combined factor

related to the secondary nucleation mechanism.

Recently, the isothermal melt crystallization kinetics of

s-PP was investigated with DSC. [6]

The overall

crystallization kinetics was estimated by the direct fitting

of the experimental data with the Avrami model. The

temperature dependence of a

k was described by eq. (4)

over the c

T ranging of 40-90 oC:

6

26 1807.1 1.39 107.08 10 exp exp

273 441.6c c

a

c

kT T T f

(5)

wherea

k has a unit of min-1

and f is c

2T / Tc+441.8 . In

addition, the spherulitic growth kinetics of s-PP in grime

III (over the c

T range of 45-110 oC) was used to

approximate the temperature dependence of G over the

whole investigated c

T range:

5

8 754.8 3.6 109.1 10 exp exp

273 441.8c c c

GT T T f

(6)

where G has a unit of µm/min and f is c2T / Tc+441.8 .

The crystallization kinetics of polymer, the

Avrami rate constant a

k is associated with the overall

rate of crystallization including the effect of nuclei

density and crystal growth rate. This study proposes the

average Avrami rate constant as a measure of the overall

rate of crystallization under temperature gradient.

The average of ak =

1

x

ak T x dx

X

(7)

where a

k T x depends on crystallization temperature

in the horizontal axis T x and X is a length of

simulation area (800 x 800 pixel).

3. Simulation of morphological development and

overall crystallization kinetics

Figure 1 shows the simplified algorithm for our

simulation. In this work, we investigated crystallization

of s-PP under various temperature gradients in x

direction. Both nucleation and growth process during

crystalllization are simulated using Monte Carlo model

in 800×800 square unit cells. 25 unit cells are assumed to

be equivalent with 1 µm2 and 1 steptime is 6 seconds.

Initially, all 640,000 cells are considered as an

amorphous entity before the crystallization occurs. The

heterogeneous nucleation (i.e., all nuclei occur

instantaneously at the corresponding temperature) is

assumed in the model. The nuclei density N(T) at

specific temperature can be calculated using Equation

(2), (5) and (6). The nuclei are located randomly and

assumed to occupy the area of 1 unit cell.

Each nucleus is followed by the subsequent radial

growth. Crystal growth rate can be calculated using

Equation (6). Each crystallite can grow until it impinges

with adjacent crystallites. Crystallinity, which is defined

as the crystalline volume fraction, and morphology are

recorded at each time step.


163

Figure 1. Simplified algorithm for our simulation

Four case studies with different temperature

gradients are summarized below and shown in Figure 2.

Case 1: Study effect of temperature gradient by fixing

temperature at a center line as 70 oC (i.e., 60 to 80

oC, 50

to 90 oC and 40 to 100

oC).


left at 40 oC (i.e., 40 to 60

oC, 40 to 80

oC and 40 to 100

oC).

Both cases follow the simple linear equation:

( 1)1

R L

L

T TT x T x

X

(8)

whereR

T and L

T are the temperatures in boundary.


left and right boundary temperature as 40 oC and 90

oC,

respectively,with a rapid cooling on the right hand side.


left and right boundary temperature as40 oC and 90

oC,

respectively, with a rapid cooling on the left hand side.

Cases 3 and 4 follow equation below:

2TT

t

(10)

Figure 2. Temperature profiles of 4 case studies


4.1 Crystallization kinetics

Figure 3 shows the relationship of G , tot

N and A

k

as a function of crystallization temperature for s-PP in

the range 40-100 °C; G and Ak exhibit typical bell-shape

dependence withc

T and the maximum values at about 70

and 60 ° C, respectively. tot

N decreases with an increase

in c

T .

START

Specify temperature gradient

t=t0

Caculate N , ka and G from Eqs. (2), (5)

and (6)

Growth

RandomlN in space area by probability

END

t=tfinal

YES

No


164

Figure 3. The relationship of G ,

totN and

Ak as a

function of c

T for s-PP

Figure 4 shows the relative crystallinity as a

function of time for all case studies. Table 1-4

summarize average Avrami rate constant for each case.

In case 1 as shown in Figure 4(a), the

crystallization of T1_60/80 has the fastest crystallization

rate because it is closed to the maximum of crystal

growth rate that at 70 oC. Because most nuclei occupy at

low temperature, the time required to reach the complete

crystallization increase when the temperature range is

broadened.

For case 2 where the temperature at left boundary

is equal to 40 oC (see Figure 4(b)), the relative

crystallinity of T2_40/60 and T2_40/80 are very close. It

was found that the crystallization of T2_40/60 is slightly

more than T2_40/80. Both conditions have a large

number of nuclei and a high crystal growth rate that can

make the crystallization very fast.

Figure 4(c) shows the result for case 3. The

crystallization of T3_c3 is formed at high temperature

when compared with others but the result shows the fast

crystallization at early time. However, T3_c3 requires

more time to complete crystallization at the later stage

when compared with other conditions because the large

of nuclei is occurred at the low temperature area and the

crystallites in this area grow very slowly to the high

temperature area.

In the last case (see Figure 4(d)), the

crystallization of T4_h1 has the highest crystallization

because it has a highest temperature which lead to small

amount of nuclei and slow growth. In all case studies, the

average Avrami rate was found to be the adequate

indicator of overall crystallization rate at the early stage

of crystallization (see Table 1-4).

Figure 4. The relative crystallinity as a function of time

for 4 case studies.

Table 1. Average Avrami rate constant for case 1

Temperature Average of ka (min2)

T1_60/80 0.1141

T1_50/90 0.1024

T1_40/100 0.0807



T2_40/60 0.12170

T2_40/80 0.11860

T2_40/100 0.08077


165 Table 3. Average Avrami rate constant for case 3


T3_c1 0.0670

T3_c2 0.0820

T3_c3 0.0999



T4_h1 0.0580

T4_h2 0.0386

T4_h3 0.0230

4.2 Morphological development

Figure 5 and 6 shows the simulated

morphological development during crystallization for

case 1 and case 3, respectively. The results captured the

crystallites growth and impingement until the complete

crystallization. It was observed that the crystallites tend

to have broader size distribution with an increase in the

temperature gradient. Because most nuclei occur at low

temperature, the simulated morphology showed the high

nuclei density on the left hand side.[3]

These results

expresses a good agreement with the crystallization

kinetics.

Figure 5. Simulated morphological development for

case1

Figure 6. Simulated morphological development for

case3

4. Conclusion

Effect of temperature gradient on crystallization

kinetics and morphological development of s-PP was

investigated under various temperature profiles using

Monte Carlo simulation. The results can be explained

using the relationships between nuclei density, crystal

growth rate and temperature. Average Avrami rate

constant over the temperature field was found to be an

adequate indicator of the crystallization rate, especially

at the early stage of crystallization

5. Acknowledge

The authors would like to thank Graduate

School and Faculty of Engineering of Kasetsart

University for their kind financial supports.

6. References

[1] Strobl, G., “Crystallization and melting of bulk

polymers: New observations, conclusion, and a

thermodynamics scheme”, Prog. Polym. Sci., 31:

398-442 (2006).

[2] Long, Y., Shanks, R.A., and Stachursili, Z.H.

“Kinetics of polymer crystallization”, Prog. Polym.

Sci., 20: 651-701 (1995).

[3] Pawlak, A., and Piorkowska, E., “Crystallization of

isotactic polypropylene in a temperature gradient”,

Colloid. Polym. Sci., 279:939-946 (2001).


166 [4] Avrami, M., “Kinetics of Phase Change. I. General

Theory”, J. Chem. Phys., 7: 1103-1112 (1939).

[5] Avrami, M., “Kinetics of Phase Change. II.

Transformation-time relations for random

distribution of nuclei”, J. Chem. Phys., 8:212-224

(1940).

[6] Supaphol, P. and Sprueill, J.E., “Thermal properties

and Isothermal crystallization of syndiotactic

polypropylene: Differential scanning calorimetry

and overall crystallization kinetics”, J. App. Polym.

Sci., 75:44-59 (2000).


167 PROCO-04

Pressure Slips Casting Using a Porous Plastic Mold: Effect of Pressure and Time on

Green Articles

Kittiya Jitklang1,a

and Dr.Somchoke Sontikaew 1,b

1King Mongkut’s University of Technology Thonburi (KMUTT)

126 Pracha Uthit Rd., Bang Mod, Thung Khru, Bangkok 10140, Thailand

Abstract

In this work, a porous plastic mold was developed for pressure slip casting of green ceramic articles.

PMMA microspheres with average diameter of 375 m and 30 m provided the porous plastic with average

pore size of 14-20 m. The porous mould was used to investigate effects of casting pressure and casting time on

weight, thickness, diameter, linear shrinkage, and green density of green ceramic disks. It was found that the wet

and dry weights of the green samples were pressure dependent at low casting pressure and became pressure

independent when the casting pressure was beyond 1 MPa. The casting time slightly affected the wet and dry

weights. The content of residual water in the green samples was independent of casting pressure and casting

time. The final content of residual water in the green samples was 50% by weight of water in a slip used. The

diameter, linear shrinkage, and green density of the green samples were merely pressure dependent.

Keywords: Pressure slip casting, Porous materials, Ceramic casting

1. Introduction

For several decades, high-pressure slip casting

process has shown a remarkable increase in productivity

and a considerable improvement in quality of ceramic

sanitary and tableware articles as compared to

conventional slip casting process using plaster of Paris as

a mold material [1]. In the former process, porous

polymer material has been used to make a mould to

extract water from slip through pores under high

pressures of 1-4 MPa [2, 3]. According to Sacmi

Whiteware resin moulds, the matrix of the porous mould

materials with average pore size of 20-25 micrometers

and % open porosity of 27-30 frequently is polymethyl

methacrylate (PMMA). The porous materials show

20 N/mm2 and 8 N/mm

2 in compressive strength and

blending tensile strength, respectively. The service life of

the mould material is around 20,000-40,000 casts. There

are a few works reported in the literatures about the

preparation, morphology and properties of porous

materials [4-6]. However, in these literatures,

investigation of using the porous polymer materials to

produce ceramic articles via the pressure slip casting

process has not been publicly reported yet.

In this work, PMMA porous plastic was

developed to investigate morphology and thickness-

forming capacity of the porous plastic. The porous

plastic obtained was used to make a pressure slip casting

mould to study effects of pressure and time on weight,

thickness, diameter, linear shrinkage, and green density

of green ceramic articles.

2. Experimental

2.1 Materials

The methyl methacrylate monomer (MMA)

with inhibitor of dimethyl-6-tertiary-butylphenol and a

minimum purity of 99.9% was purchased from Thai

MMA Co., Ltd. The polymerization initiator benzoyl

peroxide (BPO) was supplied by Sigma-Aldrich.

Ethylene Oxide/Propylene Oxide Copolymers (EO/PO),

Tergitol XD from DOW chemical, was used as a non-

ionic surfactant for emulsion polymerization of PMMA

porous material. Large and small PMMA beads with

diameter of 375 m and 30 m, respectively, were

kindly supplied by Diapolyacrylate Co., Ltd. (Mitsubishi

Rayon group). Distilled water purchased from RCI

labscan limited was used for suspension polymerization

and water-in-oil emulsion polymerization of PMMA. A

clay slip with a density of 1.49 (1.39) g/cm3

and pH value

of 8 and solid content of 49% used was supplied by Sin

Fa industrial, Thailand.


168 2.2 Porous PMMA synthesis

The PMMA porous plastics were produced by

water-in-oil (W/O) emulsion polymerization. Initially,

water and oil phases had to be prepared. In the water

phase, 40 wt% of fine PMMA microspheres and 24 wt%

of water/surfactant solution with ratio of 100/1 were

blended together. In the oil phase, there were 18 wt % of

coarse microspheres and 18 wt% of MMA. The oil and

water phases were then mixed together and were stirred

to form W/O emulsions. The emulsion was poured in an

aluminum mould. After the polymerization reaction was

complete, the porous resins were removed from the

mould and washed with warm water several times and

dried. The porous mould obtained consisted of two parts,

an upper part and a lower part, as shown in Fig. 1. The

lower part was cavity of disk shape with 36 mm in

diameter and 5 mm in height.

2.3 Pressure slip casting experiment

The pressure casting experiment was carried on

using the laboratory pressure slip casting equipment as

shown in Fig. 1. The casting equipment was placed in a

20 tons hydraulic machine. From this unit, diameter and

thickness of a mould cavity were 36 mm and 5 mm,

respectively.

The slip pressure levels used were 0.5, 1.0, 1.5 and

2.0 MPa. The casting times were 2 and 4 minutes.

The casting time was recorded when the casting pressure

reached the desired level. The sample was, then,

removed manually.


Scanning electron microscopy (SEM) was used

to investigate diameter of PMMA microspheres and pore

morphology of the porous materials. The fractured, cut,

and polished surfaces of the porous resin for SEM were

prepared. Thickness-forming capacity is used to measure

the filtration ability of the porous resin compared to that

of the plaster of Paris. To measure thickness-forming

capacity, cake samples were prepared from a plaster

mold and a porous resin mold. A clay slip was poured in

the plaster or porous resin molds at atmospheric pressure

for 60 minutes to form the cylindrical cakes. The cakes

were removed to measure their thicknesses which were

used to calculate the thickness-forming capacity. The

thickness of the cakes from the plaster mold was used as

the reference to calculate % difference in the thickness of

the other cake. The % thickness difference obtained was

defined as the thickness-forming capacity of the porous

resin while the thickness-forming capacity of the plaster

was 100. The % water content in the cake samples were

examined by measuring the weight of the samples before

and after drying in an oven at 120C for 24 hours. The

relative linear shrinkage in diameter of the dried green

samples was examined. The green bulk density was

calculated from the weight and volume of the dried green

cakes.

3. Result

3.1 Morphology

SEM images in Fig. 2 shows the coarse particles

with average diameter of between 375 µm and the fine

particles with average diameter of 30 µm. Both acrylic

powders were used to prepare acrylic porous materials.

The images of the fracture, cut, and polished surfaces of

SEM samples were shown in Fig. 3(a), 3(b) and 3(c)

respectively. From the image of the fracture surface,

acrylic particles and micro-pores with diameter of less

than 14-20 m were seen. For the image of the cutting

surface, pore connectivity and pore distribution were

observed. In the image of polished surface the pore

connectivity and the pore distribution were clearly seen.

The average size of the long pores was 60 m. The

percentage of porosity was about 16% (cm2/cm

2),

Fig. 1 Diagram of pressure casting unit


169 evaluated from the area of the pores in the SEM image of

the polished samples.

3.3Thickness-forming capacity

Thickness-forming capacity is used to compare

the filtration ability of the porous resin and that of plaster

at atmospheric pressure for 60 minutes. The thickness-

forming capacity of plaster is defined a value of 100 and

was used as the reference number for values assigned to

the other materials. Porous materials with a larger pore

size than that of plaster provided the small value of

thickness-forming capacity. Table 1 shows the thickness-

forming capacity of the home-made porous sample (PR-

4) compared to various commercial porous materials

(PR-1 to PR-3). The commercial materials with pore size

of 5-10 m exhibited 27-28% of plaster-capacity

whereas the porous material in this work showed the

thickness-forming capacity of 22.7%.

3.4Wet and dry weight

In Fig. 4(a) and 4(b), the weights of the wet and

dry samples increased with increasing pressure from 0.5

to 1 MPa and were unchanged when the casting pressure

varied from 1 to 2 MPa. At the casting pressure of 0.5 to

1.5 MPa, the weights for the casting time of 4 minutes

were slightly higher than those for the casting time of 2

minutes. At higher casting pressures (1.5 and 2 MPa)

applied, both weights for the long casting time were

insignificantly higher than those for the short casting

time. It indicated that the wet and dry weights were the

time-dependent at the low casting pressure and became

the time-independent at the high casting pressure.

3.5 Content of residual water

In Fig. 4(c), the content of residual water in the

green cast samples varies from 24.5% to 26% and was

independent of the casting pressure and time. In this

work, solid and water content in the slip used were 49%

and 51% by weight, respectively. The residual water in

the green samples was about 50% of the water content of

the slip. Previous work showed that advanced ceramic

slip with water content of 30% by weight provided

16.4 % of residual water in pressure slip casting samples

[8]. This indicated that the residual water in a pressure

casting sample was about 50% of the water content in a

slip used.

(a)

(b)

(c)

Fig. 4 Wet weight (a) and dry weight (b), and water

content (c) as a function of the casting pressure

(a) (b) (c)

Fig. 3 (a) fractured (b) cut (c) polished surfaces of the

porous resin

(a) (b)

Fig. 2 SEM images of acrylic powder with average

diameter of (a) 375 m and (b) 30 m


170

Table 1: Properties of the commercial and home-made porous resins

Porous resin

types

Avg. pore

(m)

Total absolute open porosity (%

by cm3/cm

3)

Thickness forming

capacity (%)

PR-1 [7] 1.0 45-50 100-110

PR-2 [7] 5-10 27-28 30

PR-3 [7] 20-25 30-32 5-10

PR-4 14-20 16% (by

cm2/cm

2)

22.7

3.6 Thickness and diameter

The thickness of the green cast samples was

independent of the casting time and pressure as shown in

Fig. 5(a). This was due to the fixed mould geometry. In

Fig. 5(b), the diameter of the dried green samples was

time dependent but increased with increasing pressure,

indicating that the diameter was pressure dependent. The

expansion of samples in the radius direction can be

attributed to the separation of cast particles under the

applied pressure.

3.7 Linear shrinkage and green bulk density

The linear shrinkage decreased as the pressure

is increased due to the expansion in diameter of the

samples as seen in Fig. 6(a). Similarly to the diameter,

the shrinkage was time independent. In Fig. 6(b), the

green density gradually increased as the casting pressure

increased from 0.5 to 1.5 MPa. The green density

significantly decreased when the casting pressure beyond

1.5 MPa. The decrease in the green density was due to

the increase in the diameter of the samples (see Fig. 5(b))

while mass was unchanged as seen in Fig. 4(b). For the

effect of the casting time, the green density increased

with increasing the casting time.

(a)

(b) Fig. 6 shows linear shrinkage (a) and bulk density (b) as a

function of the casting pressure

(a)

(b) Fig. 5 shows thickness (a) and diameter (b) as a function

of the casting pressure


171 4. Summary

PMMA microspheres with average diameter of

375 m and 30 m were used to produce porous

materials. The porous material obtained provided

average pore size of 14-20 m, porosity of 16% by area,

and thickness-forming capacity of 23%. Porous plastic

material was produced for making a pressure casting

mould to form the green disk samples. The results

showed that the wet and dry weights of the green

samples increased with increasing the pressure and

became unchanged when the casting pressure was

applied beyond 1 MPa. At a casting pressure, increasing

the casting time led to slight increase in the wet and dry

weight. The time dependent of the dry weight led to the

increase in bulk density with increasing casting time. The

content of residual water in the green samples did not

affected by the casting pressure and time. The final

content of residual water in the green samples was 50%

by weight of water in a slip used. Increasing the casting

pressure greatly decreased the diameter and linear

shrinkage which were not affected by the casting time.

At pressure of 2 MPa, the volume of a green sample

increased while its dry weight did not changed. This led

to the decrease in the bulk density at a high casting

pressure applied.

5. Acknowledgements

We would like to thank Assoc. Prof. Dilok

Sriprapai for professional guidance and

recommendations on this project. as well as Pathanasuk

Capital CO.,LTD., for their help in collecting the plant

data and technical support. This work has been funded

by the Research and Researcher for Industry (RRI) under

the Thailand Research Fund (TRF).

6. References

[1] Gregory D. Wallis: The development and

application of porous plastic molds for the casting of

sanitary ware and dinnerware, Ceramic engineering

& science proceedings Vol. 15(1994), p. 113-117.

[2] Mazzanti V.: Process Eng DKG 2002;79(1–2):E11–

2.

[3] Alfred Kaiser, Roel van Loo, Josef Kraus, and

Andreas Hajduk: Comparison of Different Shaping

Technologies for Advanced Ceramics Production,

Process Engineering, cfi/Ber. DKG 86 (2009) No. 4

[4] E. Jimenez Pique, L.J.M.G. Dortmans, G. de With:

Fictitious crack modeling of polymethyl

methacrylate porous material, Materials Science and

Engineering A335 (2002) p.217–227.

[5] Y. Ergun, C. Dirier, M. Tanoglu: Polymethyl

methacrylate based open-cell porous plastics for

high-pressure ceramic casting, Materials Science

and Engineering A 385 (2004) p.279–285.

[6] Metin Tanoglu, Yelda Ergu: Porous nanocomposites

prepared from layered clay and PMMA

[poly(methyl methacrylate)], Composites: Part A 38

(2007) p.318–322

[7] Information on http://www.sama-online.com/System

/00/01/06/10699/633571798146718750_1.pdf

[8] Alfred Kaiser, Roel van Loo, Josef Kraus, and

Andreas Hajduk: Comparison of Different Shaping

Technologies for Advanced Ceramics Production,

cfi/Ber. DKG 86 (2009) No. 4, E41-E48


172

PROCO-05

Comb-shaped Polycarboxylate based Copolymers with Benzaldehyde

Derivative for Molecular Model of Antimicrobial Superplasticizer

Nalinthip Chanthaset1, Hiroharu Ajiro

2, Mitsuru Akashi

2 and Chantiga Choochottiros

1*

1 Department of Materials Science, Kasetsart University, 50 Lat Yao, Chatuchak, Bangkok 10900, Thailand.

2 Department of Applied Chemistry, Graduate School of Engineering, Osaka University,

2-1 Yamada-oka, Suita, Osaka 565-0871, Japan.

Abstract

A novel polycarboxylate superplasticizers (PCs) based on methacrylic acid (MAA) and 2-aminoethyl

methacrylamide (AMA) modified with benzaldehyde derivative were synthesized wtih two diferent molar ratio

of starting materials as copolymer MAA-AMA(BZ)1 and MAA-AMA(BZ)2. Zeta potential of MAA-

AMA(BZ)1 and MAA-AMA(BZ)2 were -29±1.6 and -19±0.4 mV, respectively. The antimicrobial test showed

the antifungal property in basic solution and the optimum efficiency was up to 25% via dual culture method.

These molecular model polymers were able to dissolve in basic solution and suitable for cement application.

Keywords: comb-shaped polycarboxylate, superplasticizers, antimicrobials

1. Introduction

Superplasticizers are chemical admixtures which

are used for enhancing well dispersion. They are also

known as water reducers. In the past, latex particles were

well-known to solve the flocculation of cement particles

by neutralizing charge on cement surface and provide

good dispersion of cement particles1,2

. Development in

chemical compounds and structures of the superplas-

ticizers have been reported as linosulfonates (LS),

sulfonated melamine formaldehydes (SMF), sulfonated

naphthalene formaldehyde (SNF) and the latest one was

comb-shaped polycarboxylate derivatives (PC)3. The

comb-shaped polycarboxylic acid-based superplas-

ticizers perform excellent in reducing water to cement

ratio, enhancing flow ability because of their structure

that perform negative charge from carboxylate group in

order to neutralize charge on the cement surface and side

chain of copolymers provide steric hindrance to prevent

agglomeration of cement particles4,5

.

Generally there are many environmental problems and

infection caused by microorganisms, which occur on the

cement/concrete surface as discoloration. Method to control

microbial infections is a point of issue.

Biocidal polymers which are polymers with active

functional group. They have ability to kill microorganisms

by sterilizing ions or molecules6. Profits of polymeric

antimicrobial agents are chemically stable, non-volatile,

and hard to enchance to the human body. From many

researches revealed that functional groups which show

antimicrobial properties are molecules compose of positive

charge such as N-alamine groups, nitric oxide, phenol,

quaternary ammonium including neutral molecule such as

benzaldehyde derivatives7-9

. The benzaldehyde derivative

that has reactive antimicrobial activity such as phenols

derivative, is one of antimicrobial agents that interact with

cell membrane on surface of cell and lead to cell death

through disintegration of the cell membrane and release of

intracellular constituents10

.

Due to the composition in cement, there are mainly

silica and aluminate compound, which are tricalcium

silicate (C3S), dicalcium silicate (C2S), tricalcium alu-

minate (C3A) and tetracalcium aluminoferrite (C4AF)11

.

On cement surface layer are fulfilled with ions when added

water to cement, the reaction between C3A, C2S and water

provide cations such as Ca+. They diffuse faster than anions

which are products of cement hydration reaction. In

common, the hydration of cement paste also give large

amount of hydroxide ions that lead slurry of cement paste

to be basic around pH 10-12. On the surface of the cement

particle, the concentration of anions is higher than that

cations forming a negative charged layer. Typically, on

cement grains shows the distribution charges both positive

and negative and easily agglomerate.


173 In this work, we modified active functional group for

antimicrobial activity to molecular structure of

superplasticizer. The comb-shaped antimicrobial

copolymers based on polycarboxylate and 2-aminoethyl

methacrylacrylate hydrochloric were synthesized and

benzaldehyde derivatives, 2,4-dihydroxybenzaldehyde,

were immobilized at amine-terminated functional group.

Antimicrobial activity were consequently determined by

well agar diffusion and dual culture method against to

different types of fungal and bacteria including Aspergillus

Nigar, Aspergillus Flavous, Aspergillus Fumigatus and

Bacillus Cereus.

2. Experimental section

2.1 Materials

2-Aminoethyl methacrylate hydrochloride was

purchased from Sigma Aldrich, Japan. 2,4-

Dihydroxybenzaldehyde was purchased from Merck.

Methacrylic acid monomer stabilized with MEHQ (99%)

was brought from TCI, Japan and was purified by

distillation before used. 2,2’-Azobis(isobutyronitrile)

(AIBN) was purchased from Wako Pure Chemical

Industries (Osaka, Japan) and N,N-Dimehtylformamide

anhydrous 99.8% (DMF) was brought from Sigma Aldrich,

Japan. Other chemicals and solvent were used without

purification.

2.2 Synthesis of copolymers

MAA-AMA(BZ)1 was copolymeried via free radical

polymerization by using AIBN as initiator (0.09g,

0.05mmol). Two monomers were 2-Aminoethyl

methacrylate hydrochloride (AMA: 1.0 g, 11.6 mmol).

Reaction temperature was 70°C under nitrogen atmosphere

in anhydrous DMF for 24 h. Product of MAA-AMA1

(2.33g, 80% yield) was precipitated in acetone to hexane

ratio of 1:1 and washed with acetone. Product was

collected by filtration. Then, MAA-AMA1 (1g) were

functionalized with 3-fold of 2,4-dihydroxybenzaldehyde

(BZ: 1.61g, 34.8 mmol), they were dissolved in anhydrous

DMF and poured in three-necked flask. The second step of

reaction was stirred for 48 hours at 90°C under nitrogen

atmosphere. A heterogeneous solution were observed and

filtrated with hexane and acetone for departing unreacted

benzaldehyde. Finally, the product of MAA-AMA(BZ)1

was recovered as yellow-green powder and dried under

vacuum (50% yield).

Two products with different molar ratio (mmol) were

synthesized and label as shown in Table1.

2.3 Measurements

1H NMR with 32 scans (D2O with 40%NaOD,

400MHz) were measured by JEOUL JNM-GSX400 and

fourier-transform infrared (FTIR) spectra were recorded

from 100 FTIR spectrometer (Perkin-Elemer). The number

of scans were 64 times at resolution 4 cm-1

. Molecular

weights and molecular weight distribution were

characterized by gel permeation chromatography (Tosoh

system HLC-8120GPC). PMMA was used as standards at

35°C and 0.2MNaCl as an eluent. Zeta potentials was

measured by dynamic light scattering (DLS) method using

a Zetasizer Nano ZS (Malven Instruments, UK).

Table 1 Summary of the used quantity of monomer, benzaldehyde

derivative and yield (%)

Scheme 1 Synthesis of copolymer AMA-MAA(BZ) which

immobilized 2,4-dihydroxybenzaldehyde


174 2.4 Antimicrobial test

Fungi were kept on Sabouroud’s agar slopes and

bacteria were grown on nutrient agar. All fungi and

bacteria were authenticated by the Department of

Microbiology, Faculty of Science, Kasetsart University,

Thailand. There were Aspergillus Nigar, Aspergillus

Flavous, Aspergillus Fumigatus and Bacillus Cereus.

To determine efficiency of inhibition fungal growth by

comparison of the growth diameter, the dual culture

method was done. Agar were prepared by dissolving 39 g

of Potato Dextrose Agar (PDA: Becton, Dickinson and

company, USA) in 1000 ml boiled water stirred for

homogenous solution and cooled down. Then, poured the

liquid into plastic plates (8.6 cm diameter) approximately

20 mL of each plate. All plates were blown and sterilized

under Ultra-violet light for 2 hours. The culture was

touched lightly on the center of agar surface by a loop and

sample solution was cross over by loop as 2 line in parallel.

Solutions of MAA-AMA(BZ)2 with concentration of 10,

20 and 30 mg/mL and 2,4-dihydroxybenzaldehyde were

compared with control in same equivalent mole, so three

replicates were made.


Solubility of copolymer. Copolymers, MAA-AMA1

and MAA-AMA2 were able to dissolve in water due to

polarity of side chain that are terminated-amine and

carboxylic acid. In contrast, polymers which were

immobilized with benzaldehyde derivative, MAA-

AMA(BZ)1 and MAA-AMA(BZ)2, could dissolved in

basic solution (pH12) and chemical structure was changed

as shown in scheme 2.

Characterization. The structure of the AMA-MAA1,

AMA-MAA2, AMA(BZ)-MAA1 and AMA(BZ)-MAA2

were characterized by 1H NMR as shown in Figure 3 and

4. The 1H NMR spectra (Figure 3) shows peak shift

comparing to two copolymers due to their structure, the

chemical shifts (ppm) at 3.65 (CH2, cx), 3.24 (CH2, dz),

2.80 (CH2, cz), 2.54 (CH2, dx), 1.29 (CH2, bx,y,z), 0.56 (CH3,

ax,y,z) of MAA-AMA1 and at 3.89 (CH2, cx), 3.51 (CH2, dz),

3.25 (CH2, cz), 2.78 (CH2, dx), 1.54 (CH2, bx,y,z), 0.82 (CH3,

ax,y,z) of MAA-AMA2. Furthermore, modified copolymer

found chemical shift

Solubility of copolymer. Copolymers, MAA-AMA1

and MAA-AMA2 were able to dissolve in water due to

polarity of side chain that are terminated-amine and

carboxylic acid. In contrast, polymers which were

immobilized with benzaldehyde derivative, MAA-

AMA(BZ)1 and MAA-AMA(BZ)2, could dissolved in

basic solution (pH12) and chemical structure was changed

as shown in scheme 2.

Scheme 2 The structure of copolymer AMA-MAA(BZ) in pH12


175

Characterization. The structure of the AMA-

MAA1, AMA-MAA2, AMA(BZ)-MAA1 and

AMA(BZ)-MAA2 were characterized by 1H NMR as

shown in Figure 3 and 4. The 1H NMR spectra

(Figure 3) shows peak shift comparing to two

copolymers due to their structure, the chemical shifts

(ppm) at 3.65 (CH2, cx), 3.24 (CH2, dz), 2.80 (CH2,

cz), 2.54 (CH2, dx), 1.29 (CH2, bx,y,z), 0.56 (CH3, ax,y,z)

of MAA-AMA1 and at 3.89 (CH2, cx), 3.51 (CH2, dz),

3.25 (CH2, cz), 2.78 (CH2, dx), 1.54 (CH2, bx,y,z), 0.82

(CH3, ax,y,z) of MAA-AMA2. Furthermore, modified

copolymer found chemical shift (Figure 4) at 9.41

(CH, et), 7.01 (CH, gt), 5.95 (CH, ft, ht), 3.61 (CH2,

ct), 3.45 (CH2, dr), 2.84 (CH2, cr), 2.52 (CH2, dt), 1.29

(CH2, bt,s,r), 0.56 (CH3, at,s,r) of MAA-AMA(BZ)1 and

at 9.04 (CH, et), 7.22 (CH, gt), 6.14 (CH, ft, ht), 3.85

(CH2, ct), 3.47 (CH2, dr), 3.22 (CH2, cr), 2.73 (CH2,

dt), 1.50 (CH2, bt,s,r), 0.77 (CH3, at,s,r) of MAA-

AMA(BZ)2. The analysis from FTIR appeared Schiff

base v(C=N) of MAA-AMA(BZ)1 at 1,624 cm-1

and

MAA-AMA(BZ)2 at 1,619 cm-1

as showed in Table

2.

In order to synthesize the copolymer that had

a plenty of negative charges of carboxylate group

(COO-), we increased molar ratio of MAA : AMA

monomer from 1:1 and 2:1 (Table 1) and the number

of molecular weight (Mn) of MAA-AMA(BZ)2 was

larger than MAA-AMA(BZ)1.

The repeating units of copolymer were

relative to the number of molecular weight (Mn,

GPC) and the integral area of 1H NMR peak showed

in Table 2. Moreover, zeta potential (ζ-potential) of

MAA-AMA(BZ)1 and MAA-AMA(BZ)2 were -29

±1.6 mV and -19±0.4 mV. This implied remaining of

negative charge from COO- in basic solution.

Antimicrobial activity. Due to the abundant

of COO- group along polymer chain, the MAA-

AMA(BZ)2 were dissolved in water at pH12.

Polymer solutions with concentrations of 5, 10 and 30

mg/ml were lined across on the PDA surface. After 5

days, diameter of fungal and bacteria growth were

measured. We reported ability of antimicrobial

activity by comparing diameter of cell growth of

sample plate and control plate and reported in percent

inhibition. For A. Fumigatus inhibition test, the

MAA-AMA(BZ)2 concentration of 5 mg/ml could

inhibit 1.4% , 10 mg/ml as 2.7% and the optimum

concentration 30 mg/ml showed the best inhibition as

25.1%. Besides, we observed that polymer at 30

mg/ml concentration was able to inhibit all fungi

which were A. Nigar, A. Flavous, and A. Fumigatus

for 3.2%, 25.1% and 15.0%, respectively. But it was

unable to inhibit the Bacillus Cereus.

Figure 4. 1H NMR of MAA-AMA(BZ)1 and MAA-AMA(BZ)2.

//

Table 2 GPC, Zeta potential, and FTIR data of synthetic copolymers


176 Conclusions

Two new comb-shaped copolymer base on

polycarboxylate modified 2,4-

hydroxybenzaldehyde were synthesized via free

radical polymerization of methacrylic acid (MAA)

and 2-aminoethyl methacryla-mide (AMA) were

synthesized. The antimicrobial activity against the

fungi; A. Nigar, A. Flavous, and A. Fumigatus were

investigated. Hence, the outstanding performance

of MAA-AMA(BZ)2 copolymer was able to

dissolve in pH 12 and could act as antifungal

polymer for Aspergillus Fumigatus. So, the MAA-

AMA(BZ)2 appropriate to be molecular model for

antifungal superplasticizer.

Acknowledgement

This research is supported by Research

and Researcher for Industry (RRI: MAG, Grant

number: MSD5710077). The author would like to

thank Dr. Yaovapa Taprab, Department of

Microbiology, Kasetsart University, Thailand for

antimicrobial testing.

References

1. S.H. Lu, G. Liu, Y.F. Ma, F. Li, Synthesis and

application of a new vinyl copolymer

superplasticizer. J. Appl. Polym. Sci 117 (2010)

273-280.

2. J. Plank, Z.M. Dai, H. Keller, H. Keller, H. Hossle,

W. Seidl, Fundamental mechanisms for

polycarboxylate interaction into C3A hydrate

phases and the role of sulfate present in cement.

Cem., Concr. Res. 40 (2010) 45-57.

3. T. Hirata, A cement dispersant, in: JP Patent

84,2022, 1981, S59-018338.

4. H.Uchikawa, S. Hanehara, D. Sawaki, The role of

steric replusive force in the dispersion of cement

particles in fresh paste prepared with organic

admixture. Cem. Concr. Res. 27 (1997) 37.

5. A. Zingg, F. Winnefeld, L. Holzer, J. Pakusch, S.

Becker, R. Figi, L. Gauckler, Adsorption of

polyelectrolytes and its influence on the rheology,

zeta potential, and microstructure of various cement

and hydrate phases. J. Colloid Interface Sci. 323

(2008) 301-312.

6. A. Ahemd, J.N. Hay, J.N. Wardell, G. Cavalli,

Biocidal polymers (I): Preparation and biological

activity of some novel biocidal polymers based on

uramil and its azo-dyes. React. Funct. Polym. 68

(2008) 248-260.

7. E. Kenawy, S.D. Worley, R. Broughton, The

chemistry and applications of antimicrobial

polymers: A state-of-the-art review.

Biomacromolecules 57 (2007) 1359-1384.

8. A. Munoz-Bonilla, M.Fernanzdez-Garcia,

Polymeric materials with antimicrobial activity.

Prog. Polym. Sci. 37 (2012) 281-339.

9. ES. Park, WS. Moon, MJ. Song, MN. Kim, KH.

Chung, YS. Yoon, Antimicrobial activity of phenol

and benzoic acid derivatives. Int. Biodeter. Biodegr.

47 (2001) 209-214.

10. E. Subramanyam, S. Mohandoss, HW. Shin,

Synthesis, characterization, and evaluation of

antifouling polymers of 4-acryloyloxybenzaldehyde

with methyl methacrylate. J. Appl. Polym. Sci. 112

(2009) 2741-2749.

11. H.F.W. Taylor, Cement Chemistry. Academic Press

1997.


177

PROCP-02

Properties and Stability of Poly(Vinyl Acetate-Co-Alkyl- Acrylate) Latex Modified with

Carbonyl- and Hydroxyl-containing Monomers Piyabutree Thongsook

1, Varawut Tangpasuthadol

1,2*

1Program of Petrochemical and Polymer Science, Faculty of Science, Chulalongkorn University, Bangkok

10330 2Department of Chemistry, Faculty of Science, Chulalongkorn University, Bangkok 10330

Abstract

The aim of this work was to study the effect of functional group types on the adhesion property and stability

of pressure sensitive adhesive (PSA) based on poly(vinyl acetate-co-alkyl acrylate) latex. The latex was

synthesized by the miniemulsion polymerization in semi-continuous process having vinyl acetate (10%wt) and

2-ethyl hexyl acrylate (90%wt) as major ingredients. The modification was carried out by adding 0-1.5% of one

or more monomer types that contained –COOH, i.e. -carboxyethyl acrylate, methacrylic acid, and acrylic acid,

or –OH group, i.e. 2-hydroxyl ethyl acrylate and 2-hydroxy ethyl methacrylate. Each final adhesive was tested

on stainless-steel substrate with adhesive film thickness of 20 m. The synthesized copolymer latexes with high

stability have the zeta potential values of -30 to -36 mV and, conductivity of 1.4510-3

to 3.1310-2

mS/cm. The

modified latex could be stored for as long as 6 months at ambient temperature. Adhesion properties of the

synthesized copolymer had the peel strength of > 0.7 kg/in and cohesion time of > 24 hr. These results indicated

that the improved PSAs gave the stronger adhesive bond and were more stable than the copolymers without an

addition of –COOH monomer.

Keywords: Pressure sensitive adhesive, Emulsion polymerization, Functional monomer, Peel strength and

Cohesion

1. Introduction

Adhesive is the bonding agent for joining the

surface of two solid materials. The strength of the

adhesive is affected by duration, heating, compression

which depends on the type of surface, surface energy,

contact angle, and type of the adhesive. There are 2 types

of the adhesive; permanent and transient types.

Industrial adhesives are latexes synthesized by

the emulsion polymerization of monomers, such as vinyl

acetate, styrene, or acrylate. The polymer synthesized

from each monomer is suitable for different applications.

This study focuses on vinyl acetate–alkyl acrylate

copolymer which is pressure sensitive adhesive or PSA

by which the materials will be adhered just applying

gentle press. PSAs are applied in many applications such

as tape, sticker, OPP tape, double side tape, etc. The

carboxyl and hydroxyl groups are the key functional

groups that affect the tack and the peel properties of the

latex, caused by the interaction from hydrogen bonds

between the two materials [1].

Each application needs different latex

properties. The latex particle surface can be improved by

using the monomers containing functional groups, for

instance, carboxyl, hydroxyl, amine, and epoxy. These

monomers can provide crosslinking and surface

modification of the latex particles. Carboxyl and

hydroxyl are interesting functional groups. Carboxyl

groups in, for example, acrylic acid and methyl acrylic

acid can improve the stability of the colloid, mechanical

properties, resistance to shear, and increase the hardness

of the film. The hydroxyl groups in, for example, 2-

hydroxy ethyl acrylate and 2-hydroxy ethyl methacrylate

can improve the stability of colloid and heat resistance.

Both functional groups improved adhesion of latex

particles due to their additional hydrogen bonding and

dipole-dipole attraction force [1, 3].

Therefore, in this work, the functional groups

on latex particles were modified by incorporating

additional monomers containing carboxylic (-COOH)

and hydroxyl (-OH) groups into the latex. The adhesion

property, colloidal property, and thermal property were

investigated. It was hypothesized that a new PSA would

have an improved adhesion property and would be more

stable than the commercial available products.


178 2. Experimental

2.1 Material

2-Ethyl hexyl acrylate (2-EHA) and vinyl

acetate (VA) monomer were purified by conventional

methods (washed three times with dilute sodium

hydroxide solution and water was removed by vacuum

distillation and dried using calcium chloride). The

initiator in the process was ammonium persulfate.

Sodium dodecylbenzene sulfonate and nonylphenol

ethoxylate were used as surfactant. The chemical

structures of β-carboxyethyl acrylate (BETA-C),

methacrylic acid (MAA), acrylic acid (AA), methyl

methacrylate (MAA), 2-hydroxy ethyl acrylate (2-HEA),

2-hydroxy ethyl methacrylate (2-HEMA) are shown in

Fig. 1. Chain transfer agent was n-dodecyl mercaptan.

Buffer solution was sodium bicarbonate used to adjust

pH (~4). Sodium formaldehyde sulfoxylate and tertiary

butyl hydroperoxide were used to eliminate residue

monomers.

2.2 Synthesis of PSA latex

The adhesives were synthesized in a glass

reactor equipped with a reflux condenser, 4-blade

turbine, N2 purge, and thermometer. The surfactant

system dissolved in DW and the chain transfer agent

dissolved in the monomer were mixed under stirring. The

initiator was added and then heated to 70oC. The pre-

seeding was added continuously to the reactor using

calibrated addition pump for 15 min. After the semi-

continuous addition, the pre-seeding was allowed to be

added continuously for another 4 hr whereupon the

reaction was allowed for further 1 hr, and then sodium

formaldehyde sulfoxylate solution and tertiary butyl

hydroperoxide solution were added. The reaction was

allowed subsequently for another 1 hr. The synthesized

copolymer was cooled to rt. The formulation used in this

study are given in Table 1.

2.3 Dispersion characterization

The particle size and zeta potential were

determined by the static laser scattering technique on the

Malvern instrument zeta-sizer version 6.01.

2.4 Thermal analysis

The samples of synthesized polymer were

coated on glass plate and dried in room temperature for

12 hr. DSC analyzing technique using NETZSCH DSC

204 F1 Phoenix 240 was employed to determine the Tg

of adhesives. Standard 5-10 µg alumina DSC pans with

perforated lids were used and the samples were subjected

to two heating-cooling cycles from -120–50oC at heating

rate of 10oC/min along with an empty reference pan in a

DSC furnace. The second cycle was used for

determination of Tg while the first cycle was used only to

remove previous thermal history.

2.5 Adhesion properties

Adhesion properties were measured on the

adhesive coating on micrometer adjustable film

applicator and backing OPP film (Oriented

polypropylene film). The OPP films were corona-treated

by air. The coating weight of the dry adhesive was in the

range from 20-22 g/m2, drying time 3 min at 70

oC and

aging time 7 days at rt.

Fig 1. Monomers and functional monomers used in

copolymer synthesis


179 Table 1 Monomer feed content in each latex formula

No

.

Monomer main

2-EHA/VA

Functional

Monomer(Ratio) Experimental

Carbonyl

group

Hydroxyl

group

1 90/10 1 0 PSA-1 to PSA-5

2 90/10 0 1 PSA-6 to PSA-8

3 90/10 1 0.5 PSA-9 to PSA-16

4 90/10 1.5 0.5 PSA-17 to PSA-

20

No.

Functional Monomer

BETA-

C MAA AA

MM

A

2-

HEA 2-HEMA

PSA_1 1.25 - - - - -

PSA_2 - 1.25 - - - -

PSA_3 - - 1.25 - - -

PSA_4 - - - 1.25 - -

PSA_5 0.31 0.31 0.31 0.31 - -

PSA_6 - - - - 1.25 -

PSA_7 - - - - - 1.25

PSA_8 - - - - 0.63 0.63

PSA_9 0.84 - - - 0.41 -

PSA_10 - 0.84 - - 0.41 -

PSA_11 - - 0.84 - 0.41 -

PSA_12 - - - 0.84 0.41 -

PSA_13 0.84 - - - - 0.41

PSA_14 - 0.84 - - - 0.41

PSA_15 - - 0.84 - - 0.41

PSA_16 - - - 0.84 - 0.41

PSA_17 - - 0.93 - 0.32 -

PSA_18 - - 0.93 - - 0.32

PSA_19 0.30 0.30 0.30 0.30 0.05 -

PSA_20 0.30 0.30 0.30 0.30 - 0.05

2.6 Fourier transfer infrared analysis

The Fourier transform infrared (FTIR) spectra

of the polymer were recorded using dry sample films on

a Perkin Elmer FTIR 1600 instrument for 16 scans from

600 to 4,000 cm-1

.

2.7 Stability test condition

Three conditions were carried out: stored at

room temperature for 6 months, freeze-thaw condition (-

20oC for 24 hr and room temperature for 6 hr for 15

cycles) and accelerated condition (60oC for 24 hr and

room temperature for 6 hr for 15 cycles)

2.8 Adhesion and cohesion test method

Five test methods were used: the peel adhesion

at 180o (FINAT test method 1), the probe test (Polyken

test) the peel adhesion at 90o (FINAT test method 2),

Initial ball tack (Pressure Sensitive Tape Council test

method 6), the shear resistance (FINAT test method 8)

and loop tack (FINAT test method 9).


3.1 Characteristics of the copolymers

The characteristic FTIR bands of the carbonyl

group at 3,410-3,420 cm-1

(O-H stretch), 1,730- 1,870

cm-1

(C=O stretch), 1325-1,205 cm-1

(C–O stretch),

1,440-1,395 cm-1

and 950-910 cm-1

(O–H bend) were

observed. The characteristics of hydroxyl group was

3,550-3,200 cm-1

(-OH stretch). The bands at 2,800-

2,900 cm-1

(-CH2, -CH3 bend) 1,560 and 1,406-1,410 cm-

1 related to carboxylate group stretching of acrylate were

observed.

3.2 Tg Measurement

In general, PSA adhesion properties showed a

strong dependence on Tg. The Tg of the polymer depends

on the ratio of monomer in formula. In this work, it was

clear that adding just one type of the –COOH or –OH

containing monomers did not significantly affect their Tg

values (-53 to -50ºC) (Table 2; sample 1-7). The slight

reduction of Tg (to about -55 ºC) in sample 8-12 could be

caused by the length of pendent group in 2-HEA that

prevented closed chain packing. Adding more than one

monomer (sample 8-20) somewhat affected the Tg but

rather insignificantly.


180

3.3 Particle size and zeta potential measurement

The size of the copolymer prepared by

miniemulsion polymerization was 200 - 400 nm. Zeta

potential is the indicator of the stability of the colloidal

copolymer particles in the latex system. Since the zeta

potential governs the electrostatic stabilization and steric

stabilization of colloid. The more negative value of zeta

potential, the more stable of the colloid will be. From

this study, the zeta potential of latex no. 3, 4, 7, 16, 19

and 20 was acceptably lower than -30mV (Table 2).

Table 2 Testing results for PSA_1 to PSA_20

No.

%

Con-

versio

n

Tg (oC)

Z-Avg.

(nm)

Zeta

potential

(mV)

Conduc-

tivity

(mS/cm)

PSA_1 98.57 -49.8 318.5 0.9 6.42 x 10-4

PSA_2 96.92 -51.8 400.3 2.4 9.29 x 10-3

PSA_3 97.39 -51.3 277.8 -35.0 2.78 x 10-2

PSA_4 97.81 -52.9 284.0 -35.3 3.13 x 10-2

PSA_5 98.01 -52.8 368.8 4.3 6.25 x 10-4

PSA_6 97.58 -52.4 300.0 1.7 8.56 x 10-4

PSA_7 96.75 -51.5 366.2 -31.2 5.84 x 10-3

PSA_8 98.38 -56.2 390.1 0.6 4.08 x 10-3

PSA_9 96.54 -55.1 318.5 1.4 9.37 x 10-4

PSA_10 98.48 -55.4 315.2 1.8 9.02 x 10-4

PSA_11 98.03 -55.0 311.4 -33.5 2.98 x 10-2

PSA_12 97.19 -55.7 348.4 -32.4 3.12 x 10-2

PSA_13 98.15 -50.9 347.7 0.8 5.16 x 10-4

PSA_14 99.15 -50.4 389.1 2.7 8.57 x 10-4

PSA_15 98.47 -50.3 317.5 -19.8 1.97 x 10-2

PSA_16 98.15 -50.9 347.4 -35.4 3.10 x 10-3

PSA_17 96.40 -54.4 412.0 -10.2 9.81 x 10-4

PSA_18 97.88 -50.8 295.7 -20.4 2.47 x 10-2

PSA_19 98.17 -51.5 337.0 -34.0 2.72 x 10-3

PSA_20 99.45 -51.4 347.8 -36.2 1.45 x 10-3

3.4 Adhesion test

Tack, adhesion and cohesion properties can tell

the difference of the PSA. According to a results (Table

3), PSA_15 and PSA_16 have the highest tackiness since

they consist of AA (-COOH) or MMA (carbonyl in ester)

and hydroxyl (2-HEMA) in the ratio of 2:1, providing

the additional hydrogen bonding attraction between the

polymer chains. Furthermore, 2-HEMA structure was

steric at the pendent group and could initiate higher

tackiness more than the other formulas. However, the

others ratio were not significantly different.

From adhesion properties, i.e. 90º and

180º

peel strength,

the adhesive bond cannot be improved by either the

carboxylic or hydroxyl group (PSA_15 and PSA_16).

For cohesion properties, high cohesion properties were

achieved in the formula with both functional group types

that consist of AA and 2-HEA in the ratio of 2:1

(PSA_11), due to their intermolecular force as well as

the flexibility of the 2-HEA monomer. The results are

shown in Table 3.

3.5 Stability for PSA

The copolymers that were synthesized from monomers

containing functional groups as carbonyl (BETA-C/

MAA/ AA/ MMA) and hydroxyl (2-HEA/ 2-HEMA)

improved the stability of PSA_3, PSA_4, PSA_7,

PSA_16, PSA_19, and PSA_20, as shown in Table 3.

Zeta potential was -30 to -36 mV and conductivity was

1.4510-3

to 3.1310-2

mS/cm. However, the copolymers

PSA_1, PSA_2 and PSA_6 containing BETA-C, MAA,

2-HEMA were not stable (see Table 3). It was possible

that the monomers were not reacted and not incorporated

into the backbone of the copolymers.

4. Conclusions

This work showed that carbonyl and hydroxyl functional

group in the monomer can affect the adhesion and

stability properties of the latex. It was found in this work

that the –COOH was more effective in improving

adhesion than the –OH. The carboxylic group has

stronger dipole-dipole force due to the presence of both

carbonyl (C=O) and hydroxyl group (-OH). This and the


181

soft monomer (2-HEA) used in the formula helped

increase the elasticity of the polymers.

The stability improvement was achieved in the

latex formula containing both functional monomers,

carboxyl and hydroxyl groups (PSA_3, 4, 7, 16, 19, and

20), as also confirmed by their highly negative charge in

zeta potential analysis of the prepared latex. The authors

suggest that the sample PSA_1, 3, 4, 16, 19, and 20 are

suitable for used in adhesive in tape, label, and sticker;

PSA_6 to 7 can be used in protective tapes.

References

[1] Severtson, S., Guo, J., Xu, H., “Properties of water-

based acrylic pressure sensitive adhesive films in

aqueous environments ”, Dep. of Bioproducts and

Biosystem Engineering., University of Minnesota.

[2] Lovell, P. A., El-Aasser, M. S., “Emulsion

polymerization and emulsion polymer” 1st ed.

Wiley, England, 1997.

[3] Lili, Q., Marc, A. D.,“Manipulation of chain transfer

agent and cross-linker concentration to modify latex

micro-structure for pressure-sensitive adhesives”,

European Polymer Journal, volume 46 : 1225-

1236(2010).

Table 3 Results of application and stability test, where X= Not Pass, = Pass

No.

Application results

Zeta

potential (mV)

Conductivity

(mS/cm)

Stability method

Loop tack

(lb/in2)

Initial ball tack

(cm)

90o Peel strength

(kg/in)

180o Peel strength

(kg/in)

Cohesion

(h)

Freeze

thaw Accelerate

Room

Temperature

PSA_1 2.79 4.0 0.51 0.60 70.0 0.9 6.42 x 10-4 X X X

PSA_2 1.88 10.0 0.54 0.50 69.8 2.4 9.29 x 10-3 X X X

PSA_3 1.68 6.5 0.57 0.50 40.0 -35.0 2.78 x 10-2

PSA_4 2.22 8.0 0.51 0.60 30.0 -35.3 3.13 x 10-2

PSA_6 1.93 15.0 0.2 0.30 3.1 1.7 8.56 x 10-4 X X X

PSA_7 2.5 7.0 0.44 0.54 4.2 -31.2 5.84 x 10-3

PSA_8 2.04 15.0 0.48 0.50 63.4 0.6 4.08 x 10-3 X X X

PSA_15 3.50 5.0 0.74 0.79 35.0 -19.8 1.97 x 10-2

PSA_16 3.20 5.1 0.71 0.73 42.0 -35.4 3.10 x 10-3

PSA_19 2.31 6.0 0.74 0.52 32.0 -34.0 2.72 x 10-3

PSA_20 2.14 7.0 0.61 0.62 45.0 -36.2 1.45 x 10-3


182

PROCP-04

Comparative Study on Shear Flow of Olefin Multiblock Copolymers

Anan Sookbanthoeng, Phornsuda Phanjamnonk and Chantima Deeprasertkul*

School of Polymer Engineering, Institute of Engineering, Suranaree University of Technology,

Nakhon Ratchasima 30000

Abstract

Shear flow of olefin multiblock copolymers (OBCs) was investigated using capillary and parallel-plate

rheometers. Viscometric and linear viscoelastic properties of three different melt index OBCs were compared. It

was found that viscometric, after Bagley and Rabinowitsch corrections, and linear viscoelastic properties were

similar. This follows Cox-Merz rule. The entrance pressure loss though seemed to be negligible to all samples.

Time-temperature superposition (TTS) on linear viscoelastic properties failed at temperature of 220C. This

could suggest the presence of the mesophase structure in the melts.

Keywords: olefin multiblock copolymer, Bagley correction, Cox-Merz rule.

1. Introduction

Block copolymers are polymers comprising of two

or more monomers in a chain. Di- and tri-block

copolymers are typically obtained and have been widely

studied. Here the focus is on the olefin multiblock

copolymers (OBCs) which have recently been

synthesized1. These OBCs are ethylene-octene

copolymers which consist of crystallizable blocks (very

low comonomer content) alternating with amorphous

blocks (high comonomer content). With its unique chain

structure, many research on structure and property

relationship have been widely investigated. Interestingly,

mesophase separation transition occurring in these OBCs

has been observed using many techniques. Rheological

measurements was also applied and the presence of the

mesophase transition in the OBC melts have been

reported even at temperature well above the melting

point2,3

.

Many research work have focused on the structure

and property relationship which quite limits in the range

of low shear rates. While polymer processing techniques,

e.g. injection, extrusion, are typically in the high shear

rate range, the high shear behavior of the materials is

inevitable needed. In this work, shear results as obtained

from steady (viscometric properties) and small amplitude

oscillatory (linear viscoelastic properties) of three OBCs

with different melt indices were compared. Cox-Merz

rule was thus applied. By varying the aspect ratio of the

capillary die, entrance pressure loss was studied. Also,

effect of mesophase transition in the melt (if presence) in

the strong shear flow was also of interest.


2.1 Materials

Three olefin multiblock copolymers (OBCs) with

the same density of 0.866 g/cm3 and different melt flow

indices were used. Melt flow index was measured

according to ASTM D1238 (2.16 kg, 190C). The

samples were designated to OBCx, where x is its MFI.

These Dow Chemicals polyolefins were purchased from

Chemical Innovation (Thailand).

2.2 Thermal analysis

Thermal behavior of OBCs was examined using

differential scanning calorimetry (Perkin Elmer DSC

Pyris Diamond). Nitrogen gas was purged throughout the

measurements. The sample was heated from 40 to 160C

(1st heating) and held for 5 min at 160C. Subsequently,

the sample was cooled to 40C (cooling) and heated

again to 160C (2nd

heating). The rate of heating/cooling

was 5C /min. Melting temperature (Tm) of OBCs was

determined.

2.3 Rheological measurements

Rheological measurements were conducted using

both capillary and parallel-plate rheometers. Capillary

shear flow (viscometric) was performed using Gottfert

RG20 rheometer equipped with circular die of 1 mm in


183 diameter. Three dies with aspect ratio (L/D) of 10, 20,

and 30 were used. All measurements were done at

190C. Bagley and Rabinowitsch corrections were

applied to obtain true viscosity. Rheometer (TA

Instrument AR-G2) equipped with 25 mm parallel-plate

geometry was used to obtain linear viscoelastic

properties. Small amplitude oscillatory shear (SAOS),

frequency sweep in a range of 0.1-100 rad/s within 2%

strain was carried out at temperatures 135, 150, 170, 190,

and 220C under nitrogen atmosphere. Master curve was

constructed at reference temperature 190C.


3.1 Thermal properties

DSC thermograms of the samples are depicted in

Figure 1. Melting (from the 2nd

heating) and

crystallization temperatures determined at peak are listed

in Table 1. As seen, the samples melts completely at

below 130C. Rheological studies were conducted at

temperature ranging from 135 to 220C above their

melting temperatures.

Table 1 Physical Properties of OBCs used.

Samples MFI

(g/10min)

(g/cm3)

Tm1

(C)

Tm2

(C)

Tc

(C)

OBC1 1 0.866 - 119.4 90.1

OBC5 5 0.866 118.3 122.5 100.7

OBC15 15 0.866 118.3 122.5 101.8

3.2 Rheological properties

Using capillary rheometer equipped with die of

different L/D ratios (10, 20, and 30) and entrance angle

of 180, for all samples, significant difference in flow

was not noticed using different L/D. Bagley and

Rabinowitsch corrections were performed on the

apparent data. The true and the apparent viscosities were

compared as shown in Figure 2. As seen, the true data

did not much deviate from that apparent results for all

the samples. Though among them OBC1 may show

slightly more difference. The results suggested that there

is only minute entrance pressure loss for these OBC

samples. Bagley plots were shown in Figure 3. If the

mesophase was present in the melts, one would say the

high shear viscosity was not affected.

Figure 1. DSC thermograms of (a) OBC1 (b) OBC5 and

(c) OBC15 where the 1st heating, cooling and 2

nd heating

are bottom, middle and top curves.


184

Figure 2. True (open symbols) versus apparent (closed

symbols) viscosities of (a) OBC1 (b) OBC5 and (c)

OBC15 at 190C with three L/D ratios.

Linear viscoelastic properties were obtained from

small amplitude oscillatory shear (SAOS) measurements

using a rheometer equipped with parallel-plate geometry.

Plots of storage (G) and loss (G) moduli versus angular

frequency () of OBC1 at each temperature are shown in

Figure 4. In the frequency range studied, all G except at

220C decreased as approaching terminal region at low

frequency. Instead, plateau was observed at 220C and

this was also found in OBC5 and OBC15. Using time-

temperature superposition (TTS), master curve at

reference temperature of 190C of OBC1 was

constructed as presented in Figure 5. Typical terminal

region for linear flexible chain is observed. TTS fails to

shift the data at 220C. This could indicate the presence

of mesophase transition in the melts. Further study on

this mesophase in the melts is being conducted.

Figure 3. Bagley plots of (a) OBC1 (b) OBC5 and

(c) OBC15 at 190C using dies with L/D of 10, 20, and

30.

Viscometric and linear viscoelastic properties are

equal as predicted by Cox-Merz rule i.e.

where . True (corrected) shear viscosity ( )

and reduced complex viscosity ( *) at 190C were

compared in Figure 6. As shown, the results follow Cox-

Merz rule very well. This suggests that these OBCs

behave similar to homogeneous linear flexible polymer

melts.


185

Figure 4. Plots of G and G moduli versus of OBC1

at various temperatures.

Figure 5. Master curve of OBC1 at the reference

temperature of 190C.

Conclusions

Shear flow properties of OBCs as measured in

steady and dynamic tests was investigated. According to

Cox-Merz rule steady shear and linear viscoelastic

properties were the same in the homogeneous melts.

Using the circular die of three aspect ratios, relatively

small entrance pressure loss was observed in these

OBCs. The presence of the mesophase structure in the

melts was noticed.

Figure 6. True shear viscosity ( ) and reduced

complex viscosity ( *) of OBCs at 190C

References

[1] Arriola D.J., Carnahan E.M., Hustad P.D., Kuhlman

R.L., Wenzel T.T. Science 2006, 312, 714-719.

[2] Park H.E., Dealy J.M., Marchand G.R., Wang J., Li

S., Register R.A. Macromolecues 2010, 43, 6789-

6799.

[3] He P., Shen W., Yu W., Zhou C. Macromolecules

2014, 47, 807-820.


Natural and Synthetic Rubbers


187

KN-RUBBER-1

New Focus on Rubber Science and Technology

Yuko Ikeda

Kyoto Institute of Technology, Matsugasaki, Sakyo, Kyoto 606-8585, JAPAN

Phone +81 724 7558, Fax +81 724 7558, * E-Mail: [email protected]

Abstract The sulfur cross-linking reaction of rubber, i.e., vulcanization is one

of the most important reactive processes in polymer technology. However, the

details for the reactions have not yet been conclusively clarified, because of the

complicated chemical reactions between rubber and cross-linking reagents at

each processing step. An important key to control the network formation by

sulfur cross-linking is still sought for the development of the rubber industry.

Ikeda et al. was inspired to clarify the vulcanization mechanism by smallangle

neutron scattering techniques, and reported the effects of the combination and

composition of the sulfur cross-linking reagents on rubber network

formations.1 The combination and composition of zinc oxide (ZnO) with other

reagents were found to be crucial to control the structural network

inhomogeneity in the N-(1,3-benzothiazol- 2-ylsulfanyl)cyclohexanamine

(CBS) accelerated vulcanization of isoprene rubber (IR) as shown in Fig. 1. The

mesh size (ξ) in a two-phase inhomogeneous structure of sulfur cross-linked

isoprene rubber was unexpectedly found to be controlled by the amounts of

ZnO and stearic acid (StH) used. A time-resolved zinc K-edge X-ray absorption

fine structure (Zn K-edge XAFS) spectroscopy supported the two-phase

network formation.2

The observations are very important for rubber science and

technology, but it has remained unclear why the combination of ZnO and StH

can be a key for controlling the CBS accelerated vulcanization. Generally, it is

thought that StH can be reacted with ZnO to form zinc stearate (ZnSt2) as an

essential cure activator. Zinc 1,3-benzothiazole-2-thiolate with ligands has long

been postulated to be generated from ZnO, StH, and CBS. It has been accepted

as an active intermediate, where stearate was thought to be one of the ligands.

However, the in situ details of the generated zinc salt of StH for sulfur cross-

linking during the vulcanization have not been well understood. Most studies

have focused on materials isolated from the vulcanization reactions. What is the

role of the zinc salt of StH in the vulcanization reaction? On the way to

determine this, the formation of a specific structural complex generated from

ZnO and StH at a high temperature was recently found by a combination of

time-resolved Zn K-edge XAFS spectroscopy and time-resolved infrared

spectroscopy.3,4 The structure is dominantly a bridging bidentate zinc/stearate

complex, the molar ratio of the zinc ion to stearate and the coordination number

of which are unexpectedly one and four, respectively. Combination with a

density functional calculation for identifying the intermediate predominantly

suggests that its most possible structure is (Zn2(μ-O2CC17H35)2)

2+(OH−)2·XY, where X and Y are water and/or a rubber segment. This

intermediate has been unknown despite the long history of rubber science and

technology. The newly observed zinc/stearate complex may play a role to

accelerate the sulfur cross-linking reaction of rubber like an enzyme because of

the high Lewis activity of the zinc ion.

A filler network, on the other hand, is also an important topic to

reveal a reinforcement effect by filler mixing for rubber materials. However, a

selective formation of clear filler network structure could not be achieved due

to various kinds of filler aggregation by the mechanical mixing. Recently, a

combination of in situ silica filling in the natural rubber (NR) latex and solution

casting was found to be a good method to produce the model nanocomposites

providing a filler network.5,6 Characteristic morphology conferred good

dynamic mechanical properties on the composites. Furthermore, the filler

network structure of the in situ silica-filled NR composites showed a unique

stepwise strain-induced crystallization behavior.7 Pure rubber phases in the

filler network were found to afford highly oriented amorphous segments and

oriented crystallites upon stretching.

These novel observations may be useful for resulting improvements

in performance of rubber materials such as tires in order to construct a green

sustainable society on the Earth.

Yuko Ikeda

Kyoto Institute of Technology

Faculty of Molecular Chemistry and

Engineering Hashigami-cho, Mastugasaki, Sakyo-ku, Kyoto 606-8585,

Japan

Education and Academic Career:

1984 Bachelor of Engineering from Kyoto

Institute of Technology

1986 Master of Engineering from the Graduate School, Kyoto Institute of

Technology 1986 Enrollment to the PhD Course at the

Graduate School of Agriculture Science,

Nagoya University 1988 Left the School

1988 Assistant Professor of Kyoto Institute

of Technology 1991 Doctor of Engineering from the School

of Engineering, Kyoto University. The

title is “Studies on Blood-Compatible Polyurethanes with Triblock Polyether

Soft Segments”

1991 Visiting Researcher at Hyogo Prefecture Technology Center (to 1997)

1993 Visiting Researcher at University of

Bayreuth, Germany Under Prof. Dr. Manfred Schmidt (for six months)

1996 Visiting Researcher at Texas Christian

University of U. S. A. Under Prof. C. D. Gutsche (for six months)

2007 Associate Professor of Kyoto Institute

of Technology 2015 Professor of Kyoto Institute of

Technology

Awards:

1986 Excellent Paper Award from Society

of Rubber Industry, Japan

1996 Excellent Paper Award from Society of Rubber Industry, Japan

2008 The 3rd Excellent Presentation Award

from Society of Rubber Industry, Japan 2014 The 29th Oenslager Award from

Society of Rubber Industry, Japan on

“Fundamental study on cross-linking of rubber”

Major Research Topics:

(1) Sulfur cross-linking of rubber (2) Synthesis, properties and morphology of

physically cross-linked elastomers

(3) Reinforcement of rubber (4) Preparation of functionality elastomeric

materials

Publications:

Original papers, 124; Books, 42; Reviews,

60; Patents, 21; Essays and the others, 26


188

RUBBERO-01

The Use of Modified Palm Oil as Processing Oils in Tyre Tread Applications

Vorapot Thongplod1

, Pongdhorn Sae-oui 2

and Chakrit Sirisinha1, 3

1Department of Chemistry, Faculty of Science, Mahidol University, Salaya Campus, Nakhon Pathom, 73170

2National Metal and Materials Technology Center, Thailand Science Park, Pathumthani, 12120

3Rubber Technology Research Centre (RTEC), Faculty of Science, Mahidol University, Salaya Campus,

Nakhon Pathom, 73170

Abstract

The present work focused on a preparation of Modified Palm Oil (MPO) as benzyl esters of fatty acids.

The MPO was then used as an alternative to petroleum-based Distillate Aromatic Extract (DAE). Fatty acids

were first prepared by a hydrolysis reaction of palm oils and thereafter esterified with benzyl alcohol in the

presence of sulphuric acid as a catalyst. The reaction based on fatty acid: benzyl alcohol: sulphuric acid molar

ratio of 1.5:1.0:0.04 gave high yield of benzyl esters (>70%). The characterisation of MPO was conducted by

FT-IR and 1H-NMR techniques. Blended oils with various ratios of MPO to DAE as well as MPO to TDAE

were prepared and used as Rubber Process Oils (RPOs). Properties of rubber compounds and vulcanisates were

measured. Referred to the overall results gained, the MPO was capable of partly substituting the petroleum-

based DAE in rubber tyre tread applications.

Keyword: Rubber process oil, Palm oil, Modification, Rubber compounding, Mechanical properties

1. Introduction

Rubber process oils (RPOs) are generally added

to rubber compounds for improving processability and

state-of-mix in some circumstances. Petroleum-based

oils are typically major sources of RPOs [1]. Distillate

aromatic extracts (DAE) with high content of polycyclic

aromatic hydrocarbons (PAHs) are widely used as

effective rubber process oils (RPOs) for rubber

manufacturing especially tyre tread applications. This is

because of their high compatibility between rubber

matrix (mostly SBR for tyre tread of passenger car tyres)

and RPOs leading to enhanced product performance.

However, the European legislation classified the DAE as

a carcinogenic substance [2, 3]. Non-carcinogenic

alternatives including treated distillate aromatic extract

(TDAE) and mild extraction solvate (MES) [4] have

been developed to replace DAE in rubber products

especially tyre tread applications. Nonetheless, both

TDAE and MES are still petroleum-based products

which are not renewable.

Palm oil as renewable oil or green oil has

therefore gained attention due to its environmental

friendliness. Recently, attempts to use the palm oil as

RPO have been carried out by modifying the palm oil via

chemical processes. In this work, the modification of as-

received palm oil with benzyl ester to prepare the

Modified Palm Oil (MPO) having low PAH was

conducted. Success of MPO utilisation as RPO by partly

substituting the petroleum-based RPOs was revealed and

discussed.

2. Experimental

2.1 Materials

Palm oil, DAE, TDAE were used as the rubber

process oils (IRPC Plc. Co., Ltd., Thailand).

Chemicals used for modification of palm oil

were as follows: NaOH, HCl and H2SO4 as supplied by

RCI Labscan Ltd., Thailand; benzyl alcohol and Na2CO3

anhydrous as supplied by Ajax Finechem Pty., Ltd.,

Australia.

Compounding ingredients were as follows:

styrene butadiene rubber (SBR 1502; styrene content of

23.5%, JSR Co., Ltd., Thailand), butadiene rubber (BR;

Kelton 2340A, The East Asiatic Plc. Co., Ltd.,

Thailand), carbon black (N330; Thai Carbon Product

Co., Ltd., Thailand), stearic acid (Petch Thai Chemical

Co., Ltd., Thailand), zinc oxide (Kitpiboon Chemical

Part Co., Ltd., Thailand), paraffin wax (Kitpiboon

Chemical Part Co., Ltd., Thailand) and N-t-butyl-2

benzothiazolesulfenamide (TBBS, (Kitpiboon Chemical


189 Part Co., Ltd., Thailand), N,N’-Diphenylguanidine

(DPG, Reliance Technochem Co., Ltd., Thailand), N-

(1,3-dimethylbuthyl)-N’-phenyl-p-phenylenediamine

(6PPD, Flexsys, Belgium), sulphur (Chemmin Co., Ltd.,

Thailand). All compounding ingredients were used as

received.

2.2. Modification and characterisation of palm oils

2.2.1 Preparation of modified palm oils (MPO)

According to Figure 1, fatty acid was prepared

by hydrolysed palm oil with NaOH at 70°C for 9 hrs at a

molar ratio of oil: NaOH of 1:6. Afterwards, saturated

NaCl solution was added into the mixture giving a 2-

layered separation. Hydrochloric acid (4 M) was then

incorporated into the upper layer to perform an

acidification. Subsequently, the solution was neutralised

with water, and the fatty acid product obtained was

heated at 100 ºC to evaporate the trace water before

performing the esterification reaction [5].

Figure 1 Chemical reaction diagram of fatty acid

preparation

For the esterification reaction shown in Figure

2, a fixed amount of benzyl alcohol and H2SO4 as a

catalyst were added into the prepared fatty acid, and then

refluxed at 100ºC for 3 hrs. The molar ratio of fatty acid:

benzyl alcohol: catalyst was fixed at 1.5:1:0.04 [5].

Thereafter, the solution was neutralised with water.

Finally, the residual fatty acid in the mixture was

eliminated by extracting with 10% w/v Na2CO3 solution,

leading to an occurrence of desired benzyl ester.

Figure 2 Chemical reaction diagram of the esterification

reaction with benzyl alcohol.

2.2.2 Characterisations of MPO

The MPO prepared was characterised using: (i)

a Fourier transform infrared spectroscopy (FTIR; Bruker,

Equinox 55) with Attenuated Total Reflectance (ATR)

mode and (ii) 300 MHz proton nuclear magnetic

resonance spectroscopy (1H-NMR; Bruker, AM400) with

deuterated chloroform (CDCl3) as a solvent and

tetramethylsilane (TMS) as an internal standard.

2.3 The use of MPO as RPOs in rubber compounds

2.3.1 Rubber compound preparation

Mixing for preparing rubber materbatch

(denoted as Compound A) was performed using

Brabender Plasticorder with rotor speed, fill factor and

initial mixing chamber temperature of 50 rpm, 0.75 and

50°C, respectively. Sulphur and accelerators were mixed

with Compound A on a two-roll mill. The RPOs with

various DAE/MPO ratios of 20/0, 15/5, 10/10, 5/15 and

0/20, and various TDAE/MPO ratios of 20/0, 15/5,

10/10, 5/15 and 0/20 ratios were used at a given total

RPO loading of 20 phr.

2.3.2 Rheological properties

Mooney viscosity (ML1+4 @ 100 °C) of rubber

compounds was measured using the Mooney viscometer

(TechPro viscTECH+, USA). Rheological properties of

rubber compounds were monitored using the Rubber

Process Analyser (Alpha Technologies model RPA2000,

USA).

O

O

O

O

O

O

OH

O OH

HO

OH

OH

O

O

O

HO

3 NaOH +

Palm oil (triglyceride)

Fatty acid Glycerol

- Palm oil : NaOH (by mole) = 1:6

at 70°C for 9 hrs. - NaCl (saturated)

- HCl (4M)

- Water

Benzyl ester as a product

Benzyl alcohol Fatty acid - Fatty acid : Benzyl alcohol : H

2SO

4

(by mole) = 1.5:1:0.04, refluxing for 3 hrs.

- Neutralizing with water

- Adding 10% w/v Na2CO

3 solution

+

+ 3

http://en.wikipedia.org/wiki/Nuclear_magnetic_resonance_spectroscopy

http://en.wikipedia.org/wiki/Nuclear_magnetic_resonance_spectroscopy


190

Table 1 Compounding recipe used in this work.

*phr = part per hundred of rubber

2.3.3 Cure characteristics

Cure behaviour was determined with the use of

moving die rheometer (TechPro MDR, USA) at 160 °C.

2.3.4 Mechanical properties

To prepare rubber vulcanisates, the compounds

were compression moulded under clamping pressure of

14 MPa at 160 °C for the optimum cure time (tc90) as

pre-determined from the MDR.

Hardness was measured using Shore A

durometer (H17A, Cogenix Wallace) according to

ASTM D2240. Abrasion resistance was determined

using abrasion tester (Zwick model 6120) as per DIN

53516.

2.3.5 Determination of heat build-up

The flexometer (BF Goodrich flexometer Model

II) was used to assess the temperature rise in

vulcanisates, or the so-called heat build-up (HBU) as per

ASTM D623.


3.1 Characterisation of MPO

The comparative FTIR spectra of unmodified

palm oil (i) and MPO prepated (ii) are shown in Figure

3. The additional absorption band at a wavenumber of

700 cm−1

is observed in the MPO representing the out-

of-plane bending of C–H on the mono-benzene. Also, the

MPO displays the absorption bands at 3,400 cm−1

due to

the stretching vibration of O–H in residue free fatty acid

structure of MPO.

Figure 3 FT-IR spectra of unmodified palm oil (i) and

MPO (ii).

Figure 4 shows 1H-NMR spectra of unmodified

and modified palm oils. After modification, the MPO

exhibits the chemical shift at 7.4 ppm attributed to the

proton in aromatic ring of benzyl alcohol, suggesting a

complete esterification. Other signals at 4.7 - 5.1 ppm are

assigned to the hydroxyl proton in residue fatty acid.

Figure 4

1H-NMR spectra of unmodified and modified

palm oils.

Ingredients Amount

(phr*)

SBR 1502 70

BR 01 30

ZnO 3

Stearic acid 1

N330 60

RPOs 20

6PPD 3

Paraffin wax 1

TBBS 2

DPG 0.5

Sulphur 2

Blends

oils

(T)DAE / MPO

20/0

15/5

10/0 5/15

0/20


191 3.2 Rheological properties

Figure 5 shows Mooney viscosity of compounds

with various RPO blend ratios. Evidently, at any given

blend ratio, the TDAE based offers lower compound

viscosity. Furthermore, an increased portion of MPO is

capable of lowering the compound viscosity, implying

the processability improvement. Figure 6 represents the

effect of RPO blend ratio on Payne effect of rubber

compounds and thus degree of filler dispersion. In

theory, the greater the G’, the poorer the dispersion

level [6, 7]. It is obvious that, the increased proportion of

MPO in DAE/MPO blends reduces the magnitude of

Payne effect, suggesting the improved degree of filler

dispersion in rubber compounds. However, no significant

difference in the TDAE/ MPO blend system is observed.

Figure 5 Money viscosity of SBR/BR blends with

various RPO blend ratios.

Figure 6 Payne effect of SBR/BR blends with various

RPO blend ratios.

3.3 Cure characteristics

Figure 7 illustrates scorch time (ts2) of

compounds incorporated with various blend ratios of

(T)DAE/MPO. With increasing MPO proportion in

blended RPO, the scorch time decreases suggesting the

improved cure efficiency with MPO. Such cure

promotion might be caused by the existence of benzyl

ester acting as a cure activator [8].

Figure 7 Scorch time (ts2) of SBR/BR blends with

various RPO blend ratios.

Figure 8 reveals torque difference (MH – ML) as

an indication of the crosslink density of vulcanisates. It is

evident that the crosslink density of vulcanisates

increases (i.e., cure promotion) with increased proportion

of MPO in blended RPO. This could be attributed to the

presence of fatty acid ester as reported elsewhere [8].

Figure 8 Torque difference (MH-ML) of SBR/BR blends

incorporated with various RPO blend ratios.


Figure 9 exhibits that the hardness of vulcanisates

is independent of MPO proportion in blended RPOs.

1

2

3

4

5

100/0 75/25 50/50 25/75 0/100

t s2

(m

in.)

Blend ratio (%)

DAE/MPO

TDAE/MPO


192

Figure 9 Hardness of SBR/BR blends with various RPO

blend ratios.

Figure 10 reveals the decreases in abrasion loss

of vulcanisates with increasing MPO proportion in

blended RPO. Such improvement in resistance to

abrasion is attributed to the increased crosslink density

and enhanced degree of filler dispersion as illustrated in

Figures 6 and 8.

Figure 10 Abrasion loss of SBR/BR blends with various

RPO blend ratios.

Figure 11 demonstrates the decrease in heat

build-up (HBU) of vulcanisates with increasing MPO

portion in blended RPO. Again, the increased in

crosslink density and state-of-mix are responsible [9].

Figure 11 Heat build-up of SBR/BR blends with various

RPO blend ratios.

4. Conclusion

The modified palm oil (MPO) was successfully

prepared and characterised by FT-IR and 1H-NMR

techniques. The RPO blends prepared from MPO and

distillate aromatic extracts (DAE) or treated distillate

aromatic extracts (TDAE) are capable of functioning as

rubber process oil (RPO) effectively. The improvement

of processability and crosslink density are found in the

blended RPOs with increased MPO portion. Superiority

in heat build-up and abrasion resistance are resulted in

blended RPO having relatively high MPO proportion.

5. Acknowledgment

The authors would like to express their gratitude

to the Thailand Research Fund-Research and Researcher

for Industry (TRF-RRI) and IRPC Plc. Co., Ltd. for the

financial support of this research.

6. References

[1] Lansdown, A. R., “Lubrication and Lubricant

Selection: A Practical Guide, 3rd

Edition”, Tribology

in Practice Series, United Kingdom, Amer Society

of Mechanical Engineers (2003).

[2] Dasgupta, S., Agrawal, S. L., Bandyopadhyay, S.,

Chakraborty, S., Mukhopadhyay, R., Malkani, R. K.

and Ameta, S. C., “Characterization of Eco-Friendly

Processing Aids for Rubber Compound”, Polymer

Testing, 26(4): 489-500 (2007).

40

50

60

70

80

90

100

100/0 75/25 50/50 25/75 0/100

Ab

rasi

on

loss

(m

m3)

Blend ratio (%)

DAE/MPO

TDAE/MPO


193 [3] Sahakaro, K. and Beraheng, A., “Epoxidized Natural

Oils as the Alternative Safe Process Oils in Rubber

Compounds”, Rubber Chemistry and Technology,

84(2): 200-214 (2011).

[4] Null, V., “Safe Process Oils for Tires with Low

Environmental Impact”, Raw Meterials and

Applications, KGK 52: 799-805 (1999).

[5] Boontawee, H., Nakason, C., Kaesaman, A.,

Thitithammawong, A. and Chewchanwuttiwong, S.,

“Application of Benzyl Ester of Modified Vegetable

Oils as Rubber Processing Oils”, Advanced

Materials Research, 415-417: 1164-1167 (2012).

[6] Takino, H., Iwama, S., Yamada, Y. and Kohjiya, S.,

“Effect of Processing Additives on Carbon Black

Dispersion and Grip Property of High-Performance

Tire Tread Compound”, Rubber Chemistry and

Technology, 70(1): 15-24 (1997).

[7] Byers, J. T., “Fillers for Balancing Passenger Tire

Tread Properties”, Rubber Chemistry and

Technology, 75(3): 527-548 (2002).

[8] Barlow, F. W., “Rubber Compounding, 2nd

Edition”,

Principles, Materials and Techniques, New York,

Marcel Dekker (1988).

[9] Sombatsompop, N. and Kumnuantip, C., “Rheology,

Cure Characteristics, Physical and Mechanical

Properties of Tire Tread Reclaimed Rubber/Natural

Rubber Compounds”, Journal of Applied Polymer

Science, 87(10): 1723-1731 (2003).


194 RUBBERO-02

Thermoplastic Elastomers Based on Graft Copolymers of Natural Rubber and

Poly(diacetone acrylamide)/Polyamide-12

Gosalee Phersalaeh1 Bencha Thongnuanchan

1*, Anoma Titithammawong

1 and Charoen Nakason

2

1Department of Rubber Technology and Polymer Science, Faculty of Science and Technology, Prince of

Songkla University, Pattani, 94000, Thailand 2Faculty of Science and Industrial Technology, Prince of Songkla University, Surat Thani, 84000, Thailand

Phone +66 07331 2213, *E-Mail: [email protected]

Abstract

The aim of the present study was to improve the compatibility in blends of natural rubber (NR) and

polyamide-12 (PA-12) by grafting hydrophilic monomer, diacetone acrylamide (DAAM), onto NR backbone.

Graft copolymers of NR and poly(DAAM) prepared using 10 wt% of DAAM (NR-g-PDAAM10) was first

synthesized. Blends of NR-g-PDAAM10/PA-12 was then prepared at a blend ratio of 60/40 (wt%) by simple

blend and dynamic vulcanization techniques. The mechanical and rheological properties of the resulting blends

were subsequently investigated and compared to those of the corresponding blends based on unmodified NR.

The results reveal that dynamic vulcanization led to a significant increase in both mechanical and rheological

properties of the blends. The size of vulcanized rubber particles in thermoplastic vulcanizate (TPV) based on the

NR-g-PDAAM10/PA-12 blend was found to be smaller than those in the NR/PA-12 TPV, which is probably

due to the compatibilizing effect of DAAM groups present in NR-g-PDAAM10 molecule. It was also observed

that the former exhibited higher tensile strength and elongation at break than the latter. These results indicate

that the interfacial adhesion between NR and PA-12 phases was improved by the presence of DAAM groups in

NR molecule.

Keywords: Thermoplastic Elastomers, Natural Rubber, Diacetone Acrylamide.

1. Introduction

Thermoplastic elastomers (TPEs) are materials

which possess the characteristics of both thermoplastics

and elastomers. These materials exhibit properties

similar to those of vulcanized rubbers at ambient

temperatures and they can be processed in the molten

state as thermoplastics. TPEs can be broadly classified

into two main groups: TPEs based on block copolymers

and those derived from rubber/thermoplastic blends

(blended TPEs). Thermoplastic natural rubber (TPNR) is

one of blended TPEs which have gained considerable

interest in recent years due to its relative ease of

preparation. TPNR can typically be made by blending

natural rubber (NR) with thermoplastics at temperatures

above the melting point for semi-crystalline polymers or

above the glass transition temperature for amorphous

polymers [1].

Polyamides (PAs) are semi-crystalline polymers

which contain the amide (-CONH-) linkage in their

backbone. The presence of polar amide groups along the

PAs backbones gives rise to inter-chain hydrogen bonds,

which account for good mechanical properties of PAs

[2]. However, the degree of compatibility in blending

PAs with hydrophobic NR is relatively low. This would

lead to the formation of incompatible blends with poor

mechanical properties due to dissimilarity in polarity.

Thus, the modified forms of NR, that bear functional

groups capable of interacting with the functional groups

of PAs (i.e., -NH2, -COOH, -NH-(C=O)-), is preferably

used for blending with PAs. Blending of epoxidized

natural rubber (ENR) with PAs (i.e., polyamide 6 [3] and

PA-12 [4]) is the most widely studied blending system.

Diacetone acrylamide is a functional monomer

containing polar amide and ketone groups in its

molecule. The hypothesis of the present study is that

grafting of hydrophilic DAAM monomer onto the NR

backbone would not only increase its polarity, but also it

would improve interfacial adhesion between PAs and

NR. It is expected that the polar functional groups in the

grafted poly(diacetone acrylamide), PDAAM, chains

would participate in hydrogen bonding with amide

groups in PAs chains (see Figure 1), resulting in better

blend compatibility.


195

N CH2 CH2 CH2 C

O

H

. .

n9

CH C

O

N C

CH3

CH3

CH2 C

O

CH3

H

CH2

.

NR

NCH2CH2CH2C

O H

..

9 m

Nylon-12

Nylon-12

NR-g-PDAAM

Figure 1. The hydrogen bonds formed between NR-g-

PDAAM and PA-12 chains.

In the present study, blends of NR/PA-12 at

a fixed blend ratio of 60/40 (wt%) were prepared by

simple blend and dynamic vulcanization techniques.

Effect of types of natural rubber (i.e., unmodified NR

and graft copolymers of natural rubber and

poly(diacetone acrylamide), NR-g-PDAAM) on the

properties of the blends were investigated.

2. Experimental Section

2.1 Materials

Two types of natural rubber were used in the

present study: air dired sheet (ADS) and NR-g-PDAAM

prepared using 10 wt% of DAAM (NR-g-PDAAM10).

ADS was manufactured by a local factory operated by

Khuan Pan Tae Farmer Cooperation (Phattalung,

Thailand). Diacetone acrylamide (DAAM) with purity

99% was manufactured by Sigma-Aldrich Chemicals

(Steinheim, Germany). NR-g-PDAAM10 was prepared

via seeded emulsion polymerization. The details of

preparation and characterization procedures of the NR-g-

PDAAM10 have been described elsewhere [5]. The

amount of grafted PDAAM in the NR-g-PDAAM as

determined by 1H-NMR is about 5.47 wt%. Injection

molding PA-12 grade (Grilamid L20G) with a melting

temperature of 178C was used in this work. It was

manufactured by EMS-Grivory GmbH, Gross-Umstadt,

Germany. All chemicals were used as received.

2.2 Preparation of rubber/PA-12 blends

In the present study, the blend ratio of rubber/

PA-12 was fixed at 60/40 (%wt/wt). The blends were

prepared in an internal mixer with a mixing chamber of

50 cm3 (Brabender® GmbH & Co.KG, Germany), with

a rotor speed of 60 rpm. Both the simple blend and

dynamic vulcanization techniques were prepared using

a two-step mixing process. In the simple blend

technique, rubber (i.e., NR or NR-g-PDAAM10) was

first masticated for 1 min at 50°C before it was mixed

with

an antioxidant (i.e., 6PPD) for 1 min. The compounding

formulation is shown in Table 1. When the mixing step

was completed, the mixture was removed from the

mixing chamber. Simple blends were then prepared by

melt mixing the pre-compounded rubber with pre-dried

PA-12 at 160°C for 6 min. It is important to note that the

temperature in the mixing chamber was gradually raised

to about 200°C due to the frictional heat before the blend

was dumped out of the mixer.

Table 1. Compounding formulations for simple blends

and dynamic vulcanized blends

Ingredients Quantities (phr)

Simple blend Dynamic vulcanization

Rubber 100 100

6PPD 1 1

HRJ-10518 - 7

For dynamic vulcanization technique, the pre-

compounded rubber was first prepared by mixing rubbers

with 6PPD and HRJ-10518 (i.e., a phenolic curing agent)

at 50°C for 1 min, using the compounding formulation as

shown in Table 1. The rubber was masticated for 1 min

before antioxidant and curing agent were sequentially

added into the mixing chamber. Thermoplastic

vulcanizate was prepared by melt-mixing of the pre-

compounded rubber with PA-12, using the same mixing

condition as the simple blend technique.

3. Characterizations

3.1 Contact angle measurement

Contact angles of different test liquids on the

surfaces of NR, NR-g-PDAAM10, and PA-12 were

determined by the sessile drop method at ambient

humidity using a contact angle meter (DM300, Kyoma

Interface Science Co., Japan). Three test liquids used for

the contact-angle measurements were distilled water,

ethylene glycol, and formamide.


196 3.2 Mechanical testing

Tensile properties of rubber/PA-12 blends were

measured according to ASTM D412 using dumbbell-

shaped specimens. The test was conducted at 252°C at

a crosshead speed of 500 mm/min using Hounsfield

Tensometer H 10 KS (the Hounsfield Test Equipment

Co., Ltd, U.K.).

3.3 Morphological characterization

Leo scanning electron microscope (model VP

1450, Leo, UK) was employed to examine morphologies

of cryogenically fractured surfaces of rubber/PA-12

blends. After the blends were fractured under liquid

nitrogen, one of phases was preferentially extracted prior

to SEM analyses.

3.4 Rheological characterization

Shear flow properties of rubber/PA-12 blends in

term of relationship between apparent shear viscosity

with apparent shear rate were studied using a capillary

rheometer, G ttfert Rheo-Tester 2000, (Werkstoff-

Pr fmaschinen GmbH, Germany). A capillary die has

a diameter (D) of 1.0 mm, a length (L) of 20 mm (i.e., an

aspect ratio L/D=20/1), and an entrance angle of 180°.

The measurements were conducted over a wide range of

shear rates (i.e., 5–1000 s−1

) at 180°C.


4.1 Contact angles

Contact angle can be defined as the angle formed

between a substrate surface and a tangent line drawn

from the contact point of liquid droplet along the liquid-

vapor interface [6]. In the present work, the contact angle

of water on the surfaces of NR, NR-g-PDAAM and PA-

12 were measured to characterize their hydrophilicity.

The average of at least five different measurements was

taken as the value of contact angle.

The results in Table 2 show that the water contact

angle for NR was 104.8°. After NR was modified by

grafting with 10 wt% of DAAM, its water contact angle

was decreased to 86.7°. The reduction of the water

contact angle is due to the presence of polar groups (i.e.,

amide and ketone) in the NR-g-PDAAM10 molecules.

This allows a water drop to partially wet its surface and

minimizes the water contact angle, as evident from

Figure 2. In the case of PA-12, the average water contact

angle was found to be 89.3°. Furthermore, it can also be

seen form Table 2 that the values of contact angles for

a particular liquid on the same polymer surface were

different. This is attributed to the difference in surface

tension of different test liquids.

Table 2. Values of contact angles for different test

liquids against the surfaces of NR, NR-g-PDAAM10,

and PA-12

Materials Contact angle (degree)

distilled water ethylene glycol Formamide

NR 104.8 87.6 63.2

NR-g-PDAAM10 86.7 68.5 79.7

PA-12 89.3 68.4 70.1

Figure 2. Water droplets formed on the surfaces of (a)

NR, (b) PA-12 and (c) NR-g-PDAAM10.

Fowkes proposed that the surface free energy of

both solid and liquid can be assumed to be the sum of

two major components (i.e., polar components, P, and

dispersive components, D) [7]. The values of contact

angles measured for different test liquids on a particular

surface can be used to estimate the surface free energy

(). In this work, a video-based software SCA 20

(Dataphysics Instruments GmbH, Germany) was used to

estimate surface free energy from contact angle results.

The estimated values of surface free energies for NR,

NR-g-PDAAM10, and PA-12 are presented in Table 3.

The results reveal that the polarity and hydrophilicity of


197 NR increased after grafting with PDAAM. This is

evident from a dramatic increase in the value of its polar

component from 1.16 to 7.49 mJ.m-2

. Moreover, it is also

observed that the difference in surface energies between

NR-g-PDAAM10 and PA-12 was smaller than that

observed between NR and PA-12. A small value of

surface energy difference between polymer pairs is

expected to provide good adhesion between the two

phases. Hence, these results suggest that PA-12 is more

likely to be compatible with NR-g-PDAAM10 than

unmodified NR in a blend.

Table 3. Values of surface energies and its components

for NR, NR-g-PDAAM10, and PA-12

Materials Surface energy (mJ.m-2)

Total Dispersion Polar

NR 43.78 42.62 1.16

NR-g-PDAAM10 26.23 18.74 7.49

PA-12 29.53 24.24 5.29

4.2 Mechanical and morphological properties

Figure 3 shows stress-strain curves for simple

blends of NR/PA-12 and NR-g-PDAAM10/PA-12

compared with their dynamically vulcanized

counterparts. It is clearly seen that NR-g-PDAAM10/PA-

12 blend exhibited higher tensile strength and elongation

at break than NR/PA-12 blend. This is because the polar

functional groups present in the NR-g-PDAAM

molecules facilitate compatibility between the rubber and

PA-12 phases.

It can also be observed that when the rubber phase

(i.e., NR or NR-g-PDAAM10) in the simple blends was

dynamically cured during its melt mixing with

PA-12, the stiffness and toughness of these materials

referred to as thermoplastic vulcanizates (TPVs)

noticeably increased. It is evident from the observation

that the initial slopes of the stress-strain curves for both

types of TPVs were steeper than those of the

corresponding simple blends. It was also observed that

these TPVs showed a higher tensile strength and

elongation at break as compared with their simple

blends. This observation can be explained on the basis of

morphological change occurred during dynamic

vulcanization.

Figure 3. Stress-strain curves of simple blends and

thermoplastic vulcanizates (TPV) based on NR/PA-12

and NR-g-PDAAM10/PA-12 blends at a blend ratio of

rubber/PA-12 = 60/40 wt%.

Figure 4. SEM micrographs of 60/40 rubber/PA-12

blends based on NR and NR-g-PDAAM10: (a) NR/PA-

12 simple blend (×500), (b) NR-g-PDAAM10/PA-12

simple blend (×500), (c) NR/PA-12 TPV (×10,000), (d)

NR-g-PDAAM10/PA-12 TPV (×10,000).

Morphological studies of the simple blends were

then carried out by examination of cryo-fractured

surfaces using SEM. Rubber phases (i.e., NR or NR-g-

PDAAM) were etched with a solvent mixture of toluene

and methyl ethyl ketone (50/50 v/v) for 3 days at

ambient temperature prior to SEM analyses. It can be

0

4

8

12

16

20

24

0 20 40 60 80 100 120 140 160 180 200

NR/PA-12 TPVNR-g-PDAAM/PA-12 simple blend

NR-g-PDAAM/PA-12 TPV

NR/PA-12 simple blend

Str

ess

(MP

a)

Strain (%)


198 seen from Figures 4a and 4b that the simple blends based

on blends of NR/PA-12 and NR-g-PDAAM10/PA-12

exhibited co-continuous phase morphology. However,

the latter exhibited finer grain morphology than the

former, which corroborated that PA-12 was more

compatible with NR-g-PDAAM than NR.

Table 4. Mechanical properties of rubber/PA-12 blends

based on unmodified NR and NR-g-PDAAM10

Rubber/PA-12

blends

Mechanical properties

Tensile strength

(MPa)

Elongation

at break (%)

Tension set

(%)

NR/PA-12 blend 5.1 0.7 153 16 458

NR-g-PDAAM10

/PA-12 blend 7.4 1.2 166 20 404

NR/PA-12 TPV 14.4 1.4 146 17 322

NR-g-PDAAM10

/PA-12 TPV 17.1 0.6 196 14 282

During the dynamic curing process, the in-situ

vulcanization of rubber phase occurs during melt

blending with molten thermoplastic. As vulcanization

proceeds, the viscosity of rubber phase increases

significantly due to an increase in the degree of

crosslinking. Consequently, the crosslinked rubber phase

will break up into fine rubber particles, which are

dispersed in the thermoplastics matrix, due to the shear

stress. This leads to the formation of dispersed-phase

morphology. The morphological analysis of the NR/PA-

12 and NR-g-PDAAM10/PA-12 TPVs was carried out

using SEM. The PA-12 phase was selectively extracted

from the TPVs with hot N,N’-dimethylformamide

(~100°C) prior to SEM analysis. As expected, both types

of the TPVs show dispersed-phase morphology where

the size of vulcanized rubber particles dispersed in the

PA-12 matrix was within 0.8-3 m range (see Figures 4c

and 4d).

It becomes obvious that dynamic vulcanization

causes the change in phase morphology of the simple

blends from co-continuous phase morphology to

crosslinked rubber particles dispersed in a continuous

PA-12 matrix. Hence, the strength of semi-crystalline

PA-12 matrix is expected to provide a significant

increase in the mechanical properties of TPVs. This

would result in better mechanical properties of the TPVs

when compared with the simple blends. Moreover, the

crosslinking of rubber phase during melt mixing

substantially improved elasticity of the simple blends,

which was reflected in the lower value of tension set, as

presented in Table 4.

In addition, it is also observed that

PDAAM10/PA-12 TPV showed smaller particles of

vulcanized rubber dispersed in the PA-12 matrix as

compared to those observed for NR/PA-12 TPV. This

result is consistent with the observation that tensile

strength and elongation at break of the former were

higher than those of the latter. This is because the

mechanical properties of TPVs depend strongly on the

size of vulcanized rubber particles [8]. Smaller size of

vulcanized rubber particles are expected to provide better

mechanical properties of TPVs as they would facilitate

greater stress distribution between the thermoplastic

matrix and the dispersed vulcanized rubber.

Figure 5. Plots of apparent shear viscosity versus

apparent shear rate of simple blends and TPVs prepared

from NR/PA-12 and NR-g-PDAAM10/PA-12 blends

with a blend ratio of 60/40 wt% at 180°C.

4.3 Rheological properties

Figure 5 displays the apparent shear viscosity at

various apparent shear rates for the simple blends of

NR/PA-12 and NR-g-PDAAM10/PA-12 compared with

their dynamically vulcanized counterparts. In all cases,

a shear-thinning behaviour (i.e., pseudo-plasticity) was

observed as the apparent shear viscosity decreased with

increasing the shear rate. Furthermore, it was found that,

1E+2

1E+3

1E+4

1E+5

10 100 1000 104

NR-g-PDAAM/PA-12 TPVNR/PA-12 TPVNR-g-PDAAM/PA-12 simple blendNR/PA-12 simple blend

Ap

pea

ren

t sh

ear v

iscosi

ty (

Pa.s

)

Apprarent shear rate (1/s)


199 at a given shear rate, the rubber/PA-12 TPVs clearly

exhibited higher apparent shear viscosity than those of

the corresponding simple blends. This indicates that the

vulcanized rubber particles dispersed in the PA-12

matrix restrict the movement of polymer chains in the

direction of shearing flow, resulting in higher values of

apparent shear viscosity.

Additionally, it is also seen that the apparent shear

viscosity at low shear rates of the NR-g-PDAAM10/PA-

12 TPV were much higher than those of NR/nylon-12

TPV. This can be related to the presence of polar groups

in the NR-g-PDAAM molecule which make its more

compatible with nylon-12 than a non-polar NR.

Therefore, the NR-g-PDAAM/nylon-12 TPV is expected

to have better interfacial adhesion between polymeric

phases, leading to a higher melt viscosity. These results

provide supportive evidence that the DAAM groups

present in the NR-g-PDAAM10 molecules enhance the

blend compatibilization between NR and PA-12 phases

and, subsequently, improve properties of the blends.

5. Conclusion

In the present study, an attempt was made to

demonstrate that the grafting of NR with hydrophilic

DAAM improved its compatibility with PA-12. TPEs

based on blends of either unmodified NR or NR-g-

PDAAM10 with PA-12 were then prepared at a 60/40

(wt%) blend ratio of rubber/PA-12 by simple blend and

dynamic vulcanization techniques. Dynamic vulca-

nization of the rubber phase was shown to markedly

enhance the mechanical and rheological properties of the

blends. Smaller size of vulcanized rubber particles

dispersed in the PA-12 matrix was seen in the NR-g-

PDAAM10/PA-12 TPV. This is probably responsible for

the higher tensile properties of the NR-g-PDAAM10/PA-

12 TPV when compared with the NR /PA-12 TPV. It

was also observed that the apparent shear viscosities in

low shear regions of the NR-g-PDAAM10/PA-12 TPV

was much higher than those of the NR/PA-12 TPV.

These results indicate that the interfacial adhesion

between NR and PA-12 was improved by the presence of

DAAM groups in NR molecule.

Acknowledgments

This work was supported by the Higher Education

Research Promotion and National Research University

Project of Thailand, Office of the Higher Education

Commission.

References

[1] Ibrahim, A. and Dahlan, M. Prog. Polym. Sci. 1998,

23, 665-706.

[2] Baker, A-M. and Mead, J.L., “Thermoplstics”, In:

Harper, C. A. (ed) Handbook of Plastics Technologies,

New York, McGrawHill: 2.11-2.14 (2006).

[3] Narathichat, ., Kummerl we, C., Vennemann, N.

and Nakason, C. J. Appl. Polym. Sci. 2011, 12, 805–814.

[4] Tanrattanakul, V., Sungthong, N. and Raksa, P.

Polym. Test. 2008, 27, 794–800.

[5] Thongnuanchan, B., Ninjan, N., Kaesaman, A. and

Nakason, C. Polym. Bull. 2015, 72, 135–155.

[6] Yuan, Y. and Lee, T. R., “Contact Angle and

Wetting Properties”, In Bracco, G., and Holst, B. (ed.s)

Surface Science Techniques, New York ,Springer-Verlag

Berlin Heidelberg : 3-5 (2013).

[7] Fowke, F.M. Ind. Eng. Chem. 1964, 56, 40-52.

[8] Coran, A.Y. and Patel, R. Rubber Chem. Technol.

1983, 56, 1045-1060.

http://link.springer.com/search?facet-author=%22Bodil+Holst%22


200

RUBBERO-03

Thermoplastic Vulcanizates Based on Natural Rubber/Propylene-Ethylene Copolymer

Blends: Influence of Viscosity and Ethylene Contents of the Copolymer on the Properties

T. Wohmang B. Thongnuanchan

and A. Kaesaman

*

Department of Rubber Technology and Polymer Science, Faculty of Science and Technology,

Prince of Songkla University, Pattani, 94000

Phone +6689 658 5892, Fax +66 73331099, *E-mail: [email protected]

Abstract

This work investigates the effect of viscosity and ethylene content of propylene-ethylene copolymer

(PEC) on the properties of thermoplastic vulcanizates (TPVs) based on natural rubber (NR) and PEC blends.

Blends of NR and PECs with different viscosities and ethylene contents were prepared at a blend ratio of

NR/PEC = 40/60 wt% by dynamic vulcanization. The results reveal that the properties of NR/PEC TPVs are

significantly influenced by the melt flow index (MFI), and ethylene content of the PECs. Tensile strength of the

TPVs decreases slightly but elongation at break increases with increasing the viscosity of PECs. It is also

observed that the apparent shear viscosity at various shear rates of the TPVs increases with a decrease in the

MFI of the PEC matrix. Increasing the ethylene content in the PEC results in a decreases in the tensile strength

of the TPVs whereas the elongation at break increases. The NR/PEC TPVs can be reprocessed, but deterioration

in their mechanical properties is observed.

Keywords: Thermoplastic vulcanizates (TPVs), Propylene – ethylene copolymer (PEC), Natural rubber (NR),

Dynamic vulcanization.

1. Introduction

In recent years, themoplastic vulcanizates (TPVs)

are materials which have drawn remarkable interests by

both academia and industry as they combine the ease of

processing of thermoplastics with the characteristics of

vulcanized rubber. The TPVs are normally phase

separated systems, in which one phase is soft and

rubbery at room temperature while the other is hard and

solid. The advantages of TPVs over rubbers are their

processing properties as they can be processed like the

conventional thermoplastic and so they can be recycled.

The materials are likely to be used to replace the

conventional vulcanizated rubber products wherever they

are applicable. TPVs can be applied for a variety of

products such as automotive parts, electronic devices,

medical equipment and industrial buildings, etc. [1].

TPVs based on blends of natural rubber (NR) and

polyolefins (i.e., polypropylene, PP, and polyethylene,

PE) is one of the most widely studied systems in the field

of thermoplastic elastomers. Polyolefins are widely used

for the preparation of TPVs due to their good mechanical

properties, chemical resistance, and relatively low cost

[2-3]. PP is often considered as the polymer of choice for

blending with NR as it has a higher softening

temperature (i.e.,105°C) than PE, resulting in a higher

service temperature for NR/PP blends. TPVs based on

NR/PP blends with different mechanical properties can

be typically prepared by varying the blend compositions.

Additionally, copolymers of propylene and ethylene

(PEC) are often used in place of PP homopolymer for

blending with NR when greater impact strength and

ductility at low temperatures are required. This is

because PP has relatively low impact strength at a

temperature below 0ºC. In the present work, TPVs based

on blends of NR and different grades of PEC were

prepared at a blend ratio of NR/PEC = 40/60 wt% by

dynamic vulcanization. PEC (Versify®, Dow Chemical)

is selected to blend with NR in order to produce TPVs.

PECs with different viscosities and ethylene contents

were investigated for their influence on the properties of

resulting TPVs

2. Experimental Section

2.1 Materials

Natural rubber (STR 5L) was supplied by Teck

Bee Hang, Thailand. PEC (Versify®) of various grades

(i.e., 2300, 3000, 3200, 3300, 3401 and 4301) was

supplied by Dow Chemical Company, USA and some of

mailto:[email protected]


201 their properties are shown in Table 1. Brominated

dimethylol phenolic resin (SP-1055) as crosslinking

agent was supplied by Schenectady International Inc.,

USA. Trimethyl-dihydroquinoline (TMQ) as anti-oxidant

was supplied by Flexsys Inc., USA.

Table 1. Some properties of virgin Versify®.

Versify®

grade

Melt flow

index*

[dg/min]

Ethylene

content

[%]

Hardness

[Shore A]

Denoted

Sample

codes

V-2300 2 12 88

M2E12**

V-3000 8 5 96 M8E5

V-3200 8 9 94 M8E9

V-3300 8 12 85 M8E12

V-3401 8 15 72 M 8E15

V-4301 25 12 84 M25E12

* Melt flow index (MFI) was tested at 230°C, 2.16 kg.

** M and E stand for melt flow index and ethylene content,

respectively.

Table 2. TPVs composition at NR/PEC 40:60 blend

ratio.

Components

Quantity

(% by weight)

NR (STR 5L) 40

Versify®* 60

SP-1055 5.2

TMQ 2

* At varied viscosity and ethylene content as detailed in Table 1.

2.2 Preparation of TPVs

STR5L and SP-1055 were first compounded for 6

min in an internal mixer (Chareon Tut Co. Ltd.,

Thailand) and then stored for 24 hrs prior to being used

to prepare TPVs. The compound formulations of the

NR/PEC TPVs are shown in Table 2. All TPVs were

prepared by a batch process in Brabender Plasticorder

(Brabender Plastograph EC plus, Brabender GmbH &

Co. KG, Germany) having a mixing chamber volume of

50 cm3. The mixer temperature was kept at 170°C under

a constant rotor speed of 60 rpm. PEC was first

mechanically melt-mixed for 3 min, then the antioxidant

(TMQ) was added and mixed for 1 min followed by the

addition of NR compound (STR5L+SP-1055). The

mixing was continued for another 4 min to complete the

dynamic vulcanization process. After that, the resulting

TPV was immediately removed from the mixing

chamber and passed through a cold two-roll mill in order

to convert it into a 1.5 mm thick sheet. The sheet was

then cut and compressed under a compression molding

machine having a cooling system (Scientific, Labtech

Engineering Co.,Ltd., Thailand) at 180°C for 5 mins,

before cooling down under pressing for another 5 mins.

Test specimens were later die-cut from the compression

molded sheet and kept at room temperature for 24 hrs

before the tests.

2.3 Testing procedures

2.3.1 Rheological study

The rheological properties of NR/PEC TPVs were

investigated using a capillary rheometer, G ttfert Rheo-

Tester 2000, (Werkstoff -Pr fmaschinen GmbH,

Germany). A capillary die had a diameter (D) of 1.0 mm,

a length (L) of 20 mm (i.e., an aspect ratio L/D=20/1),

and an entrance angle of 180°. The measurements were

carried out over a wide range of shear rates (i.e., 5–1000

s−1

) at 170°C.

2.3.2 Mechanical properties

Tensile tests were carried out according to

ASTMD 412 on dumb-bell shaped specimens using a

universal tensile testing machine (Model H 10 KS

Hounsfied Test Equipment Co, Ltd., UK) at a constant

cross-head speed of 500 mm/min. Tension set was tested

at room temperature after stretching the samples for 10

min at 100% elongation. Three specimens were tested for

each sample and average value was reported. Hardness

of the samples was measured with a durometer (Shore A)

according to ASTM D2240.

2.3.3 Recyclability study

The moulded TPVs samples were first cut into

small pieces using a Bosco plastic grinding machine. The

sample was then put into the internal mixer and the

mixing was performed for 5 mins at 150°C. After that,

the resulting material was reprocessed based on

compression molding at 170°C as the procedure


202 described in section 2.2. The specimens were then

subjected to property testing.

3. Results and Discussion 3.1 Properties of virgin PECs

Figure 1 shows the plot of apparent shear

viscosity as functions of apparent shear rate for all virgin

PECs. The materials exhibit shear-thinning behavior, i.e.,

apparent shear viscosity decreases with increasing

apparent shear rate. The PEC grade that has the lowest

melt flow index (MFI): M2E12 shows the highest shear

viscosity whereas those the same values of MFI: M8E5,

M8E9, M8E12 and M8E15 show almost the same level

of the shear viscosities, and M25E12 with the highest

MFI displays the lowest shear viscosity.

Figure 1. Apparent shear viscosity as functions of

apparent shear rate for virgin PECs.

Figure 2 displays the stress – strain curves of all

virgin PECs. The stress-strain behaviors of the PECs

with ethylene content of 5 and 9 wt% (i.e. M8E5 and

M8E9 respectively) shows yield point, which is a

character of plastic deformation. The produced material

is called Plastomer, i.e. a polymeric material that

combines qualities of elastomers and plastics, such

as rubber-like properties with the processing ability of

plastic. The PECs with 12-15 % ethylene contents (i.e.,

M2E12, M8E12, M25E12 and M8E15) do not show a

clear necking phenomenon while being stretched and

display elastic deformation. The materials become softer

with increasing ethylene contents as observed by its

lower stress and longer strain at break.

Figure 2. Stress-strain curve of PECs virgin copolymers.

3.2 Influence of viscosity on the TPV properties

Influence of PEC’s viscosity on the properties of

TPVs based on 40/60 NR/PEC blends is shown in

Figures 3-5.

Figure 3. Apparent shear viscosity as functions of

apparent shear rate for 40:60 NR/PEC blends with PECs

of various MFI values and ethylene content of 12 %.

Figure 3 shows the plot of apparent shear

viscosity as functions of apparent shear rate for 40/60

NR/PEC blends prepared by using PECs with various

MFI. The apparent shear viscosity decreases with

increasing apparent shear rate, and at a particular shear

http://en.wikipedia.org/wiki/Plastic


203 rate, the TPVs that was produced by the PEC with higher

MFI also show lower apparent shear viscosity. The flow

property of the TPVs corresponds to the viscosity of the

PEC which becomes a matrix after dynamic

vulcanization.

Figure 4, shows the stress-strain behaviors of

NR/PEC blends, which compose of the PECs; M2E12,

M8E12 and M25E12, which have MFI 2, 8 and 25

dg/min respectively. The stress-strain curves show

similar shape especially at low strain range. Initially, the

tensile stress sharply increases but upon further

deformation in the range of 50-300 % strain, the slope of

the curve decreases and the stress almost linearly

increases with strain. However, the slope of the curve

increases sharply again when the strain is above 400 %.

All the stress-strain curves of the TPVs show the

character of soft and tough elastic material.

Figure 4. Stress-strain curve of 40:60 NR/PEC blends

using PECs at various MFI with the same ethylene

content of 12 %.

Tensile properties, both tensile strength and

elongation at break, of the NR/PEC blends with PECs of

the same ethylene content but various MFI values are

shown in Figure 5. The elongation at break of NR/PEC

blends increases, but tensile strength slightly decreases

with increasing the MFI of PEC. With the same ethylene

content, the PEC with higher MFI or lower viscosity

should be caused by its lower molecular weight and so

can be deformed easier. In this case, the elongation at

break and tensile strength of the TPVs are not only

affected by the matrix properties but also by the

dispersed rubber phase. The higher molecular weight of

the matrix and/or finer dispersed rubber domains

generally lead to better mechanical properties of the

TPVs.

Figure 5. (a) Tensile strength; (b) Elongation at break of

NR/PEC blends with PECs of MFI values.

After passing through a recycle process, the

materials show a significant drop in the elongation at

break whereas the tensile strength of the TPVs does not

change considerably, as can be seen in Figure 5. The

TPVs with PEC of higher viscosity, i.e. lower MFI, show

a larger decrease of elongation at break, which indicates

a greater extent of degradation after recycling process.

(A) (b) (b)

(a)


204 3.3 Influence of ethylene content on the TPV

properties

The ethylene content of PECs has an influence

on the properties of TPVs as shown in Figures 6 - 8.

Figure 6 shows the plot of apparent shear viscosity

versus apparent shear rate of the 40/60 NR/PEC blends

with various ethylene contents in the PEC component. It

can be seen that apparent shear viscosity decreases with

increasing apparent shear rate. However, at a particular

shear rate, all the blends with different ethylene content

in the PEC component show similar shear viscosities as

those PECs have the same MFI values.

Figure 6. Apparent shear viscosity as function of

apparent shear rate for 40:60 NR/PEC blends with PECs

of various ethylene contents and MFI of 8 dg/min.

Figure 7 shows the stress-strain curves of

NR/PEC blends, containing the PECs of various ethylene

contents but the same MFI values. The result indicates

that the presence of ethylene in the PEC enhances the

molecular chain mobility and so the material can be

deformed easier. As seen in Table 1, the PECs with

different ethylene contents have different hardness, that

is, this polymer becomes softer with increasing ethylene

content. Additionally, the crystallinity of PECs also

decreases as the ethylene content in copolymers

increases. The values of crystallinity for PECs containing

5 (M8E5), 9 (M8E9), 12 (M8E12) and 15% ethylene

content (M8E15) are 44, 30, 17, and 14 wt%,

respectively, as reported by Dow Chemical Company.

Figure 7. Stress-strain curve of 40:60 NR/PEC blends

with PECs of various ethylene contents and MFI of 8

dg/min.

Choudhury and Bhowmick [6] reported that the

interphase interactions between the NR and PP phases

increased with the addition of ethylene propylene diene

monomer (EPDM) rubber into the NR/PP blend. This is

because EPDM has some structural similarity with PP

phase. Moreover, it is amorphous and elastomeric in

nature similar to NR. Thus, it is expected that PECs with

the lowest value of crystallinity, i.e. M8E15, is more

likely to be compatible with NR in a blend than the other

types of PECs. However, the results in Figure 8 show

that increasing the ethylene content of PEC in the

NR/PEC blends from 5 to 9, 12 and 15 %, the tensile

strength tends to decrease and elongation at break

increases respectively. These results indicate that the

tensile properties of the TPVs are strongly influenced by

the properties of the matrix.

It can also be seen from Figure 8a that the

NR/PEC blends with PECs of various ethylene contents

can be reprocessed without significantly affecting their

tensile strength. However, a significant drop in

elongation at break is observed for the TPVs with PEC

having ethylene contents higher than 5 % (i.e., M8E9,

M8E12, and M8E15) after recycling process (Figure 8b).

The decrease in the elongation at break of TPVs is

mainly attributed to thermal degradation of the PEC

matrix taking place during the recycling process since

the dispersed NR domains have already been

dynamically vulcanized. Hence, these results indicate


205 that the ethylene content in PEC is one of important

factors affecting the mechanical properties of NR/PEC

TPVs after recycling.

Figure 8. (a) Tensile strength; (b) Elongation at break of

NR/PEC blends with PECs of various ethylene contents

(%).

4. Conclusions

Thermoplastic vulcanizates based on NR/PEC

blends with 40/60 blend ratio have been developed. The

properties of these TPVs are influence by the melt flow

index, i.e. viscosity and ethylene content of the PECs.

With increasing the viscosity of PECs, tensile strength of

the TPVs decreases slightly but elongation at break

increases. The ethylene content in the PEC blend

component significantly affects the stress-strain

behaviors of the TPVs in which increasing ethylene

content decreases tensile stress but increases elongation

at break. The apparent shear viscosity at various shear

rates of the TPVs is in accordance with the melt flow

index of the PEC matrix. The materials can be recycled

but deterioration of the properties is clearly observed.

5. Acknowledgement This work was supported by the Higher Education

Research Promotion and National Research University

Project of Thailand, Office of the Higher Education

Commission.

6. References

[1] Basuli U., Chaki T.K., and Naskar K. “Influence of

Engage® copolymer type on the properties of

Engage®/silicone rubber-based thermoplastic

dynamic vulcanizates” Express Polymer Letter: 2,

846–854 (2008)

[2] Nakason, C. and Kaewsakul, W. “Influence of oil

contents in dynamically cured natural rubber and

polypropylene blends” Journal of Applied Polymer

Science: 115, 540–548(2010)

[3] Naskar K. “Dynamically vulcanized PP/EPD

thermoplastic elastomers” Ph.D. Thesis, University

of Twente, The Netherlands (2004)

[4] Robert P. L. and Robert A. K. “The mechanism of

phenolic resin vulcanization of unsaturated

elastomers” Rubber Chemistry and Technology: 62,

106 – 123(1988)

[5] Utara, S. and Boochathum, P. “Novel dynamic

vulcanization of polyethylene and ozonolysed

natural rubber blends: effect of curing system and

blending ratio” Journal of Applied Polymer Science:

120, 2606–2614(2011)

[6] Choudhury, N. R. and Bhowmick, A. K. “Adhesion

between individual components and mechanical

properties of natural rubber-polypropylene

thermoplastic elastomeric blends“Journal of

Adhesion Science and Technology: 2, 167-177

(1988).

(b)

(a)


206 RUBBERP-01

Study of Rheological Behaviour Used in Quality Control of Raw Natural Rubber (NR)

via Stress Relaxation

Apichet Ratchamontree1 , Chakrit Sirisinha

1, 2

1 Department of Chemistry, Faculty of Science, Mahidol University, Salaya, Nakornprathom 73170, Thailand

2 Rubber Technology Research Centre, Faculty of Science, Mahidol University, Salaya, Nakornprathom 73170,

Thailand

Phone 0-2441-9816-20 ext. 1142 Fax. 0-2354-7151, E-Mail: [email protected]

Abstract

Presently, standard Thai rubber (STR) is controlled by a variety of scientific values such as initial

plasticity (P0), plasticity retention index (PRI), Mooney viscosity, contents of dirt, ash, nitrogen, and volatile

matter. However, these scientific values are not sufficient for controlling the mixing behaviour or production

process stability in view of tyre industry. Therefore, the characterisation method in QC process capable of

predicting the mixing behaviour or production process stability is of interest in this work.

Relaxation test is one of the effective tools for separating viscous response from elastic response of

elastomer. In the present work, Mooney stress relaxation and RPA 2000 stress relaxation of STR 10 prepared

with different drying temperature (i.e., STR 10_H and STR 10_L) are focused. The former was prepared with

high temperature while the latter was dried at low temperature. The stress relaxation results were fitted with a

power law model. The STR 10_H demonstrates greater magnitude of viscous response than STR 10_L. In other

words, the stress relaxation test is considered as one of efficient methods for a quality control of raw NR

stability.

Keywords: Stress relaxation, Natural rubber, Mooney viscosity, Quality control

1. Introduction

Typically, rubber used in most rubber applications

is divided into 2 main categories, namely, natural rubber

(NR) and synthetic rubber (SR). Although SR possesses

various properties suitable for the applications of

interest, NR is still known to be a very important

material for making products required outstanding

properties including high elasticity and strength with

minimal heat build-up. The main products made from

NR include tyres and condoms[1].

As a rule of thumb for tyre manufacturing, the

mixing and processing behaviours of rubber compounds

are very crucial in order to achieve good processability

and product quality. There are numerous factors

governing the quality of mixing. One of important

factors is the consistency of raw rubber quality. This is

even more important for the tyres made from natural

rubber (NR), because the quality of NR depends strongly

on rubber clones, tree-age [2], planting areas, seasons

[3], and processing of latex to solid rubber [4].Molecular

weight (Mw) and its distribution are known to affect

properties of solid NR, including gel content and thus

green strength of NR [5]. In addition, one of the

important properties of NR is its increase in viscosity

over storage duration which is known as “storage

hardening”. This could be found in all grades of NR even

in viscosity-constant (CV) grade [6].These factors are

acknowledge to influence significantly the mixing

behaviour of raw NR incorporated with compounding

ingredients.

Therefore, it is of interest in this work to utilise a

stress relaxation behaviour of raw NR as a tool for a QC

process of raw NR in order to ensure the production

process stability.

2. Experimental

The STR 10_H and STR 10_L were kindly

provided by Michelin Co., Ltd., Thailand. Raw

NRwascharacterised for a relaxation behaviour using

optimised test conditions of 2 different rheometers: (i)

Mooney viscometer at 100 °C with rotor speed of 2 rpm,

relaxation duration of 20 mins, andpre-heating time of 1

min; and (ii) RPA 2000 at test temperature, strain

amplitude, and pre-heating time of 120 °C, 50 %, and 3

mins, respectively.

According to ASTM D1646 the stress relaxation

results were fittedwith power law equation expressed in

Eq. (1).


207 (1)

where:

= absolute value of torque at 1 second after rotor or die

stopped

= rate of relaxation

= relaxation time

The area under stress relaxation curve (SR Area)

was calculated from Eq. (2).

(2)

where:

= area under the stress relaxation curve from 1second

to the end of the test

= total time of the stress relaxation test in seconds

, = values obtained from Eq. (1)

Finally, the percentage of torque retention at 30 s

after stopping the rotor (%Mret 30) was calculated using

Eq. (3).

(3)


Rheological test of raw NR is known to be an

effective tool for predicting the mixing behaviour in tyre

industry. At present, standard Thai rubber (STR) is

graded by contents of dirtiness and especially Mooney

viscosity value. However, Mooney viscosity measured is

a combined responses of viscous component and elastic

component simultaneously. Referred to the samples

investigated in this work, both STR 10_L and STR 10_H

demonstrate no significant different in Mooney viscosity

with in test tolerance (i.e., 80.08 ± 0.95vs. 83.46 ± 1.33

MU). Thus, the stress relaxation test was selected for

further characterisation of viscoelastic properties.

Relaxation curves from Mooney viscometer (at

high test strain) and RPA 2000 (at low test strain) are

shown in Figs. (1) and (2), respectively. According to the

steady shear flow with constant shear rate, the shear

strain imposed on rubber bulk is a function of shear

deformation as illustrated in Eq. (4).With high test strain,

the relaxation behaviour of both samples was similar, but

the results with low test strain display some differences

in relaxation behaviour. The STR 10_H, prepared with

higher drying temperature, demonstrates greater

magnitude of viscous response. Molecular weight (Mw)

and gel network are believed to be responsible for such

difference. The STR 10_H was subjected to higher level

of thermal degradation during the drying process, leading

to the higher extent of molecular chain-end. Therefore,

STR 10_H might possess lower Mw and higher gel

network. On the contrary, the higher Mw and lower gel

network might be resulted in the STR 10_L. The

relaxation behaviour at high strain test gives the results

dominated by Mw because the gel network might be

completely broken. In contrast, the results measured at

low test strain are caused by a combination of Mw and

gel network.

(4)

where:

= shear strain

= shear strain rate (s-1

)

t = shear strain time (s)

Fig. (1). Mooney stress relaxation curve of raw NR

measured at high test strain


208

Fig. (2). RPA 2000 stress relaxation curves of raw NRs

measured at low test strain

Fig. (3). Rate of relaxation of raw NR measured with

different rheometers

Fig. (4). Relax (a/b) of raw NR measured with different

rheometers

Fig. (5). SRArea of raw NR measured with different

rheometers

Fig. (6). %Mret 30 of raw NR measured with different

rheometers

According to ASTM D1646, the 4 main

relaxation parameters are used for estimating the elastic

and viscous responses. The STR 10_H shows the higher

rate of relaxation, lower Relax (a/b), lower SRArea, and

lower %Mret 30 than STR 10_L. The results suggest

higher magnitude of viscous response [7]. The

discrepancies in molecular characteristics as affected by

drying temperature are expected to play role on mixing

behaviour, properties of compounds and vulcanisates.

4. Conclusions

Stress relaxation test in the present work was

measured at 2 strain magnitudes: at high and low strains,

Mooney stress relaxation test of 2 raw NR prepared with

different drying temperatures reveals no significant

difference in viscoelastic properties because of the

molecular slippage and disruptions of gel network at


209 high strain. In other words, the results of high strain test

came from main factor as Mw. However, with low test

strain, there are some differences in viscoelastic

properties because of the combined effect of Mw and gel

network. All relaxation parameters suggest the higher

viscous response in the STR 10_H.

Acknowledgement

The author thanks the Rubber Technology

Research Centre (RTEC) and Michelin Co., Ltd.,

Thailand for test equipment and financial supports.

References

[1] Chapman A. Natural rubber and NR-based

polymers: renewable materials with unique

properties. Transport. 2007;5:8.

[2] Kovuttikulrangsie S, Sakdapipanich JT. The

molecular weight (MW) and molecular weight

distribution (MWD) of NR from different age and

clone Hevea trees. Songklanakarin J Sci Technol.

2004;27:337-42.

[3] Moreno RMB, Ferreira M, Gonçalves PdS, Mattoso

LHC. Technological properties of latex and natural

rubber of Hevea brasiliensis clones. Scientia

Agricola. 2005;62(2):122-6.

[4] Dick JS, Harmon C, Vare A. Quality assurance of

natural rubber using the rubber process analyzer.

Polymer testing. 1999;18(5):327-62.

[5] Kawahara S, Isono Y, Sakdapipanich JT, Tanaka Y,

Aik-Hwee E. Effect of gel on the green strength of

natural rubber. Rubber chemistry and technology.

2002;75(4):739-46.

[6] Yunyongwattanakorn J, Sakdapipanich JT. Physical

property changes in commercial natural rubbers

during long term storage. Rubber chemistry and

technology. 2006;79(1):72-81.

[7] aláč J. Viscosity, Relaxation and Stability of

Natural Rubber. Open Macromolecules Journal.

2009;3:41-4.


210 RUBBERP-02

The Effect of Epoxide Functionality on Basic and Compound Properties of In-House and

Industrial Epoxidized Natural Rubbers

Patompong Pummor and Wisut Kaewsakul*

Department of Materials Science and Technology, Faculty of Science,

Prince of Songkla University, Hat Yai Campus, Songkhla 90110 Thailand

Abstract

The intrinsic properties of the raw rubbers are an important factor which determines the overall properties of the

compounds and end-use vulcanizate rubber products. The epoxidized natural rubber (ENR) is one of the most

commonly used modified natural rubber materials in industry. Many research works have been carrying out on ENR-

related subjects. However, a few of those has paid an attention on the difference between raw epoxidized natural

rubbers produced from laboratories and from industries. In the present work, the ENRs from the two production

sources were studied for their basic and compound properties. Based on the results, ENRs at the same level of

epoxidation degree from both sources show the similar basic properties in terms of chemical structure, thermal

dependence, specific gravity, and cure characteristics of compounds, except for their viscosity. The epoxide groups

play a significant role on basic characters and compound properties of the epoxidized natural rubbers when compared to

the unmodified natural rubbers.

Keywords: NR; ENR; compound; properties; epoxidize

1. Introduction

Natural rubber (NR) is a renewable material since it

can be produced from Hevea brasiliensis or Pará trees.

Rubber products in a variety of forms are made from

natural rubber such as tires, conveyer belts, building- and

bridge-pads, motor mounts, latex-based products: glove

and condom, and etc. According to the statistical

information reported by IRSG in 2013, the consumption

of natural rubber in the world market was shared with

synthetic rubbers (SRs), which are produced from

petroleum oils, in a proportion of NR/SRs at 42.4/57.6

[1].

The growing concerns regarding the increase of

synthetic rubber prices, the shortage of petroleum oils as

well as the environmental issues, there have been

attempted to modify natural rubber by chemically adding

some specific functionalities onto its molecules so that

the intrinsic properties of NR can be diversified. Up-to-

date, the most practical useful modified NR is epoxidized

natural rubber (ENR).It has been successfully produced

since 1920s and there were many intensive studies on

ENR-related subjects in the 1980s-1990s, led by the

research team from the Tun Abdul Razak Research

Centre (TARRC) [2,3]. There are two grades of ENR

available in the market nowadays: ENR-25 and ENR-50

which possess epoxide contents of 25 and 50 mol%,

respectively.

Number of research works were carried out using

ENRs as the material for the subject of interest, for

instance polymer blends [4-5], thermoplastic elastomers

[6], reinforcement of rubbers with either conventional

reinforcing fillers [7,8] or nanofillers [9-12], and so on

[13-14]. Based on the literature survey, the production

sources of the ENRs employed in the ENR-related works

were from both laboratories [15] and factories [4, 5].

Although, the two types of ENRs are not totally the same,

none of those has paid an attention on the difference

between these laboratory- and industry-prepared ENRs.

With our concerns, the outcomes derived from the

previous studies may not correspond to the practical

applications. In the other words, the reproducibility of the

related works for the practical use is hardly to be

achieved.

In the present study, our aim is to explore the

difference between in-house and commercial ENRs on

their basic properties, processing properties focusing on

the viscosity, and cure behaviors of compounds. Two

forms of natural rubber were used: 1) unmodified or


211 normal NRs, i.e. Ribbed Smoked Sheet 3: RSS3, and

Standard Thai Rubber 20: STR20; and 2) Epoxidized

NRs from laboratory and industry with the epoxide

contents of approximately 25 and 50 mol%.

2. Experimental

2.1 Materials

The rubbers used in this study were unmodified

natural rubbers: Ribbed Smoked Sheet 3 or

RSS3;andStandard Thai Rubber 20 or STR20 (Rubber

Estate Organization, Thailand), and commercial modified

natural rubbers: Epoxidized Natural Rubbers with the

epoxide contents of 25 and 50 mol% or ENR-25 and

ENR-50, respectively (Muang Mai Guthrie Public Co.

Ltd., Thailand). The in-house ENR-25 and ENR-50 were

prepared and used to compare with the commercial ones

as the preparation detail is given in paragraph 2.2. High

ammonia natural rubber latex: HA, with approximately

60 wt% dry rubber content: DRC (Chana Latex Co. Ltd.,

Thailand), hydrogen peroxide (Riedel-De Haen,

Germany), formic acid (FlukaChemie, Switzerland),

alkylphenolethoxylate or Teric N30, and commercial

grade methanol (J.T. Baker, USA), were used for ENRs

modification. The compounding ingredients formulated

in the rubber compounds were: zinc oxide (Global

Chemical, Thailand); stearic acid (Imperial Chemical,

Thailand); processing aromatic oil (H&R ChemPharm

(Thailand) Ltd., Thailand); polymerized 2,2,4-trimethyl-

1,2-dihydroquinoline or TMQ; N-(1,3-

diphenylguanidineor DPG; N-tert-butyl-2-benzothiazyl

sulfenamide or TBBS (all from Flexsys, Belgium); and

sulfur(Siam Chemicals Co. Ltd., Thailand).

2.2 In-house preparation of ENRs

The ENRs were prepared using HA latex with the

DRC of approximately 60 wt% via an in-situ performic

epoxidation reaction in which the performic acid was

generated by a reaction between formic acid and

hydrogen peroxide inside the reactor. The recipe used for

this synthesis is given in Table 1. The reaction was

carried out in a continuously stirred reactor, ata stirring

speed of 60 rpm. All reactants were first added into the

reactor at room temperature. The latex was diluted to

have a DRC of approximately 20 wt%, then stirred to

remove the preservative ammonia for 10 minutes, and

stabilized against coagulation by adding a non-ionic

surfactant, i.e. Teric N30, and held for 25 minutes. After

that, formic acid and hydrogen peroxide were

simultaneously added dropwise over a time period of 30

minutes. The reactor was later on placed in 50oC water

bath, continuously stirred, and held for 2:20 and 5:30

h:min. to obtain the ENR lattices with epoxidation degree

of approximately 25 and 50 mol%, respectively. After

the modification reaction finished, the ENR latex was

coagulated with methanol, thoroughly washed with water

and then dried in a vacuum oven at 40oC for 72 hours.

Table 1 Epoxidation recipe used for ENRs preparation.

Ingredients Quantity

gram mole

60 wt% HA Latex 192.67 2.05

Distilled water 385.34 -

10 wt% Teric N30 13.00 -

94 wt% Formic acid 44.04 1.03

50 wt% Hydrogen peroxide 176.80 4.10

The 1H-NMR spectroscopic technique was used to

analyze the molecular structure of the ENRs and their

epoxide contents by following the calculation equation:

Eq. 1. In addition, the epoxide content in commercial

ENRs was again confirmed with this technique. The

ENR samples were dissolved in deuterated chloroform

(CDCl3) prior to the performing of analysis. An example

1H-NMR spectrum of ENR-25 is depicted in Figure 1.

Epoxide content (mol%) = A2.7/ (A2.7+A5.14) x100(1)

where A2.7 is the integrated peak area at 2.7 ppm which

assigns to the attached proton in oxirane ring of ENR

molecules; and A5.14 is the integrated peak area at 5.14

ppm which assigns to the olefinic proton on NR

molecules.


212

Figure 1.An example of 1H-NMR spectrum of ENR-25

prepared in-house.

Table 2 Compound formulation used in this study.

Ingredients Amounts (phr)a)

Rubberb) 100.0

Process oil 7.5

ZnO 5.0

Stearic acid 1.0

TMQ 1.0

TBBS 1.0

DPG 0.5

Sulfur 2.0 a) The phr unit is parts per hundred rubbers.

b) Six types of rubbers used were two grades of unmodified NRs:

RSS3 and STR20; and four types of modified NRs: ENR-25 and

ENR-50 in commercial and laboratory-prepared forms denoted as

ENR-25C, ENR-50C, ENR-25L, and ENR-50L, respectively.

Table 3 Mixing procedure for the compounding preparation.

Mixing procedure Time (min)

Step I: Mixing in internal mixer

-Rubber 2

-ZnO, stearic acid, and

TMQ

3

-Discharged Kept until it was

cooled down to

room temperature.

Step II: Addition of curatives on

Two roll mill

-DPG and TBBS 3

-Sulfur 2

-Sheeted out Kept for at least 16

h prior to the step of

sample preparations

2.3 Compound preparation

The rubber compounds were prepared using the

formulation shown in Table 2. Mixing process was

carried out in an internal mixer (BrabenderPlasticoder

350s, Germany) with an initial mixer temperature setting

at 60oC, a rotor speed of 60 rpm, and 2 steps of mixing

procedure as detailed in Table 3. For the incorporation of

curatives, it was performed on two roll mill so that the

mixing temperature can be minimized with an

achievement of adequate chemical dispersion.

2.4 Determination of Mooney viscosity and cure

characteristics

Mooney viscosity of the raw rubbers and their

compounds were determined using a Mooney Viscometer

(MV2000vs, Alpha Technologies, USA) according to

ASTM D1646. The test was performed with a small

rotor for raw rubbers and a large rotor for the compounds,

the test temperature at 100oC, and the preheating and

measuring times of 1 min and 4 mins., respectively.

Vulcanization behavior of the compounds was

characterized using a Moving Die Rheometer

(MDR2000, Alpha Technologies, USA) with the cure

temperature of 150oC. Scorch time (Ts2), optimum cure

time (Tc90) and rheometer cure torque are reported.

2.5 Determination of thermal dependence and specific

gravity of raw rubbers

All raw rubber materials were analyzed for their

thermal properties by using a differential scanning

calorimetry (DSC) analyzer (PerkinElmer, USA). The

thermoanalytical condition was set at the temperature in

the range of -80 to 180oC upon the heating rate speed of

10oC/min Specific gravity of the raw rubbers was

measured using a digital specific gravity analytical device

(Mettler Toledo, Switzerland). The measurement was

performed at room temperature with the specimen weight

of about 10 grams.


3.1 Basic properties of raw rubbers

The intrinsic properties of raw rubbers were

investigated in order to gain the information regarding the

chemical, physical, and thermal characteristics prior to

the study of their compound properties. Epoxide content,

glass transitiontemperature, specific gravity, and viscosity

of the rubbers are shown in Table 4.


213

In Table 4, it can be observed that the ENRs

produced from laboratory and industry contain more or

less equal quantities of epoxide functionality in their

molecules. The glass transition temperatures of the

considering ENRs are somewhat not different. The higher

the epoxide content, the higher the Tg due to the reduction

of double bonds in NR molecules when it has been

modified with epoxide functionalities [2,3]. Molecular

chains of NR are highly flexible because the proton

adjacent to the double bonds in the molecules can freely

rotate/move. Hence, lower extent of double bonds, i.e.

more epoxidationdegree, in rubber main chains leads to a

decrease in elastic properties as the transition in the

region of glassy to rubbery states of the materials

remarkably shifts to higher temperature. In addition, the

ENRs from the two sources show similar specific gravity.

The ENR with higher epoxide content has a little higherin

specific gravity due to better molecular chain packing.

The self-associated intermolecular interactions of ENR

molecules are the main reason for the close packing of

rubber polymer chains [2,3]. Although, the comparable

values in terms of epoxide content, Tg, and S.G. of

the in-house ENRs compared to those of commercial

ENRs are achieved, the raw rubber viscosities as

indicated by MS(1+4)100oC of both types of ENRs

areobviously different. This result is important for the

processing properties of these rubbers. The laboratory-

prepared ENRs significantly show higher Mooney

viscosity than that of commercial ENRs. It is attributed

to the minimization of ENR viscosity from industry,

which has been done for an appropriate fabrication

purposes. One of the difficulties in rubber processing

comes from the undesired high rubber viscosity. So, the

viscosity properties of commercial ENRs is probably

controlled by incorporation of some small amount of

peptizers and/or plasticizers, as will be proved later.

Based on the results, the ENRs prepared from

laboratory are only difference in viscosity compared to

their counterparts received from industry, thus providing

good basic information for the further compound study.

3.2 Mooney viscosity of the compounds

With regards to the difference in the viscosity of

raw rubber materials, the Mooney viscosities of the

compounds based on those rubbers show the similar trend

with the starting rubbers used (see Table 4) as depicted in

Figure 2.

Figure 2.Mooney viscosities of ENR and NR compounds.

The compounds have much lower viscosity

compared to those of the raw rubbers (the large Mooney

rotor: ML, is used when the rubber specimen is soft or

has low viscosity, but when it is too hard the small rotor:

MS, is applied) due to mastication effect and lubrication

effect arose from the addition of process aromatic oil and

Table 4 Basic properties of raw ENRs and unmodified NRs used in this study.

Properties RSS3 STR20 ENR-25L ENR-50L ENR-25C ENR-50C

Epoxide content (mol%) 0.00 0.00 25.37 51.92 25.93 52.61

Tg (๐C)

a) -68.29 -67.61 -48.11 -26.43 -47.30 -25.96

Specific gravity, S.G. 0.88 0.91 0.91 0.94 0.92 0.96

MS(1+4)100oC 36 50 63 84 38 38

a) Glass transition temperature


214 a small amount of stearic acid which can give peptizing

influence towards rubber molecules.

3.3 Cure characteristics of the compounds

Vulcanization curves of the rubber compounds

show different cure behaviors as illustrated in Figure 3.

ENR-25s produced either from laboratory or industry

give similar cure characteristics to the unmodified NRs:

RSS3 and STR20. But, the ENRs with 50 mol% epoxide

content from both sources show different behaviors

compared with the others. The distinct differing between

the two groups of cure curves is that ENR-50 compounds

exhibit plateau characters (i.e. constant lines) after

reaching the optimum cure torque. On the other hand, the

unmodified NR and ENR-25 compounds clearly give

reversion aspects. This is attributed to ENR-50s possess

less extent of double bond sites on the molecules. The

double bonds or diene conjugates are sensitive to be

broken down under thermal and/or oxidative

environments. Apart from the breakdown of rubber

molecules, the sulfidic linkage stability can also be the

cause. Based on the results, it can be presumed that

unmodified NRs and ENR-25s provide higher quantity of

polysulfidic crosslink structure than the other crosslink

types: di- and mono-sulfidic bonds. The polysulfidic

linkage has lowest bond energy compared among the

sulfidic structures generated in accelerated sulfur

vulcanization system, and hence the poorest bond

stability upon the thermal annealing circumstances.

However, the evidence to support the extents of each type

of sulfidic linkages generated in sulfur-cured NR and

ENR systems has not been reported. So, it will be worth

to further study on this concern.

Figure 3. Rheological cure curves of NR and ENR Compounds

Table 5 shows the values of cure properties

derived from the cure curves (figure 4). Depending on

the crosslink density and molecular weight, the

rheological torque differences of the rubber compounds

are variable. Unmodified NRs give comparable torque

difference compared to ENR-50 compounds. In fact,

unmodified NRs contain the most numbers of double

bond sites on the molecules and so higher crosslink

density in the compounds. In addition, normal NRs

possess higher molecular weight than the modified ones

due to the chain scission during the epoxidation

modification [2,3] as a result in higher torque or moduli

of the compounds. However, it can be seen that ENR-50

compounds show more or less equal level of torque

difference compared to RSS3- and STR20-compounds. It

is because of the significant factor of intermolecular

interactions of epoxide groups. The polar epoxide

functionality can induce the strong polar-polar or dipole-

Table 5 The values of cure properties of ENR and NR compounds.

Properties RSS3 STR20 ENR-25L ENR-50L ENR-25C ENR-50C

Minimum torque, MH(dN.m) 0.58 0.79 1.22 1.03 0.64 0.64

Maximum torque, ML (dN.m) 10.07 9.38 8.68 10.68 8.12 10.94

Torque difference, MH-ML(dN.m) 9.49 8.59 7.46 9.65 7.48 10.3

Scorch time, Ts2 (min) 1.04 1.06 0.4 0.39 0.37 0.43

Optimum cure time, Tc90 (min) 2.39 2.07 1.26 6.4 1.21 6.31

Cure rate index, CRI 29.16 31.95 60.24 14.72 63.29 14.83


215 dipole interactions between the ENR molecules. The

higher epoxide contents give the stronger intermolecular

interaction forces, and so contributing to the torque of the

rubber matrix. From the previous study [2,3], ENR

molecules can generate the molecular networking via the

self-association of opened oxirane rings [2,3], leading to

contribution to matrix moduli/torque.

Considering the rate of vulcanization reaction, the

ENR-25 compounds show fastest vulcanization rate as

indicated by highest values of cure rate index. According

to the report of Gelling et al. (1988) [2], methyl group

attached to the oxirane ring is able to boost the

vulcanization reaction because of its electron

donatability. Nevertheless, the ENR-50 compounds show

the lowest curing speed, attributed to less extent of

double bonds in ENR-50, and thus reducing in the

reaction potential between curatives and rubber

molecules.

4. Conclusions

The chemical structure, thermal dependence and

specific gravity of the raw epoxidized natural rubbers

with the same level of epoxide contents obtained from

laboratory and from industry show more or less similar

characteristics, except for the viscosities. ENRs produced

in-house have a significant higher in viscosity compared

to industrial ENRs. Their compounds show the same

trend of viscosity with the raw ENRs, but in lower values.

Vulcanization behaviors of the compounds based on

ENR-25s produced from both sources exhibit reversion

aspects like the unmodified NRs based compounds,

however ENR-25s provide a little lower in torque

difference. The ENR-50s from both production sources

show different cure behaviors and properties compared to

the ENR-25s and unmodified NRs as the cure curves tend

to be plateau after reaching the maximum cure torque.

The torque difference of the compounds prepared with

ENR-50s is comparable with the values of unmodified

NR compounds. In addition, the speeds of vulcanization

reaction of both ENR-50 compounds are rather slower

than those of ENR-25s and normal NRs based

compounds.

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217 RUBBERP-03

Processability of TiO2-Filled-Natural Rubber Threads.

Wallapa Lorcharoenraung1*

and Ploenpit Boochathum2

1,2Department of Chemistry, Faculty of Science King ongkut’s University of Technolygy Thonburi,

126 Prach-utis Road, Bangmod, Toong-kru, Bangkok 10140, Thailand

Phone +6683 905 4086,*E-Mail:[email protected]

Abstract This work studied on preparation of rubber product by injection, dipping and molding process.

Chloroacetated natural rubber (CNR) was chemically modified from natural rubber latex using formic acid and

hydrogen peroxide in the ratio of 1:0.25:0.25 by mole at 50°C for 3 hr. After the reaction, chloroacetic acid was

added when the mixture was cool down to room temperature and stirred for 1 hr. Rubber molding method can

be applied by filled the natural rubber latex compounds into the rectangular mold and control the condition of

molding at room temperature. Then, Rubber thread vulcanizates for a diameter of 1 x 2 mm2 should be cured

at33 ºC, 48 hr and cured at 120 ºC, 5 min. However, %Elongation at break and tensile strength of CNRafter

vulcanized decreasing when compare to NR after vulcanized. The tensile strength of CNR was found least than

NR because the numbers of double bonds in CNR were least than those of NR.

Keywords: Chloroacetated natural rubber, Dipping,Injection, Molding, Natural rubber thread

1. Introduction

The rubber thread industry is one of important

rubber industries. Main raw rubber which was used in

this thread industry is “Natural Rubber Latex”.

Nowadays, the textile industry are expanding both

domestically and abroad. Thailand is a major exporter by

the year 1998 with 2 export value of 197.24 billion U.S.

dollars[1].Rubber thread which typically is composed of

natural rubber has excellent resiliency and other

desirable properties such as high elongation. However,

the disadvantage for using natural rubber in cloth

products is the growing of microorganisms along the

natural rubber threads which causes the skin allergy[2,3].

Rubber thread is commonly used in a number of

products, including narrow elasticized fabric for textile

applications such as waist bands and shoulder straps of

foundation garments, for toys, and for braided

("bungee") cord. The production process of rubber latex

thread comprises the following steps; preparing

materials, filtering, extruding, forming, cleaning,

vulcanizing, drying, cooling, placing in a box and

packaging. The superfine and high-elastic rubber latex

thread is one of products of the most with wide market

prospects [4,5].Rubber thread has been made by cutting

narrow strips from sheet rubber, yielding thread with a

square cross-section. Another method involves streaming

uncoagulated latex compound through a small-aperture

nozzle or capillary into a bath of coagulant, e.g., acetic

acid, washing the coagulated thread in a water bath, and

drying and heat-curing the final product[6].

However, most applications of NR are limitation

due to the low stability when increase temperature,

oxygen, sunlight, etc. and the high solubility inmost

hydrocarbon/hydrophobic solvents includingoils.

Therefore, to improve the stability of natural rubber to be

used widely, it must modify the structure of NR latex[7].

The chemical modification of NR latex by introduction

of functional group and lead to hydrophilicity along NR

backbone such as addition of functional groups including

epoxy, hydroxyl and ester groups on NR become to

Chloroacetated natural rubber (CNR) are alternative

strategies to provide more interaction between NR and

titanium dioxide[8,9].Recently, the effect of the addition

of stabilizer such as Triton-X on latex coagulation and

results were shown that stabilizer which added into latex

would protect the appearance of latex coagulation during

preparation[10].

This research studied the appropriate rubber

thread processing maid from functionalized natural

rubber latex. Then preparation process of rubber thread

with addition of inorganic material i.e., titanium by

injection, dipping and molding were studied for

comparison.


218 2. Experimental

2.1 Materials

Natural rubber latex 60.00%DRC was perchased

from theBond Chemicals Co. Ltd.Hydrogen peroxide

35%w/w, formic acid 85%w/w, Triton-xand chloroacetic

acidsupplied by Merck were used as recieved, 50%w/v

dispersed chemicals includingWingstay-L,sulfur,zinc

diethyldithiocarbamate(ZDEC), titanium dioxide (TiO2)

and zinc oxide (ZnO) were perchased from Rubber

Research Institute of Thailand.

2.2 Preparation of chloroacetated natural rubber

Chloroacetated natural rubber (CNR)was

prepared using by NR latex(60.00% DRC) and place into

the reaction vessel, followed by epoxidation using

hydrogen peroxide (H2O2) and formic acid (HCOOH )

which were dropped slowly while the mixture was

heated at temperature of 50°C. The mole ratio of NR:

H2O2: HCOOH=1:0.25:0.25, respectively. The mixture

solution was stirred at 50°C for 1.5 and 3hr will be

ENR1.5 and ENR3. Subsequently, chloroacetic acid

(ClCH2COOH) was added when the mixture was cool

down to room temperature and stirred for 1 hr. The

functionalized NR latex obtained were washed with

alkaline solution and a plenty of water till pH of the

washing water turned to pH of about 7 and rubber was

completely dried at room temperature will be CNR1.5

and CNR3[9]. The functional groups on modified natural

rubber were tested by FTIR. Wavelength scans from 500

to 4000 cm-1

using a 4 cm-1

resolution and 32 scans in

ATR mode. Confirmation the functionality of rubber by

Glass transition temperature (Tg) of FNR and unmodified

NR were measured using Differential Scanning

Calorimeter technique (Mettler-Toledo, DSC1) under

nitrogen gas flow. The approximately 5-10 mg sample

weight was encapsulated in hermetic pan. The

temperature scan was from -100°C to 25°C with the

heating rate of 10 °C/min.

2.3 Preparation of rubber latex compound.

Compounds were prepared according to the

described in Table 1. 50%w/v of sulfur dispersion was

first added to rubber latexes after that 50%w/v ZDEC

dispersion, 50%w/v Wing stay L dispersion, 50%w/v

TiO2 dispersion and 50%w/v ZnO dispersion was added

respectively.

Table 1. The compound formulation

Ingredients Part per

hundred (phr)

NR latex or CNR 60%DRC 100.0

50%w/v Sulfur dispersion 1.5

50%w/v ZDEC dispersion 1.0

50%w/v Wing stay L dispersion 0.5

50%w/v TiO2 dispersion 5.0

50%w/v ZnO dispersion 1.0

2.4Preparation processes of rubber threads

Three methods of preparation processes for

rubber threads production were injection, dipping and

molding process.

1) Hot water temperature affects the forming

rubber threads.

Heating water in a 100 mL beaker until the

desired temperature was 85°C and then put a syringe

containing compound dipped into water. The end of the

syringe in hot water. Then gently wash compound into

the water to save the results and repeated the experiment

with a water temperature of 80, 75, 70, 65, 60, 55, 50 and

room temperature (33°C), respectively, to study the

effect of temperature. In forming an rubber threads.

2) Concentration of Triton - X.

CNR latex prepared using a solution of Triton - X

concentrated 10.00% v/v of 180 mL and 13.33% v/v of

135 mL of water and then injected into the hot water

compound. The NR latex unattended solution Triton - X,

so did not do the experiments.

3) Concentration of acetic acid.

Prepared a solution of acetic acid at the

concentrations of 15% v/v, 20% v/v and 25% v/v, then

washed compound into a solution of acetic acid at the

different concentrations for preparing rubber thread.


219 2.4.2 Dipping process

Preparation of rubber fine thread was carried out

dipping by using process. The rubber compounds were

prepared using Triton-X solution 10.00%v/v 180

mLand13.33%v/v 135 mL. The clean test tube was used

as dipping mold. The test tube molds was dipped into the

coagulant before dipping into rubber latex compound for

2 time. The dipped mold was allowed to dry at room

temperatures for 24 hr and then vulcanization was done

under the temperatures for 120 ºC, 5 min. Washed the

dipped mold with a plenty of water then the fine film was

removed from the mold.

2.4.3 Molding process

Rubber molding method can be applied by filled

the natural rubber latex compounds into the rectangular

mold which was long: 112.00 cm, wide: 0.30 cm and

deep: 0.30 cm size and control the condition of molding

at room temperature. Then, rejected rubber vulcanizates

samples from the mold and washed samples in warm

water and separated all of natural rubber vulcanizates

samples. After the process above, sample will cure at

cured at 33 ºC, 48 hr and cured at 120 ºC, 5 min.

2.5Mechanical properties of rubber vulcanizate

Tensile strength and elongation at break of rubber

thread were measured at a crosshead speed of 500

mm/min, 500 N load with regard to Thai Industrial

Standard (TIS) 2556-2554 rubber thread or ISO 2321

Rubber threads-Methods of test.


3.1 Characterization of functional natural rubbers.

Figure 1 showed FTIR spectra of the selected

chloroacetated natural rubber (CNR3) molecule and

showed the unmodified natural rubber molecule (NR).

Peaks at 3391 cm-1

for OH stretching, at 1668 cm-1

for

C=C stretching and at 841 cm-1

for =C-H out of plane

bending for NR and the extra peaks for the epoxide ring

at 876 cm-1

and 1253 cm-1

, carbonyl group at 1752 cm-1

and C-O-C stretching at 1134 cm-1

. The molecular

structure of CNR prepared composed of epoxy groups on

the main chains and chloroacetate. Pending groups

randomly locate on the rubber molecules as shown in

Figure 2.

Figure 1 FTIR spectra of NR and CNR3.

Figure 2Proposed molecular structure of CNR.

The addition of the functionality of rubber

molecules was confirmed by glass transition temperature

values measured by using DSC. The results are shown in

Figure 3, The DSC thermogram indicated that glass

transition temperature of NR(-66.61°C) was much lower

than that of CNR3(-57.3°C).It was believed that the

rigidity of CNR molecules was due to the significant

interactions among the functional groups added on the

rubber molecules. From these results, we could

calculate %epoxy of each samples and found

that %epoxy of NR equal to 0%, %epoxy of CNR1.5

equal to 10.65%and %epoxy of CNR3 equal to 12.36%

respectively. It was believed that the rigidity of CNR

molecules was due to the significant interactions among

the functional groups added on the rubber molecules.


220

Figure 3DSC spectra of NR, CNR1.5 and CNR3.

3.2 Results of forming rubber thread.

3.2.1Injection process

The injection molded rubber thread was

conducted by injecting latex compound into hot water or

acetic acid solution for coagulation of rubber threads.

The effect of concentration of surfactant was

studied by using concentration of surfactant

from10.00%v/v180 mLand13.33%v/v 135 mL. The hot

water at 65ºC was used for thread coagulation.

Figure 4 was found that both CNR1.5 and CNR3

compounds prepared by using Triton-X solution

concentration of 10.00%v/v gave the brittle gel with

short thread because the compound contains water

bubbles up during the epoxidation reaction. While those

with the Triton-X solution concentration of 13.33%v/v

gave stronger and longer thread but it was not stable.

However, these properties were not sufficient. The

Triton-X solution, which is a surfactant types Octyl-

Phenol Ethoxylate, served as Emulsifier enhance

stability, that stabilizer which added into latex would

protect the appearance of latex coagulation during

preparation

Therefore the coagulation condition was changed

to acetic acid solution with variable solution of 15 %v/v,

20 %v/v and 25 %v/v at room temperature of 33°C

which was shown in Figure 5. It was found that the

concentrations of acetic acid did not affect the properties

of all threads. It was remarkable that the CNR1.5thread

compound and CNR3thread compound showed that the

latex compounds were injected not strong and dissolved

in acetic acid quickly because the latex compounds were

not included in acetic acid. Hence the rubber gel was

injected into the formation of the gel was not strong

enough to lead vulcanization.

Figure 4 Rubber thread characteristics process by

injection process which was coagulated in hot water at

65ºC by using Triton-X concentration of a) 10.00%v/v

and b) 13.33%v/v.


injection process which was coagulated different

concentration of acetic acid include 15 %v/v, 20 %v/v

and 25 %v/v by a) CNR1.5 and b)CNR3. (Use room

temperature 33° C and Tri ton – X13.33%v/v)


221 The characteristics of different types of rubber

thread were produced from injection shown in Figure 6.

It was found that compounds from conventional natural

rubber latex (NR) cannot form into gel formation at all

temperatures study because NR no has heat-sensitive

agent. The latex compounds of CNR1.5 and CNR3 gave

longer rubber line at 65 ± 5ºC and 32 ± 2ºC while at the

higher temperature the shorter gel rubber thread because

increasing the water temperature was able to the rubber

gel was broken easily. The stability of rubber line of

compound CNR3 was better than rubber line of

compound CNR1.5 because CNR3 has more epoxy

groups.


injection process which was coagulated in hot water at

different temperature at a) 80 ± 5ºC, b) 65 ± 5ºC, c) 50 ±

5ºC and d)32 ± 2ºC.

3.2.2 Dipping process.

At the Triton-X concentration of 10.00%v/v,the

rubber film from CNR1.5latex compound showed a

plentyof air bubble on the film. In addition the stabilizer

concentrtion of 10.00% v/v was found to be not enough

tostabilize CNR3 latex compound.For the Triton-X

solution of13.33% v/v, the rubber film was smooth and

without the presence of bubles. It was obvious that the

strenght of the vulcanized film still was in sufficient.


dipping process which was coagulated on room

temperature 33°C by using Triton-X concentration of a)

10.00%v/v and b) 13.33%v/v.

3.2.3Molding process

The threads produced from the molding process

were in the dimension of the cross-section area of 2

mm2and 1120 mm

2in length, which was shown in figure

8.

The elongation at break of threads in figure 9,

produced from CNR latex was found to be slightly less

than that produced from the conventional NR latex. But

for tensile strength NR thread was higher than those of

CNR because of more numerous of the double bond in

molecular chains of NR These results confirmed more

numerous of cross-linking of NR prevail over CNR.

Figure 8 Wash chemicals that do not react out of rubber

threads and rubber threads forming


222

Figure 9 %Elongation at break and Tensile strength of

NR, CNR1.5 and CNR3.

4. Conclusion

From studying of the preparation of rubber

threads by three methods such as injection molding,

dipping and casting. It was found that the preparation of

rubber threads by injection molding and dipping could

not be applied for molding of rubber threads but the

preparation of rubber threads by casting the latex on the

mold could be applied for molding of rubber threads.

From the experimental, it can be seen that using of the

mold which had cross-section area 3 mm2 to produce

rubber threads had cross-section area 2 mm2.

Degree of cross-linking of NR was found to be

more than CNR and leaded to better tensile strength and

elongation at break than those of CNR.

References

[1] Tantiwiboonchai,N.,“A Study on the

competitiveness of Rubber Products Produced

fromConcentrated Latex of Thailand in the World

arket”. Thammasat University,211-226(2010).

[2] Nimsuwan, C., “Rubber thread”,Research and

Development Centre for Thai Rubber Industry,1-5

(2011)

[3] PatanakunW., Opanukul, W., Na Ranong, N. and

Wichitchonlachai,N.,“Extrudedlatex thread

production by pilot model machine”,Research report

of rubber, 640-657 (2010)

[4] Bosshard,A., “Rubber thread-cutting apparatus”,U.S.

Patent No. 2,567,634 (1951)

[5] Wilhelm,F.J., O'Neill,K.J., John,IIF., Maglio, R.

and Cabral, E., “Process of making rubber thread”,

U.S. Patent No. 5,679,196 (1997)

[6] ax, D., “Rubber thread and method of making

same”,U.S. Patent No. 2,149,425 (1939)

[7] Yoksan, R., “Epoxidized Natural Rubber for

Adhesive Applications”,Kasetsart Journal (Nat.

Sci.), Vol. 42 : 325-332 (2008).

[8] Na Wichian,A.,Prakaimaneewong, and P. Na

Ranong, N.,“ Preparation of Epoxidized Natural

Rubber from Field Latex”,Research report of

rubber, 1-14 (2009)

[9] Heping, Y., Sidong, L. and Zheng, P., “Preparation

and Study of Epoxidized Natural Rubber”,Journal of

Thermal Analysis and Calorimetry, Vol. 58 : 293-

299 (1999).

[10] Boochathum, P. and Rongtongaram, N.,

“Characterization of Processability and Silica-silica

Network of Silica Filled Functionalized Natural

Rubber Composite”,28th

The Polymer Processing

Society (PPS) Conference, December 11-15,

Pattaya (Thailand) (2012).


223 RUBBERP-04

Preparation and Properties of Epoxidized Natural Rubber/Carbon Nanotubes

Nanocomposites

Piyaphorn Mungmeechai *, Saowaroj Chuayjuljit

and Anyaporn Boonmahitthisud

Department of Materials Science, Faculty of Science, Chulalongkorn University, Bangkok 10330

Phone +6689239 6601, Fax +66 22185561, *E-Mail: [email protected]

Abstract

This research aimed to investigate the effects of carbon nanotubes (CNTs) on the tensile properties and

thermal stability of epoxidized natural rubber (ENR)/CNTs nanocomposites. ENRs with different levels of

epoxidation were prepared from concentrated natural rubber (NR) latex via in situ performic acid epoxidation

method using various molar ratios of formic acid and hydrogen peroxide (H2O2) at 50°C for 4 h. An ENR with

30–35 mol % epoxidation was prepared by using a molar ratio of 0.75:0.75 formic acid:H2O2, and this formic

acid:H2O2 molar ratio was further used in the preparation of ENR/CNTs nanocomposites with an inclusion of

five loadings of CNTs (0.5–2.5 parts per hundred parts of rubber) in the in situ epoxidation reaction of NR. The

tensile properties (tensile strength and modulus at 300% strain) and thermal stability of the prepared

nanocomposites were found to be improved with the inclusion of an appropriate loading of the CNTs as compared

to those of pure NR, but with a decrease in the elongation at break in a dose-dependent manner. However, the

elongation at break of the nanocomposites was largely retained, giving the values of between 591.2–691.6% as

compared with that of the pure NR (727%), since these CNTs reinforced ENR vulcanizates are soft

nanocomposites.

Keywords: Epoxidized natural rubber, Carbon nanotubes, In situ epoxidation, Nanocomposites

1. Introduction

Natural rubber (NR), one of the important renewable

natural polymers, has numerous advantage properties

such as high mechanical strength, low heat buildup and

excellent resilience and elasticity [1–3]. It is known that

NR has high tensile and tear strength due to its ability to

undergo strain crystallization [2]. However, NR has

some drawbacks such as poor ozone, thermal, weathering

and oil resistance owing to its unsaturation hydrocarbon

chain structure and non-polar nature, which limit its use

in many applications [3,4]. One of the straightforward

and convenient methods to solve these disadvantages is

the introduction of polar groups onto the NR backbones.

Epoxidation of NR is a simple reaction that can be

occurred in latex stage via in situ peracid epoxidation,

and is a well-known and effective method of improving

the oil and thermal resistance of NR [4–6]. As NR is

epoxidized, double bonds are randomly changed to

epoxide groups, leading to an increased polarity and

glass transition temperature with increasing level of

epoxidation, while the ability to strain crystallize can be

retained up to about 50% epoxidation [7,8]. The extent of

epoxidation is controlled by the peracid content, reaction

temperature and reaction time. The ENR of lower than

50 mol % epoxidation is a typical elastomer, while that

of higher epoxidation becomes harder and lower

resilience and elasticity [9]. The ENR can be vulcanized

by sulfur-cured systems similar to NR, and its

vulcanizate can retain some advantages of NR, including

high tensile and tear strength [2,8]. However, reinforcing

filler is still needed for rubber to gain appropriate

properties for specific applications, especially, nano-

sized reinforcing particles [10-12]. The present work

focused on the use of carbon nanotubes (CNTs) at a very

low loading (0.5–2.5 parts per hundred parts of rubber,

phr) to prepare ENR/CNT nanocomposites. CNTs are

valuable for nanotechnology due to their superior

properties such as very high aspect ratio (up to 104),

specific surface area, modulus (about 103 GPa) and tensile

strength (~ 50 GPa) and low density (~ 1.3 g/cm3)

[11,13]. In this study, ENR/CNT nanocomposites were

prepared via in situ performic acid epoxidation of NR

using formic acid and hydrogen peroxide (H2O2) in the

presence of CNTs. The curing characteristics and

properties of the vulcanizates in terms of the tensile

properties and thermal stability were examined.


224 2. Materials and Experimental methods

2.1 Materials

NR latex having 60% dry rubber content (DRC),

polyalcohol ethyleneoxide condensate (Terric 16A-16),

zinc oxide (ZnO), stearic acid, sulfur, n-cyclohexyl-2-

benzothaiazyl sulphenamide (CBS), Voltamol and bentonite

were obtained from Rubber Research Institute of Thailand

(Bangkok, Thailand). Formic acid (98%) and H2O2 (30%)

were purchased from Asian Scientific Co. (Bangkok,

Thailand). CNTs was supplied by EM-Power Co.

(Bangkok, Thailand). Sodium carbonate was purchased

from Merck Thailand Co. (Samutprakarn, Thailand).

Methanol was purchased from RCI Lab Scan Company

(Bangkok, Thailand). All materials were used as

supplied.

2.2 Preparation and characterization of ENR

ENR was prepared via in situ epoxidation of NR

using formic acid and H2O2. The NR latex was first diluted

to 20 % DRC with distilled water and stabilized with Terric

16A-16 under stirring for 1 h. The assigned amounts of

formic acid (0.25–1 M of isoprene unit) and H2O2 (0.25–

1 M of isoprene unit) were added into the stabilized NR

latex according to the formulation in Table 1. Formic

acid was fed drop by drop into NR latex within 10–15

min at 40°C, followed by raising the temperature to 50°C

within 15 min prior to adding the H2O2 within 5–10 min

with continuous stirring. The reaction was allowed to

proceed at 50°C for 4 h. The obtained product was

coagulated in methanol, filtered, washed with distilled

water, soaked in sodium carbonate solution (5%) for 5–10

min, washed again until neutral, pressed into thin sheet

(1.5–2 mm), and then dried at 60°C for 48 h.

Table 1. Formulation for epoxidation reaction of NR.

Formula

Ingredient 1 2 3 4 5 6 7

NR 100a 100 100 100 100 100 100

Terric

16A-16 3a 3 3 3 3 3 3

Formic

acid 0.75b 0.75 0.75 0.75 0.25 0.50 1

H2O2 0.25b 0.50 0.75 1 0.75 0.75 0.75

a dry content in g,

b content in M of isoprene unit

The extent of epoxidation was analyzed by Fourier

transform infrared spectroscopy (FT-IR) using a Nicolet

6700-FT-IR over a frequency range of 400–3400 cm-1

.

The mol % epoxidation was calculated from Eq. (1) using

data obtained from the characteristic FT-IR peaks at 870

and 835 cm-1

in according to Davey and Loadman [14].

Mol % epoxidation = [A870/(A870 + A835)] × 100 (1)

where A870 and A835 are the absorbancies at wavenumbers

870 and 835 cm-1

, respectively.

2.3 Preparation and characterization of ENR/CNTs

CNTs were first prepared in an aqueous dispersion

at 2 wt% using Voltamol (2 wt%) and bentonite (1 wt%)

as a dispersant of CNTs in distilled water by ball milling

at 200 rpm for 72 h. The ENR/CNTs nanocomposites

were then prepared via in situ epoxidation of NR with

various loadings of CNTs (0.5, 1, 1.5, 2 and 2.5 phr). The

CNTs was first added to the stabilized NR latex and the

formic acid and H2O2 were then added as previously

mentioned. The mol % epoxidation of ENR was also

examined by the FT-IR analysis and Eq. (1).

2.4 Compounding, cure characterization and

vulcanization

The NR, ENR and ENR/CNTs nanocomposites were

mixed with ZnO (5 phr), stearic acid (2 phr), CBS (1 phr)

and sulfur (2 phr) in an internal mixer (MX500-D75L90)

at 70°C for 9 min, followed by a two-roll mill for 1 min.

Curing behaviors in terms of scorch time (ts2) and

cure time (t90) were determined at 130°C according to

ASTM D2084 using a moving die rheometer (MDR)

(A0225-rheo Tech MD). At least three specimens were

determined for each composition, and the average value

was reported.

The compounded rubber was cured in a compression

molding machine (LP-S-20, LabTech Engineering) at

130°C using the t90 obtained from the MDR.

2.5 Tensile properties

The tensile test was performed on a dumbbell-shape

specimen in according to ASTM 412 (type D) using a


225 universal testing machine (T-TS01) with a load cell

capacity of 1 kN at a cross-head speed of 500 mm/min.

The values of tensile strength, modulus at 300% strain

and elongation at break were averaged and reported from

at least five specimens for each composition.

2.6 Thermogravimetric analysis (TGA)

The thermogravimetric analysis was performed on a

Mettler Toledo TGA/SDTA 851e analyzer to examine the

thermal stability in terms of the temperatures for onset

(Tonset), end set (Tend set), 50% weight loss (50%) and

maximum decomposition rate (Tmax). About 10 mg of

sample were scanned over a temperature range of 30–

1000°C at a heating rate of 20°C/min under a nitrogen

atmosphere.


3.1 FT-IR analysis and epoxide content

Representative FT-IR spectrum of NR shown in Fig.

1(a) exhibited characteristic peaks at 2860, 1650, 1450,

1375 and 835 cm-1, which are assigned to the -CH stretching,

-C=C- stretching, -CH2- deformation, C-H deformation

of CH2 and =C-H deformation, respectively. A selected

ENR showed the new characteristic peaks at 1240 with a

small peak at 870 cm-1

(Fig. 1(b)), which are assigned to the

C-O-C ring vibration of epoxide groups, which were not

found in the NR. This confirmed the formation of

epoxide rings from the reaction of performic acid with

C=C bonds on the NR backbones. By varying formic

acid:H2O2 ratios, the obtained ENRs possessed different

mol % epoxidation as calculated from the Eq. (1), using

data obtained from the FT-IR spectra (not shown here).

The calculated mol % epoxidation of ENRs is shown

in Table 2. The results showed that the mol % epoxidation

increased with increasing either formic acid or H2O2

content as a consequence of the increased performic acid

content, and was in the range of 17–46 mol %.

Figure 1. FT-IR spectra of (a) NR, (b) ENR and (c)

ENR/CNTs nanocomposite.

Table 2. Mol % epoxidation of ENR.

Formula

Character 1 2 3 4 5 6 7

Mol %

epoxidation 17 28.5 31.3 42 18.4 33.3 46

Since ENR with low mol % epoxidation has higher

strain crystallization, resilience and elasticity, but lower

oil resistance than ENR with higher mol % epoxidation

[7-9]. With the balance of these properties, the ENR with

medium mol % epoxidation (~ 31.3 mol %), prepared

from 0.75 M H2O2 and 0.75 M HCOOH (Formula 2) was

used for further blending with different loadings of CNTs

(0.5, 1, 1.5, 2 and 2.5 phr) in the latex stage.

Accordingly, this ENR was denoted as ENR-30.

Fig. 1(c) shows the FT-IR spectrum of ENR/CNTs

nanocomposite. The characteristic peaks at 1240 and 870

cm-1

also confirmed the formation of epoxide rings via in

situ epoxidation in the presence of CNTs. It was found

that the mol % epoxidation of ENR in the nanocomposites

was slightly higher than that in the neat ENR-30 and was

in the range of 33.6–34.2 mol %. This may be due to the

acidity on the surface of CNTs. Since commercial CNTs

usually possess carboxylic acid groups on their surfaces

[11].

3.2 Cure characteristic

Scorch time (ts2) and cure time (t90) obtained from

the rheographs (not shown here) of NR, ENR 30 and

ENR/CNTs nanocomposites are listed in Table 3. As can

be seen, ENR 30 had shorter ts2 and t90 than NR, which

due to the epoxide groups that activated the adjacent


226 double bonds, and thus increased the cure rate of ENR 30

[15]. However, the ts2 and t90 of the nanocomposites

increased with increasing loading of CNTs. This is

because -COOH groups on the CNTs surface absorbed the

basic accelerator species, thereby delaying the ts2 and t90

of ENR 30 [11].

Table 3. The ts2 and t90 of NR, ENR and ENR/CNTs

nanocomposites.

Sample tt2

(min)

t90

(min)

NR 8.5 17.1

ENR 30 8.1 16.6

ENR/CNTs

100/0.5 8.7 17.2

100/1.0 9.7 18.4

100/1.5 10.2 19.8

100/2.0 10.5 19.4

100/2.5 9.0 15.0


The tensile properties in terms of the tensile strength,

elongation at break and modulus at 300% strain of NR,

ENR 30 and ENR/CNTs nanocomposites are

summarized

in Table 4. As can be seen, the tensile strength of NR

was slightly higher than that of ENR 30. This was due to

the nature of NR that has higher level of strain

crystallization than ENR 30 [2,7]. However, the tensile

strength of ENR/CNTs nanocomposites at 0.5 and 1 phr

CNTs was not improved compared to that of the NR and

ENR 30, which may be due to an insufficient level of

dispersion of CNTs in the ENR matrix. As the CNTs

loading was 2 phr, the tensile strength increased up to a

maximum value, suggesting the well dispersion of CNTs

in the ENR matrix and also the better stress transfer at the

interfacial of CNTs and ENR, due to the very high

specific area and aspect ratio of the CNTs. At higher

CNTs loading (2.5 phr), a remarkable decrease in the

tensile strength was observed. This may be due to the

agglomeration of CNTs particles that restricted the strain

crystallization of ENR and allowed less stress transfer

across each phase. The elongation at break of ENR 30

and ENR/CNTs nanocomposites was found to be lower

than that of the NR. This is because the polar epoxide

groups tightly held the ENR molecules and then restricted

the mobility of rubber chains, resulting in a lower

elongation. However, the elongation at break of NR

(727%) was largely retained in ENR 30 and ENR/CNTs

nanocomposites due to the low mol % epoxidation and

the very low loading of CNTs, giving an elongation at

break range of between 591.2–695.3%. Moreover, the

modulus at 300% strain of ENR 30 and ENR/CNTs

nanocomposites slightly increased compared to that of

the NR, indicating that the CNTs reinforced ENR

vulcanizates were soft nanocomposites.

Table 4. Tensile properties of NR, ENR and ENR/CNTs

nanocomposites.

Sample Tensile

strength

(MPa)

Elongation

at break

(%)

Modulus

at 300%

strain

(MPa)

NR 24.0 727.0 1.8

ENR 30 23.3 673.8 2.1

ENR/CNTs

100/0.5 21.2 662.5 2.1

100/1.0 22.9 683.7 2.1

100/1.5 24.1 691.6 2.1

100/2.0 26.5 695.3 2.2

100/2.5 15.2 591.2 2.4

3.4 TGA analysis

TGA was performed to evaluate the thermal stability

of NR, ENR 30 and ENR/CNTs nanocomposites. The

TGA curves of samples are shown in Fig. 2, while the

values of Tonset, T50%, Tend set and Tmax are listed in Table 5.

As can be seen, TGA curves of ENR 30 and ENR/CNTs

nanocomposites exhibited a similar characteristic. From

Table 5, it was noticed that the thermal stability of NR

was improved by introducing epoxide groups onto its

backbones, since the Tonset, T50%, Tend set and Tmax of ENR

30 were all shifted to higher temperatures compared to

those of the NR. This is because the epoxide groups

caused an increase in the intermolecular attraction and


227 reduced the chain mobility, and so the thermal stability

was improved. The results also indicated that the

presence of CNTs at low weight fractions had no

significant effect on the thermal stability of ENR/CNTs

nanocomposites.

Figure 2. TGA thermograms of ENR 30 and ENR/CNTs

nanocomposite.

Table 5. TGA-derived data for the samples.

Sample Tonset

(°C)

Tend set

(°C)

T50%

(°C)

Tmax

(°C)

NR 363.7 422.6 394.7 389.6

ENR 30 375.3 427.9 405.7 399.2

ENR/CNTs

100/0.5 375.3 428.0 406.0 399.8

100/1.0 375.4 428.8 406.3 398.4

100/1.5 375.7 429.0 407.3 398.5

100/2.0 376.5 429.2 408.0 399.5

100/2.5 374.1 425.4 403.0 395.7

4. Conclusion

The preparation of ENR/CNTs nanocomposites was

simultaneously performed with the in situ epoxidation of

NR using formic acid and H2O2 in order to disperse CNTs

homogeneously in ENR latex prior to compounding with

vulcanizing ingredients. The scorch and cure time of the

ENR/CNTs nanocomposites were found to be shorter

than those of the NR, and mainly depended upon the

mol % epoxidation of ENR. Among the prepared

nanocomposites, optimum tensile strength and elongation

at break was achieved at 2 phr CNTs loading, while the

modulus at 300% strain was nearly the same. Moreover,

the addition of CNTs does not improve the thermal

stability of the ENR 30 to a great extent.

References

[1] Peng, Z., Feng, C., Luo, Y. Li, Y. and Kong, L.X. Self-

assembled natural rubber/multi-walled carbon

nanotube composites using latex compounding

techniques, Carbon, 48(15), 2010, 4497-4503.

[2] Sae-oui, P., Sirisinha, C. and Hatthapanit, K. Effect

of blend ratio on aging, oil and ozone resistance of

silica-filled chloroprene rubber/natural rubber

(CR/NR) blends, Express Polymer Letters, 1(1),

2007, 8-14.

[3] Kalkornsurapranee, E., Sahakaro, K., Kaesaman, A.

and Nakason, C. Influence of reaction volume on the

properties of natural rubber-g-methyl methacrylate,

Journal of Elastomers and Plastics, 42(1), 2010, 17-

34.

[4] Saramolee, P., Lolopattananon, N. and Sahakaro, K.

Preparation and some properties of modified natural

rubber bearing graft poly(methyl methacrylate) and

epoxide groups, European Polymer Journal, 56(1),

2014, 1-10.

[5] Zhao, Y., Huang, B., Yao, W., Cong, H., Shao, H.

and Du, A. Epoxidation of high trans-1,4-

polyisoprene and its properties, Journal of Applied

Polymer Science, 107(5), 2008, 2986-2993.

[6] Gelling, I.R. Epoxidised natural rubber, Journal of

Natural Rubber Research, 6(3), 1991, 184-205.

[7] Ismail, H. and Chia, H.H. The effects of

multifunctional additive and vulcanization systems

on silica filled epoxidized natural rubber

compounds. European Polymer Journal, 34(12),

1998, 1857-1863.

[8] Akinlabi, A.K., Okieimen, F.E. and Aignodion, A.I.

Thermal aging properties and chemical resistance of

blends of natural rubber and epoxidized low

molecular weight natural rubber, Journal of Applied

Polymer Science, 98(4), 2005, 1733-1739.

[9] Bac, N.V., Terlemezyan, L. and Mihailov, M. On the

stability and in situ epoxidation of natural rubber in

latex by performic acid, Journal of Applied Polymer


228 Science, 42(11), 1991, 2965-2973.

[10] Ranimol, S., Rosamma, A., Treesa, C., Siby, V.,

Kuruvilla, J. and Sabu, T. Rheological behavior of

nanocomposites of natural rubber and carboxylated

styrene butadiene rubber lattices and their blends,

Journal of Applied Polymer Science, 101(4), 2006,

2355-2362.

[11] Shanmugharaj, A.M. and Ryu, S.H. Influence of

aminsilane-functionalized carbon nanotubes on the

rheometric, mechanical, electrical and thermal

degradation properties of epoxidized natural rubber

nanocomposites, Polymer International, 62(10),

2013, 1433-1441.

[12] Teh, P.L., Ishak, Z.A.M., Hashim, A.S., Kocsis, J.K.

and Ishiaku, U.S. On the potential of organoclay with

respect to conventional fillers (carbon black, silica) for

epoxidized natural rubber compatibilized natural

rubber vulcanizates, Journal of Applied Polymer

Science, 94(6), 2004, 2438-2445.

[13] Zhou, X., Zhu, Y., Liang, J. and Yu, S. New fabrication

of styrene-butadiene rubber/carbon nanotubes

nanocomposite and corresponding mechanical

properties, Journal of Materials Science and

Technology, 26(12), 2010, 1127-1132.

[14] Davey, J.E. and Loadman, M.J.R. A chemical

demonstration of the randomness of epoxidation of

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and activation energy of ENR 25 compared with

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Polymer Journal, 35(4), 1999, 711-719.


Smart and Intelligent Polymers


230

KN-SMART-1

Self-Assembly of Ultralong Polyion Nanoladders Facilitated by Ionic Recognition and

Molecular Stiffness

Yun Yan*

College of Molecular Engineering, Peking University, Beijing, 100871

[email protected]

Abstract

Ionic interaction has emerged as an important driving force for the

fabrication of ionic self-assembly. Well-known examples are the

polyion assemblies formed in water by pairs of oppositely charged

polyelectrolytes, which can be forced into films and capsules with

appropriate protocols, or made into micelles and vesicles by end-

attaching uncharged water-soluble blocks to (at least one of) the

polyelectrolytes. This is also the case when one of the polyelectrolytes

was replaced by reversible or “soft” coordination polymers, namely,

supramolecular chains in which organic bisligands and transition metal

ions alternate in a regular fashion. By proper choice of the bisligands,

one can even obtain supramolecular polyelectrolytes, which in turn can

form polyion assemblies just like ordinary polyelectrolytes. Usually,

polyion assemblies formed with covalent polylectrolytes or

coordination polymers have internally random structures, since the

polymers are flexible and the ionic interactions are isotropic between

the flexible chains. In this work we show that precise alignment of

polyelectrolyte chains inside polyion assemblies can be achieved by

imparting proper stiffness to the molecular tiles.

Yun Yan

Dr Yun Yan has been an associate

professor at Peking University, China,

since 2008.She earned her bachelor

degree at Northeast Normal University

(1997), China, and obtained the PhD

in physical chemistry at Peking

University. After two postdoctoral

studies in Bayreuth University

(Germany) and Wageningen

University (the Netherlands), she

joined Peking University as an

associate professor. Her current

interest is molecular self-assembly in

solutions, including 1) surfactant in

solution; 2) Self-assembly of

amphiphiles; 3) self-assembly based

on coordination chemistry; 3) adaptive

soft materials; 4) sevelopment of

efficient self-assembling methods for

functional materials.

She was selected into the New

Century Training Program for the

Talents by the State Education

Commission of China in 2009, Best

Researchers in Colloid Science of

China (2013), Teaching Award of

Peking University (2013), winner of

Outstanding Youth Science

Foundation, Natural Science

Fundation of China (NSFC, 2014),

Excellent doctoral thesis instructor of

Peking University (2014).


231

KN-SMART-2

Self-Assembled Polymer Electrolytes for Future Electrochemical Devices

Moon Jeong Park*

Department of Chemistry, Division of Advanced Materials Science,

Pohang University of Science and Technology (POSTECH), Pohang 790-784, Korea *E-Mail: [email protected]

Abstract

The global energy crisis and an increase in environmental

pollution in the recent years have drawn the attention of the scientific

community to develop innovative ways to improve energy storage and

find more efficient methods of transporting the energy. Polymers

containing charged species have the potential to serve as electrolytes in

next-generation energy systems and achieving high ionic conductivities

from these electrolytes is the key to improving the device efficiency.

Although the synthesis and characterization of a wide variety of

polymer electrolytes have been extensively reported over the last

decade, quantitative understanding of the factors governing the ion

transport properties of these materials is in its infancy. In this seminar, I

will present the underpinning of key factors affecting the

thermodynamics, morphologies and ion transport in polymer

electrolytes by focusing on the use of block copolymers and ionic

liquids (ILs). Various strategies for accessing improved ionic

conductivity and high cation transference number from IL-containing

block copolymers are elucidated. The major accomplishment of

obtaining well-defined nanoscale morphologies for these IL-containing

block copolymers is particularly emphasized as a novel means of

controlling the transport properties. The applications of IL-containing

block copolymers in high temperature fuel cells, lithium batteries, and

electro-active actuators are also enclosed.

Keywords: Block copolymers, Ionic liquids, Self-assembly.

References

(1) Hoon Kim, Joungphil Lee, Hyungmin Ahn, Onnuri Kim, Moon

Jeong Park*, "Synthesis of Three-Dimensionally Interconnected

Sulfur-Rich Polymers for Cathode Materials of High-Rate Lithium-

Sulfur Batteries" Nature Communications, 2015 AOP.

(2) Gyuha Jo, Hongchan Jeon, Moon Jeong Park*, "Synthesis of

Polymer Electrolytes Based on Poly(ethylene oxide) and an Anion-

Stabilizing Hard Polymer for Enhancing Conductivity and Cation

Transport" ACS Macro Lett., 2015, 4, 225−230.

(3) Hyungmin Ahn, Sungyeon Kim, Onnuri Kim, Ilyoung Choi, Chang-

Hoon Lee, Ji Hoon Shim, Moon Jeong Park*, "Blue-emitting Self-

assembled Polymer Electrolytes for Fast, Sensitive, Label-free

Detection of Cu(II)-ions in Aqueous Media" ACS Nano 2013, 7(7),

6162-6169.

(4) Onnuri Kim, Tae-Ju Shin, Moon Jeong Park*, "Fast Low-Voltage

Electroactive Actuators Using Nanostructured Polymer Electrolytes"

Nature Communications, 2013, 4, 2208. DOI:

10.1038/ncomms3208.

(5) S. Y. Kim, S. Kim, Moon Jeong Park*, "Enhanced Proton

Transport in Nanostructured Block Copolymer Electrolyte / Ionic

Liquid Membrane under Water Free Conditions". Nature

Communications, 2010, 1:88.

Education

2006: Ph.D. School of Chemical and

Biological Engineering, Seoul

National University; Mentor: K.

Char

2004 BK-21 Visiting Scholar,

University of Minnesota; Mentor:

T. P. Lodge

2002: M.S. School of Chemical

Engineering, Seoul National

University

2000: B.S. School of Chemical

Engineering, Seoul National

University

Professional Positions

Present Associate Professor,

Department of Chemistry, POSTECH

Mar. 2013 – present: Editorial Advisory

Board, Journal of Applied Polymer

Science

Mar. 2015 – present: Editorial Advisory

Board, Journal of Polymer Science:

Polymer Physics

Honors and Awards

• POSCO Technology Award, POSCO,

Korea, 2015

• Young Scientist Award, John Wiley

& Sons and The Korean Polymer

Society, 2013

• Chong-Am Science Fellowship for

Young Professors, 2011

• Best Lectureship at POSTECH, 2011

• Asia Excellence Award for Young

Scientists in Polymer Science, Osaka,

Japan, 2011

Research Interests

Synthesis and characterization of model

hard/soft materials to elucidate the

mechanisms of charge and ion transport

through these nanostructured materials-

Current efforts focus on correlating

nanoscale structures with ion transport

properties to establish prospective

avenues geared towards high

temperature polymer electrolyte

membrane fuel cells and electroactive

polymer actuators.


232

KN-SMART-3

Non-ionic thermoresponsive polymers of UCST-type in water: Challenges and

perspectives

Seema Agarwal

Faculty of Biology, Chemistry and Earth Sciences, Macromolecular Chemistry II and Bayreuth

Center for Colloids and Interfaces, University of Bayreuth, Universitätsstraße 30, 95440

Bayreuth, Germany

[email protected]

Abstract

Stimuli-responsive polymers or “smart” polymers exhibit a

predictive and sharp change in properties upon only small

changes in the environment (e.g. temperature, pH, ionic strength,

radiation, mechanical stimuli etc.). One class of smart polymers

are thermoresponsive polymers, the best known example being

Poly(N,N-isopropyl acrylamide) PNiPAAm which undergoes

phase separation from aqueous solution at approximately 33 °C.

In contrast to the LCST polymers, only few polymers are known

that show an upper critical solution temperature (UCST) in water

with a sharp phase transition. In most cases the UCST is either

based on ionic interactions or hydrogen bonding. The first is

observed for some polybetaines or quaternized polyurethanes.

However, ionic interactions are disturbed in electrolyte solution.

This is why novel polymer systems showing a sharp UCST over a

wide range of concentrations and being tolerant to electrolytes are

highly desirable. Such polymers can open new application

opportunities.

This presentation will highlight polymers and architectures

showing sharp UCST in water and biological fluids. The reasons

for slow research in this field of UCST polymers in contrast to

LCST polymers will also be discussed by taking few examples.

Seema Agarwal

Universität Bayreuth

Lehrstuhl für Makromolekulare Chemie II

Universitätsstraße 30, 95447 Bayreuth

Germany

Curriculum vitae

1995 PhD Polymer Chemistry, Indian

Institute of Technology (I.I.T.), Delhi, India.

1997 Post-doc at Philipps Universität,

Marburg, Germany.

2007 Habilitation for Macromolecular

Chemistry, Department of Chemistry,

Philipps-Universität Marburg, Germany.

Present Academic Director and Professor at

University of Bayreuth, Germany

Awards and Other Responsibilities

Since 2010 Guest Professor- Jiangxi Normal

University, Nanchang, China

1997–1999 Alexander von Humboldt

Fellowship, Germany

2009 Hermann-Schnell Award (GDCH),

Germany

2009–2011 Associate Editor: Polymers for

Advanced Technologies (Wiley)

Since 2013 Chief-Editor e-Polymers

(De Gruyter publisher)

Since 2013 Managing committee member and

scientific coordinator of a working group

COST 1206 (Electrospun Nano-fibres for bio

inspired composite materials and innovative

industrial applications

Research profile and projects

Functional biodegradable polymers for

biorelevant applications, antibacterial

polymers, thermoresponsive and

photoresponsive polymers, Special fiber

morphologies and properties by bicomponent

electrospinning, nanocomposites. Publications

Peer Reviewed journals: >150

mailto:[email protected]


233

KN-SMART-4

Polymer-Based Smart Devices: Electronics on Paper, Plastic and Textile

Teerakiat Kerdcharoen

Department of Physics and NANOTEC’s Center of Excellence, Faculty of science,

Mahidol University, Bangkok 10400, Thailand


Abstract

Polymer has become basic material for infrastructure

development of human civilization in the past century. Recently,

polymer has been further upgraded to the next version, “Polymer

2.0”, in which advanced functionalities such as electronics and

logical functions can be embedded. In this lecture, we will

demonstrate various polymer-based smart devices based on the

paper, plastic and fabric substrates. Our devices are mostly

wearable including interactive data glove, wearable electronic

nose, gait monitoring and sniffing shoes and smelling shirts. As

focused on the gas sensors and electronic nose technology, we

will present the new concept of human health monitoring based

on the human body and wearable device.

Keywords: Printed electronics; gas sensor; electronic nose;

wearable intelligence

Teerakiat Kerdcharoen

Education:

1986-1989 B.Sc. (Chemistry)

Chulalongkorn University, Thailand

1990-1991 M.Sc. (Physical Chemistry)

Chulalongkorn University, Thailand

1992-1995 Dr.rer.nat (Physical

Chemistry) University of Innsbruck,

Austria

1999-2000 Postdoc (Materials Science &

Engineering) Technical University of

Munich, Germany

Awards :

1999-2000 DFG Postdoctoral Fellow,

Technical University of Munich

2001 Cherry L. Emerson Fellow, Emory

University, USA

2001 Thailand Young Scientist Award

2005 Thailand Toray Science &

Technology Aid Prize

2008 Thailand Research Fund’s

Outstanding Research Prize

2008 NECTEC’s Top Ten Outstanding

Research

2009 Mahidol University Publication

Prize

2014 1st Prize, Innovation for Crime

Combating Contest 2014

(Department of Special Investigation)

2015 Invention Award, National

Research Council of Thailand

Research Interests:

Hardware and software development of

smart devices

Chemical sensors, Electronic Nose

Precision Farming, Smart Farm


234

SMARTO-02

Preparation of Microcapsules Containing Citronellal Oil And Galangal Extract

Kankamon Sinpaksa and Arunee Kongdee Aldred

Program in Applied Chemistry, Faculty of Science, Maejo University, ChiangMai 50210

Abstract The microcapsules have been prepared by in-situ suspension polymerization. In this method, microcapsules of

melamine-formaldehyde (MF), urea-formaldehyde (UF) containing citronellal oil and galangal extract were prepared.

Morphology of microcapsules were examined by using Scanning Electron Microscope (SEM). Functional groups of

microcapsules were analyzed by using Fourier transform infrared spectrometer (FT-IR). The sizes of microcapsules

were investigated by using particle size analyzer (PSA). Thermal analysis of microcapsules was analyzed by using

Thermogravimetric analyzer (TGA). The results showed that microcapsules were spherical, but different in shape

depended on chemicals added in the reaction. The absorption band corresponding to functional groups in citronellal oil

and galangal extract did not appear in FT-IR spectra of microcapsules loaded with citronellal oil and galangal extract,

even though microcapsules were coloured. MF microcapsules trapped citronellal oil as the evidence of evaporation

temperature of citronellal oil in MF at 125.4 °C in TGA thermogram, while UF microcapsules could not trap.

Keywords: microcapsules, melamine-formaldehyde, urea-formaldehyde, citronellal oil and galangal extract

1. Introduction

Microencapsulations is a process in which tiny

particles or a coating to give a small size of

microcapsules. The microcapsules have been widely

applied in many applications such as paper industries,

food stuff, pharmaceuticals, cosmetics and coating

materials to protect substrates from their surrounding

environment (Haghi et al., 2014). It has been prepared by

in-situ suspension polymerization. The microcapsules are

small sphere including of wall around core. The core

material can be a spherical or irregular particle, liquid-

phase suspended solid, solid matrix or liquid forms.

The classification of microcapsules can be

divided into three basic categories according to their

morphology as mononuclear, poly-nuclear and matrix

types. Mononuclear microcapsules contain the wall

around the core, while poly-nuclear microcapsules have

many core encapsulated in the wall. In matrix type, the

core material is distributed homogeneously into the wall

material. In addition to these three basic morphologies,

microcapsules can be mononuclear with multiple walls

(Umer et al., 2011).

In this paper, the single-wall microcapsules of

melamine-formaldehyde (MF), urea-formaldehyde (UF)

containing citronellal oil and galangal extract were

prepared by using in-situ suspension polymerization. The

reaction of MF-prepolymer is shown in Fig. 1.

Fig. 1. The reaction of melamine-formaldehyde pre-

polymer.

The MF microcapsules are produced by

dropping the emulsion of citronellal oil and galangal

extract into the prepolymer. The reaction mechanism is

shown in Fig.2.

MF prepolymer MF prepolymer

Fig. 2. The reaction of melamine-formaldehyde

polymerization (Zhang and Wang, 2009).

The UF microcapsules were included urea and

formaldehyde material as a wall. They were prepared by

using in-situ suspension polymerization. The reaction of

UF-prepolymer and UF condensation polymerization are

shown in Fig. 3. and Fig. 4., respectively.


235

Fig. 3.The reaction of urea-formaldehyde prepolymer

(Yuan et al., 2006).

Fig. 4.The reaction of urea-formaldehyde condensation

polymerization (Yuan et al., 2006).

Citronellal oil is extracted by many methods

such as solvents or microwave extraction. Citronellal oil

has been reported to possess a mosquito repellent action.

A chemical structure of citronellal oil is shown in Fig. 5.

(Hwang et al, 2006).

Fig. 5.The chemical structure of citronellal oil.

The galangal extract was extracted from

cymbopogonnardus. It has been found to contain strong

bioactive compounds as cineole, eugenol, L’8-cineole,

D,L’-acetoxychavicol acetate, p-coumarydiacetate and

palmitic acid etc (Oonmetta-aree et al., 2006).

Citronellal oil and galangal extract were

encapsulated in MF and UF microcapsules. Morphology

and particle size and functional group of microcapsules

were analyzed. Encapsulation capability of citronellal oil

was studied using TGA. The microcapsules containing

citronellal oil and galangal extract will be further used

for textile finishing to impart mosquito repellent action.

2. Experimental

2.1. Preparation of MF microcapsules

containing citronellal oil and galangal extract.

First step, preparation of citronellal oil and

galangal extract emulsion. 7 % (w/v) of sodium dodecyl

sulfate (SDS) was dissolved in 100 ml of water and 5 ml

citronellal oil/galangal extract was added into the SDS

solution the emulsion was stirred at a speed of 300 rpm

at 50 °C.

The second, preparation of MF-prepolymer.

36 % (v/v) formaldehyde solution was dissolved in 2.5 %

(w/v) melamine solution which its pH was adjusted to 8-

9 with 10 % (w/v) Na2CO3 and the condensation was

carried out at 70 °C under stirring for 1 hr.

Finally, preparation of melamine-formaldehyde

microcapsule containing citronellal oil and galangal

extract. The emulsion from the first step was dropped

into the MF pre-polymer and stirred for 15 minutes.

After that 0.03 % (w/v) poly(vinyl alcohol, PVA) was

added and the pH of solution was adjusted to 4 with

10 % (v/v) CH3COOH. The solution was stirred at 70 °C

by mechanical stirrer at a speed of 1000 rpm for 1 hr.

2.2. Preparation of UF microcapsules containing

citronellal oil/galangal extract.

5.0 g urea, 0.5 g ammonium chloride, 0.5 g

resorcinol and 10 ml of 5 % (w/v) PVA solution were

dissolved in 260 ml of water. The pH of solution was

adjusted to 3.5 with 10 % (v/v) of hydrochloric acid and

then 12.60 ml of 36 % (v/v) of formaldehyde was added

in the solution. This solution was stirred at 1000 rpm,

55 °C for 4 hrs. The microcapsules were received, they

were filtered and washed by 10 % (v/v) ethanol for

microcapsules containing citronellal oil but

microcapsules containing galangal extract, was washed

by distilled water.

2.3. Analyses of microcapsules

Microcapsules were examined by using a

scanning electron microscope (SEM) and they were

analyzed by using a Fourier transform Infrared

spectrometer (FT-IR). The microcapsules were mixed

with KBr and their pellets were prepared. Spectra of wall

and core material i.e. with citronellal oil and


236 microcapsules with and without citronellal oil and

galangal extract were recorded by Fourier Transform

Infrared spectrometer over the range of 400-4000 cm-1

.

The microcapsules size was investigated using Particle

Size Analyzer (PSA). Thermogravimetric analysis

(TGA) was used to studyencapsulation capability of

microcapsules at a heating rate of 10 °C/min in nitrogen

atmosphere.

3. Result and Discussion

MF and UF microcapsules without citronellal

oil and galangal extract are white while microcapsules

containing citronellal oil are yellow and those containing

galangal extract are brown. Morphology of MF

microcapsules was examined by using SEM, they are

spheres as shown in Fig. 6. Concentration of SDS affects

to size of MF microcapsules. Size of MF microcapsules

was decreased when SDS was added in the reaction.

The surface of MF microcapsules without SDS

is smooth but that of MF microcapsules with SDS are

coarse.

(a) (b)

Fig. 6. SEM images of MF microcapsules with (a) 0%

(w/v) and ( b) 7% (w/v) SDS (5000x magnification).

The size of UF microcapsules was decreased

when concentration of PVAwas added in the reaction as

shown in Fig. 7.

(a) (b)

Fig. 7. SEM images of UF microcapsules with varied

concentration of PVA (a) 0% (w/v) (b) 5% (w/v) PVA

(5000x magnification).

SEM images of MF microcapsules and UF

microcapsules containing galangal extract areshown in

Fig 8. (a) and (b), respectively. MF and UF

microcapsules with galangal extract are spherical, but

agglomeration of particles is found in MF microcapsules.

(a) (b)

Fig 8. SEM images of (a) MF and (b) UF microcapsules

with galangal extract (5000x magnification).

Fig. 9.shows FT-IR spectra of citronellal oil and

galangal extract. They show absorption bands of O-H


, C-H stretching at 2915 cm-1

,

C=O stretching at 1728 cm-1

and C=C stretching at 1644

cm-1

and absorption bands of C-H stretching at 2978 cm-

1, C=O stretching at 1716 cm

-1, C=C stretching at 1640

cm-1

,C-O bending at 1048 cm-1

and aromatic bending at

878 cm-1

.

Fig. 9. FT-IR spectra of (a) citronellal oil and (b)

galangal extract.

Fig. 10. shows FT-IR spectra of (a) citronellal

oil (b) MF microcapsules without and (c) MF

microcapsules with citronellal oil, It show absorption

bands of C-H stretching at 2961cm-1

and 2929 cm-1

, N-H

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2013/8/21

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237

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06

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00

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and 1570 cm-1

and C-N stretching

at 1349 cm-1

and 1347 cm-1

.

The spectra of (a) galangal extract (b) UF

microcapsules without and (c) UF microcapsules

containing galangal extract are shown in Fig. 11. They

show absorption bands of C-H stretching at 2971cm-1

and 2975 cm-1

, N-H stretching at 1656 cm-1

and 1654

cm-1

respectively and C-N stretching at 1252 cm-1

and

1254 cm-1

respectively.

Fig. 10. FT-IR spectra of (a) citronellal oil (b) MF

microcapsules without and (c) MF microcapsules with

citronellal oil.

Fig. 11. FT-IR spectra of (a) galangal extract (b) UF

microcapsules without and (c) UF microcapsules with

galangal extract.

The absorption bands of citronellal oil and

galangal extract are not present in spectra of MF

microcapsules, this may cause from low amount of

citronellal oil trapped by microcapsules. The

encapsulation capability of MF microcapsules and UF

microcapsules containing citronellal oil was analyzed by

TGA which will be discussed later.

Microcapsules size was analyzed using PSA.

Fig. 12.shows the particle size distribution of MF

microcapsules (a) without (b) with citronellal oil. Size

distribution of MF microcapsules without citronellal oil

is larger than that of with citronellal oil which confirmed

by SEM image in Fig. 6 (b). The highest number of

particles falls in 94.67 µm and 29.18 µmfor MF

microcapsules without and with citronellal oil,

respectively.

Fig. 12.shows the particle size distribution of MF

microcapsules (a) without and (b) with citronellal oil.

Fig.13. shows the particle size distribution of

UF microcapsules (a) without and (b) with citronellal oil.

The particle size distribution of MF microcapsules

containing citronellal oil was larger than MF

microcapsules without citronellal oil. The highest

number of particles falls in approximately 29.18 µm and

59.87 µm for UF microcapsules without and with

citronellal oil and galangal extract,respectively. This

result may attribute to encapsulation capability of

citronellal oil by UF.

Fig.13.shows the particle size distribution of UF

microcapsules (a) without and (b)with citronellal oil.

0

1

2

3

4

5

6

7

8

9

10

0.01 0.10 1.00 10.00 100.00 1000.00

Vo

lum

e (

%)

Particle size (µm)

0

1

2

3

4

5

6

7

8

9

10

0.01 0.1 1 10 100 1000

Vo

lum

e (

%)

Particle size (µm)

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(a)

(b)

(c)

(b)

(c)

(a)


238 In Fig.14., TGA thermogram of citronellal oil

shows weight loss at 128.0 °C, meaning evaporation of

the oil. Thermogramof Fig. 14. (b) MF microcapsules

without citronellal oil shows weight loss of 12.31% at

54.4 °C that corresponds to evaporation of moisture in

microcapsules (Park et al., 2001). The degradation of MF

with citronellal oil as 11.28% is observed at 159.5 °C as

26.95% at 328.1 °C and 40.57% at 418.3 °C. For

thermogram of MF microcapsulescontaining citronellal

oil in Fig. 14. (c),moisture is found to evaporate at

50.0°C with 5.01% weight loss, evaporation of

citronellal oil is at 125.4 °C and degradation of MF

microcapsules at 270.1 °C and 388.6 °C with 40.29%

and 21.5% weight losses.

Fig.14.Thermogram of (a) citronellal oil (b) MF

microcapsules without and (c) with citronellal oil

Fig. 15.Thermogram of (a) citronellal oil (b) UF

microcapsules without and (c) with citronellal oil.

In Fig. 15. (b), thermogram of UF

microcapsules shows weight loss of evaporatedmoisture

at 44.5 °C for 13.66%, weight loss of UF microcapsules

is at 266.8 °C for 74.09%. For UF microcapsules with

citronellal oil in Fig. 15. (c), weight loss of 99.04% is

found at 255.0 °C. No evaporation of citronellal oil is

observed in TGA thermogram.

4. Conclusion

MF microcapsules and UF microcapsules were

spherical and different in shape. MF microcapsules

prepared by addition of SDS were small and had coarse

surface, while UF microcapsules prepared by addition of

PVA were small.Microcapsules containing citronellal oil

were yellow while those containing galangal extract were

brown color. FT-IR spectra of UFmicrocapsules

containing and without citronellal oil and galangal

extract were similar as low amount of citronellal oil and

galangal extract, probably lower than detection limit of

FT-IR. Using PSA, the sizes of microcapsuleswith

citronellal oil were larger than those without citronellal

oil. MF microcapsules could trap citronellal oil as weight

loss of 19.96% was found at 125.4 °C in TGA

thermogram. Fabric finishing with MF microcapsules

containing citronellal oil will be further prepared.

Mosquito repellency of the fabric is expected.

5. Acknowledgement

The authors greatly appreciate Program in

Chemistry, Faculty of Science, Maejo University,

ChiangMai, Thailand for financial support.

6. Reference

[1] Haghi E. H., Mirabedini S. M., Imani M., Farnood

R. R., 2014. Preparation and Characterization of

Pre-silane Modified Ethyl Cellulose- based

Microcapsules Containing Linseed Oil, Colloids and

Surfaces A: Physicochemical and Engineering

Aspects, pp. 1-30.

[2] Hwang. J. S., Kim. J. N., Wee. Y. J., Yun. J. S.,

Jang.H. G., Kim. S. H. and Ryu. H. W., 2006.

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[3] Oonmetta-aree J., Suzuki T., Gasaluck P. and

Eumke G., 2006, Antimicrobial properties and

action of galangal (Alpiniagalanga Linn.) on

(a) citronellal oil (b) MF

i(a) citronellal oil (b) MF without

citronellal oil (c) MF with

citronellal oil i54.4°C

F159.5°C

t125.4°C

270.1°C

o128.0 °C

i328.1°C

l388.6°C

418.3°C

-120

-100

-80

-60

-40

-20

0

0 50 100 150 200 250 300 350 400 450 500

We

igh

t (%

)

Temperature (°C)

(a) ci

ronellal

il

b) UF witho

t citronellal oil (c) UF with citronellal oil

(a)

(b)

(b.)

(

a

)

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270.

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388.6°C

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44.5°

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266.8°C

255.0 °C

(a.)


239 Staphylococcus aureus, LWT-Food Science and

Technology, vol. 39, pp.1214-1220.

[4] Park J. S., Shin S. Y., and Lee R. J., 2001,

Preparation and characterization of microcapsules

containing lemon oil, Colloid and Interface Science,

vol 241, pp. 502-508.

[5] Umer H., Nigam H., Tamboli M. A., Nainar M. S.

M., 2011. Microencapsulation: Process, Techiques

and Applications, Pharmaceutical and Biomedical

Sciences, vol 2, pp. 474-481.

[6] Yuan L., Liang G., Xie Q. J., Li L. and Guo J., 2006,

Preparation and characterization of poly(urea-

formaldehyde) microcapsules filled with epoxy

resins, Polymer, vol. 47, pp. 5338-5349.

[7] Zhang H., Wang X., 2009, Fabrication and

performances of microencapsulated phase change

materials based on n-octadecane core and resorcinol-

modified melamine–formaldehyde shell, Colloids

and Surfaces A: Physicochemical and Engineering

Aspects, vol. 332, p. 129-138.


Polymer Research in Industry Sector


241

PTT

Development of bioplastics-based lamination application

Narin Kaabbuathong

Process Technology Research Department

PTT Research and Technology Institute

PTT Public Company Limited

71 M. 2, Phahonyothin Rd., Sanubtub, Wangnoi,

Ayutthaya 13170 THAILAND


Abstract

Recently, plastic-based lamination packaging products are

widely used in our daily-life due to their various advantages noted as;

ease of processability, good mechanical properties, good gas barrier and

cost competitiveness. Thus, we can easily find them in a broad range of

packaging product. These laminated plastics can be fabricated into a

simple disposal packaging such as paper cup or paper food tray or a

complex multilayer flexible packaging. Unfortunately, these products

are typically produced by laminating the conventional plastic on the

desired substrate. After their end of life, these non-biodegradable

products are littered and still remained in the nature for hundreds of

year causing of a waste pollution.

To relieve the above concern, the concept of replacing

conventional plastics with bioplastics into the products was introduced

by PTT Group since 2012. As known, bioplastics have several benefits

over conventional plastics in term of biodegradability, sustainability of

raw material and better LCA. However, some limitations of applying

bioplastics to the conventional lamination machines are needed to

solve. The most common technical issues are known to be; an

instability of molten bioplastic curtain during the lamination process,

high percent of neck-in and low adhesion strength between bioplastic

film and substrate.

Hence, this talk will focus on the key success factors of bioplastic

lamination technology together with the corresponding R&D activities

and/or tool which are applied to overcome all of those relative technical

issues.

Narin Kaabbuathong

Education

Jan 2000 - March 2003: Ph.D., Material

Engineering, Department of Science and

Technology, University of Rome “Tor Vergata”, Rome, Italy

May 1997 - April 1999: M.Sc., Polymer Technology, Department of Polymer

Technology, The Petroleum and

Petrochemical College, Chulalongkorn University, Bangkok, Thailand

May 1994 - March 1997: B.Sc., Polymer Science, Department of Material Science,

Chulalongkorn University, Bangkok,

Thailand

Professional Experiences

2012 - Present: Head of Polymer and Advanced Material, PTT Research and

Technology Institute (PTT RTI), PTT

Public Company Limited

May 2006 – 2012: Researcher of Polymer

and Advance Material Research Group,

Process Technology Research Department, PTT Research and Technology Institute

(PTT RTI), PTT Public Company Limited

October 2003 - April 2006: Western

Digital (Hard Disk Drive Manufacturer)

Staff Engineer/ Department of Research and Development

Academic activities

July 2007 - 2013: Committee of The Royal

Institute Thailand

September 2004 - Present: Invited lecturer and Thesis Co-Advisor (M.Sc. Student)

December 2003 - Present: Committee of Thai Institute of Physics, Thai Institute of

Physics (TIP), Bangkok, Thailand

August 2002 - September 2002: Committee of international conference,

International Conference on electroceramic society IX in association

with University of Rome, Rome, Italy


242

PTTGC

From Bioscience to Polymer Science: Our Sustainable Prospects

Sukhgij Ysothonsreekul

Vice President for Scientific Research, Science and Innovation, PTT Global Chemical

PTT Global Chemical Public Company Limited

Abstract

It is indisputable that Thailand harbors abundant agricultural

resources waiting to be converted to precious molecules to

further chemically modified for polymer technology. PTT

Global Chemical Plc. (“PTTGC”) envisions such advantage

situation of Thailand, we have set path of our future and

sustainable growths, especially the polymer products, through

the integrated technology of biological science, chemistry,

and polymer science. Utilizing sustainable biomass combined

with the cutting-edge technology, PTTGC has moved forward

to become a leader in biopolymer business and confirmed to

create a sustainable future.

From molecular biology to genetically modified microbes,

fermentation technology to produce biomolecules of interest,

and chemical modification, to polymer science, our scientists

dedicated to produce novel/ high specialty materials to be

further used in variety of product applications. Not only the

Bio-based products such as proteins, carbohydrates,

polyunsaturated fatty acids, but also the chemical building

blocks such as succinic acid, lactic acid, muconic acid, adipic

acid, and ethanol can be produced via biotechnological route.

Through our acquired the US-based biotechnology company,

Myriant, and lessons-learned from the commercial scale

production of bio-succinic acid, PTTGC has become a front-

runner in committing to produce products in green business.

We have moved towards the using of our abundant biomass

available in Thailand as feedstock in the fermentation process

by the genetically engineered microbes from Myriant’s

sophisticated technology. A novel development of microbial

system should enable us to convert sugars contained in the

biomass into high value, high purity renewable chemicals

efficiently at commercial scale. Moreover, PTTGC has

enhanced our sustainability by using renewable bio-based

chemicals as monomers to produce biodegradable plastics and

high value specialty materials. Our research focused on

polymerization, compounding, and processing enables us to

produce biomaterials and bioplastics for film application, and

the future possibilities abound. With sustainability and social

responsibility mindsets, it is our commitment to be frontier in

utilization of our abundant agricultural biomass to produce

the environmental friendly products for society.

Ph.D. (Biochemistry) Kansas State University

Academic Positions:

1999-2003 Assistant Research Professor, University of

California, San Diego, USA 2004-2005 Associate Research Professor, University of

California, San Diego, USA

2006-2013 Associate Professor in Biochemistry, Naresuan University

2008-present Adjunct Professor, Graduate School of

Biomedical Sciences, University of North Texas Health Science Center, Fort Worth, Texas, USA

2011-present Invited Professor, University of Franche-

Comte, Besancon, France

Administrative Positions:

2006-2009 Associate Dean for Research and Graduate

Studies, Faculty of Medical Science, Naresuan University, Phitsanulok, Thailand

2006-2009 Director of Center for Forensic Sciences,

Faculty of Medical Science, Naresuan University, Phitsanulok, Thailand

2008-present Member of the Forensic Science Research

Board, Central Institute of Forensic Sciences, Ministry of Justice, Thailand

2008-present Genetic Specialist to Forensic Genetic

Laboratory, Royal Police Department 2010-2012 Vice President for Research and External

Relations, Naresuan University, Phitsanulok, Thailand

2010-2012 Chairman of Naresuan University Research Development Board

2010-2012 President of Toastmaster Club of Naresuan

University 2010-2013 Chairman of Naresuan University

International Development Board

2010-2013 Member of Naresuan University Human Resource Management Board

2010-2013 Member of Naresuan University

International College Board 2012-2013 Vice President for International Affairs,

Naresuan University, Phitsanulok, Thailand

2012-2013 Advisor to Naresuan University Research

Development Board

2012-2013 Advisor to Logistic and Supply Chain Curriculum Development, Naresuan University for the

Office of Higher Education Commission

2013-present Vice President for Scientific Research, Science and Innovation, PTT Global Chemical


243

AGILENT

Optical characterization of thin films using a new Universal Measurement Accessory

Heng Soo Chin

Applications Engineer, Agilent Technologies, Inc.

Abstract

Traditional approaches to thin film characterization have relied

on a single angle or a small set of angles, often measured with relative

reflectance accessories. This left thin film designers the task of

correcting results to absolute values or extrapolating data from a limited

angular set out to the angles of interest to estimate thin film response.

Also, the limited or lack of transmission data has resulted in

assumptions being made about the end product. The fine angular

control and automation of the new approach enable research to capture

both absolute reflectance and transmission at the designed angle,

removing the guess work and allowing for precise and detailed

validation of thin film designs.

Heng Soo Chin

Ms. Heng Soo Chin is currently an

Applications Chemist for Molecular

Spectroscopy with Agilent

Technologies, Singapore, covering

Singapore and ASEAN regions. She

addresses applications for FTIR

Spectroscopy, FTIR Microscopy, UV-

Vis, and UV-Vis-NIR systems. She

holds a BSc (Hons) in Chemistry and

MSc in Advanced Chemical

Engineering from Imperial College,

London. Prior to Agilent, Soo Chin

gained R&D experience in the aerospace

industry and technical experience from

the analytical industry.


244

SCG

High Performance Composites for Industry

Nopphawan Phonthammachai

Composite Research, R&D Center, SCG Chemicals Co., Ltd

10-I1 Road, Map Ta Phut Industrial Estate, Muang District

Rayong 21150, Thailand

Tel: +66 3891 2826, Fax: +66 3868 4676, Email: [email protected]

Abstract

Due to the tailor-made properties of composite materials, they

have been paid attentions by industries for advanced and innovative

product development satisfying customer’s requirements and global

trend. Examples on the replacement of steel and conventional materials

by composite have been shown more and more from aerospace to

automotive, infrastructure until electronic and consumer products;

driving the trend of composite development toward more specific

properties. With these movements, the concerns on processability, cost,

and relevant certified tests of composite materials have also been

raised.

Nopphawan Phonthammachai

Education 2001-2005: Doctor of Philosophy in

Polymer Science, The Petroleum and

Petrochemical College, Chulalongkorn University, Bangkok, Thailand in academic

partnership with Case Western Reserve

University, US 1997-2001: Bachelor of Science in

Chemistry, Naresuan University,

Phitsanulok, Thailand

Professional Experiences

2012 – Present: Senior Researcher Head

of Composite Research Group, SCG Chemicals, Thailand

2008 –2012: Scientist at Institute of

Materials Research and Engineering, Singapore

2005- 2008: Research Fellow at School

of Material Science and Engineering, Nanyang Technological University,

Singapore

Research Interest

1. Designing and tailoring the properties

of active nano-fillers for polymer-based composites

2. Development of low-cost, convenience

and effective methods to produce nano-fillers and composites

3. Surface modification and interfacial

study of nano-fillers and composites 4. Characterizations of nano-fillers and

composites

5. Correlated study between composition, process and property of composite

materials (deformation/failure mechanism,

interfacial interaction, screw design and processing condition, specific property)

6. Applications of composite materials

Awards and Honors

2004: Research Scholarship Award at the

Department of Macromoleular Science, Case Western Reserve University, Ohio,

USA

2004: The Outstanding Student Award of Chulalongkorn University

2001: Royal Golden Jubilee Scholarship

from Thai Research Found 2000: Research Scholarship Award at the

Department of Chemical Engineering,

University of Melbourne, Australia


245

SYNCHROTRONS

Synchrotron Light for innovative polymer

Prae Chirawatkul

Beamline Scientist

[email protected]

I have got a Ph.D. in Physics from the University of Bath, UK. I

started working at the Synchrotron Light Research Institute, SLRI, Thailand, in

2011. Between 2013-2014, I was working as a post-doctoral researcher for the

9A-USAXS (Ultra small angle x-ray scattering) beamline at the Pohang

Accelerator Laboratory (PAL), South Korea. Now I am a beamline scientist at

the SAXS (small angle x-ray scattering) beamline at the SLRI. Together with

the SAXS team, we provide technical support to our users. We help users to

plan their experiments and help them to familiarize themselves with the control

software and measurement protocols in the experimental station. My area of

interest is beamline instrumentation. I am also leading a project to construct a

new beamline which supports wide angle x-ray scattering technique in a

grazing incidence geometry and a coupling measurement between x-ray

absorption and x-ray scattering techniques.

Wanwisa Limphirat

(Pattanasiriwisawa)

[email protected],

[email protected]

Phone: +66-44-217 040

ext. 1480

Fax: +66-44-217 047

Address

Synchrotron Light Research Institute (Public Organization),

111 University Avenue, Muaeng, Nakhon Ratchasima 30000, Thailand

Education

1995-1998 Bachelor of Science (1st class honor), Kasetsart University,

Bangkok, Thailand

1999-2004 Doctor of Philosophy (Physics),

Suranaree University of Technology, Nakhon Ratchasima, Thailand

Research Interests

X-ray Imaging, X-ray Absorption Spectroscopy (XAS), Infrared

Microspectroscopy

Professional Experience

2005-present Researcher: Synchrotron Light Research Institute (SLRI),

Nakhon Ratchasima, Thailand

2010-2011 Beamline Manager (Beamline IR): Synchrotron Light Research

Institute (SLRI), Nakhon Ratchasima, Thailand

2008 Visiting Scholar: SOLEIL Synchrotron, Gis sur Yvette C_edex, France,

September 2008

2007 Visiting Scholar: High Energy Accelerator Organization (KEK),

Tsukuba, Ibaraki, Japan, 10-16 May 2007

2000-2002 Research Assistant: High Energy Accelerator Organization (KEK),

Tsukuba, Ibaraki, Japan

Scholarships/Award

1996-1999 Recipient of the Development and Promotion of Science and

Technology Talents Project (DPST) from The Institute for the Promotion of

Teaching Science and Technology

2000-2004 Recipient of the Royal Golden Jubilee (RGJ) scholarship from the

Thailand Research Found

2002 Recipient of the AIEJ scholarship from the Association of International

Education, Japan

1999 Recipient of Professor Dr. Tab award

1996-1998 Recipient of Kasetsart outstanding academic performance award

(every year)


Instruments session


Bara Scientific Co. Ltd. AXIS SUPRA, Cutting edge XPS technology for Polymer

Tan Teck Beng Assistant Sales Manager, Shimadzu Asia Pacific Pte Ltd

Abstract

Polymer characterization is done with most

of the basic techniques such as FTIR, GCMS or

many instruments. Especially polymers find usages

on to many fields especially coatings & coatings

with very thin films are always challenge to

characterize.

X-ray Photoelectron Spectroscopy (XPS)

is the most powerful technique for chemical

compound analysis on material surface. Polymers

are insulating materials & getting them analyzed

with good charg-neutralizer is key to obtain great

significant, dependable data. AXIS SUPRA is a

newly designed cutting edge technology from

KRATOS ANALYTICAL gives superb signal and

resolution on Polymer materials.

Tan Teck Beng

Assistant Sales Manager

Shimadzu Asia Pacific Pte Ltd

Tan Teck Beng is the Application Specialist in Shimadzu Asia

Pacific for XRF, XRD, ICP, SEM, SPM and ESCA. He has conducted Application Seminars and Workshops on RoHS at

Panasonic, Sharp, Sony, Pioneer, National Metal and Technology

Center (Thailand), Intertek (Hong Kong) and Hong Kong PSB. The countries where he has conducted Shimadzu RoHS Seminars

include Malaysia, Indonesia, Thailand, India, Philippines and

Singapore. Shimadzu is currently providing Application and Technical Support for their EDX for RoHS Screening to some 200

companies in Asia Pacific Region. He has an MSc in Materials

Science from National University of Singapore.


Crest Nanosolution (Thailand) Ltd.

Simultaneous Measurement of Thermogravimetric Differential Thermal Analysis and

Photoionisation Mass Spectrometry Complexed Through Unique Skimmer Interface System,

TG-DTA-PIMS

Tadashi Arii1 and Taisuke Yoshiki

2*

1 SBU-TA, Thermal Analysis Division, Rigaku Corporation, Tokyo 196-8666, Japan

2 X-ray Diffractometer & Thermal Analyzer, Rigaku Asia and Pacific Limited, Pathumthani 12130, Thailand

Phone +81-3-3479-0618, Fax +81-3-3479-6112, *E-Mail: [email protected]

Abstract

Although thermal analysis has wide range applications to understand

thermophysical and chemical changes at a macro-molecular level, it is necessary to

perform complex measurements such as hyphenated methodology combined with other spectroscopic methods to obtain specific micro-molecular information on reaction

products. Nowadays, one of the powerful complex techniques is well known as TG-DTA-

MS, which is a simultaneous measurement technique composed of thermogravimetric differential thermal analysis (TG-DTA) combined with mass spectrometry (MS) through

an interface system. It is suitable for the qualitative analysis of the different gases evolved

in response to heating a polymeric material in TG-DTA process. This presentation aims to

propose a novel thermoanalystical method that integrates a “skimmer-type interface” and

a “photoionization method”, in order to overcome the essential problems and the

limitations of conventional TG-DTA-MS. The traditional style of interface system in TG-DTA-MS employs a capillary

type which consists of a relatively long narrow tube (about 2m length) connecting both

devices, and it adopts the principle that the injection tube is maintained at a constant

temperature (< 300℃). The keeping temperature of the interface is determined by

considering re-condensation and transformation of the injected gases as well as the user’s safety. The gaseous products re-condense within the interface and are trapped internally if

the boiling points of the gasification products exceed the interface temperature, adversely

if the interface temperature is too high the gasification products cause the gas transformation due to secondary reactions between the activated pyrolysates, which often

interferes with the measurement results especially in the degradation analysis of

polymers. The proposed skimmer type interface consists of two concentric ceramic tubes

with orifices of the different diameters located into the furnace of the TG-DTA.[1] The

interface connects directly the two devices, one at atmospheric pressure of the sample in the TG-DTA and the other at high vacuum MS chamber. The sample position and two

orifices of skimmer interface as well as the MS ion source are arranged in a straight line.

The gas injection port of the skimmer interface is located in close proximity to the sample and is thermally programmed under the same environment up to maximum operation

temperature of the instrument. In this way, high-precision gas analysis becomes

fundamentally possible because the interface length is negligibly minimized. In the conventional MS, although a gaseous molecule is ionized by colliding with

an accelerated electron by the electron ionization (EI) method of 70eV, a part of the

generated molecule ion further decomposes, and consequently the molecule ion is observed simultaneously with the fragmented ions. The analyses of multiple complex

gases formed by the pyrolysis of polymeric materials using conventional TG-DTA-

(EI)MS become more difficult because many kinds of pyrolysates evolve simultaneously or continuously during heating. This means that the fragment ions generated as a result of

the higher ionization potential of EI often obstruct the identification of the gases formed

by heating process. In order to overcome such complex evolved gas analysis, one feasible approach is the use of a selective and soft fragment-less ionization method.

Single-photon ionization with a vacuum ultraviolet (VUV) light source with the

ionization potential of 10.2eV is a particularly soft and selective ionization method, suited well for detection of both aromatic and aliphatic organic species. The photoinization (PI)

is the simplest electron transfer reaction induced by photoabsorption. Consequently,

because only molecular (parent) ions of the gases are observed in the resulting fragment-less mass spectrum, it is possible to directly differentiate multiple gases evolved by using

the discrete information on their molecular ions.[2] With above-mentioned potential advantages, use of both of the skimmer type

interface and the PI method for the TG-DTA-MS analysis greatly enhances instrument

adaptability to broader classes of organic compounds including polymer resins and the features permit a better understanding of the thermal behavior and precise pyrolysates of

polymers. To demonstrate the effectiveness of the technique, the results of its application

to the TG-DTA-MS analysis of typical polymeric materials are presented with the direct characterization of polymer degradation products by focusing on the minute structural

difference between the samples.

References

[1] T. Arii J. Mass Spectrom. Soc. Jpn, 53, 211-216 (2005)

[2] T. Arii and S. Otake, J. Therm. Anal. Cal., 91, 419-426 (2008)

Tadashi Arii

Education:

1984 March B.E. in Chemistry,

Department of Materials Science and

Technology, Nagaoka University of Technology.

1997 March Ph.D. in Chemistry,

Department of Energy and Environment Science, Nagaoka University of Technology.

Awards:

2007 Research Award of the Mass

Spectrometry Society of Japan

Experience:

1984-Present Thermal Analysis Division, Rigaku corporation

Been in charge of Strategic Business

Unit Manager

Skills:

Thermal Analysis, Mass Spectrometry


JAIMA. Introduction of JAIMA and Investigation Project

for Analytical Instruments Related Industries

Goto Ryozo Japan Analytical Instruments Manufactures’ Association

Abstract JAIMA (Japan Analytical Instruments

Manufactures’ Association) consists of analytical

instruments related companies. JAIMA was founded

with 18 members in 1960, and now around 200

companies join the association at present.

Our menbers separated 2 group , one of A

(made and developed in Japan) and other B ( not

made in Japan but sale in Japan). Various

committees are organized within JAIMA to respond

to several requests from outside, and to submit the

advice and action proposals to the board of directors.

One of most important activity for JAIMA

is to held “JASIS”. JASIS is one of the Largest

Exhibitions in Asia for Analytical and Scientific

Instruments. JAIMA held JASIS every year (on

September, Makuhari near Tokyo). Now we research

“ the way of systematic human resource

development for Instrumental Analysis and

Analytical Science in ASEAN “ Today I introduce

our An organization and activity.

Graduate the HIROSHIMA University.

The doctor's degree (engineering section) is given.

The winners Prize in the Ion-Chromatography

technology (2009)

Received the Technical Lifetime for FIA

Achievement Award(2013)

Adviser for Technical Affairs Committee ,JAIMA

Research field: Ion Chromatography,

Electrochemical Analysis , Environmental Analysis ,

etc.

E-mail : [email protected]

SM Chemical

Abstract

Size exclusion chromatography (SEC) is a convenient and

reproducible method to analyze and determine relative molecular mass

and molecular mass distribution of polymers. It is widely used for R&D

and QC in polymer industry. We, Tosoh Corporation have a long

history to develop SEC columns and equipment since 40 years ago and

have been perusing high-throughput, reproducible and easy-to-use SEC

system. One of the key factor for especially high-throughput is column

technology and we developed polystyrene/divinylbenzene packing

materials with 3 micron as particle size and packed in semi-micro

columns (4.6 mm I.D. x 15 cm) few years ago. The columns achieved

high-throughput analysis for half measurement time of conventional

column (30 cm length) with the same or better resolution. Recently we

employed multi-pore size distribution technology on

polystyrene/divinylbenzene packing materials with 3 - 6 micron as

particle size and packed in semi-micro columns (4.6 mm I.D. x 15 cm)

and commercialized under the brand name of TSKgel

SuperMultiporeHZ columns. This newly employed technology enables

wide separation range of polystyrene calibration curves in

tetrahydrofuran with good linearity on identical packing materials and

no distortion on the chromatograms which are sometimes observed as

mixed-bed type columns or series of individual columns are used.

Using this column reproducible SEC was achieved in addition to high-

throughput and high resolution results.

Ivan Lim

Ivan Lim is the Regional Sales

Executive in Tosoh Asia (Bioscience

Separation Division). Tosoh is the leader in the development and production of gel

permeation chromatography system and

columns that are used extensively for analysis of polymers. Ivan was graduated with

Bachelor Degree in Applied Science

(Industrial Chemistry) from University of Science Malaysia and has been in the

chemical industry for 6 years in various

capacities. Currently, he is in charge of the chromatography product sales in Thailand,

India, Singapore, Malaysia, Indonesia,

Australia and Vietnam.


Chiang Mai University Research

Research Laboratory for the Production of High Quality Resorbable Polymers,

Chiang Mai University

Winita Punyodom

1*, Robert Molloy

1,2, Puttinan Meepowpan

1, Kanarat Nalampang

1, Runglawan Sonsunan

1,

Kiattikhun Manokruang1, Patnarin Worajittiphon

1, Nawee Kungwan

1, Wathuka Booncharoen

1, Donraporn

Daranarong1, Wootichai Khotasen

1, Vivan Thammongkol

3,

Narin Kaabbuathong3, Tuspon Thanpitcha

3, Chaiyapruk Katepetch

3

1

Polymer Research Laboratory, Department of Chemistry, Faculty of Science,

Chiang Mai University, Chiang Mai, 50200 2

Materials Science Research Center, Faculty of Science, Chiang Mai University, Chiang Mai, 50200 3PTT Research and Technology Institute, Wang Noi, Ayutthaya 13170, Thailand

Phone: +66-5394-3345, Fax : +66-5389-2277, *E-mail: [email protected]

Abstract

Bioplastics industry in Thailand can be regarded as the country’s

new-wave business due to the fact that there has not yet been the full-

cycle supporting the industry and the products have not been distributed

widely in the market. Chiang Mai University’s innovation aims to utilize

Thailand’s abundant raw materials to strengthen the country’s Bioplastics

industry in order to become a leader in the region. This research work can

be divided into 3 main areas, namely: (1) polymers for use in biomedical

applications, (2) polymers for environmental applications, and (3) novel

initiators/catalysts for the controlled ring-opening polymerization (ROP)

of cyclic esters. The polymers that are featured in this work are, for the

most part, what can be termed as “specialty polymers” which have been

purpose-designed to meet the specific requirements of various

applications. As such, they are value-added materials which, if they can

be commercialized, will reduce the country’s dependency on expensive

imported products. In its wider context, this research is aimed at

developing new materials for which there is a definite need in Thai

society and at an affordable cost. In the case of specialty polymers such

as these, their added value (as compared to commodity plastics) comes

mainly from the technological know-how involved in producing them

rather than from the cost of the materials themselves. This know-how is

gradually being developed through a combination of pure and applied

research together with close collaborations with end-users and industry.

Keywords: Specialty polymers; Bioplastics; Biomedical Applications;

High Quality Resorbable Polymers

Winita Punyodom

Education

1996-2000 PhD (Polymer Physics) IRC in

Polymer Science and Technology,

University of Leeds, Leeds, United

Kingdom 1990-1994 M.S. (Chemistry)

Chiang Mai University, Chiang Mai,

Thailand 1986-1990 B.S. (Chemistry)

Chiang Mai University, Chiang Mai,

Thailand

Fields of speciallization

• Polymer Chemistry

• Polymer Synthesis, Characterisations and

Testing

• Biodegradable Polymers for Use in

Biomedical Applications and as Bioplastics

List of Sponsors and Exhibitors

Gold Sponsors BASF (Thai) Limited SCG Chemicals CO., LTD. PTT Global Chemical Public Company Limited

Silver Sponsors Thailand Convention and Exhibition Bureau (TCEB)

Bronze Sponsors Bruker Biospin AG

Exhibitors Agilent Technologies (Thailand) Ltd. Bara Scientific Co., Ltd. Coax Group Corporation Ltd. Crest Nanosolution (Thailand) Limited HORIBA (Thailand) Limited LMS Instruments Co., Ltd. RI Technologies Ltd. SM Chemical Supplies Co., Ltd. Ulvac (Thailand) Ltd. Japan Analytical Instruments Manufacturers' Association (JAIMA) Chiang Mai University- Research Laboratory for the Production of High Quality Resorbable Polymers Synchrotron Light Research Institute

Cover-Proceedings-create outline - Polymer Society of Thailand

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