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Polymer 54 (2013) 5230e5240

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Polymer

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The role of non-covalent interactions and matrix viscosity onthe dispersion and properties of LLDPE/MWCNT nanocomposites

Alexandros A. Vasileiou a, Aristides Docoslis a, Marianna Kontopoulou a,*, Peng Xiang b,Zhibin Ye b

aDepartment of Chemical Engineering, Queen’s University, Kingston, Ontario K7L 3N6, CanadabBharti School of Engineering, Laurentian University, Sudbury, Ontario P3E 2C6, Canada

a r t i c l e i n f o

Article history:Received 16 April 2013Received in revised form2 July 2013Accepted 12 July 2013Available online 19 July 2013

Keywords:Polyethylene nanocompositesCarbon nanotubesFunctionalization

* Corresponding author. Tel./fax: þ1 613 533 3079.E-mail addresses: kontopm@queensu.ca,

chee.queensu.ca (M. Kontopoulou).

0032-3861/$ e see front matter � 2013 Elsevier Ltd.http://dx.doi.org/10.1016/j.polymer.2013.07.034

a b s t r a c t

Linear low density polyethylene (LLDPE)/multi-walled carbon nanotube (MWCNT) composites wereprepared by melt compounding, following two different compatibilization strategies that involved non-covalent interactions between the matrix and the filler. The first approach involved grafting pyridinearomatic moieties on the maleated polyolefin backbone, which are able to interact by pep stacking withthe surface of the nanotubes. The second method implemented non-covalent/non-specific surfacefunctionalization of the MWCNTs with a hyperbranched polyethylene (HBPE). The enhanced interfacialinteractions established in the composites containing LLDPE functionalized with pyridine graftsimproved the dispersion of the nanotubes within the polymer matrix. Dispersion was also favoured byhigher matrix viscosity. Composites containing finely dispersed MWCNTs exhibited an increase in therheological and electrical percolation thresholds, and a significant improvement in mechanical proper-ties. On the contrary the composites based on the low viscosity matrix contained large amounts of ag-gregates, which promoted lower percolation thresholds. Manipulation of matrix viscosity andcompatibilization resulted in composites with good mechanical properties, and low percolationthresholds.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Polymer/multi-walled carbon nanotube (MWCNT) nano-composites are receiving steadily increasing attention from thescientific community [1]. The incorporation of graphitic carbonswithin commodity polymers, such as polyolefins, is of great interest[2], because it can result in electrically and thermally conductivematerials suitable for applications such as electromagnetic-reflective materials, electrostatic charge-dissipative materials, andsemiconductor layers in high voltage cables [3,4].

Achieving efficient dispersion of the MWCNTs within polyolefinmatrices remains a challenge, given that they are thermodynami-cally driven to aggregate by van der Waals forces and pep in-teractions. This has prevented polyolefin/MWCNT nanocompositesfrom reaching their full potential on an industrially relevant scale.Covalent modifications of MWCNTs are frequently used in order toimprove the weak interfacial adhesion [1]. A notable drawback ofcovalent functionalization is the disruption of the extended p

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conjugation in nanotubes, which has a profound effect on theirelectrical properties. Non-covalent functionalization of the carbonnanotubes (CNT), characterized by the existence of weak CHepand/or pep interactions, is an alternative method for tuning theinterfacial properties of nanotubes without compromising theirconductivity [5]. This strategy provides efficient dispersion ofMWCNTs in a polyolefin matrix, without compromising theirinherent conducting properties [6].

Polymer functionalization may further improve the compati-bility of the matrix with the functionalized carbon nanotubes, byestablishing interfacial interactions between the functional groupsof the compatibilizer and surface of the nanotubes. Maleic anhy-dride grafts [7,8] have been employed to improve the dispersion andproperties of MWCNT-containing composites with polypropyleneand polyethylene. Furthermore coupling agents, such graft or blockcopolymers, are commonly used to improve thewettingor adhesioncharacteristics of the nanotubes, thus reducing their tendency toagglomerate in CNT/polymer composites [7e9]. As an example,maleic anhydride grafted polystyrene-b-poly(ethylene-co-butylene)-b-styrene (SEBS-g-MA), a triblock copolymer with sty-rene blocks at both ends and an ethylene/butylene block in themiddle, has been used to promote interfacial adhesion and

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e5240 5231

dispersion of the MWCNTs in a polyolefin matrix. The two endscontaining aromatic rings exhibit compatibility with the MWCNTs,while the rubbery ethylene/butylene block is compatible with thepolyolefin matrix, such as polypropylene [7]. Thomassin et al. [9]functionalized a polypropylene matrix through the introduction ofaromatic moieties onto the polymer backbone, which are able tointeract by pep stacking with the surface of the carbon nanotubes.This approach avoids the need for covalent modification of thenanotubes, thus allowing the extended p conjugation in the nano-tubes, which is primarily responsible for their characteristic electricproperties, to remain intact. Furthermore, the aromaticmoieties caninteract with hydroxyl and carboxyl groups that may be present onthe surface of the nanotubes, due to their synthetic process [7,8].

In this studywe have applied two compatibilization strategies toestablish non-covalent interactions between the polyolefin matrixand the MWCNTs in composites prepared by melt compounding.The first method involves grafting aromatic pyridine groups ontothe backbone of a maleated linear low density polyethylene(LLDPE-g-MAn). The second approach employs the non-covalentfunctionalization of MWCNTs with a hyperbranched polyethylene(HBPE) [6,10]. The LLDPE-g-MAn matrix was chosen because of itsindustrial importance [11,12] and its amenability to functionaliza-tion reactions. The effect of matrix viscosity and compatibilizationon the microstructure, as well as on the electrical and mechanicalproperties of the composites is investigated in detail.

2. Experimental

2.1. Materials

MWCNTs (purity > 95%, diameter 30 � 15 nm and length 1e5 mm) were purchased from Nanolab Inc. (Massachusetts, USA) andused as received. The specific surface area (SSA) of the nanotubeswas 300 m2/g, as determined by BrunauereEmmetteTeller (BET)characterization. MWCNT samples weighing 0.5e1.0 g were firstdegassed at 110 �C for 24 h and then subjected to a multipoint BETphysisorption analysis (Autosorb-1 Quantachrome, USA) for nitro-gen relative vapor pressures in the range 0.1e0.3 at 77 K.Fusabond� E439 and E528, which are both maleic anhydride graf-ted LLDPEs (LLDPE-g-MAn) with a grafting degree of 0.5e1.0 wt.%and MFI (190 �C/2.16 kg) of 2.7 g/10 min and 6.7 g/10 min respec-tively, were supplied from E.I. DuPont Canada. These are denoted asLLDPE-g-MAn(H) and LLDPE-g-MAn(L), representing respectivelyhigh and low viscosity. 4-Aminomethylpyridine (AMP, 98% purity)was supplied from Aldrich. HBPE was synthesized from ethyleneusing a chainwalking Pdediimine catalyst, as described in detail byYe and Li [5]. HBPE has a complex and irregular dendritic structure,containing a large number of branches of various lengths (frommethyl to hexyl and higher) and extensive branch-on-branchstructures [5]. All other materials and solvents used were ofanalytical grade and were used without further purification.

2.2. Synthesis of LLDPE-graft-aminomethylpyridine (LLDPE-g-Py)

LLDPE-graft-aminomethylpyridine (LLDPE-g-Py) was synthesizedin a Haake Rheomix E3000 by adding a molar excess (compared tomaleic anhydride functions) of 4-aminomethylpyridine withinLLDPE-g-MAn at 190 �C for 10 min. The resulting polymer was solu-bilized in xylene at 120 �C and precipitated in methanol in order toremove the excess of amine bearing molecules. The grafting reactionwas confirmed by FTIR spectra obtained using a Vertex 70 FTIRspectrometer (Bruker Optics). Samples were thin films, prepared in ahydraulic hot press at 190 �C and 50 MPa. The resolution for eachspectrumwas 2 cm�1 and the number of co-added scans was 24. The

spectra presented were baseline corrected and converted to theabsorbance mode.

2.3. Functionalization of MWCNTs

Mixtures of HBPE and MWCNTs in tetrahydrofuran (THF) with amass ratio of 1.0 were prepared by adding HBPE into dispersionscontaining 2.0 mg MWCNT/mL of THF. The resulting mixtures weresonicated for 1 h and then stirred overnight. The supernatant so-lutions were vacuum filtered drop-wise through Cu-Formvar/Carbon film with pore size of 0.22 mm. After washing with THF,the modified nanotubes were dried in a vacuum oven overnight atroom temperature, according to the procedures described by Petrieet al. [10]. The amount of adsorbed HBPE on the MWCNTs wasabout 23 wt.%, as estimated by thermogravimetric analysis [10].

2.4. Melt compounding

LLDPE nanocomposites containing pristine or modifiedMWCNTs were prepared by melt compounding using a DSMResearch 5 mL Micro-Compounder (DSM Resolve, GeleenNetherlands), at a temperature of 190 �C, screw speed of 60 rpm,and a mixing time of 10 min. A masterbatch containing 20 wt.%MWCNTs was prepared first and was subsequently redispersed inthe respective polymeric matrices, to obtain compositions rangingfrom 0.5 to 12.5 wt. %. This corresponded to about 0.3 and 7.4 vol. %,respectively, calculated using Eq. (1):

f ¼ dm �Wf�dm � df

��Wf þ df

(1)

where f is the volume fraction, Wf is the filler’s weight fraction, dmis the density of LLDPE (0.92 g/cm3), and df is the average density ofthe MWCNTs (1.65 g/cm3). For the MWCNTs modified with HBPE,the weight of the neat MWCNTs was taken into account whenpreparing the composites, to ensure comparable results. Prior tomelt mixing, the nanotubes were dried by heating in a vacuumoven at 120 �C for 24 h. The two components were physicallypremixed before being fed in the micro-compounder. During themixing period, the melt temperature and torque were continuouslyrecorded. After preparation, the nanocomposites were placed in adesiccator to prevent moisture absorption prior to further testing.

2.5. Composite characterization

Scanning Electron Microscopy (SEM) was carried out using aJEOL JMS-840A scanning microscope equipped with an energy-dispersive X-ray (EDX) Oxford ISIS 300 micro-analytical system.Surfaces were fractured in liquid nitrogen and subsequently coatedwith gold. TEM images were obtained for ultra-thin films of thepolymer nanocomposites prepared using a Leica ultra-microtome.Images of the nanocomposites were obtained using a FEI Tecnai20 instrument (FEI Co., Eindhoven, Netherlands). The compositemelts were also observed using an Olympus BX 51 optical micro-scope (Tokyo, Japan). Nanocomposite films were loaded on aLinkam SCC 450 Hot Stage (Surrey, UK) at 160 �C and pressed to athickness of 25 mm. Images were recorded after 5 min, once thedesired thickness was reached, using transmitted light.

Raman studies were performed using a Jobin-Yvon/Horiba mi-cro-Raman Spectrometer (Model: LabRAM) equipped with a632 nm He/Ne laser source, 1800 1/nm grating and an OlympusBX41microscope system. The laser power at the samplewas kept at0.17 mW. Collection of the spectra was performed in the back-scattered mode with the use of a filter at room temperature underthe following conditions: �100 microscope objective, 100 mm

Scheme 1. Reaction between LLDPE-g-MAn and 4-aminomethylpyridine during meltmixing.

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e52405232

pinhole size, 300 mm slit width, and 30 min exposure time. Eachspectrum represents the average of two measurements.

Tensile properties were measured using an Instron 3369 uni-versal tester, at crosshead speeds of 50 mm min�1. Dumbbell-shaped specimens were cut with a type-V die according to ASTMD 638 from sheets with average thickness of 3.0 mm, which wereprepared by compression moulding of the compounded samples at190 �C and 10 MPa using a Carver press, which were prepared bycompression moulding of the compounded samples at 190 �C and10 MPa using a Carver press for 5 min.

Thermal analysis of the nanocomposites was performed using adifferential scanning calorimeter (TA Instruments Q100 Series DSC)calibratedwith Indium and Zinc standards. For eachmeasurement, asample of about 5 � 0.1 mg was placed in a sealed aluminium pan,and heated to 200 �C at a scanning rate of 20 C min�1. From thesescans the melting temperature (Tm) and the heat of fusion (DHm) ofthe nanocomposites weremeasured. The crystallinity of the sampleswas calculated using the DHm values with a heat of fusion of 293 J/gfor fully crystalline polyethylene [3], after normalization of thenanocomposites’ DHm to the proper LLDPE amount contained. Thecrystallization temperature (Tc) of the melt was also recorded bycooling the samples from200 to�30 �Cata cooling rateof 10Cmin�1.

Rheological characterization was carried out on a stress-controlled Reologica ViscoTech oscillatory rheometer using 20 mmparallel plate fixtures, with a gap of 1 mm at 190 �C. Compressionmoulded disks with a diameter of 20 mmwere prepared using theCarver press as described above. Stress sweep experiments werecarried out from 1 to 103 Pa at a frequency of 0.1 Hz and temper-ature of 190 �C to identify the limits of linear viscoelasticity. Timesweeps, conducted at 0.1 Hz and 1% strain revealed that the sam-ples were stable for times up to 30 min, except for the LLDPE-g-MAn(L)/MWCNT samples containing higher than 5 wt% MWCNT,which showed some time dependency. The rheometer was oper-ated in the dynamic oscillatory mode within the linear viscoelas-ticity region at a frequency range of 0.01e25.1 Hz. The complexviscosity (h*), elastic modulus (G0) and tan d were measured as afunction of angular frequency (u).

Volume resistivity was measured under DC current at roomtemperature. Samples were prepared by compressionmoulding themelt-compounded composites in a Carver press at 190 �C and a forceof 10MPa to get a thinfilm of 0.8mm. The thin compositefilmwith adiameter of 6 cm was put into the measuring chamber (Keithley8009 Resistivity Test Fixture) of the Keithley 6517B Electrometer/High Resistance Meter (Keithley Instruments, Inc., Cleveland, Ohio,USA) for an electrification time of 1 min. For more conductivesamples (lower than 107 U cm), an Agilent 34401A 6½ Digit multi-meterwas used tomeasure resistance. Sample preparation includedcutting test specimens with a dimension of 1 cm2 and thickness of0.04 cm and gold sputtering the surface to reduce contact resistancebetween the sample and the electrodes of themultimeter. The edgesof the sampleswere trimmed after gold sputtering to prevent short-circuit during themeasurement of the resistance. The resistivitywasthen calculated based on the standard relation:

r ¼ R� Al

(2)

where r is the resistance in Ohms, A the contact surface area in cm2

and l the average sample thickness in cm.

3. Results

3.1. Grafting reaction of LLDPE-g-MAn with aminopyridine

The reaction between the succinic anhydride groups and theamino functional group proceeds through a nucleophilic attack of

the amino group to the anhydride moiety, shown in Scheme 1. Thefirst stage of the reaction leads to the formation of an acid-amide.From this stage the reaction can proceed in two ways: either thefree carboxylic acid further reacts with another amine moiety,resulting in a diamide, or there is further condensation between thecarboxylic group and amide formed, leading to the formation of animide, with the dehydration water being eliminated during theprocess due to the high temperature [13]. Whether the first route orthe second is followed is determined by the reaction parameters,such as temperature, reaction time, amine concentration, etc.Higher temperatures and reaction times favour the formation of thecyclic imide. At the high reaction temperature used in the presentwork (190 �C) it is expected that the imide is formed exclusively,leading to complete formation of LLDPE grafted with 1-(pyridin-4-ylmethyl)-1H-pyrrole-2,5-dione (Py) groups (LLDPE-g-Py).

The reaction between the maleic anhydride groups (MAn) and4-aminomethylpyridine was confirmed by the FTIR spectra shownin Fig. 1. LLDPE-g-MAn exhibits absorptions at 1866 cm-1 (asym-metric stretching of the carbonyl groups of maleic anhydride,weak), at 1791 cm-1 (symmetric stretching of the carbonyl groups ofmaleic anhydride, strong), and a small one at around 1712 cm-1, dueto the symmetric stretching of the carbonyl groups of the carboxylicacid groups formed by a very small portion of hydrolysed maleicanhydride groups. The presence of imide functional groups inLLDPE-g-Py is evidenced by the characteristic absorption bands ofthe imide ring at 1777 cm-1 (asymmetric stretching of the carbonylgroups of the imide ring, weak) and 1714 cm-1 (symmetricstretching of the carbonyl groups of the imide ring, strong) [14].

The functionalization with pyridine groups did not affectsignificantly the rheological properties of the polymers. Fig. 2shows that the complex viscosities of the maleic anhydride andpyridine-grafted (Py) LLDPEs are very similar. The slight decrease incomplex viscosity observed in the LLDPE-g-Py (H) may be attrib-uted to the increase of free space due to the introduction of the

Fig. 1. FTIR spectra of LLDPE-g-MAn(H) and LLDPE-g-Py(H).

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e5240 5233

bulky aromatic pyridine groups and/or due to limited degradationof the matrix during processing. Thus, we do not expect any sub-stantial viscosity-related effect on the dispersion of the nanotubesstemming from the functionalization.

3.2. Effect of polymer functionalization and viscosity onnanocomposite properties

3.2.1. Raman spectroscopyFunctionalization with Py aims at establishing non-covalent in-

teractions between the aromatic moieties and the surface of theMWCNTs by pep stacking. The interactions between the nanotubesand the aromatic side groups in the compounded nanocompositescan be evidenced using Raman spectroscopy. Representative highfrequency Raman spectra of the neat MWCNTs and compositescontaining 5 wt. % MWCNTs are presented in Fig. 3. MWCNTs havetwo characteristic peaks: The first band around 1590 cm�1 isassigned to the tangential mode (G-band) and is associatedwith the

Fig. 2. Complex viscosity as a function of frequency of LLDPE-g-MAn and LLDPE-g-Pyat 190 �C.

ordered sp2-hybridized carbon network. The band located around1340 cm�1 is the disorder-induced D-band, which is related to localdefects that originate from the structural imperfections and asso-ciated with the disordered sp3-hybridized carbon network [15].

In the LLDPE-g-MAn/MWCNT composite, the G-band waslocated at 1587 cm�1, while the D-band was centered at 1336 cm�1.Generally the shifting of the G- and D- band peaks to higherwavenumbers in MWCNTs containing composites is attributed tothe polymer penetration into the MWCNT bundles during meltmixing, resulting in their disentanglement, and to the generation ofstrong compressive forces between the polymer chains andMWCNTs [16]. In the present work, further shifts of 20 cm�1 and9 cm�1 are noted in the G-band and D-band peak, respectively, inthe presence of the pyridine side-group. This provides evidence ofnon-covalent interfacial interactions between the polymer matrixand the MWCNTs.]

Stronger interfacial interactions between the polymer and thematrix are expected to result in more efficient stress transfer to theMWCNTs during compounding, thus breaking up more effectivelythe filler aggregates. Stress transfer is also strongly influenced bythe viscosity of the matrix [17,18]. Both factors are investigated indetail below.

3.2.2. MWCNT dispersionThe effects of matrix viscosity and compatibilization on the

MWCNT dispersion within the LLDPE matrices are clearly seen inthe optical microscope (OM) images presented in Fig. 4. The dis-tribution of MWCNTs in the composites based on the high viscositypolymer matrix (Fig. 4a and c) appears homogeneous, whereasmicron-size aggregates, and areas devoid of nanofiller, are seen inthe low viscosity samples (Fig. 4b and d). This is due to the lowershear stresses generated during compounding, which are not suf-ficient for efficient break-up of the aggregates, thus resulting inworse dispersion [18].

The effect of viscosity on dispersion is also depicted in the TEMimages (Fig. 5a and b). The LLDPE-g-MAn(H) matrix (Fig. 5a) con-tains individually dispersed nanotubes, together with some resid-ual aggregates that are around 500 nm in size, whereas extensiveaggregation is present in the LLDPE-g-MAn(L) matrix (Fig. 5b). Itshould be noted that in all cases the MWCNTs are generallyrandomly distributed throughout the matrix, with no preferential

Fig. 3. Raman spectra of high viscosity LLDPE-g-MAn and LLDPE-g-Py samples con-taining 5.0 wt.% MWCNTs and pure MWCNTs.

Fig. 4. Representative OM images of (a) LLDPE-g-MAn(H), (b) LLDPE-g-MAn(L) (c) LLDPE-g-Py(H) and (d) LLDPE-g-Py(L) with 3.0 wt.% MWCNTs samples.

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e52405234

orientation. Moreover the nanotubes appear worm-like, with asmaller “effective aspect ratio” than if they were fully elongated.This microstructure is attributed to the preparation method,involving compression moulding.

Even though in the present work it is hard to quantify the lengthof the MWCNTs because of the difficulty in isolating the nanotubesby dissolving the polyolefin matrix without affecting their proper-ties, a variation of the MWCNT length is obvious when comparingsamples based on different matrices. Inspection of the magnifiedTEM images (Fig. 5e and f) reveals that the MWCNTs appear to belonger andmore entangled in the unmodified lowviscosity samples(Fig. 5e), compared to the high viscosity samples (Fig. 5f). The effectsof processing on MWCNT breakage in the presence of high shearforces and their consequence on the nanocomposite properties arewell documented [18]. Higher matrix viscosity generates highershear forces during compounding, which may result in a higherincidence of nanotube breakage, as suggested by the TEM images.The reduced aspect ratio may affect the conductivity and physicalproperties of the composites, as discussed in the following sections.

The effectiveness of the Py functionality in further improvingMWCNT dispersion can be assessed by examining the OM and TEMimages. For both matrices, a reduction in the size and number ofaggregates is observed in the presence of the pyridine side group(Fig. 4c and d). The TEM images confirmed that the enhancedinterfacial interactions between LLDPE-g-Py and MWCNTs resultedin improved MWCNT dispersion, as evidenced by the presence of agreater amount of isolated nanotubes inside the matrix (Fig. 5c andd), compared to the LLDPE-g-MAn counterparts. This effect can bealso seen in the SEM images of Fig. 6, which demonstrate thereduction of the micron-size agglomerates to aggregates that areabout 500 nm in size.

From these results it is evident that higher viscosity and func-tionalization with Py result in improved MWCNT dispersion. Thedifferences in the microstructure of the composites affect theirproperties profoundly, as discussed in the following sections.

3.2.3. Electrical and rheological percolation thresholdAs shown in Fig. 7, the addition of MWCNTs to the LLDPE matrix

resulted in significant increases in conductivity, by almost 13 orders

of magnitude compared to the unfilled polymer. The maximumconductivities achievedwere in the order of 10�4 Sm�1, in the rangeof static dissipative materials [19,20]. Even though the maximumconductivity was the same for all composites, the percolationthreshold, defined as the value of the filler content above which theelectrical properties increase in an exponential manner [21], variedsignificantly depending on the type of the matrix.

The percolation threshold concentration for each composite wasestimated by fitting the conductivity data, using power-law re-lations above and below the critical concentration for percolation,4c [18,22]:

s ¼ smatrix

�4c � 4

4c

��s

4 < 4c (3a)

s ¼ m�4� 4c1� 4c

�t

zmð4� 4cÞt 4 > 4c (3b)

where s is the electrical conductivity of the composite, smatrix is theconductivity of the matrix, 4 is volume fraction of the filler, 4c iscritical volume fraction at percolation, s and t are the critical ex-ponents below and above percolation respectively and m is a con-stant. The best fits obtained by fitting Eqs. (3a) and (3b) to theexperimental data are summarized in Table 1. The values ofpercolation threshold observed were generally higher than thosereported in the literature for polyolefin/MWCNTcomposites, whichare typically around 2.5 wt.% [16,23].

The data summarized in Table 1 show that the electricalpercolation threshold is the lowest in the LLDPE-g-MAn matrix,which exhibited the worse dispersion (Fig. 5b) and the largestnumber of aggregates (Fig. 6a). Functionalization with Py resultedin an increase in the percolation threshold in all matrices. The effectof functionalization is relatively small in the low viscosity samplesas it shifts the electrical percolation by only 0.7 wt.%, but it is verydramatic in the high viscosity system, which experiences an almosttripling of the percolation threshold, from 3.2 to 8.6 wt.% in theLLDPE-g-Py(H)/MWCNT composites; these exhibited the bestMWCNT dispersion (Fig. 5c). The percolation threshold for the

Fig. 5. TEM images of (a) LLDPE-g-MAn(H), (b) LLDPE-g-MAn(L), (c) LLDPE-g-Py(H) and (d) LLDPE-g-Py(L) containing 5 wt. % (2.9 vol. %) MWCNT. (e) and (f) are higher magni-fications of the areas marked in (a) and (b), respectively.

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e5240 5235

transition from insulating to conductive materials is associatedwith the formation of conductive pathways across the material dueto the presence of interconnected nanotubes [21] and consequentlyit is influenced by the microstructure of the composites. Improveddispersion in the Py-functionalized composites results in the

Fig. 6. SEM images of cryo-fractured surfaces of (a) LLDPE-g-MAn(L) and (b) LLDPE-

presence of individual nanotubes, with smaller interparticle dis-tances. In spite of the smaller interparticle distance, these well-dispersed nanotubes exist as individual particles, without inter-connectivity; thus a larger amount of nanotubes is needed to ach-ieve percolation, as evidenced by the results shown in Table 1.

g-Py(L) with 5 wt.% MWCNTs (representative nanotube aggregates are circled).

Fig. 7. Volume conductivity as a function of MWCNT concentration. Solid lines denotefits obtained by using Eqs. (3a) and (3b). Fig. 8. Reduced storage modulus vs. MWCNT concentration at 0.1 rad/s and 190 �C.

Solid lines denote fits obtained by using Eqs. (5a) and (5b).

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e52405236

The differences in dispersion are also reflected on the rheolog-ical properties of the composites. Fig. 8 depicts the compositiondependence of the reduced storage modulus, G0

r, defined as,

G0r ¼ G0

G00

(4)

where G00 is the storage modulus of the matrix and G0 is the

modulus of the composite. Both are evaluated at a fixed frequencyof 0.1 rad/s.

At low loadings, the MWCNTs act as isolated objects and thecomposition dependence of the reduced storage modulus is rela-tivelyweak, without significant differences in the ratio between thelow and high viscosity composites. In all composites, substantialincreases in the storage modulus are noted when the MWCNTcontent increases above a certain loading, consistent with thepresence of a rheological (geometrical) percolation threshold.

In order to quantify the rheological percolation threshold, fg,equations of a similar form as the ones used for the electricalpercolation thresholds can be used, by replacing conductivity pa-rameters with the corresponding reduced storage modulus ones:

G0r ¼

4g � 4

4g

!�a

4 < 4g (5a)

G0rzn

�4� 4g

�b4 > 4g (5b)

Table 1Electrical and geometrical percolation thresholds obtained by fitting Eqs. (3) and (5)on the experimental data of conductivity and storage modulus.

Sample Electricalpercolation

Geometricalpercolation

fc

(vol.%)wc

(wt.%)fg

(vol. %)wg

(wt.%)

LLDPE-g-MAn(H) þ MWCNT 1.8 3.2 1.7 3.0LLDPE-g-Py(H) þ MWCNT 5.0 8.6 4.7 8.2LLDPE-g-MAn(L) þ MWCNT 1.5 2.7 1.3 2.3LLDPE-g-Py(L) þ MWCNT 1.9 3.4 1.9 3.3LLDPE-g-MAn(H) þ (HBPE-MWCNT) 4.0 7.0 3.7 6.5

The fitting of the equations to the data is shown in Fig. 8 and theresulting geometrical percolation thresholds obtained are sum-marized in Table 1. The LLDPE-g-MAn(H), LLDPEg-MAn(L) andLLDPE-g-Py(L) matrices exhibit percolations between 2.3 and3.3 wt%. A comparison with the respective TEM images (Fig. 5a, band d) reveals that these composites contain large numbers ofaggregates, approximately 500 nm in size, that are interconnectedwith individual entangled MWCNTs.

Furthermore the comparison between the LLDPE-g-MAn(H) andLLDPEg-MAn(L)matrices (Fig. 8) reveals that the formerhas a higherpercolation threshold, presumably due to the smaller aggregate sizeand better dispersed MWCNTs. The smaller length of the MWCNTsmay also contribute to this, as it offers less chance for entangle-ments. On the other hand, the LLDPE-g-Py(H)/5 wt.% MWCNTcomposite (Fig. 6c) is still below the percolation threshold, giventhat the nanotubes are finely dispersed and not interconnected.

As shown in Table 1, in all cases the rheological percolationthresholds are lower than the electrical ones. Generally the rheo-logical percolation threshold is attributed to interactions betweenthe filler particles [24]. The filler structure restrains the long-rangemotion of the polymer chains when the average distance betweenthe nanotubes reaches a critical value between the polymer’sentanglement distance and twice the radius of gyration, namely,from 10 to 100 nm [25]. Geometrical percolation may therefore betaking place at a slightly lower concentration, because it does notrequire direct contact between nanotubes, contrary to electricalpercolation which needs direct contact or proximity of the nano-tubes, to allow for electron hopping across the nanotubes.

The effect of flocculation in the low viscosity samples containingamounts of nanotubes higher than 5 wt. % during rheologicalcharacterization (see Fig. S1 in supplementary information) mayalso be a reason for the slightly lower geometrical percolation timesthatwere noted in the lowviscosity samples, since it would result inincreasing values of the elastic modulus during measurement [26].

3.2.4. Thermal and mechanical propertiesGenerally the inclusion of conductive fillers in a polymer matrix

has a reinforcing and stiffening effect, which is usually counter-acted by a significant loss in ductility [27]. One motivation forfunctionalization is to overcome this problem by improving the

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e5240 5237

interfacial adhesion and compatibility between polymer and filler[28]. The improved dispersion afforded by this approach avoids theformation of aggregates, which act as stress concentrators and ul-timate failure points during a tensile test. The effectiveness of ourcompatibilization approach on the tensile properties of the com-posites is investigated in this section.

As expected, substantial improvements in the Young’s modulusare observed upon introduction of the MWCNT (Fig. 9a). Thecompatibilized samples exhibited higher moduli than their non-compatibilized counterparts at filler loadings above 1 wt.%. Thehighest value was observed for LLDPE-g-Py(H) containing 5.0 wt.%MWCNTs, which had a 62% improvement in Young’s moduluscompared to the pure polymer, even though the unfilled LLDPE-g-Py sample had a lower modulus compared to LLDPE-g-MAn. Thesechanges are attributed solely to the presence of the filler, as therewas no significant effect of the MWCNTs on the crystallinity andthermal properties of the polymer. According to the data presentedin Table 2, there is only a slight increase in the crystallizationtemperature, Tc, and a minor reduction in the melting temperature,Tm, attributed to the nucleation effect of the nanotubes and a smallpossible effect on the crystalline structure, respectively.

Significant modulus improvements are also noted in the lowviscosity samples, with a 44% improvement in the modulus inLLDPE-g-Py/MWCNT 5 wt.% composite compared to the unfilledpolymer. A more noticeable increase in crystallinity, as shown inTable 2, may be a contributing factor to this change. The tensilestrength at break follows different trends, depending on the type ofmatrix (Fig. 9b). Increases in tensile strength at concentrations ofMWCNT up to 2 wt.% can be seen when the LLDPE-g-MAn(H) ma-trix is used, followed by a decrease at concentrations beyond the

Fig. 9. (a) Young’s modulus, (b) stress at the break point and (c) el

percolation threshold. This is associated with the presence ofMWCNT aggregates at high contents, which function as stressconcentrators. A continuous improvement in strength is evident inthe LLDPE-g-Py(H)/MWCNT composites, reaching to a maximum of22% increase compared to the neat polymer. This is attributed to thefiner dispersion and the improved interfacial interactions betweenpolymer and nanoparticles. On the contrary, a decreasing trend isnoted upon introduction of MWCNT in LLDPE-g-MAn(L), due to theextensive MWCNT aggregation, whereas the decrease is less pro-nounced in the respective Py-derivative, which had improveddispersion.

Filler aggregation resulted in a reduction of the elongation atbreak. This was more pronounced at higher concentrations wherethe size and number of the aggregates became larger. However, dueto the better dispersion of the MWCNT and the smaller number ofaggregates, this decrease was less significant in the compatibilizedsamples, which retained better elongation at break compared totheir unmodified counterparts. While the elongation at the breakpoint for the LLDPE-g-MAn(H)/5 wt.% MWCNT composite wasdecreased by 50%, this reduction was only 23% when using thepyridine derivative. In spite of the reduction, this compositemaintains values of the elongation at break of above 900% for allcompositions, which make it suitable for engineering applicationsthat require ductility. In the low viscosity samples, due to thepresence of larger nanotube aggregates, the decrease was morepronounced, reaching 65% in the case of the LLDPE-g-MAn(L)/5 wt.% MWCNT composite. However, the compatibilization signif-icantly constrained this deterioration, limiting it to about 40%whenthe pyridine derivative was used. It is therefore obvious that effi-cient dispersion, which results in higher interfacial area between

ongation at the break point of the LLDPE/MWCNT composites.

Table 2Thermal properties of LLDPE/MWCNT composites.

Sample Tm (�C) Tc (�C) Xm (%) Sample Tm (�C) Tc (�C) Xm (%)

LLDPE-g-MAn(H) 117.9 97.1 31.7 LLDPE-g-MAn(L) 118.7 98.2 37.3þ 1.0 wt.% MWCNT 117.4 97.5 32.3 þ 1.0 wt.% MWCNT 118.4 98.8 39.2þ 3.0 wt.% MWCNT 115.8 99.1 33.3 þ 3.0 wt.% MWCNT 117.4 100.0 44.5þ 5.0 wt.% MWCNT 115.0 101.5 33.6 þ 5.0 wt.% MWCNT 116.9 99.2 38.8

LLDPE-g- Py(H) 115.2 95.7 29.5 LLDPE-g- Py(L) 119.1 99.7 35.4þ 1.0 wt.% MWCNT 114.7 98.2 30.6 þ 1.0 wt.% MWCNT 118.7 100.1 37.6þ 3.0 wt.% MWCNT 112.8 101.1 31.6 þ 3.0 wt.% MWCNT 117.2 102.2 44.1þ 5.0 wt.% MWCNT 111.5 102.8 31.9 þ 5.0 wt.% MWCNT 116.0 101.8 41.6

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the filler andmatrix [21] is the determining factor for improvementin mechanical properties seen in the Py functionalized composites.

3.3. Effect of non-covalent functionalization of MWCNTs onnanocomposite properties

Given that Py functionalization compromised significantly theconductivity of the high viscosity material and led to an increase inthe percolation threshold due to the improved dispersion, as analternative we functionalized the MWCNTs using a hyperbranchedpolyethylene (HBPE). Given that the surface coverage of nanotubesby HBPE is only about 30% the inherent conducting properties ofthe nanotubes are not compromised by this modification [6,10].Coating the MWCNTs with HBPE serves to disrupt the fillerefillerinteractions, which consist mainly of van der Waals forces betweennanotubes [6,10]. HBPE would also act to prevent secondary ag-gregation during compounding, given that the interactions be-tween nanotubes become weaker, thus reducing propensity foraggregation. Observation of the macro- and nano-dispersion of thecomposites (Fig. 10a and b respectively) revealed a drop in size andnumber of MWCNT agglomerates and a better dispersion overallwith the use of HBPE, compared to the respective compositespresented in Figs. 4a and 6a.

However, the effect on percolation threshold was less pro-nounced compared to the introduction of the aromatic moietiesonto the polymer’s backbone, presumably because of the less effi-cient dispersion. As seen in Figs. 7 and 8, Table 1 and Fig. 11, theHBPE-modified composites had properties that were in-betweenthe corresponding ones based on LLDPE-g-MAn(H) and LLDPE-g-Py(H). For example, the electrical percolation threshold was 7.0 wt.% (compared to 3.2 wt.% and 8.6 wt.% for LLDPE-g-MAn(H)/MWCNTand LLDPE-g-Py(H)/MWCNT, respectively).

The same “in-between” behaviour can also be seen in the me-chanical properties of the composites with HBPE/MWCNTs (Fig. 11).Furthermore, the modification of the MWCNTs with HBPE did not

Fig. 10. Representative (a) TEM of LLDPE-g-MAn(H)/5.0 wt.% (2.9 vol%) HBPE-M

result in significant variations of the thermal properties of thecorresponding composites.

4. Discussion

It is well-known in the literature that carbon nanotube disper-sion is themain factor influencing the conductivity andproperties ofpolymer/MWCNTcomposites. In recent years it has been recognizedthat a certain level of nanotube aggregation is required to achieveconductivity [22,29e31]; however this comes at the expense ofmechanical properties. Therefore, a fine balance exists between therequirement for improved conductivity and the need to maintainphysical properties that are sufficient for engineering applications.

Nanofiller dispersion can be controlled mainly by mechanicalmeans, i.e. compounding procedures, intensity of mixing, stresstransfer due to matrix viscosity, and chemical means, including thevarious methods of compatibilization and functionalization of thepolymer and/or matrix. Obviously, depending on the target use ofthe composites, nanofiller dispersion must be controlled in order toachieve the desirable balance between conductive and mechanicalproperties; this will dictate the choice of approach for the prepa-ration procedure of the nanocomposites.

In this work, we examined both techniques: The first employedtwo compatibilization strategies, involving modification of thematrix and the MWCNTs, aiming at the least disruption of theinherent conducting properties of the nanotubes. The secondinvolved the dispersion of the nanotubes by mechanical means,through manipulation of the viscosity of the matrix.

Both compatibilization strategies resulted in enhanced disper-sion of the nanotubes obtained by melt compounding, due to theincreased interfacial interactions. This resulted in significantlyimproved mechanical properties compared to the MAn-functionalized composites. However the electrical and geomet-rical percolation thresholds increased. This confirms that somedegree of aggregation is desirable, because it facilitates the

WCNTs and (b) OM images of LLDPE-g-MAn(H)/3.0 wt.% (HBPE-MWCNTs).

Fig. 11. (a) Young’s modulus and (b) stress at the break point of the prepared nanocomposites with HBPE/MWCNTs compared to samples with unmodified MWCNTs.

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formation of sample-spanning continuous conductive pathwayswithin the polymer matrix, as depicted by the schematic of Fig. 12a,which contains a conceptual illustration of aggregates inter-connected by individual entangled nanotubes. On the other handthere is no continuous network in composites that contain thesame amount of well-dispersed nanotubes (Fig. 12b); these com-posites would therefore require significantly higher amounts offiller to achieve the same effect. Indeed the percolation thresholdincreased more than two-fold in the case of the high viscositymatrices, which exhibited the best dispersion.

The degree of aggregation can be further manipulated throughthe choice of matrix viscosity [18] (or potentially compoundingtemperature, which would also influence the viscosity). This pro-vides the interesting prospect of utilizing the viscosity of thematrixas an additional variable to manipulate dispersion, while simulta-neously ripping the benefits of better interfacial interactions andimproved mechanical properties that the compatibilizations canafford, without suffering significant losses of conductivity. Forexample, high Young’s modulus can be achieved by LLDPE-g-Py(L)/MWCNTcomposites. These composites still maintain elongations atbreak above 600%, while the electrical percolation threshold is only3 wt.%. Additionally, non-covalent functionalization of theMWCNTs can be exploited, because it provides a balance betweendispersion and mechanical properties. The results presented inTable 1 and Figs. 9 and 11 provide therefore some guidelines for the

Fig. 12. (a) Network formation consisting of MWCNT aggregates (circled in red) interconnectmatrix. (For interpretation of the references to color in this figure legend, the reader is refe

selection of matrix viscosity and compatibilization method,depending on the desired end-product use.

Processing by melt compounding may pose an inherent limi-tation in the level of conductivity that can be achieved withinpolyolefins. Mechanical forces during themelt mixing process, suchas shear and normal forces induced by the rotors, can causebreakage of the nanotubes. In this work, TEM imaging suggests thathigher viscosity of the matrix polymer resulted in more pro-nounced shortening of the nanotubes, presumably because of theincrease in the shear forces exerted by the higher viscosity matrix.Furthermore, increasing the amount of nanotubes leads toincreased viscosity of the composites, which would result ineven higher shear forces during mixing and, therefore shorternanotubes [32].

Given that the percolation threshold depends on the aspectratio, a considerable reduction of the aspect ratio of the nanotubeswill lead to higher percolation threshold concentrations [18,32,33].Breakage is expected to happen along structural defects present onthe walls of the nanotubes, when shear forces exceed the strengthat the locations of the defects. The length reduction during prep-aration of the polymer/MWCNT composites is a significant issuethat needs to be overcome if the aim is to improve the performanceof these materials. Nanotubes of high quality and purity, containingless amorphous material, would be desirable to mitigate this effect.Mild mixing conditions, involving shorter mixing times and lower

ed with individual MWCNTs and (b) randomly distributed MWCNTs within the polymerrred to the web version of this article.)

A.A. Vasileiou et al. / Polymer 54 (2013) 5230e52405240

rotation speeds, may also be more favourable and may contributetowards achieving composites with a better balance of properties.

5. Conclusions

Reactive compounding of a maleated LLDPE with 4-amino-methylpyridine resulted in the grafting of pyridine (Py) aromaticmoieties on the polymer backbone, which are able to interact byp�p stacking with the surface of MWCNTs. The presence of non-covalent interactions between the Py grafts and MWCNT inLLDPE/MWCNT composites prepared by melt compounding wasevidenced by Raman spectroscopy.

The Py-functionalized composites exhibited improved disper-sion and consequently better mechanical properties compared tothe non-compatibilized counterparts. These improvements werecounteracted by increases in the rheological and electrical perco-lation threshold concentration. Similar results were obtained byimplementing non-covalent surface functionalization of theMWCNT using a hyperbranched polyethylene, which serves tointerrupt filler/filler interactions, therefore preventing aggregation.

On the contrary, in the non-compatibilized samples, increasedphase segregation of the filler driven by intrinsic incompatibilityled to percolation at smaller concentrations of nanotubes and anensuing deterioration in the mechanical properties.

High matrix viscosity favoured dispersion, because of increasedstress transfer to the nanotube aggregates during compounding.Consequently the compatibilized composites based on a high vis-cosity matrix exhibited the highest electrical and geometricalpercolation thresholds.

The results of the present work suggest that non-covalent in-teractions, which serve to improve the interfacial compatibilitybetween polymer and MWCNT, and low viscosity matrices thatallow for some degree of secondary aggregation, can provide a goodcombination of electrical and mechanical properties, suitable fortargeted engineering applications.

Acknowledgements

Financial support from the Natural Sciences and EngineeringResearch Council of Canada (NSERC) through the Discovery andAccelerator Supplement programs is gratefully acknowledged. Thepolymers used in this study were donated by E.I. DuPont Canada.

Appendix A. Supplementary material

Supplementary material related to this article can be found athttp://dx.doi.org/10.1016/j.polymer.2013.07.034.

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