EFFECT OF GRAPHITE AND NbC ON MECHANICAL ... - USM

EFFECT OF GRAPHITE AND NbC ON MECHANICAL

PROPERTIES OF AISI304 BINDED WC

TRAN BAO TRUNG

UNIVERSITI SAINS MALAYSIA

2013

EFFECT OF GRAPHITE AND NbC ON MECHANICAL

PROPERTIES OF AISI304 BINDED WC

by

TRAN BAO TRUNG

Thesis submitted in fulfilment of the

requirements for the degree

of Doctor of Philosophy

May 2013

ii

ACKNOWLEDGEMENTS

I would like to express my most sincere thanks that come deeply from my

heart to those who in one way or another contributed to make this research study

possible.

Firstly, I am deeply grateful to my main supervisor, Assoc. Prof. Dr.

Zuhailawati Hussain for her supervision and advice from the early to the final stage

of this research work. She has always been patient and encouraging in times of new

ideas and difficulties. She has listened to my ideas and discussed frequently with me

which led to the key insights of my work. She is a true scientist and a dedicated

teacher that I want to be. Above all, she made me feel a friend, which I appreciate

from my heart.

I would like to express my sincere gratitude to my co-supervisor Prof. Zainal

Arifin Ahmad for his helps in all the time of this research. I have never forgotten his

supports, for his patience, motivation, enthusiasm, and immense knowledge. He is a

great teacher that I have met. I keep in my mind his impressive saying to me “Take a

seat and talk to me as a friend”.

I would like to deliver my thanks to my Japanese advisor Prof. Ishihara N.

Keiichi, the School of Energy Science, Kyoto University, for his guidance and advice

on my research. Special thanks to Prof. Hideyuki Okumura and Dr. Eiji Yamasue for

their supports during the time of my study in Japan.

I am grateful to Professor Dr. Hanafi Ismail as the Dean and the staffs of the

School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia,

for their kindness and support.

This work would not have been possible without the financial support and

cares from AUN-SEED Net/JICA. I would like to deliver my sincere thanks and

iii

deepest gratitude for their generosity in giving me this opportunity to pursue my

Ph.D.’s degree.

My special thanks and appreciation are also extended to those people who, in

one way or another, helped me accomplish this research:

To Mdm. Fong, Mr. Kemuridan, Mr Shahid, Mr. Farid, Mr. Khairi, Mr.

Rashid, Mr. Zaini, Mr. Fadzil, and Mr. Fujimoto Shoji for their kindness, help, and

assistance.

To all my friends: Duong, Long, Bang, Viet, Mahani, Anny, Nini, Sunisa,

Zahir, Macara, Endo, Siba, Luong, Luyen, etc.

Last but not the least, I would like to take this opportunity to express my

gratitude to my family for their love.

Thank you to all of you!

iv

TABLE OF CONTENTS

Page

ACKNOWLEDGEMENTS ii

TABLE OF CONTENTS iv

LIST OF TABLES x

LIST OF FIGURES xi

LIST OF ABBREVIATIONS xvi

LIST OF SYMBOL xvii

ABSTRAK xviii

ABSTRACT xix

CHAPTER 1: INTRODUCTION

1.1 Introduction 1

1.2 Problem statements 4

1.3 Objectives of study 6

1.4 Research scope 7

CHAPTER 2: LITERATURE REVIEW

2.1 History of tungsten carbide and hard metals 10

2.2 Binder phase 13

2.2.1 Cobalt binder 14

2.2.2 Nickel binder 16

2.2.3 Iron binder 18

2.2.4 Ni-Fe and Co-Ni-Fe binders 19

v

2.2.5 Fe-Cr-Ni binders and stainless steel binder 20

2.3 - Phase 21

2.4 Graphite 24

2.5 Consolidation tungsten carbide hard metal powders 25

2.5.1 Green consolidation 25

2.5.2 Sintering process 27

2.5.2.1 Solid state sintering 27

2.5.2.2 Liquid phase sintering (LPS) 33

2.5.3 Grain growth 36

2.5.4 Grain growth inhibitors 38

2.6 Sintering method 42

2.6.1 Vacuum sintering 43

2.6.2 Hot pressing 43

2.6.3 Hot isostatic pressing 44

2.6.4 Pseudo hot isostatic pressing 44

2.7 Mechanical properties of WC-hard metals 45

2.7.1 Hardness of WC-hardmetals 46

2.7.2 Fracture toughness of WC-hardmetals 49

2.7.3 Other mechanical properties of WC-based hardmetals 52

2.8 Mechanical alloying 53

2.8.1 Introduction of mechanical alloying 54

2.8.2 Mechanical alloying mechanisms 55

2.9 Summary 57

vi

CHAPTER 3: RAW MATERIALS AND METHODOLOGY

3.1 Raw Materials 59

3.1.1 Tungsten carbide (WC) powder 59

3.1.2 Stainless steel powder 60

3.1.3 Niobium carbide (NbC) powder 60

3.1.4 Graphite powder 61

3.2 Research methodology 61

3.2.1 Mixing of WC-FeCrNi hardmetal powders by planetary

ball milling

63

3.2.2 Green body compaction 64

3.2.3 Sintering in a vacuum furnace 64

3.2.4 Sintering by PHIP process 65

3.3 Data analysis 67

3.3.1 Phase identification 67

3.3.2 Microstructure observation 68

3.3.3 Density measurement 69

3.3.4 Hardness testing 70

3.3.5 Fracture toughness measurement 71

CHAPTER 4: RESULTS AND DISCUSSION

4.1 Raw material analysis 73

4.1.1 Tungsten carbide (WC) powder 73

4.1.2 FeCrNi powder 74

4.1.3 Graphite (Cgr) powder 76

vii

4.1.4 Niobium carbide (NbC) powder 78

4.2 Effect of milling time on microstructure and mechanical properties

of WC-10FeCrNi hardmetals

79

4.2.1 Phase identification and microstructure of as-milled

powder

80

4.2.2 Phase identification and microstructure of sintered samples 83


4.2.4 Hardness and fracture toughness 88

4.3 Effect of sintering temperature on microstructure and mechanical

properties of WC-10FeCrNi hardmetals

90


4.3.2 Density measurement of sintered samples 93


4.4 Effect of sintering time on the microstructure and mechanical


96


4.4.2 Density of sintered samples 101

4.4.3 Mechanical properties of sintered samples 102

4.5 Role of binder phase composition 103

4.5.1 The role of binder in phase and microstructure of sintered

samples

104


4.5.3 Hardness and fracture toughness of sintered samples 107

4.6 Effect of graphite addition on microstructure and mechanical

properties WC-10FeCrNi hardmetals

109

viii

4.6.1 Phase identification and microstructures 109


4.6.3 Vickers hardness and fracture toughness 115

4.6.4 Summary 116

4.7 Role of NbC addition on vacuum sintered WC-10FeCrNi-2Cgr

hardmetals

117

4.7.1 Phase identification and microstructures 117



4.7.4 Summary 124

4.8 The role of NbC addition as WC grain growth inhibitor for PHIP

sintered WC-10FeCrNi-2graphite hardmetals

125

4.8.1 Phase identification and microstucture 125



4.8.4 Summary 136

4.9 Effect of temperature on microstructure and mechanical properties

of samples sintered by PHIP

137

4.9.1 Phase identification and microstructure 137


4.9.3 Hardness and fracture toughness of sintered samples 142

4.9.4 Summary 143

4.10 Effect of pressure on the microstructure and mechanical properties

of PHIP sintered samples

144

4.10.1 Phase identification and microstructure 144

ix



4.10.4 Summary 149

CHAPTER 5: CONCLUSION AND RECOMMENDATIONS

5.1 Conclusion 150

5.2 Recommendations 152

REFERENCES 153

APPENDIX

174

x

LIST OF TABLES Page

Table 2.1 Physical and mechanical properties of some WC-Co hardmetals 15

Table 2.2 Characteristic stages of liquid phase sintering 35

Table 2.3 The relationship between hardness and microstructural

parameters

48

Table 2.4 Expressions for fracture toughness (KIC) calculation from

Vickers indentation crack systems

52

Table 3.1 Physical properties of WC 59

Table 3.2 Properties of alloy AISI304 60

Table 3.3 Physical properties of niobium 61

Table 3.4 Physical properties of graphite 61

Table 4.1 -phase fraction formed with vs. sintering temperatures 92

Table 4.2 Composition of samples 104

Table 4.3 Composition of samples with graphite addition 109

Table 4.4 Composition of sample with NbC addition 117

Table 4.5 WC grain size of vacuum sintered samples with NbC addition 122

Table 4.6 Composition of sample with NbC addition sintered by PHIP 125

Table 4.7 Amount of -phase of PHIP-sintered samples various with NbC contents

129

Table 4.8 WC grain size of PHIP sintered samples with NbC addition 131

Table 4.9 Relative densities of sintered samples 132

xi

LIST OF FIGURES

Page

Fig. 2.1 W-C phase diagram 10

Fig. 2.2 The atomic structure of WC crystal 11

Fig. 2.3 Vertical section of W-C-Co calculated at 10 wt.% Co 15

Fig. 2.4 Vertical section of W-C-Ni calculated at 10 wt.% Ni 18

Fig. 2.5 Vertical section of W-C-Ni calculated at 10 wt.% Fe 19

Fig. 2.6 Figure 2.6 W-Co-C isothermal sections: a) at 1000oC and b) 1400oC

23

Fig. 2.7 Crystal structure of Fe3W3C 24

Fig. 2.8 Crystal of -graphite 25

Fig. 2.9 Compressibility curves of the ball milled WC–10Co powders 26

Fig. 2.10 Stages of densification of a range of WC-Co alloys; Stage I below

eutectic temperature; Stage II densification at eutectic temperature

and Stage III subsequent densification

28

Fig.2.11 SEM images of morphology evolution of WC-10Co_10 nm

powders when heated at different temperatures: (a) as-milled

power, (b) 800oC, (c) 1000oC, (d) 1100o C (e) 1200oC, and (f)

1300oC

30

Fig. 2.12 Schematic representation of WC particles bounded by the Co

binder phase

31

Fig. 2.13 Schematic of the hardmetal solid state sintering mechanism (a, b,

c, d, e) (Silva et al., 2001) and (f) Cobalt spreading on WC plate at

1050oC under argon, 1-WC, 2-WC/Co region and 3-Co partice

33

Fig. 2.14 SEM images of a) WC-30Co and b) WC-30Co-1VC sintered at

1400oC for 1h in a vacuum furnace

40

Fig. 2.15 The relationship between Vickers hardness and volume percent Co 47

Fig. 2.16 (a) The correlation of the fracture toughness KIC with the mean

linear path in binder phase λ, and (b) with the carbide crystals

contiguity G

49

Fig. 2.17 The correlation between hardness and fracture toughness of

WC-Co and with added cubic carbides

50

xii

Fig. 2.18 Schematic of Vickers indentation cracks and Palmqvist cracks: d-

diagonal of the indentation left in the surface; l- Palmqvist crack

length

51

Fig. 2.19 Ball-powder-ball collision of powder mixture during mechanical

alloying

56

Fig. 3.1 Flow chart of the experimental procedure 62

Fig. 3.2 Sintering diagram of samples in the vacuum furnace 65

Fig. 3.3 Heating diagram of PHIP process 66

Fig. 3.4 Schematic diagram of experiment set up for PHIP. (1) Die, (2) Heating coil, (3) Specimen, (4) Upper punch, (5) Lower punch, (6) Thermocouple, (7) Electric circuit, and (8) Silica sand powder

66

Fig. 3.5 PHIP system and sample fabrication process: 1 – Heating coil, 2 – Sample, 3 – Stainless steel mould, 4 – Silica sand, 5 – Themorcouple, 6 – Upper punch, 7 – Putting in pressing chamber and connecting with pressing system, 8 – Pressure gauge, 9 – Pressing chamber, 10 – System controlling

67

Fig. 3.6 Schematic of indentation mark in Vickers hardness measurement 71

Fig. 3.7 FESEM image of indentation on sample WC-10FeCrNi-2Cgr-1NbC

sintered by PHIP at 1300oC for 45 min including 15 min pressing

at 20 MPa and the crack lengths at the corners of the indentation

(L1-L4)

72

Fig. 4.1 XRD patterns of WC raw powders 73

Fig. 4.2 a) FESEM image and b) EDX result of WC raw powder 74

Fig. 4.3 XRD patterns of FeCrNi powders 75

Fig. 4.4 a) FESEM image and b) EDX result of FeCrNi powders 76

Fig. 4.5 XRD patterns of graphite powders 77

Fig. 4.6 a) FESEM image and b) EDX result of Cgr powders 77

Fig. 4.7 XRD patterns of NbC powders 78

Fig. 4.8 a) SEM image and b) EDX result of NbC powders 79

Fig. 4.9 XRD patterns of WC-10FeCrNi powders at different milling time 80

Fig. 4.10 WC crystallite size at different milling time 82

Fig. 4.11 Back scattered electron FESEM images of WC-10FeCrNi at

various milling time: a) 5 h, b) 10 h, c) 15 h and d) 20 h

83

Fig. 4.12 a) XRD patterns of sintered samples at 1300oC and b) magnification at 40-50 of 2 degree

84

xiii

Fig. 4.13 -phase fraction vs. milling time of sintered samples at 1300oC 85

Fig. 4.14 Back scattered electron FESEM images of sintered samples at different milling time; arrows show the pores in the microstructures

87

Fig. 4.15 Relative densities of pre-compacted and sintered samples at

1300oC

88

Fig. 4.16 Vickers hardness and fracture toughness vs. milling time 89

Fig. 4.17 a) XRD patterns of sintered WC-10FeCrNi at different temperatures and b) magnification at 35-55 of 2

81

Fig. 4.18 FESEM back scattered electron images of sintered WC-10FeCrNi

at different temperatures:a) 1250oC, b)1300oC, d) 1350oC for 1 h

and EDX result of point X in Fig. 4.18a

93

Fig. 4.19 Relative density of WC-FeCrNi at different sintering temperature 94

Fig. 4.20 Vicker hardness and fracture toughness of sintered samples vs.

sintering temperatures

96

Fig. 4.21 a) XRD patterns of sintered samples at different sintering time; 15,

30, 45 and 60 min, and b) magnification at 35-55degree of 2

97

Fig. 4.22 -phase fraction various with sintering time 98

Fig. 4.23 Back scattered electron images of sintered WC-10FeCrNi at different sintering time: a, 15min; b) 30min; c) 45min and d) 60min, and EDX analysis of X, Y, Z and W points in SEM images

100

Fig. 4.24 Effect of sintering time on the density of sintered sample 101

Fig. 4.25 Vickers hardness and fracture toughness of samples at various

sintering time

103

Fig. 4.26 XRD patterns of samples various with binder content sintered at

1350oC

105

Fig. 4.27 -phase fraction vs. initial binder contents 105

Fig. 4.28 SEM back-scatter electron images of samples sintered at 1350oC

with different binder contents

106

Fig. 4.29 Density and relative density of samples vs. initial binder contents 107

Fig. 4.30 Vickers hardness and fracture toughness of sintered samples with

different binder content

108

Fig. 4.31 XRD patterns of mixed powders with graphite addition 110

Fig. 4.32 XRD patterns of sintered samples vs. graphite addition 111

Fig. 4.33 FESEM backscattered electron images of samples after sintering at 113

xiv

1350oC in vacuum furnace various with graphite contents: 0 wt.% HC0, 1 wt.% HC1, 1.5 wt.% HC2, 2 wt.% HC3, 2.5wt.% HC4, 3 wt.% HC5

Fig. 4.34 Relative densities of sintered samples vs. graphite contents 114

Fig. 4.35 Vickers hardness, HV30, and fracture toughness of sintered samples

116

Fig. 4.36 a) XRD patterns of milled samples various with NbC contents and

b) magnification at 40-50 of 2

118

Fig. 4.37 XRD patterns of vacuum sintered samples with vs. NbC contents 119

Fig. 4.38 Figure 4.38 FESEM back scattered electron images of sintered

samples with different NbC contents

120

Fig. 4.39 EDX result of point X in Fig. 4.30 (HN5) 122

Fig. 4.40 Density of sintered sample vs. NbC content 123

Fig. 4.41 Vickers hardness of as-sintered samples with NbC addition 124

Fig. 4.42 XRD patterns of samples HS1 and HS6 after sintering 126

Fig. 4.43 FESEM back-scattered electron images of as-sintered samples: (a) HS1 and (b) HS6

127

Fig. 4.44 XRD patterns of as-sintered samples with various amounts of NbC; 0 wt.% NbC (HS1), 1 wt.% NbC (HS2), 1.5 wt.% NbC (HS3), 2 wt.% NbC (HS4) and 5 wt.% NbC (HS5)

129

Fig. 4.45 FESEM back-scattered images of as-sintered samples with: (a) 1 wt.% NbC (HS2), (b) 1.5 wt.% NbC (HS3), (c) 2 wt.% NbC (HS4) and (d) 5 wt.% NbC (HS5)

130

Fig. 4.46 EDX analysis of a) point X in Fig. 4.45c and b) point Y in Fig. 4.45d

131

Fig. 4.47 Hardness and fracture toughness of PHIP-sintered samples vs with

NbC addition

133

Fig. 4.48 (a)-Hardness and (b)-fracture toughness of this work compared

with references

135

Fig. 4.49 XRD patterns of WC/10FeCrNi-1NbC sintered by PHIP at

different temperature

138

Fig. 4.50 FESEM back-scattered images of PHIPed-samples at: (a) 1150oC, (b) 1200oC, (c) 1250oC and (d) 1300oC

139

Fig. 4.51 Density of PHIPed samples at different temperature 141

Fig. 4.52 (a) Vickers hardness and (b) fracture toughness of

WC-10FeCrNi-2graphite-1NbC sintered by HIP and

WC-10FeCrNi sintered in vacuum furnace

142

xv

Fig. 4.53 XRD patterns of PHIPed samples vs. applied pressure 144

Fig. 4.54 -phase fraction vs. applied pressure in PHIP process 145

Fig. 4.55 FESEM back scattered electron images of samples vs. pressure in PHIP

146

Fig. 4.56 Density vs. applied pressure of samples sintered by PHIP 147

Fig. 4.57 Hardness and fracture toughness vs. pressure of PHIPed samples 148

xvi

LIST OF ABBREVIATIONS

Ar Argon gas

bcc Body centered cubic

BPR Ball to powder ratio

EDX Energy dispersion X-ray spectroscopy

fcc Face centered cubic

FESEM Field emission scanning electron microscopy

FeCrNi AISI304 stainless steel

hcp Hexagonal closed pack

HIP Hot isostatic pressing

LPS Liquid phase sintering

MA Mechanical alloying

PHIP Pseudo hot isostatic pressing

rpm Rotation per minute

TRS Transverse rupture strength

XRD X-ray diffraction

xvii

LIST OF SYMBOLS

έ Shrinkage rate

Plastic strain

Ferrite structure

Austenite structure

Yield strength

Density

Average atomic volume

C Contiguity

d Half diagonal indentation

D Grain size

E Young’s Modulus

H Hardness

HV Vickers hardness

KIC Fracture toughness

L Palmqvist crack length

<x> Mean radius

xviii

ABSTRAK

Kajian ini mengkaji peranan keluli nirkarat AISI304 untuk menggantikan

Co sebagai bahan pengikat untuk logam keras berasaskan WC. Serbuk WC-AISI304

dihasilkan dari serbuk bahan mentah (WC dan AISI304) menggunakan kaedah

pengaloian mekanikal. Serbuk logam keras kemudiannya disinter menggunakan dua

kaedah pensinteran; pensinteran vakum dan pensinteran PHIP. Untuk memperbaiki

sifat-sifat mekanikal sampel tersinter, grafit (Cgr) dan NbC telah ditambahkan

sebelum pengisaran. Keputusan menunjukkan bahawa fasa- (Fe3W3C) terbentuk di

dalam sampel-sampel tersinter ketika proses pensinteran. Penambahan Cgr telah

mengurangkan pembentukan fasa-. Disebabkan fasa ini dihapuskan, kedua-dua

kekerasan dan kekuatan patah sampel tersinter telah meningkat. Tumbesaran butir

WC boleh direncatkan dengan penambahan NbC. Peningkatan kandungan NbC

menyebabkan peningkatan kekerasan tetapi mengurangkan kekuatan patah sampel

tersinter. Di samping itu, kajian ini juga menunjukkan potensi yang tinggi PHIP

dalam menjanakan ketumpatan sampel yang lebih tinggi berbanding pensiteran

vakum. Oleh itu, kaedah ini boleh meningkatkan sifat-sifat mekanikal logam keras

WC-AISI304. Kekerasan Vickers logam keras WC-AISI304-2Cgr-xNbC (x = 1 - 5)

yang dihasilkan ialah dalam julat 1600 ke 1660 kg/mm2 dan kekuatan patah, KIC, dari

8.7 ke 8.3 MPa.m1/2 apabila menggunakan pensinteran vakum. Walau bagaimanapun,

sampel yang sama yang dihasilkan melalui pensinteran PHIP memberikan kekerasan

Vickers dari 1640 ke 1820 kg/mm2 dan kekuatan patah, KIC, dari 10 ke 7.3 MPa.m1/2.

Nilai-nilai kekerasan dan kekuatan patah ini adalah dalam julat pertengahan

berbanding dengan sistem yang disebutkan dalam tinjauan persuratan. Keputusan ini

menunjukkan bahawa AISI304 boleh dicadangkan untuk menggantikan bahan

pengikat dari Co sebagai usaha untuk menghasilkan alat pemotong.

xix

ABSTRACT

This work studies the role of AISI304 stainless steel as a Co replacement

binder for WC-based hardmetals. WC-AISI304 hardmetal powders were produced

from raw powders (WC and AISI304) by mechanical alloying technique. The

hardmetal powders were then sintered by two sintering methods; vacuum sintering

and PHIP sintering. To improve mechanical properties of sintered samples, graphite

(Cgr) and NbC were added prior to milling. The results show that -phase (Fe3W3C)

formed in the sintered samples during sintering. Cgr addition has enabled to reduce

the formation of -phase. As this phase was eliminated, both hardness and fracture

toughness of sintered sample were improved. WC grain growth can be inhibited by

the addition of NbC. Increasing NbC content led to an increase of hardness but

reduce fracture toughness of sintered samples. Besides that, this work also show a

higher potential of PHIP in generating higher density of sintered samples compared

to vacuum sintering, and hence, improving mechanical properties of WC-AISI304

hardmetals. The Vickers hardness of WC-10AISI304-2Cgr-xNbC (x = 1 - 5)

hardmetals produced is in range of 1600 to1660 kg/mm2 and fracture toughness, KIC,

from 8.7 to 8.3 MPa.m1/2 by vacuum sintering. However, the same samples produced

via PHIP sintering gave the Vickers hardness from 1640 to 1820 kg/mm2 and

fracture toughness, KIC, from 10 to 7.3 MPa.m1/2. These values of hardness and

fracture toughness are in the intermediate range compared to other systems provided

by literatures. The results indicate that AISI304 could be proposed to replace Co

binder in order to fabricate cutting inserts.

1

CHAPTER 1

INTRODUCTION

1.1 Introduction

WC-based hardmetals have been developed since 1920s and widely used in

cutting tool industries. The combination between the high hardness of WC particles

and high ductility of the binders (Co, Ni or Fe) produces a hardmetal with high

hardness, high fracture toughness and good wear resistance. These properties are

very important in industries such as cutting edges in turning machines, drilling

screws in mining equipments, and die industries and so on (Jorn, 1985; Ettmayer,

1989; Hanyaloglu et al., 2001; Fernandes and Senos, 2011). WC-based carbide

hardmetals dominate about 95% of cemented carbide cutting tools in the market (Yao

et al., 1999). The annual report of W for use in WC worldwide shows that in 2008,

50,000 tons of W were consumed which account for nearly 60% of the world’s W

consumption including recycled materials (Schubert et al., 2010).

In recent decades, many researches have been carried out to improve the

mechanical properties of WC-Co hardmetals by controlling their microstructures.

Fine and ultrafine grained WC hardmetals are more important today for high

performance of cutting tools, chipless formation or other applications. As a rule, the

importance of finer grain size of WC is derived from the understanding that hardness

and wear resistance increase with decreasing WC grain size (Kim et al., 1997; Niels,

2004; Huang et al., 2008b). However, sintering WC-Co process is often performed at

temperatures ranging from 1300-1500oC in a vacuum furnace, thus grain growth of

WC takes place and consequently, the mechanical strength and hardness of the tools

are limited.

2

Mechanical alloying or other synthesis methods could be used to gain nano

particles or ultrafine WC powders effectively but abnormal grain growth could occur

and affect the mechanical properties of hardmetals. In this case, a small amount of

some other carbides such as VC, NbC, TiC, TaC or Cr2C3 have been added to work

as the grain growth inhibitors during sintering process (Huang et al., 2007; Huang et

al., 2008a; Huang et al., 2008b; Barbatti et al., 2009; Mahmoodan et al., 2009;

Weidow et al., 2009a). The advance sintering methods to limit WC grain growth

have been investigated using a variety of techniques including liquid phase sintering,

hot isostatic pressing (HIP) (Soares et al., 2012), spark plasma sintering (SPS)

(Zhang et al., 2004b; Sivaprahasam et al., 2007), microwave sintering (Breval et al.,

2005), high frequency induction-heated and pulse plasma sintering (Kim et al., 2004;

Kim et al., 2007; Shon et al., 2009).

Since the first patent of WC-based hardmetals was issued around 1923 by

Schröter in the German company "Osram Studiengesellschaft", Co is the best choice

as the binder of WC-based hardmetals (Schroter, 1925). WC-Co hardmetals have

excellent mechanical properties such as high hardness, high strength and fracture

toughness, and high wear resistance (Viswanadham and Lindquist, 1987; Ettmayer,

1989; Hanyaloglu et al., 2001; Weidow et al., 2009a; Schubert et al., 2010). However,

recently, there are an increasing suspects of work diseases concerning with Co and

WC-Co containing dust. Several reports have shown that occupational exposure to

Co-containing dust has been associated with the development of different pulmonary

diseases including fibrosing alveolitis and lung cancer (David et al., 1990; Lison et

al., 1995; De Boeck et al., 2003; Koutsospyros et al., 2006). The excessive chronic

inhalation of hardmetal particles can be associated with the occurrence of different

lung diseases including an excess of lung cancers. The elective toxicity of hardmetals

3

is based on a physico-chemical interaction between Co metal and WC particles to

produce activated oxygen species resulting in induction of DNA damage (De Boeck

et al., 2003).

Besides that due to the high cost and depleting of Co, and the need to

improve some properties of WC-based hardmetals such as corrosion resistance and

oxidation resistance (Penrice, 1987; Gille et al., 2000; Fernandes and Senos, 2011),

many studies have been carried out to find new binders to replace Co in WC-based

hardmetals. Some transition metals have been investigated to replace Co with other

elements such as Ni, Fe, Cr and Mn and show that it is possible to obtain high

density after sintering at temperatures near to those used for WC/Co cemented

carbides. The use of Ni binder has been shown to increase corrosion resistance of

WC-based hardmetals compared to WC-Co (Tracey, 1992). The addition of Cr or Cr

combined with Ni in the binder phase has been reported not only to improve

oxidation and corrosion resistance but also to act as WC grain growth inhibitor in

WC-based hardmetals (Penrice, 1987; Tracey, 1992).

The replacement of Fe-Mn alloys also make improvement in WC hardmetal

properties such as hardness and toughness, however, the optimum composition has

not been pointed out (Hanyaloglu et al., 2001). The development of Fe-Ni binders

also can be seen in literature (Penrice, 1987; Uhrenius et al., 1997). Fe-rich alloys

have been preferred as binders, not only because they are inexpensive but also

because they improve mechanical properties and enhance sinterability due to their

good wettability with WC (Uhrenius et al., 1997).

Recently, Fe rich alloys and stainless steel have been reported to have a

potential to replace Co in WC based cemented carbides. One research team have

4

intensively used sputtering technique to coat WC powder with stainless steel and Fe

rich alloys and discovered some advantages of these alloys to replace Co in

densification (Fernandes et al., 2003b; Fernandes et al., 2007; Fernandes, 2008;

Fernandes et al., 2009a; Fernandes et al., 2009b). They reported good sintering

characteristics of the resultant composite powders and an improvement in the

hardness of those hardmetals (Fernandes et al., 2003b; Fernandes et al., 2009b).

As known, AISI304 stainless steel is a basic grade of austenite stainless steel

containing 18 wt.% of Cr and 8-10 wt.% of Ni. It possesses high strength, and high

oxidation and corrosion resistance in comparison to plain carbon steel. The presence

of Cr and Ni in this composition can combine their benefits in the binder phase for

WC-based hardmetals (Penrice, 1987; Tracey, 1992; Uhrenius et al., 1997).

Considering the potential of stainless steel as a binder replacement to Co, this

research was carried out using AISI304 stainless steel to work as metal binder phase

in WC-based hardmetal. This research investigated the role of binder phase in terms

of binder content, sintering time, sintering temperature as well as sintering methods

on the microstructure and mechanical properties of sintered samples such as hardness

and fracture toughness. Beside that the effect of grain growth inhibitor and

microstructure controlling by additional components such as NbC and graphite was

also studied.

1.2 Problem statements

Literatures show a good sintering characteristic of WC/Fe-alloys or stainless

steel hardmetals and good wetting properties of WC with iron-alloys. The reports

show that liquid binder phase was formed at eutectic temperature about 1150oC for

5

WC-Fe rich binder phase (Fernandes et al., 2003b; Fernandes et al., 2007), which

make this temperature suitable for liquid state sintering. However, similar to WC-Co

system, the formation of -phase ((M,W)6C, M = metal) has been reported during

sintering of WC-Fe rich alloys hardmetals (Fernandes et al., 2003b; Fernandes et al.,

2007; Fernandes et al., 2009a; Fernandes et al., 2009b). The formation of -phase in

the microstructure has been attributed to decarburization of WC during sintering or to

a C content in the initial composition that is below the critical minimum necessary to

prevent the formation of -phase (Fernandes et al., 2007; Fernandes et al., 2009b).

-phase were reported to form during sintering WC coated with stainless steel at low

temperature, of about 750oC and increases with temperature up to 1100oC (Fernandes

et al., 2007). They also have reported that the -phase, in the temperature ranging

from 1200-1325oC, can be represented approximately as (Fe2.3Ni0.3)(Cr0.6W2.8)C. The

presence of this brittle phase leads to a decrease in mechanical properties,

particularly in fracture toughness (Uhrenius et al., 1997; Yao et al., 1998; Upadhyaya

et al., 2001).

Thus, eliminating the formation of this phase becomes an important point in

fabrication of WC-based hardmetals with new binders. It is reported that 2.5 -3 wt.%

graphite has an ability to eliminate the formation of -phase for sintering WC-10Fe

rich alloys hardmetals (Fernandes et al., 2007; Fernandes et al., 2009a). So the role of

graphite content addition on the microstructure and mechanical properties of the new

WC-based hardmetals need clarification in particular its hardness and fracture

toughness.

Although several researches have been reported on sintering WC/Fe-alloys or

stainless steel binders (Hanyaloglu et al., 2001; Fernandes et al., 2003b; Fernandes et

6

al., 2007), the grain growth of WC in these kinds of binder phases is not well known.

Controlling grain growth of WC in iron-alloys or stainless steel binders through the

addition of grain growth inhibitor and the use of different sintering temperatures are

necessary to be investigated. It is known that NbC has been used up to 5 wt.% as WC

grain growth inhibitor for WC-Co system during sintering (Huang et al., 2008a).

Thus, it is necessary to investigate the ability of NbC as the WC grain growth

inhibitor in WC/Fe-alloy system.

In general, WC-based hardmetals could be sintered in a vacuum furnace which is

sufficient for sintering large number of products and requirement equipments but the

full densification is difficult to attain. In this case advanced sintering methods,

including hot isostatic pressing (HIP) and spark plasma sintering (SPS), are the

most suitable in order to achieve higher density of sintered samples. However, since

the typical hot isostatic pressing is too expensive, low cost and simple isostatic

pressing technique, such as pseudo hot isostatic pressing (PHIP) which uses sand as a

pressure delivery medium, is proposed. Thus, an investigation on densification of

WC in stainless steel binder during PHIP is required because information about this

process is limited. PHIP is an advanced technique to produce high densification of

sintered samples (Park et al., 1996; Xu et al., 2003; Zhang et al., 2010). Considering

PHIP promises high mechanical properties of sintered samples, PHIP process was

also used in comparison to vacuum sintering method to sinter the samples.

1.3 Research objectives

To solve the stated problem statements and in order to investigate the ability

of stainless steel as a binder to WC-based hardmetals, this research aimed to fabricate

7

WC/AISI304-stainless steel hardmetals using different sintering methods (in a

vacuum furnace and PHIP sintering) with improved mechanical properties

The objectives are listed as following:

1. To study the effect of milling time, sintering time, temperature and binder

content on the microstructure and mechanical properties of WC/AISI304

hardmetals sintered in a vacuum furnace.

2. To control the formation of -phase during vacuum sintering process by adding

graphite in order to compensate C lost.

3. To investigate the ability of NbC to inhibit WC grain growth during sintering

for mechanical properties improvement.

4. To study the effects of sintering temperature and pressure on the microstructure

and mechanical properties of WC/AISI304 hardmetals sintered by PHIP.

1.4 Research scope

AISI304 stainless steel (namely as FeCrNi) was used as the binder in

WC-based hardmetal. WC was mixed with FeCrNi powders by a planet ball milling

to produce WC-AISI304 hardmetal powders. The effect of milling to produce high

homogeneity of WC and FeCrNi powders and the role of binder phase was studied

according to the following flow:

In order to optimization of milling process in term of milling time:

WC-10FeCrNi (wt.%) powders were fabricated at different milling time (5, 10, 15

and 20h). The pre-compacted samples were sintered at 1300oC in a vacuum furnace

8

and then, investigated microstructure and mechanical properties (hardness and

fracture toughness).

The effect of sintering temperature: WC-10FeCrNi (wt.%) powders milled at

15h were pre-compacted and sintered 1h at different temperature (1250, 1300 and

1350oC) in the vacuum furnace.

The role of sintering time: WC-10FeCrNi (wt.%) powders milled at 15h were

pre-compacted and sintered at 1350oC under vacuum with different sintering time; 15,

30, 45 and 60 min.

The role of binder content: 8, 10, 12 and 15 wt.% of FeCrNi were mixed with

WC by ball milling at 15h. The pre-compacted powders were then sintered at 1350oC

for 1h in the vacuum furnace.

1, 1.5, 2, 2.5 and 3 wt.% of Cgr were added in WC-10FeCrNi before milling for

15h to eliminate the formation of -phase during sintering at 1350oC for 1 h in the

vacuum furnace.

1, 1.5, 2, 5 wt.% of NbC were added in WC-10FeCrNi-Cgr to investigate ability

of NbC to inhibit WC grain growth during sintering at 1300oC for 1 h in the vacuum

furnace.

To study the role of PHIP sintering method, 1, 1.5, 2, 5 wt.% of NbC were

added in WC-10FeCrNi-Cgr to investigate ability of NbC to inhibit WC grain growth

during sintering. PHIP was done at 1300oC for 45 min including 15 min pressing at

20 MPa, and as well as the effect of PHIP method on the microstructure and

mechanical properties of WC-10FeCrNi hardmetals.

9

The effect of sintering temperature by PHIP: WC-10FeCrNi- Cgr-1NbC were

sintered at 1200, 1250 and 1300oC for 45 min including 15 min pressing at 20 MPa.

The effect of loading pressure by PHIP: WC-10FeCrNi-Cgr-1NbC was sintered

at 1300oC for 45 min including 15 min pressing at 0, 5, 15, 20 and 25 MPa.

10

CHAPTER 2

LITERATURE REVIEW

2.1 History of tungsten carbide and hard metals

The discovery of tungsten carbides was marked to the invention of W2C in

1896 by Moissan, and of WC in 1898 by Williams, working at Moissan’s laboratory

at the school of Pharmacy at the University of Paris (Yao et al., 1998; Yao et al.,

1999). Fig. 2.1 shows the advance W-rich part of the binary W-C equilibrium

diagram (Kurlov and Gusev, 2006). Three stoichiometries have been found: -W2C,

-WC1-x, and -WC.

Figure 2.1 W-C phase diagram (Kurlov and Gusev, 2006)

W2C has hexagonal structure with three modifications: the PbO2, Fe2N, and

CdI2 types, denoted , ’, and ”, respectively. These polymorphs are stable at

different temperatures. W2C phase results from the eutectoidal reaction between

elemental W and -WC at 1250oC and melts congruently at approximately

11

278510oC and forms eutectic melts with the W solid solution at 22155oC and with

-WC1-x at approximately 2755oC, and it exhibits a comparatively wide homogeneity

range 25.4-34 at% C at 27155oC. The cubic sub-carbide WC1-x (where 1x0.5)

crystallizing in the NaCl type structure presented , and the hexagonal WC denoted .

This phase is originated from a eutectoidal reaction between and at 2516oC and

melts at approximately 27855oC.

The technical importance of -WC is the only binary phase stable at room

temperature and has almost no solid solubility up to 2384oC but may become carbon

deficient between this temperature and its incongruent melting point. The -WC has

a crystal structure of simple hexagonal (space group P6m2) with lattice parameters a

= 0.2906 nm and c = 0.28375 nm (c/a ratio of 0.9764). The carbon atoms take the

asymmetric position of (1/3, 2/3, 1/2) in the unit cell as shown in Fig. 2.2 and this

asymmetric occupation of carbon atoms divides the prismatic planes into two

different families of planes with different atom arrangements. These two families of

prismatic planes can have a different affinity to carbon because W atoms on each

plane have a different number of W–C bonds. The planes with high affinity to carbon

grow preferentially in the saturated carbon conditions and disappear finally leaving

triangular prism in shape (Kim et al., 2003).

Figure 2.2 The atomic structure of WC crystal (Kim et al., 2003)

12

The first sintered WC products were produced in 1914 for using in drawing

dies and rock drills. Powdered WCs or mixed with Mo2C were pressed and sintered

just below the melting temperature of the pure WC. However, the sintered products

were very brittle and unsuccessful in industrial applications (Yao et al., 1998; Yao et

al., 1999).

WC hardmetals, as sometimes called cemented WCs, were originated in the

USA and in Germany denoted to the electronic incandescent lamp industry. Since the

need of a replacement to the expensive diamond dies for being used in the

wire-drawing of fine W filaments, manufacturers have been searching for other

materials to make the dies. Because of their extreme hardness, WCs are the subject

of a potential substantial research and development effort on the part of incandescent

lamp industry and their suppliers, for more than two decades (Yao et al., 1998; Yao

et al., 1999). The year of 1923 made an important milestone to the invention of WC

hardmetals as submitted by Schröter, Germany (Schroter, 1925). In this patent, WC

comprising carbon content in the range of 3 – 10 wt.% was mixed with not more than

10 wt.% of metal binders, namely Fe, Ni and Co, and then, the mixture was sintered

at the temperature lower than WC melting point, about 1500-1600oC. This patent,

then, caused the revolution of metal cutting tool materials and no one could imagine

the enormous breakthrough for this material in the tooling industry. When the new

tools, made from sintered WC-Co, were first placed in the market in 1927, they

caused a sensation in the machine tool industry, by allowing cutting speeds three to

five times faster than the best high-speed steel tools in use at that time (Yao et al.,

1998; Yao et al., 1999). Few years later, Schwarzkopf discovered that the solid

solutions of more than one carbide, particularly TaC, TiC, NbC, VC and Mo2C, have

superior mechanical properties to the individual carbides in cemented hardmetals

13

(Upadhyaya, 1998; Fernandes and Senos, 2011).

Another important advance in the development of WC hardmetals happened

in the late 1960s and early 1970s with the application of coating technique.

Hardmetal tools were coated with titanium carbide (TiC), titanium nitride (TiN),

titanium carbonitride (TiCN) or alumina (Al2O3), which are extremely hard, thus

increase the abrasion resistance of hard metals (Yao et al., 1998; Yao et al., 1999;

Fernandes and Senos, 2011).

Nowadays, the application of WC hardmetal products has become widely

used in various industry sectors including metal cutting, machining of wood, plastics,

composites, soft ceramics, chipless forming, mining and construction, structural parts,

wear parts, and military components. There has been a continuous expansion in the

consumption, from an annual world total of 10 tons in 1930 to about 50,000 tons in

2008. This shows an import role of hardmetals in the world’s economy (Schubert et

al., 2010).

2.2 Binder phase

The manufacturing process of hardmetals, until now, was based on powder

metallurgy (PM) processing including liquid phase sintering. Principle steps of

Schröter (Schroter, 1925) invention consist of sintering a mixture of 90 wt.% WC

and 10 wt.% binder phase (Fe, Ni, Co) at suitable temperatures at which the binder is

liquid and complete consolidation of the compact occurs. The binder phase plays a

very important role in sintering of hardmetals. It is responsible for densification

through wetting, spreading and formation of agglomerates. The hard phase has a

passive role, since the WC particles do not sinter together but are moved by the

14

binder (Silva et al., 2001). So far, the metallic binder can be a variety of elements,

alone or in combination such as Co, Ni, Fe, and Mo or it can also contain other

materials, such as stainless steel, superalloys, Ti, etc. In addition, the contents of

metal binder have been verified from the original work of Schröter, ranging from 3 to

30 wt.%, even higher (Penrice, 1987; Gille et al., 2000).

2.2.1 Cobalt binder

Since the appearance of WC hardmetals, cobalt has been the optimal choice

for the metal binder. Cobalt metal has two allotropic modifications, a close-packed

hexagonal (cph) structure, , stable at temperatures below approximately 400oC, and

a face centered cubic (fcc) structure, , stable at higher temperatures (Upadhyaya,

2001). This metal possesses high hardness, yield stress, toughness and strength.

Nowadays, more than 90% of all WC-based hardmetals have been reported using Co

as the preferred binder metal with contents ranging from 3 to 30 wt.%. The

dominance of Co binder relative to other metal binders is concerned with its superior

wettability with WC, higher solubility of WC in Co at sintering temperature and

providing excellent mechanical properties to hard metals (Penrice, 1987; Fernandes

and Senos, 2011).

The understanding of the phase diagram of cemented carbides is an

important tool to predict phase composition after sintering step and to select the

adequate sintering conditions. The vertical section of the W-C-Co phase diagram

calculated at 10 wt.% Co is shown in Fig. 2.3 (Guillermet, 1989a). The best

properties of WC-Co system have been obtained within the two phase region, fcc-Co

and WC phase. The points denoted as a and b in Fig. 2.3 define, respectively,

15

minimum and maximum carbon contents of alloys which are in two-phase state of

fcc Co and WC just after the equilibrium solidification. Due to the important role of

WC-Co hard metals, several investigations concerning the W-C-Co system have

been carried out over the last century (Pollock and Stadelmaier, 1970; Guillermet,

1989a; Markström et al., 2005). Some physical and mechanical properties of WC-Co

obtained from those works are listed in Table 2.1.

Figure 2.3 Vertical section of W-C-Co calculated at 10 wt.% Co: M6C

(-phase,Co3W3C), FCC (-Co), gr (graphite), P (peritectic point), E (eutectic point)

(Guillermet, 1989a)

16

Table 2.1 Physical and mechanical properties of some WC-Co hardmetals

Nominal

composition Average

grain size

(m)

Density

(g/cm3)

Hardness

(HV) Fracture

toughness

(MPa.m1/2)

Kc

(N/mm) Reference

94WC+6Co 0.1 14.83 2280 - 370

Gille et al., 2000 90WC+10Co 0.1 14.46 2043 - 530

88WC+12Co 0.1 14.30 1910 - 570

85WC+15Co 0.1 13.95 1700 - 725

85WC+15Co 0.258 14.53 1992 11.9 - Kim et al., 2004

90WC+10Co 0.38 14.79 1756 11.6 - Kim et al., 2007

WC+2.9Co 0.94 14.51 2014 6.5 -

WC+12Co 1 - 1748 11.4 - Deorsola et al., 2010

Besides the aim to get two-phase state of WC-Co hard metals after sintering,

the tendency to produce fine and ultrafine or nano microstructure of WC-Co

hardmetals arises from the understanding that the mechanical properties such as

hardness and wear resistance increase with the decrease of WC grain size. To fulfill

the requirements of the hardmetal industry and the trend toward finer grain size tools,

several studies have been done (Gille et al., 2000 Kim et al., 2004; Kim et al., 2007;

Deorsola et al., 2010).

2.2.2 Nickel binder

Ni has been proposed in many researches as the binder phase to WC-based

hardmetals, mainly as a substitute for Co during periods of Co scarcity. Ni retains its

fcc structure at all temperatures below its melting point and a similarly lattice

parameter of fcc Ni compared to fcc Co. The melting point of Ni at 1453oC is

appreciably lower than Co at 1495oC, however it is necessary to use increased

17

sintering time and temperature to gain satisfactory densification (Penrice, 1987;

Fernandes and Senos, 2011). Vertical section of W-C-Ni phase diagram calculated at

10 wt.% Ni is shown in Fig. 2.4 (Guillermet, 1989b). Comparing with the

corresponding section of W-C-Co in Fig. 2.3, the width of two-phase state region

remain essentially similar, however the range of favorable C contents moves

backwards to lower value as comparison with the stoichiometric composition. In

addition, the change from W-C-Co to W-C-Ni involves an appreciable increase in the

equilibrium temperature of eutectic (E) and peritectic (P) points (Figs. 2.3 and 2.4)

(Raghavan, 2007).

The partial or complete replacement of Co by Ni leads to a decrease of

hardness of hardmetals. The decrease of hardness is probably the principal reason

why Ni has not been accepted widely to replace Co in WC hard metal industry,

exceptionally in very significant quantities of wear applications where require higher

corrosion and erosion resistance or higher oxidation (Penrice, 1987; Tracey, 1992;

Voitovich et al., 1996). In some cases, mixtures of Co-Ni were used as the binder to

improve toughness but not significantly reducing hardness for economic purpose or

some applications such as in mining and hot rolling equipments (Voitovich et al.,

1996; Zhang and Sun, 1996; Aristizabal et al., 2012).

18

Figure 2.4 Vertical section of W-C-Ni calculated at 10 wt.% Ni: M6C

(-phase,Ni3W3C), FCC (-Ni), gr (graphite), P (peritectic point), E (eutectic point)

(Guillermet, 1989b)

2.2.3 Iron binder

Iron is a soft metal that can dissolve a small amount of C (0.021 wt.%) up to

910oC with a magnetic bcc crystal structure, called -Fe (ferrite). At the higher

temperature, up to 1400oC, and in the presence of higher C contents, -Fe undergoes

a phase transition from bcc to fcc crystal structure, also called -Fe (austenite). This

phase is metallic and non-magnetic and dissolve considerably more C, 2.04 wt.% at

1146oC (Fernandes and Senos, 2011).

The W-C-Fe system has been subjected to several experiments and

theoretical investigation over the years Bergström, 1977; Gustafson, 1988;

Guillermet, 1989b. A vertical section of W-C-Fe phase diagram calculated at 10

wt.% Fe can be seen in Fig. 2.5 (Guillermet, 1989b). The two-phase state region

moves to higher content of C. Fe forms a ternary eutectic melt at only 1143oC and it

19

also a possible replacement for Co (Jia et al., 1998). However, complete substitute of

Co by Fe in WC-based hardmetal did not hold much promise. Fe, unlike Co, is a

carbide former, thus, a C-rich composition is needed and even to control material

within two-phase state is difficult (Penrice, 1987).

Figure 2.5 Vertical section of W-C-Ni calculated at 10 wt.% Fe (Guillermet, 1989b)

2.2.4 Ni-Fe and Co-Ni-Fe binders

The complete substitution of Ni-Fe binder to Co in WC-10 wt.% Co has

been reported having better properties in comparison with the substitution by either

Fe or Ni alone (Upadhyaya and Bhaumik, 1988). This study also shows that the grain

size of WC increased with increasing of Ni content in the binder but resulting in a

decrease of oxidation resistance. In other study, the hardness of WC-(Fe,Ni) with

Fe-rich binder depends on the C contents and lower than WC-Co one (Uhrenius et al.,

1997). The study also shows that the Ni-rich Ni-Fe or Ni-Fe-Co binders have no

tendency to increase hardness for composition close to or within the two-phase

region of WC and fcc metal. In general, C content plays an important role in

20

controlling phase and final density, and consequently, influences on mechanical

properties of WC-(Fe,Ni,C) hardmetals (Viswanadham and Lindquist, 1987;

González et al., 1995; Uhrenius et al., 1997). Controlling the phase formation of

WC-(Fe,Ni,C) by heat treatment has been also carried out to get higher mechanical

properties of WC-(Fe,22Ni,C) (Viswanadham and Lindquist, 1987). The study also

suggested that the hardness and fracture toughness improved at the suitable content

of C after quenching in liquid nitrogen.

Partially replacement of Co binder by a Co-rich binder (60Co-20Ni-20Fe)

and a Fe-rich binder (75Fe-15Ni-10Co) have been done in WC-Co hardmetals with

2.5 m of WC grain size (Gille et al., 2000). The results show that the hardmetals

with Fe-rich binder are higher in hardness but lower in toughness in comparison to

Co-rich binder hardmetals.

2.2.5 Fe-Cr-Ni binders and stainless steel binder

Different compositions of Fe-Cr-Ni have been used as the replace binder to

Co in WC-10Co hardmetals using magnetron sputtering method to coat WC powders

and sintering in a vacuum furnace (Fernandes et al., 2007; Fernandes et al., 2009a).

The thermal reactivity between WC and Fe/Cr/Ni was investigated (Fernandes et al.,

2009a). The results show that no -phase occurs for WC-10Fe after sintering at

1400oC and the substitution of half percent of the Fe content by Ni stabilized the

austenite -Fe. The introduction of Cr in the binder induced the formation of Cr2C

and limited the formation of -phase. Decreasing Cr content leads to the formation

of -phase. The addition of graphite has enabled to reduce -phase formation in

WC-Fe/Ni/Cr hardmetals as well (Fernandes et al., 2007; Fernandes et al., 2009a).

21

The use of AISI304 stainless steel has been also proposed as binder to

replace Co (Fernandes et al., 2003a; Fernandes et al., 2006). Particle surface

investigation revealed that a high uniformity of the coating distribution on WC

particles was attained by sputtering technique, enabling to complete surface coverage

at low binder content (1 wt.%) (Fernandes et al., 2006). The sputtering technique is

shown to be efficient in promoting densification of hardmetals, which has been

attributed to the high uniformity of coated powder. High densities (of approximately

95 %) were obtained at a relatively low sintering temperature (1325oC) with only 6

wt.% of binder phase (Fernandes et al., 2003a).

2.3 - Phase

Decarburization of WC or deficiency of C comparing to stoichiometric leads

to the formation of ternary compounds of W, Co and C, called -phase that exists in

two types: M6C and M12C (M=W and Co or Fe). M12C phase is substantially

constant composition. On the other hand, M6C can vary within the range of

M3.2W2.8C and M2W4C. The M6C type of -phase is in equilibrium with the liquid

phase and can nucleate and grow during the sintering process. The M12C type is

formed in the solid state (during cooling) with small grains distributed throughout the

matrix and is therefore effectively less embrittling (Yao et al., 1998; Yao et al., 1999)

than M6C. However, M6C was reported to be the most high-temperature stable

carbide (Pollock and Stadelmaier, 1970).

The M6C carbide, which has many isomorphs among ternary carbides,

contains at least two types of metal atoms. Its formation is favoured by the

combination of one weak and other strong carbide formers, e.g. Fe and W

22

(Bergström, 1977). All the examined -phases (Fe3W3C, Fe6W6C, Co3W3C,

Co6W6C) adopt the cubic symmetry with the space group Fd3m and Z = 16 (for

Fe3W3C and Co3W3C) and Z = 8 (for Fe6W6C and Co6W6C), (Z: the number of C

atoms). The M6C carbide has a fcc structure (cF112) containing 96 atoms per unit

cell. The W atoms occupy the 48 sites; Fe and Co are placed in two non-equivalent

32e and 16d sites, where as C is located in the 16c. The only difference between

M3W3C and M6W6C is that these phase contain 16 and 8 C atoms (per cell),

respectively (Suetin et al., 2009, Ramnath and Jayaraman, 1987).

- phase formation moves to low C contents in the phase diagram. Fig. 2.6

shows the ternary diagrams of W-Co-C calculated at 1000oC and 1400oC (Pollock

and Stadelmaier, 1970). While the homogeneity range of M12 was found to remain

small, M6C spread out to include Co3W3C and Co2W4C at 1400oC and narrowed

back down to a small range around Co2W4C at 1000oC. Their evaluation of the -

phase in W-Co-C system based on XRD patterns of as-cast alloys revealed two -

phases, Co6W6C and Co2W4C with lattice constants near 10.90 and 11.20 Å,

respectively.

23

Figure 2.6 W-Co-C isothermal sections: a) at 1000oC and b) 1400oC

(Pollock and Stadelmaier, 1970)

The ideal structure of M6C is quite complicated and consists of eight regular

octahedral of W atoms centered in a diamond cubic lattice and eight regular

tetrahedral of Fe (or Co) atoms centered in the second diamond cubic lattice that

interpenetrates the first through the 1/2, 1/2, 1/2 unit cell translation. Sixteen

additional Fe (or Co) atoms are tetrahedrally coordinated around the Fe (or Co)

tetrahedral and sixteen C atoms surround the W octahedral in tetrahedral

coordination (Suetin et al., 2009). The model crystal structure of Fe3W3C is shown in

Fig. 2.7 (Fernandes, 2008).

24

-phase is an undesired phase in microstructure of sintered products because

it results in degradation of mechanical properties and cutting performance (Yao et al.,

1998). This brittle phase was attributed to the decrease of fracture toughness of

hardmetals (González et al., 1995; Uhrenius et al., 1997; Upadhyaya, 1998;

Upadhyaya et al., 2001). The transverse rupture strength was achieved highest value

without the presence of -phase in WC-6Co (Upadhyaya et al., 2001).

Figure 2.7. Crystal structure of Fe3W3C; Carbon Tungsten Iron

(Fernandes, 2008)

2.4 Graphite

Graphite, Cgr, is an allotrope of C. There are two forms of Cgr in nature:

α-Cgr with hexagonal structure and β-Cgr having rhombohedra structure. Most of Cgr

is α-Cgr and it possesses a layer structure in which each C is directly bound to three

other C atoms at a distance of 1.415Å. The layer planes are stacked parallel to each

other at the distance of 3.354 Å. And, the layers of atoms are arranged in an

ABABAB... repeat fashion. The only difference of the β-form (rhombohedral) is the

layers in the arrangement of ABCABCABC... although the C-C distances and the

25

interlayer spacing remains the same as in the α-form (Norley, 2011; Webelements,

2012). Crystal structure of -Cgr is shown in Fig. 2.8.

From a C control point of view, the most suitable composition are those

where neither -phase nor Cgr are formed. The lack of C leads to the formation of

-phase and the high C content leads to the precipitation of Cgr in the microstructure

of sintered products. The presence of Cgr led to a decrease of hardness (Uhrenius et

al., 1997) and a lower transverse rupture strength (Upadhyaya et al., 2001).

Figure 2.8 Crystal of -Cgr (Norley, 2011)

2.5 Consolidation tungsten carbide hard metal powders

2.5.1 Green consolidation

Green consolidation is a process that produces a shape with certain

dimensions by pressing loose powder mass by an external pressure. In general,

compaction pressure in the range of 21-42 kg/mm2 is used to produce sufficient

26

green strength compacts (Upadhyaya, 1998). After compaction, a green density of

about 60% of theoretical density can be generally attained for a given hardmetal

composition. Despite of any value of compact’s green density; hardmetal parts obtain

almost complete densification upon liquid phase sintering. However, with less than

60% theoretical density, shrinkage is high and dimensional control is difficult. The

relationship between pressure and relative density of WC-10Co compacted powder at

different milling time is shown in Fig. 2.9 (Hewitt and Kibble, 2009).

The hardmetal powders do not exhibit plastic flow even at high pressure. So,

to improve the strength of green compact, a plasticizer is used which has a low yield

stress. The common plasticisers used are polyvinyl alcohol, poly ethylene glycol,

starch solution, paraffin wax and synthetic resins. And hence, pre-sintering is needed

in some cases to de-waxing or de-binding the plasticisers before sintering process

(Upadhyaya, 1998; Petersson, 2004; Fang et al., 2005; Liu et al., 2006).

Figure 2.9. Compressibility curves of the ball milled WC–10Co powders

(Hewitt and Kibble, 2009)

27

2.5.2 Sintering process

The green compacts of powders were undergone heat treatment steps to

consolidate the power particles into coherent structure via mass transport on the

atomic scale. This process leads to an improvement of mechanical properties, like

strength, fracture toughness and wear resistance. Generally, sintering theory is more

accurate for single phase powder sintering by solid state diffusion. However, in

systems consist of more than two phases such as WC-Co hardmetals, a part of liquid

phase formed during sintering process. Therefore, the mechanism of sintering

process becomes more complex.

WC based hardmetals are mostly sintered to full densification by the

formation of a liquid phase, known as the binder phase. Although, there are many

researches investigating liquid phase sintering mechanism of WC based hardmetal, it

is still not fully understood (Ettmayer, 1989). In general, the powder compact is

heated to a temperature above the eutectic point in order to attain dense materials.

During heating period, the mechanisms of solid state transformation sintering also

take place such as changes in density, C content or WC-binder diffusions. Hence, the

sintering mechanism must include solid state sintering mechanisms.

2.5.2.1 Solid state sintering

There are some theories reported concerning with the solid state sintering

mechanisms of WC-Co hardmetals (Meredith and Milner, 1976; Allibert, 2001; Silva

et al., 2001; Soares et al., 2012). The reports have revealed that high densification of

aggregate particles occurs at the temperature below eutectic point (the formation of

liquid phase) as shown in Stage I in Fig. 2.10 (Meredith and Milner, 1976). The

28

eutectic point changes with the volume fraction of Co (approximately 1300oC for Co

volume fractions using by their work as seen in Fig. 2.10).

Figure 2.10 Stages of densification of a range of WC-Co alloys; Stage I below

eutectic temperature; Stage II densification at eutectic temperature and Stage III

subsequent densification (Meredith and Milner, 1976)

The densification of the carbide particles, which takes place as the results of

enhanced surface and interfacial diffusing forming close-packed boundaries, will

first occur in Co-rich regions around Co particles leading to localized densification

or aggregation of carbide particles. And the initial size of such aggregates depends

on the size and distribution of Co (Meredith and Milner, 1976). The high shrinkage

rates were reported to depend on the initial size of WC particle and over 90% of

densification of the submicron or nanocrystalline WC–Co powder can be achieved at

solid state temperatures during heat-up prior to the formation of a liquid phase (Arató

et al., 1998; Gille et al., 2002; Fang et al., 2005). Contents of Co have been also

reported to effect of the solid state sintering. High Co content grades could be

29

successfully sintered at solid state to full densification (higher than 90% with 12

wt.% Co), with excellent microstructural quality and an effective control of the grain

growth. However, at lower Co contents, 3.5 wt.% Co, about 84% of densification

was reported (Soares et al., 2012). Full densification can be achieved by solid state

sintering with the application of external pressure (Huang et al., 2008c).

The investigations of microstructure evolution during solid state sintering

have been carried out somewhere else (Silva et al., 2001; Wang et al., 2008). Fig.

2.11 shows the changes of WC grains in WC-10Co sintering at different

temperatures (Wang et al., 2008). Fig. 2.11a and b shows a little change in WC

grains from initial powders to sintered grains at 800oC. With increase of temperature

from 800oC to 1000oC, the size of primary particles within aggregates increased

significantly (Fig. 2.11c). Between 1000oC and 1100oC (Fig. 2.11d), the aggregates

developed a multi-faceted surface morphology that has the layered structure of

crystal platelets, suggesting a preferred orientation of particles during aggregation.

At 1200oC (Fig. 2.11e), the aggregates have become much larger grains with angular

shapes and smooth surfaces. At 1300oC, the grains have grown to several hundred

nanometers (Fig. 2.11f). At this temperature, the rapid grain growth was stated that

liquid phase has occurred (Wang et al., 2008).

30

Figure 2.11 SEM images of morphology evolution of WC-10Co_10 nm

powders when heated at different temperatures: (a) as-milled power, (b) 800oC, (c)

1000oC, (d) 1100o C (e) 1200oC, and (f) 1300oC.

Complete densification of WC-Co occurs when liquid phase of the binder is

formed. The stages II (intermediate stage) and III (final stage) in Fig. 2.9 occur

during liquid phase sintering (LPS). Stage II corresponds to collapse and filling of

voids at the eutectic temperature when the Co dissolves carbide and forms eutectic

liquid that flows into the void space. In this stage rapid densification occurs by the

high shrinkage volume fraction. Further sintering occurs more slowly in stage III by

diffusion leading to densification of the aggregates formed in stage I (Meredith and

Milner, 1976).

31

Recently, a general model was proposed to analyze densification kinetics

during solid state sintering of either a single phase material, or a binary mixture of a

hard phase with a minor soft phase (Missiaen and Roure, 1998). The shrinkage rate

of WC-Co system, ἐ, can be expressed as a function of temperature, physical

parameters of the material, and geometrical parameters of the microstructure, when

the dominant mechanism for densification is the volume diffusion in the Co binder

phase as in equation 2.1 (Missiaen and Roure, 2000):

έ

(

) (2.1)

where is the Co surface tension, Co the average atomic volume in the binder phase,

DV the volume diffusion coefficient of the diffusion-limiting atom in the binder

phase, k - the Boltzmann constant and T the temperature. SV (WC:Co) and SV (P:Co)

are, respectively, the WC-Co and pore-Co specific surface areas per unit volume. <

H (P:Co) > is the mean curvature of the pore-Co surfaces. < L > is the mean

centre-to-centre distance between WC particles which are bounded by the binder

phase, and < x > is the mean radius of a slice of the binder phase as shown in Fig.

2.12.

Figure 2.12 Schematic representation of WC particles bounded by the Co binder phase (Missiaen and Roure, 1998).

32

The evolution of stereological parameters in densification during solid state

sintering has also been proposed (Missiaen and Roure, 2000). The decrease of the

pore surface area with density is practically linear. The pore-Co surface area

variation is homothetic, with a proportionality constant equal to the Co-volume

fraction in the solid mixture, which indicates a uniform rearrangement of the two

solid phases around pores. The WC-Co surface area evolution is particular of the

binary mixture and corresponds to a minimization of the surface energy by solid state

“wetting” of WC grains by the binder phase. This surface area first increases rapidly,

then slowly, as WC grains are gradually covered. The total mean pore curvature is

the sum of the pore-WC and pore-Co curvatures. It remains negative over the whole

density range, due to the convex shape of grains of the major WC phase. On the

contrary, the pore-Co mean curvature is positive over the whole density range, which

is evidence of a good “wetting” of WC grains by the Co-binder phase.

A schematic of the hard metal solid-state sintering mechanism is also

proposed as seen in Fig. 2.13a,b,a,d,e (Silva et al., 2001). In every site where a Co

particle locates, there is a formation of a WC-Co agglomerate. A group of WC

particles formed around each Co particle (Fig. 2.13a). A thin Co layer spreads over

the WC particles and links the WC particles around it (Fig. 2.13b). Since Co spreads

further, the WC-Co agglomerate is formed (Fig. 2.13c) and dense WC-Co

agglomerates are produced (Fig. 2.13d). When the WC-Co agglomerates link

together (Fig. 2.13e), the agglomerates grow in size and a network is formed. In the

process, the movement of WC particles to the center of the agglomerate forms a large

peripheral porosity around each agglomeration. The larger the agglomerates the

larger are the peripheral porosity. The size of the agglomerates is co-determined by

the WC particle size. The closure of the peripheral porosity depends on the

33

deformation of the agglomerates and on the enlargement of the contact area between

them. The larger the peripheral porosity the slower this stage proceeds (Silva et al.,

2001). The formation of WC-Co agglomerates and the spreading of Co binder during

solid state have been also confirmed later by sintering WC-(Co, Ni, Fe) at 1050oC

under argon as shown in Fig. 2.13f, the spreading of Co on WC plate at 1050oC (de

Macedo et al., 2003).

Figure 2.13. Schematic of the hardmetal solid state sintering mechanism (a, b, c, d, e)

(Silva et al., 2001) and (f) Cobalt spreading on WC plate at 1050oC under argon,

1-WC, 2-WC/Co region and 3-Co partice (de Macedo et al., 2003)

2.5.2.2 Liquid phase sintering (LPS)

In industrial practice, conventional WC-Co pressed powders are sintered by

liquid phase sintering at temperatures between 1300°C and 1500°C ensuring nearly

100% dense and strong products (Soares et al., 2012). From the technical point of

view, LPS is very attractive sintering method as it provides faster sintering and

complete densification without the need of any external pressure. Higher sinterability

is due to the enhanced atomic diffusion in the presence of liquid phase which

ultimately facilitates material transport. The minimum criteria for successful liquid

(f)

34

phase sintering are: (i) a low temperature liquid; (ii) high solubility of solid in liquid

binder, and (iii) good liquid wetting of the solid grains (Bhaumik et al., 1996;

Upadhyaya, 2001).

The major densification is achieved during the first stage of liquid phase

sintering. The good densification characteristics of WC-Co are due to the initial

solution of the carbide by Co, which results in an extensive densification even before

the first liquid is formed (Bhaumik et al., 1996; Silva et al., 2001). When the liquid is

formed (stage II in Fig. 2.10), further formation and spreading over the grains is fast

due to surface energy reason. As soon as the liquid penetrates the particle boundaries,

a capillary force is developed which leads to particle rearrangement for closer

packing in the presence of liquid phase. Although solution and reprecipitation of

solid occurs concurrently with rearrangement, the rearrangement events dominate the

early stage of densification (Bhaumik et al., 1996). Solubility of the solid in the

liquid further aids rearrangement for closer packing in the presence of a liquid phase.

In WC-Co hard metals, the WC particles are mostly irregular in shape. As the

capillary force acts to bring the particles closer, the misalignment of the center of

gravity gives a torque. This torque leads to rapid particle rearrangement, bringing flat

surfaces into contact. This is observed for a low volume fraction of liquid where the

neighboring contacts do not merge (Upadhyaya, 2001; German et al., 2005).

Once the liquid has spread, this allows repacking of rounded grains and

allows for densification relatively quickly. Liquid spreading rates can be in the

micrometers per second range, making the densification events rapid (German et al.,

2005). Grain growth, shape accommodation, and densification occur simultaneously.

Pore filling takes place and leads to a decrease of the amount and size of the pores as

35

a result of grain size increase process. At the end of this stage, pores have either been

eliminated or stabilized by a trapped internal atmosphere (German et al., 2005).

The maximum amount of densification attainable due to rearrangement is

influenced by some important factors, such as the amount of liquid present, particle

size, solubility of the solid in liquid and contact angle (Froschauer and Fulrath, 1976;

Bhaumik et al., 1996; Upadhyaya, 2001; Soares et al., 2012). Table 2.2 shows the

characteristic stages of liquid phase sintering (German et al., 2005).

Table 2.2. Characteristic stages of liquid phase sintering (German et al., 2005)

Phase Volume of liquid for

full densification (%)

Time frame

Rearrangement 25-35 Spontaneous

Contact dissolution 18-25 Seconds

Solution-reprecipitation 5-15 Minutes

Solid skeleton sintering 0.1-5 Hours

In general, nearly full densification of WC-Co is achieved before the final

stage of liquid phase starts, and further holding time does not lead to a significant

changing in densification (German et al., 2005). However, microstructural changes

of practical importance take place during the final stage. Grain size and size

distribution, grain shape, and as well as the binder phase distribution occur.

Abnormal grain growth may appear during this stage caused by the small amount of

coarse carbides acting as seeds for rapid grain coarsening. These changes effect on

properties like wear resistance, strength, fracture toughness, magnetic properties and

ductility. The improvement in density is highly dependent on the characteristics of

pores and any internal gases trapped in the pores.

36

2.5.3 Grain growth

Mechanical properties of hardmetals rely on their microstructure, including

WC grain size and distribution of components. The final grain size distribution of the

sintered product is affected at almost every stage in the production from the purchase

of raw materials to the final sintering of the materials. During the sintering step some

of the processing parameters that influence the final grain size distribution are: initial

grain size distribution; binder content, sintering time, cooling rate, sintering

temperature and C/W ratio in the binder phase (Chabretou et al., 1999).

The grain-growth process starts already at the preliminary solid-state

sintering stages, through adjusting and matching the grains orientations and grain

boundaries migration by thin film migration. Then, grains of different sizes fuse into

a single grain in a continuous process of directional growth and grain shaping. Later,

at liquid phase sintering stage, grain growth intensifies through an Ostwald-ripening

process. Smaller particles dissolve preferably due to their higher chemical potential,

and then precipitate in the neighboring of coarser particles reducing the system

interface area (Upadhyaya, 1998; Upadhyaya et al., 2001; Wittmann et al., 2002;

Wang et al., 2008; Soares et al., 2012).

The driving force for the growth of WC grains in a Co-based liquid is the

overall interface energy decrease. Obviously, the smaller the initial WC grain size,

the larger the interface area and the larger the driving force (Chabretou et al., 2003).

The contributions of the WC/Co and WC/WC interfaces for the several WC crystal

planes could vary and induce a slightly different growth process depending on the C

content (Wang et al., 2002).

37

To obtain microstructure refinement and increase mechanical properties, the

use of initial powders of ultrafine or nano powders is frequent (Spriggs, 1995;

El-Eskandarany et al., 2000; Gille et al., 2002; Zawrah, 2007; Xi et al., 2009).

However, during sintering, materials prepared from submicron size or nano size WC

powders undergo abnormal grain growth (Park et al., 1996) that affect mechanical

properties. Hence, small amount of grain growth inhibitors like VC, NbC, TaC,

Cr2O3 or other carbides are commonly added to control microstructure (Kim et al.,

1997; Xueming et al., 1998; Lee et al., 2006; Huang et al., 2007; Soleimanpour et al.,

2012). The mean grain size of WC increases with increasing of sintering time and

sintering temperature, even with added grain growth inhibitor (Xueming et al., 1998;

Choi et al., 2000; Wang et al., 2002; Chabretou et al., 2003).

Several model of WC grain growth have been proposed. A model of

continuous (normal) grain growth and discontinuous (abnormal) grain growth of WC

in WC-Co has been developed based on the Monte Carlo computer simulation

technique (Kishino et al., 2002). The results of these simulations qualitatively agree

with their experimental ones and suggest that distribution of liquid phase and

WC/WC grain boundary energy, as well as contamination by coarse grains are

important factors controlling discontinuous grain growth in cemented carbides. A

theoretical calculation and experiment of WC grain growth in WC–Co cemented

carbide have been studied (Sun et al., 2007). The results have showed that the grain

size increases with increasing temperature and sintering time in experiment and in

agreement with calculation model. Experiments and simulations of abnormal grain

growth have been also studied for WC-10Co system (Mannesson et al., 2011).

However, the simulations in their study give only a satisfactory prediction for short

sintering times similar to the ones used in industrial practice. For long sintering time,

38

the model cannot predict the effect of the large grains and the growth is strongly

underestimated.

The growth rate of the WC in WC–Co hardmetals is remarkably increased in

high C alloys (stoichiometric and over-stoichiometric alloys) compared with low C

alloys (sub-stoichiometric alloys) (Wittmann et al., 2002). The morphology and

growth of the WC grains are also affected by the C/W ratio in the binder phase (Park

et al., 1996; Wang et al., 2002; Chabretou et al., 2003). The growth rate of the WC

grains is observed to increase with the C content in the field {WC + liquid} in phase

diagrams. This trend is not significantly modified in the domains where graphite or

M6C are present (Wang et al., 2002). The WC average grain size is larger in the

C-rich side (Chabretou et al., 2003; Lay et al., 2008). When the binder is enriched in

C, WC grains are faceted and grain growth is slightly smaller in alloys with W-rich

binder, and significantly reduced in alloys containing -phase and/or V or Cr. The

growth rate and the morphology of the WC/interfaces, different in C-rich or W-rich

alloys, suggest that the steps controlling the WC nucleation and growth should

change with the binder composition (Wang et al., 2002; Chabretou et al., 2003).

2.5.4 Grain growth inhibitors

In general, the finer of WC grain size and the better of WC grain size

distribution are often expected to get the higher mechanical properties of hard metals.

The conventional sintering method (vacuum sintering) takes long time to get full

densification, therefore accelerates the grain growth of WC and consequently,

reducing the mechanical properties. Advanced sintering processes have been

reported being used to control WC grain growth such as hot pressed (Jia et al., 2007),

hot isostatic pressed sintering (HIP) (Azcona et al., 2002; Soares et al., 2012), spark

39

plasma sintering process (SPS) (Eriksson et al.; Zhang et al., 2004b; Sivaprahasam et

al., 2007; Zhao et al., 2009), microwave sintering (Breval et al., 2005), high

frequency induction-heated sintering (Kim et al., 2004; Kim et al., 2007) or

combined sintering processes (Sánchez et al., 2005). These methods are efficient to

produce finer microstructure of WC-hardmetals with shorter sintering time. However,

they also require complicated equipments and are high price and hence, they are not

commonly used as comparison to conventional method (liquid phase sintering in a

vacuum furnace).

The use of appropriate grain growth inhibitors has commonly used to prevent

WC grain growth during sintering. Small amount of cubic carbides such as VC, NbC,

TiC, Cr2C3, TaC has been used as the grain growth inhibitor to WC. These additives

reduce the growth rate of WC grains by lowering the solubility of WC in Co-rich

phase during liquid phase sintering. Higher solubility of an inhibitor in Co liquid

phase causes less solubility of WC and hence, leads to a finer grain structure (Seo et

al., 2003; Mahmoodan et al., 2009). Growth inhibition is more pronounced with

increasing amounts of inhibitor carbide. However, there is a critical amount of added

carbide for optimal effect on microstructure and so the mechanical properties of

WC-Co hardmetals. Higher than this amount, no further grain growth inhibition

occurs. This level is considered to correspond to the maximum solubility of the

carbide phase in liquid of Co (Yao et al., 1998).

VC has been found to be the most effective grain growth inhibitor and

followed by the order of VC>Mo2C>Cr2C3>NbC>TaC>TiC>Zr/HfC (Kim et al.,

1997; Yao et al., 1998; Lin et al., 2004; Morton et al., 2005; Huang et al., 2007). The

content of VC in WC-Co is usually kept at 0.7 wt.% and not more than 1wt.%, which

is regarded as the practical upper limit, in order to avoid embrittlement due to

40

(V,W)C precipitation at the interface of WC-Co (Hashe et al., 2007). The addition of

VC has been reported to increase the edge energy of WC crystals in comparison with

pure WC-Co and consequently, the coarsening process of WC grains is greatly

inhibited (Lee et al., 2003). While, a thin VC layer deposited on various WC crystal

faces with an epitaxial relationship during early sintering process has been attributed

to the effectiveness of VC to inhibit WC grain growth (Lay et al., 2003). The

formation of such VC film on WC/Co interfaces were recently modeled (Johansson

and Wahnström, 2010).

The use of VC in WC-Co hardmetals has been also reported to reduce

-phase formed by the formation of (V,W)C, which at the same time favors an

increase in the contiguity of the carbide network up to an approximately 50 wt% WC

substitution (Arenas et al., 1999). The effect of 1wt.% VC as grain growth inhibitor

to WC-30Co sintering in a vacuum furnace for 1h at 1400oC as shown in Fig. 2.14

(Choi et al., 2000). The presence of 1 wt.% VC has enabled to inhibit WC grain

growth and at the same time reduced the abnormal grain growth (coarsening) of WC.

Thus, the microstructure of WC-Co sample with 1 wt.% addition exhibited a more

uniform and finer microstructure.

Figure 2.14. SEM images of a) WC-30Co and b) WC-30Co-1VC sintered at 1400oC for 1h in a vacuum furnace (Choi et al., 2000)

a b

41

VC has also been considered being the most effective grain growth inhibitor

in WC-Ni hardmetals, followed by TaC, Cr3C2, TiC and ZrC. It was also reported for

WC-Ni alloys a significant growth inhibition in low C (W-rich binder environment),

in particular in alloys with -phase formation (Wittmann et al., 2002).

The effect of NbC on WC grain growth has been investigated by several

studies (Da Silva et al., 2000; Acchar et al., 2004; Huang et al., 2007; Huang et al.,

2008a; Huang et al., 2008b; Weidow et al., 2009b; Xiao et al., 2009). NbC has been

shown to play a role as an inhibitor to WC grain growth up to 5 wt.%, resulting in an

increase of room temperature hardness, toughness and flexural strength compared to

plain WC-12Co hardmetal. At higher NbC contents, the hardness, toughness as well

as bending strength of WC-NbC-Co hardmetals decreases with increasing NbC

addition up to 60wt.% due to the formation of larger size brittle (Nb,W)C grains

(Huang et al., 2008a; Huang et al., 2008b).

The use of individual TaC or combined with other carbides such as VC, TiC,

NbC to work as a grain growth inhibitor have been investigated (Kim et al., 1997;

Morton et al., 2005; Barbatti et al., 2009; Mahmoodan et al., 2009; Weidow et al.,

2009b). Individual TaC has less effect on WC grain growth inhibition than VC and

NbC. However, the combination of TaC with VC, NbC or TiC leads to a stronger

decreasing of grain size and hence, resulting in higher hardness and toughness of

WC-10Co, (Kim et al., 1997; Cha et al., 2001a; Soleimanpour et al., 2012).

The addition of other carbides or combined carbides such as Cr2C3, Mo2C,

TiC and Zr/HfC for WC grain growth inhibition has also reported in several studies

(Cha et al., 2001b; Lay et al., 2002; Hashe et al., 2007; Guo et al., 2008). The

additives of rare-earths such as Y2O3 or La2O3 in WC-Co have been carried out (Liu

42

et al., 2006; Zhang et al., 2008). The results have showed a promising of these

compounds to work as a grain growth inhibitor to WC. With small amount of

phosphorus (P) doped in nanocrystalline WC–Co matrix, the sinter activity of

WC–Co nanocomposites was enhanced, the low temperature shrinkage and

densification process was promoted, and the sintering temperature was decreased and

the grain growth of the WC–Co cermets could be inhibited to some extent (Zhang et

al., 2004a). The effect of cubic boron nitride (cBN) addition on densification, grain

growth inhibition and properties of WC-Co has been studied (Wang et al., 2012).

The ability of cBN as a grain growth inhibitor has explained due to the decrease of

grain boundary mobility with the addition of cBN.

2.6 Sintering method

Sintering is very important step to consolidate WC-based hardmetals into a

high-strength structure. After pre-compaction to produce the green bodies of mixture

powder of WC/binder, the samples are heated up at high temperature (1300-1600oC)

in order to obtain dense products. In history of WC-based hardmetals, two basic

methods have been used to consolidate these materials (Fernandes, 2008).

a) Hydrogen sintering uses a hydrogen-based atmosphere at atmospheric pressure

to dynamically control composition.

b) Vacuum sintering uses a vacuum or reduced-pressure environment for the same

effect.

Until now, many advanced techniques have been applied to improve the

densification of sintered WC-based hardmetals. Hot pressing, hot isostatic pressing

43

(HIP), spark plasma sintering (SPS) or high frequency induction-heated has been

used in order to lower sintering temperature and higher density of sintered WC-based

hardmetals and therefore improves mechanical properties of products such as

hardness and toughness.

2.6.1 Vacuum sintering

Sintering atmosphere also has an important effect on sintering behavior of

hardmetals. Vacuum sintering method has several advantages in comparison with

hydrogen atmosphere sintering such as: (i) allowing superior control of product

composition; (ii) the exchange rate of C and oxygen between the atmosphere and the

cemented carbides is very low and (iii) the main factor controlling composition is the

amount of oxygen of the carbide powder, not the rate of reaction with the atmosphere

(Bhaumik et al., 1996; Fernandes, 2008). A higher density and less residual porosity

have been reported to obtain in WC-Co doped TiC hardmetals by sintering in a

vacuum condition than by hydrogen sintering (Bhaumik et al., 1996).

So far, vacuum sintering is still the most popular sintering method in WC-based

hardmetal processing because of these advantages and the economic aspect as well as

the advantage in production of large numbers.

2.6.2 Hot pressing

Hot pressed sintering is a method to sinter samples in which the temperature

and pressure are applied at the same time. This method is to enhance the density and

shorter the sintering time and therefore improves the mechanical properties of

sintered samples (El-Eskandarany et al., 2000; Jia et al., 2007). The mixture powders

are consolidated directly in a hot pressed mold or they are pre-compacted into a

44

green body and then, the green body is consolidated by hot press equipment. Heating

methods include electric resistance furnace, spark plasma sintering or pulsed current

activated sintering (Kim et al., 2007, Eriksson et al.; Lee et al., 2003; Kim et al.,

2004) under vacuum or reduction environments.

2.6.3 Hot isostatic pressing

Hot isostatic pressed (HIP) sintering is an advanced method to consolidate

WC-based hardmetals. To achieve high performance of this method, the powder need

to be pre-compacted in high density or put in a gas-tight encapsulation to inhibit the

gas penetration, and subjecting the system to the desired temperature and pressure

(Upadhyaya, 2010). Previous studies have shown that nearly full density of WC-Co

hardmetals can be achieved at low sintering temperature and shorter sintering time

by HIP process (Azcona et al., 2002; Soares et al., 2012), consequently provide an

improvement in microstructure and mechanical properties of hardmetals. In recent

years, HIP has been more rapidly expanding in research and industries despite of the

high cost of equipments.

2.6.4 Pseudo hot isostatic pressing

Pseudo hot isostatic pressing (PHIP) has been reported being used as an

assisted sintering method. This method is less expensive and simpler in processing

than HIP method. In general, this method is similar to hot pressing method. The

pressure is transferred directly into the sample by upper punch or lower punch of the

hot pressed mold in hot pressing method, however, the pressing force is transferred

to sample through a pressing medium (silica sand) in PHIP method. The moving of

45

silica sand contacts with sample by pressure similar to the moving of gas in HIP

process and therefore the sample is pressed, so called PHIP. Until now, this method

is applied to produce high density intermetallic compounds (Ishihara and Shingu,

1990; Park et al., 1997). Ti2AlC compound has been fabricated approximately 97%

of theoretical density after using PHIP at 160 MPa for 5s pressing at 1600oC

(Xinghong et al., 2002) and even higher, 97.5% at 420 MPa for 15s pressing (Bai et

al., 2012). Ti-46Al-(Cr, Nb, W, B) alloy was fabricated by PHIP process at

temperature higher than 1100oC (Zhang et al., 2010). The results show that high

quality compressed samples without cracks can be obtained. However, no report of

using PHIP to sinter WC-based hardmetals has been done. Considering as an

effective method to sinter WC-based hardmetals, in present work, PHIP was used to

sinter several WC-AISI304 compositions in order to evaluate and compare the

microstructure and mechanical properties of hardmetals with vacuum sintering

method.

2.7 Mechanical properties of WC- based hardmetals

WC-based hardmetals are widely used for cutting applications such as

cutting inserts or drilling tools where the tools undergo heavy conditions, for

example high wear and erosion. Thus, the requirements for these materials are high

hardness, good wear resistance and high fracture toughness, etc. Hardness and

fracture toughness are the most interesting of the WC-based hardmetal mechanical

properties. In this section, hardness, fracture toughness and other mechanical

properties of WC-based hardmetals are discussed.

46

2.7.1 Hardness of WC-based hardmetals

Hardness and fracture toughness are the two most important mechanical

properties of WC- based hardmetals and other cermets. Other mechanical properties,

such as flexural strength, wear resistance and impact resistance, are fundamentally

dependent on the hardness and fracture toughness. The value of hardness also

provides a measure of the resistance of the ceramic to deformation, densification, and

cracking or fracture. Sintered hardmetal hardness is generally tested using either

Vickers diamond pyramid indentation, HV, or the Rockwell A-scale (HRA)

diamond-cone indentation (Kang et al., 2000; Morton et al., 2005). The hardness can

be effected by the content of binder phase and microstructural parameters, such as,

porosity, carbide grain size and contiguity/binder mean free path (Kang et al., 2000;

Srivatsan et al., 2002; Daoush et al., 2009). Increasing the binder content leads to

decreasing hardness values, because the hard WC phase is replaced by ductility

binder phase (Daoush et al., 2009, Ettmayer, 1989; Deng et al., 2001) as seen in Fig.

2.15.

47

0 5 10 15 20 25 30

0

500

1000

1500

2000

2500

MeasuredH

ard

nes

s (H

V -

kg

/mm

2)

Co content (Vol.%)

Published range

Figure 2.15. The relationship between Vickers hardness and volume percent Co (Daoush et al., 2009)

The hardness of hard metals increases with decreasing the carbide grain size

and the binder mean free path, and increasing the contiguity of the carbide phase

(Kim et al., 1997; Gille et al., 2000; Carpinteri et al., 2009).

Several models relating to the hardness and the microstructure parmeters of

WC-Co system have been reported as seen in Table 2.3. A phenomenological grain

size-hardness relationship for WC-10Co cemented carbides is formulated (Eq. 2.2)

(Cha et al., 2001b). The hardness has been calculated based on the in-situ hardness of

the hard phase (WC), binder phase (Co) and their volume fractions for WC-Co

hardmetal (Eq. 2.3) (Lee and Gurland, 1978). The two in-situ hardness values (Eq.s

2.4 and 2.5) were assumed to obey the Hall–Petch relations (Niels, 2004) based on

the grain size (d), and the mean free binder path (λ), respectively. This expression is

the result of a data fit based on hardness measurements rather than on a theoretical

model. Another theoretical model has been also proposed in order to find a

relationship among hardness HV30, and the microstructure parameters, WC mean

grain size, dWC, and Co volume fraction, VCo (Eq. 2.6) (Gille et al., 2000). According

48

to this model one straight line correlates hardness HV30, dWC and VCo for 120

hardmetal samples within a wide spread range of parameters values (dWC = 0.10-7.80

m, VCo = 0.10-0.38).

Physical and empirical expressions of the relationship between HV30 and

dWC in WC-Co with different Co content have been proposed (Roebuck, 2006). The

calculations were based on the Hall-Petch relation (physical based, Eq.s 2.7 and

2.89) and a logarithmic (empirical, Eq.s 2.9 and 2.10) dependence. The results

indicated that the logarithmic expressions tended to give a better fit at fine and coarse

WC grain sizes, although all expressions were reasonable for the middle of the grain

size range.

Table 2.3 The relationship between hardness and microstructural parameters

Model expressions Eq. Parameters References

HV = 550 +

√ 2.2 dWC WC grain size (mm) (Cha et al., 2001b)

HC = HWCVWCC+HCo(1-VWCC)

HW C=1382+23.1 (kg/mm2)

HCo=304+12.7 (kg/mm2)

2.3

2.4

2.5

C - contiguity

VWC- volume fraction of WC

HC – hardmetal hardness

HWC –hardness of WC phase

HCo – hardness of Co phase

dCo – Co grain size (mm)

(Lee and Gurland,

1978)

HV30=1824(1-1.65VCo+0.92

) (dWC)-0.194

2.6 VCo -volume fraction of Co (Gille et al., 2000)

6% Co HV30 = 970 + 540

10% Co

HV30 = 850 + 485

6% Co HV30=1538+742 dWC

10% Co HV30 = 1391+ 598log10 dWC

2.7 2.8 2.9 2.10

(Roebuck, 2006)

49

2.7.2 Fracture toughness of WC-based hardmetals

The fracture toughness or the resistance to crack propagation of a material is

measured by the critical intensity factor, KIC. Fracture toughness of WC-hardmetals

depends on the microstructure parameters such as the contiguity of the WC skeleton,

volume fraction of binder phase, the mean free path length of the binder and the

grain size of WC. In general, fracture toughness increase with the increased

contiguity, the volume fraction and the mean free path of binder phase or the

decreased contiguity of WC phase (Chermant and Osterstock, 1976; Ettmayer, 1989;

Yao et al., 1998; Deng et al., 2001; Shatov et al., 2008). Fig. 2.16 shows the effect of

mean linear path in binder phase () and with contiguity of carbide phase at different

WC mean diameters on the fracture toughness of WC-5Co and WC-7Co. And for a

given volume fraction of binder phase, geometrical arrangement of the ductile binder

as a continuous matrix phase is beneficial for high toughness while retaining high

strength.

Figure 2.16 (a) The correlation of the fracture toughness KIC with the mean linear path in binder phase λ, and (b) with the carbide crystals contiguity G (Shatov et al.,

2008)

50

Generally, the higher hardness will be achieved with the finer WC grain size,

however, the fracture toughness also decreases with the increase of hardness as

shown in Fig. 2.17 (Vaßen and Stöver, 1999). At a given hardness, the finer of WC

grain size leads to a higher of fracture toughness (Chermant and Osterstock, 1976).

Figure 2.17. The correlation between hardness and fracture toughness of WC-Co and

with added cubic carbides (Vaßen and Stöver, 1999)

The measurement of toughness may use the three-point bending method or

calculated based on the values of load (P) and crack length generated during hardness

testing. When the indentation process produces the cracks at the corners of the

indentation, the fracture toughness is measured relied on the applied load and the

crack length, l, called the Palmqvist type of cracks as seen in Fig. 2.18 (Niihara,

1983).

51

Figure 2.18 Schematic of Vickers indentation cracks and Palmqvist cracks: d-

diagonal of the indentation left in the surface; l- Palmqvist crack length (Niihara,

1983)

The Palmqvist method is recently used to calculate the fracture toughness

because of the simple method itself and the useful method in application to small

samples (Peters, 1979). Since the method was invented, several studies have been

carried out to find the optimal relationship between fracture toughness and the other

parameters of materials (Table 2.4). The expressions of Palmqvist cracks method

have been developed toward the correlation of fracture toughness, Vickers hardness,

Young’s modulus (E), Palmqvist crack length (l), and indent half diagonal (a) as seen

in Eq.s 2.11 and 2.12 (Niihara et al., 1982; Niihara, 1983). Later, more simplified

expressions (Eq.s 2.13 and 2.14 ) have been done by using the impirical curve-fitting

technique (Shetty et al., 1985; Ponton and Rawlings, 1989). Since, Vickers hardness

testing have been often carried out at the load of 30 kg for WC-Co hardmetals, the

fracture toughness has been calculated by Eq. 2.15 (Schubert et al., 1998).

52

Table 2.4. Expressions for fracture toughness (KIC) calculation from Vickers indentation crack systems

Expression Eq. Parameters References

(KIC./HV.a1/2)(HV/E. )2/5= 0.035(l/a)-1/2 2.11 KIC - Fracture toughness HV-Vickers hardness

E - Young’s modulus

a - indent half diagonal

l - Palmqvist crack

length

(Niihara et al., 1982)

(KIC./HV.a1/2)(H/E. )2/5= 0.048(l/a)-1/2 2.12 (Niihara, 1983)

KIC=0.089(HV.P/4l)1/2 2.13 P - applied load (Shetty et al., 1985)

KIC=0.0937(HV.W)1/2 2.14 W= P/LT LT - total length of cracks

(Ponton and Rawlings, 1989)

KIC=0.15(HV30/LT)1/2 2.15 (Schubert et al., 1998)

2.7.3 Other mechanical properties of WC-based hardmetals

Transverse rupture strength (TRS) has been an important property to

investigate the quality of WC-based hardmetals since the very beginning of

hardmetal technology. This is defined as the ratio of the bending moment to the

resisting moment of a rectangular bar that is loaded midway between two sintered

carbide supported cylinders with slowly force increasing until the fracture occurs.

This property is very sensitive to flows and imperfections in the test samples and

therefore is a good indicator for general “housekeeping standards” during the

manufacturing (Ettmayer, 1989). Since the transverse rupture strength is less well

suited for evaluating the inherent strength of experimental materials, the fracture

toughness (or the critical stress intensity factor, KIC) is in this case having more

indicative of the potential strength of the hardmetals (Ettmayer, 1989).

53

Compressive strength is one of the unique properties of WC-based

hardmetals. A uniaxial force is applied to compress on the straight cylindrical sample

or on cylinders having reduced diameters in the middle to localize the fracture.

WC-based hardmetals possess high compressive strength at room temperature

however it also decreases with the increasing of temperature; the rate of decrease

depends on the binder volume fraction and the microstructure.

The wear resistance property of WC-based hardmetals is a functional

property that depends on the fracture toughness and hardness of the hardmetals. The

wear environment and mechanism also effect on the wear resistance of the materials.

It is very important to note that the wear mechanisms have a significant role in wear

tests. Different wear applications (wear environments) have different wear

mechanisms. The wear resistance of WC-based hardmetals varies with the effects of

composition, microstructure and the real applications (Gee et al., 2007).

2.8 Mechanical alloying

Mechanical alloying (MA) is a powder metallurgy processing technique

involving cold welding, fracturing, and re-welding of powder particles in a

high-energy ball mill. In history of WC-based hardmetals, MA is the most popular

technique to produce mixture powders of WC-based hardmetals from raw materials

(WC and binders) before sintering. During the development of powder metallurgy

technology, MA has been developed in order to fabricate fine, ultrafine powder and

nano-powders of materials with high homogeneity. Nano-WC was produced with the

particle size less than 5 nm by high energy ball-mill for 120 h and ball to powder

ratio 10:1 (El-Eskandarany et al., 2000). 10 nm WC-6Co composite powders was

attained by mixing in a planetary ball mill for 100 h under argon atmosphere

54

(Xueming et al., 1998). In this research, MA was used to produce fine WC-AISI304

powders with good distribution between WC and AISI304 metal powders. This

section will introduce some basic information about MA process.

2.8.1. Introduction of mechanical alloying

Mechanical alloying (MA) has been developed since the years of 1970s

(Benjamin, 1970). This is a useful technique for producing composite powders with

controlled microstructures. In the beginning of MA history, this method was used to

produce supperalloys called oxide dispersion strengthened nickel-based alloys

(Benjamin, 1970). This solid-state powder processing technique involves repeated

welding, fracturing and re-welding of reactant mixed powder particles resulting from

the heavy deformation of the high energy ball-powder collisions (Benjamin and Volin,

1974). This technique can be divided into five stages: the initial period, the period of

welding predominance, the period of equiaxed particle formation, the start of random

welding orientation, and steady-state processing (Benjamin and Volin, 1974).

MA technique has been become an advance method to fabricate wide ranging

of materials including equilibrium and non-equilibrium alloy phases starting from

blended elemental or prealloyed powders (Suryanarayana, 2001). A number of

investigations have been carried out to synthesize a variety of stable and metastable

phases including supersaturated solid solutions, crystalline and quasicrystalline

intermediate phases, and amorphous alloys using MA method. A mixed powder of

nominal composition 40 at.% Mg gradually was used MA to convert into a

metastable supersaturared fcc Al(Mg) solid solution having approximately 23 at.%

Mg in solution (Zhang et al., 1994). Nanocrystalline (Ni,Fe)3Al intermetallic

compound was synthesized from Ni, Fe and Al powders by 80 h mixing in a

55

planetary ball milling under Ar gas (Adabavazeh et al., 2012). The microhardness of

(Ni,Fe)3Al phase produced by their work has been reported to be about 1170 HV

which was attributed to the formation of nanocrystalline (Ni,Fe)3Al intermetallic

compound. The formation of nano quasicrystalline icosahedral phase in Al86Cr14,

Al84Fe16 and Al62.5Cu25Fe12.5 alloys with 25-50 nm in size has been produced by

milling in a attritor vial with ball to powders ratio of 10:1 in weight and milling

speed of 500 rpm (Shamah et al., 2011). Nanostructured Ti6Al4V phase with a grain

size of 20-50 nm were fabricated by high energy mechanical milling at 90 h and

obtained higher hardness value of 630 HV (after heat treatment) which is higher than

those reported for Ti6Al4V alloys processed by conventional methods (Mahboubi

Soufiani et al., 2012). Nanostructures of WC-Co powders with 10 nm crystalline size

directly synthesized from pure powders, and given higher mechanical properties of

WC–6wt%Co–1wt%VC have been fabricated by prior mechanical alloying and

consolidation at different temperatures (Xueming et al., 1998).

2.8.2 Mechanical alloying mechanisms

During MA process, the powder particles are undergone a repeated sequence

phenomenon’s such as flattening, fracturing and rewelding. Whenever two hard balls

collide, small amount of powder is trapped in between them (Gilman and Benjamin,

1983; Suryanarayana, 2001) (see Fig. 2.19). The force of the impact plastically

deforms the powder particles, creates new surfaces, and enables the particles to weld

together, and hence, leading to an increase in particle size of the powders. At early

stages of milling, the soft particles have a tendency to weld together and agglomerate,

and forming large particles. The very wide range of particle sizes develops, with

several large as three times bigger than the starting particles. The composite powders

56

have a layered structure characteristic containing various combinations of the starting

components. Continuing the deformation, the particles become work hardened and

fracture by fatigue failure mechanism and by fragmentation. Fragmentation

generated by this mechanism may continue to reduce in size with the absence of

strong agglomerating force. In this stage, fracture dominates of the whole process.

Continuing impact of milling balls, the structure of the particles is steadily refined

but the particle size continues to be remaining. Consequently, the spacing of

interlayer decreases, while the layers increases in a particles. After milling for a

certain length of time, steady-state equilibrium is attained when the rate of welding

and the rate of fracturing achieve a balance. The average particle size obtained at this

stage depends on the relative ease with which agglomerates can be formed by cold

welding, fatigue and fracture strength of composite particles, and resistance of

particles to deformation.

Figure 2.19. Ball-powder-ball collision of powder mixture during mechanical

alloying (Gilman and Benjamin, 1983)

57

In WC-Co system, the quality of mixture powder depends on the parameters

such as milling time, ball to powder ratio (BPR), rotation speed, particle size of

starting powder, etc. The increase of BPR leads to decrease of particle size and

crystallite size of WC-10Co (Mahmoodan et al., 2009). This research also shows that

the 10 h milling with BPR of 5 did not produce enough energy to reduce the WC

particles to nanometer scale size. In general, the increase of milling time results in

the finer of particle size of WC and binder, however, the long of milling time leads to

the high agglomeration of WC with the binder phase that cause the sintering process

difficult to reach full densification (Torres and Shaeffer, 2010). Until now, MA is a

very useful technique to produce fine and ultrafine WC-based hardmetal powders,

even nano-powders.

2.8 Summary

WC-based hardmetals have been developed for nearly hundred years due to

their technological importance in cutting industry. Since the early year of these kinds

of hardmetals, Co has been commonly used as the binder. This is due to Co has a

good sinterability wetting with WC and WC-Co hardmetals which produced superior

mechanical properties compared to other binders. However, Co is no longer welcome

due to the depleting resources of Co, high cost and the toxicity of WC-Co containing

dust that is harmful to human body. Thus, there is an increasing attempt to find other

binders than Co. The present work was done to study the ability of ASIS304 stainless

steel to replace Co binder. WC-AISI304 hardmetal powders were produced by MA

technique. The use of MA is to coat WC particles by AISI304 binder phase with high

distribution between WC and AISI304 powders which support to sintering process

58

and result in high mechanical properties of sintered samples. Samples were sintered

using two different methods; vacuum sintering and PHIP sintering. Sintered samples

were measured for their density, mechanical properties (Vickers hardness and

fracture toughness), an investigated phase transformation and microstructure XRD

and FESEM, respectively. The research focused on the aspects such as; the effect of

MA technique, sintering time and temperature, the effect of binder phase, the role of

graphite and NbC additions to improve mechanical properties. This research also

investigated the effect of PHIP on the microstructure and mechanical properties of

sintered samples compared to vacuum sintering technique.

59

CHAPTER 3

RAW MATERIALS AND METHODOLOGY

3.1 Raw materials

WC and AISI304 stainless steel (FeCrNi) powders were the main materials

for this work. Graphite and NbC powders were used to improve the microstructure

and consequently, improve the mechanical properties of WC/FeCrNi hardmetals.

Physical, chemical and mechanical properties of these materials are also described in

this section.

3.1.1 Tungsten carbide (WC) powder

WC powder was supplied by Strem Chemical, USA with purity of 99.5%

and a particle size less than 1m. Some important physical properties of WC are

given in Table 3.1 (Stremchemical, 2012).

Table 3.1 Physical properties of WC (Stremchemical, 2012)

Properties Values

Chemical formula WC

Density (g/cm3) 15.63

Melting point (oC) 2870

Boiling point (oC) 6000

Crystal structure Hexagonal

60

3.1.2 Stainless steel powder

FeCrNi powders, grade AISI304 (FeCr18Ni10), was purchased from

Goodfellow Cambridge Co. Ltd., UK. The purity was of 99.5 % and an average

particle size of 45m. Stainless steel powder has a bright grey color and density of

7.93 g/cm3. Alloy AISI304 is a general austenitic stainless steel with a fcc structure.

Properties of AISI304 stainless steel are given in Table 3.2 (Goodfellow, 2012).

Table 3.2 Properties of alloy AISI304 (Goodfellow, 2012)

Properties Values

Chemical composition (wt.%) Fe (66.345-74.0), Cr (18-20), Ni (8-10.5),Mn (max 2), Si (max 1)

P (max 0.045), S (max 0.03) Density (g/cm3) 7.93

Melting point (oC) 1400-1455

Tensile strength (MPa) 460-1100

Modulus of Elasticity (GPa) 190-210

Rockwell B (HRB) 160-190

Electric resistivity ( m) (0.070-0.072)10-6

Thermal conductivity (Wm-1 K-1) at 23oC 16.3

Thermal expansion K-1 at 20-100oC 18

3.1.3 Niobium carbide (NbC) powder

NbC powder used was of 99.5% purity and a particle size less than 10 m.

Niobium carbide has a cubic crystal structure. It has a density of 7.82 g/cm3 and a

brown-gray colour of powder. Some important physical properties of niobium

carbide powder are given in Table 3.3.

61

Table 3.3 Physical properties of niobium (Metox, 2012)

Properties Values

Density at 20oC (g/ cm3) 7.82


Boiling point (oC) 4300

Crystal structure Cubic

3.1.3 Graphite powder

Graphite powder used was of 99.9% purity and a particle size of less than 10

m. Graphite has a black color and a density of 2.2 g/cm3. Some important physical

properties of graphite are given in Table 3.4.

Table 3.4 Physical properties of graphite (Habashi and Fathi, 1997)

Properties Values

Density at 20oC (g/ cm3) 2.2


Boiling point (oC) sublimes

Electrical resistivity ( . cm) 90.8

3.2 Research methodology

In this work, the hardmetal powders were fabricated by mixing from raw

materials; WC, FeCrNi, graphite and NbC powders, in a planetary ball milling

machine. The powders were then sintered either by a vacuum furnace or PHIP

sintering. The milled powders and sintered samples were investigated phase

components and microstructure by XRD and FESEM, respectively. Sintered samples

were measured Vickers hardness and fracture toughness using a Vickers hardness

62

tester. This section would introduce some concepts, equipments and sample

processing supporting for this work. The flow chart of sample procedure in this work

is shown in Fig.3.1.

Figure 3.1 Flow chart of the experimental procedure

63

3.2.1 Mixing of WC-FeCrNi hardmetal powders by planetary ball milling

WC and FeCrNi powders were mixed with different compositions in a high

energy planetary ball milling, model Pulverisette 5 (Fritsch Co.Ltd). In this work, the

aim of using ball milling is to coat WC particles by the binder phase and good

distribution between WC particles and binder phase. The process was performed

under argon atmosphere at room temperature with rotation speed of 200 rpm. The

250 ml of vessel and 20 mm in diameter of stainless steel balls were used for milling

process. The ball-to-powder weight ratio was selected 10:1 as previous studies (Seo

et al., 2003; Xiao et al., 2009), and in total about 30 or 40 g of samples was used. In

general, paraffin wax, polyethylene glycol or ethanol with 1 - 3 wt.% was used to

ovoid cold welding and also to work as lubricant for pre-compaction (Hanyaloglu et

al., 2001; Wittmann et al., 2002; Sánchez et al., 2005). In this work, 2 wt.% of

ethanol was added in all compositions before milling. To prevent overheat

appearance generated by the collision and friction phenomena’s between the ball to

ball and ball to jar during milling; the process was set up 30 min milling and 15 min

pausing. In addition, the effectiveness of milling is also achieved by automatic

rotating of the jar in reversed direction after each repetition running and pausing.

To study the effect of milling time on the microstructure of milled powder

and sintered samples, the mixture powder of WC-10FeCrNi was milled at 5, 10, 15

and 20 h, respectively. The milled powders were then pre-compacted at 300 MPa in a

10 mm diameter die and sintered at 1300oC in a vacuum furnace.

64

3.2.2 Green body compaction

The as-milled powders were then compacted to produce green bodies.

Stainless steel cylindrical dies with 10 mm for vacuum sintering and 18 mm

diameters for PHIP were used to make pellets of samples. Generally, the compaction

pressure in the range of 200 – 400 MPa is employed to impart sufficient green

strength to the compacts (Upadhyaya, 1998). Thus, pressure was selected at 300 MPa

(medium value) as previous studies (Huang et al., 2008b; Soleimanpour et al., 2012),

and dwelling time of 90 seconds. Pressing was done on a Specac pressing machine

with maximum 15 tons of load.

3.2.3 Sintering in a vacuum furnace

The compact samples were put in an alumina crucible and covered with

graphite powder in order to avoid decarburization at the surface of samples during

heat treatment periods. Sintering process was then performed in a vacuum furnace at

a pressure of 10 torrs. The schematic of heating process is shown in Fig. 3.2.

Generally, sintering cycle consists of pre-heated step to deplete lubricant or organic

compounds contained during milling or pressing. Pre-heated step is carried at low

temperature, ranging from 300 – 500oC (Allibert, 2001; Eso et al., 2007). In this

work, samples were pre-heated at 300oC for 1 h to deplete ethanol and other organic

compounds which may be contained in the samples during milling or pressing. Then,

the temperature was increased up to sintering temperature and holding at this

temperature for sample consolidation. After that, the samples were cooled to the

room temperature. All steps of sinter in process were controlled automatic using

programming.

65

Figure 3.2 Sintering diagram of samples in the vacuum furnace

3.2.3 Sintering by PHIP process

The aim of using PHIP is to get higher densification compared to vacuum

sintering by the applied pressure and hence, improve mechanical properties (hardness

and fracture toughness). The results would be compared with those obtained in

vacuum sintered samples. The as-milled powders were pre-compacted at 300 MPa to

produce 3 mm thick green body with 18 mm diameter mould. The green body was

then consolidated by PHIP at sintering temperature under vacuum condition. The

green body was placed in a die and covered with silica sand as pressing medium.

During the process, an electric coil element was used to heat the sample and

temperature was measured by a type-R thermocouple. The sample was PHIPed in a

vacuum condition. In the first step of heating process, sample was heated at 300oC

for 1 h to deplete all the organic matters or vapor and then heated again until the

sintering temperature. After holding at sintering temperature for 30 min, the pressing

process was carried out for 15 minutes at the applied pressure. The samples were

then cooled slowly in the furnace. The heating profile is shown in Fig. 3.3 and the

66

schematic of PHIP process as in Fig. 3.4. PHIP system using in this work is a SP-R

Pressing Machine system (ACT Electric Industry Co. LTD, Japan) in the School of

Energy Science’s laboratory, Kyoto University. Maximum pressure is 500 tons and

highest working temperature is 1500oC. PHIP system and sample fabrication process

are shown in Fig. 3.5.

Figure 3.3 Heating diagram of PHIP process

Figure 3.4. Schematic diagram of experiment set up for PHIP. (1) Die, (2) Heating coil, (3) Specimen, (4) Upper punch, (5) Lower punch, (6) Thermocouple, (7)

Electric circuit, and (8) Silica sand powder

67

Figure 3.5 PHIP system and sample fabrication process: 1 – Heating coil, 2 – Sample,

3 – Stainless steel mould, 4 – Silica sand, 5 – Themorcouple, 6 – Upper punch, 7 –

Putting in pressing chamber and connecting with pressing system, 8 – Pressure gauge,

9 – Pressing chamber, 10 – System controlling

3.3 Data analysis

3.3.1 Phase identification

X-ray diffraction (XRD) is a useful, nondestructive technique to analyze

chemical and phase structures of materials. Its application includes qualitative and

68

quantitative phase analysis, crystallography, structure and relaxation determinations,

texture and residual stress investigations, micro-diffraction and nano-materials

(Cullity and Stock, (2001)). In this research, XRD patterns of powders and sintered

samples were obtained using a Bruker D8 Advance diffractometer with CuK

radiation ( = 1.5418 Å). The 2 angle measured started from 10o and ends at 90o.

XRD patterns could be indexed and the phases present identified by comparing the

observed angular positions and intensities of the peaks with those of the standard

patterns available in the Powder Diffraction Files (ICDD files). Crystallite size was

calculated, in some cases, based on Scherrer’ equation (3.1).

(3.1)

where as: d - crystallite size, K – a shape factor of CuK equal to 0.9, λ = wavelength

of Cu Kα radiation equal to 1.5406 Å, θ = half of the diffraction angle (degree), β =

full width at half maximum (FWHM) (degree).

-phase quantification was calculated based on XRD result using Rietveld

refinement method. The data of XRD result was applied in Topas softwave with the

Rietveld refinement mode. This method is now populated to calculated phase fraction

using XRD data results.

3.3.2 Microstructure observation

The morphology of powder and microstructure of sintered samples were

investigated by Field Emission Scanning Electron Microscopy (FESEM). This work

was done on the FESEM equipment with model LEO Supra 55 VP. Both secondary

electron and back scattered electron detectors were used to study the microstructure

cosKd

69

of samples. To obtain useful information regarding to microstructure of sintered

samples, the samples were ground by SiC paper from grade 100 to 1200 and then,

polished by 0.25 m diamond paste to get high flat surface. After polishing, the

samples were cleaned by ultrasonic water bath for 10 min before being dried by a

dryer. The samples were placed on a carbon tape sticker with a holder and coating

their surfaces by gold metal in the case of using back scattered electron detector.

Energy dispersion X-ray spectroscopy (EDX) was also applied to analyze the atomic

composition and weight percentage of elements in the samples.

In this work, FESEM images of sintered samples was used to estimated WC

grain size of WC using AGI (Average grain-size intercept) method. This is a

technique used to qualify the grain – or crystal – size for a given material by drawing

a set of randomly positioned line segments on the micrograp, counting the number of

times each line segment intersects a grain boundary, and finding the ratio of

intercepts to the line length. Thus, the AGI is calculated as equation 3.2.

AGI = (number of intercepts)/(total line length) ( 3.2)

3.3.3 Density Measurement

The density of a body can be calculated based according to equation

= w1/V (Eq. 3.3)

where as , w1 and V are density, weight (in the air) and volume of the

body, respectively.

Archimedes' principle states that when a body is immersed in a fluid, there

is buoyant force acting upward on the body equal to the weight of the displaced

70

volume of fluid, which also equals to loss of weight of the body. Therefore, the loss

of weight of the body can be calculated base on Equation 3.4 as follow

m = VF.F (Eq. 3.4)

In which VF is the equal volume of displaced fluid (VF = V), F is the

density of the fluid and m = w1 – w2 is weight loss of the body (w2 is the weight of

the body in the fluid). Hence, the density of the body can be as follow equation (3.5)

=

21

1

www

.F (3.5)

In this work, the density of sintered samples was measured using

Archimedes’ principle by a Precisa gravimetrics AG equipment (D99-09-030,

Dietikon Switzerland). Distilled water was used as the fluid in equation (3.5).

3.3.4 Hardness testing

In this work, sintered samples were measured for hardness based on Vickers

hardness. Before testing, the sintered samples were ground by SiC paper and

polished by diamond paste to produce high flat surface as done for FESEM

investigation. Then, Vickers hardness, HV30, was measured using Vickers hardness

tester (Future-Tech FV-7, Japan) with diamond indenter at the load of 30kg and

dwelling time of 10 seconds. Hardness value was calculated based on the equation

(3.6).

221

2

*854.130

ddPHV (3.6)

71

in which HV is Vickers hardness, P is applied load (30 kg), and d1 and d2 are the two

diagonals of the indentation left in the surface of the material after removal of the

applied load. Fig. 3.6 shows the typical shape of the indentation after removing

applied load. Each sample was tested at least 5 times and an average value was taken

as the hardness value.

3.3.5 Fracture toughness measurement

In this work, fracture toughness or critical stress intensity factor, KIC, was

determined by the measurements of the Palmqvist radial cracks. With Vickers

hardness, HV30 (kg/mm2), fracture toughness (MNm-3/2) can be calculated using the

following equation (Schubert et al., 1998):

KIC = 0.15(HV30/LT)1/2 (3.7)

where LT= Li, the total length of cracks at the corners of Vickers hardness

indentations. The longer crack length, the lower is the fracture toughness of the

hardmetals. This simple method is useful in determining the fracture toughness of

d2

d2

Figure 3.6 Schematic of indentation mark in Vickers hardness measurement

72

ceramic materials as it allows determination of fracture toughness in small samples.

FESEM image of indentation on sample WC-10FeCrNi-2graphite-1NbC sintered by

PHIP at 1300oC for 45 min including 15 min pressing at 20 MPa and the crack

lengths at the corners of the indentation (L1-L4) are shown in Fig. 3.7.

Figure 3.7 FESEM image of indentation on sample WC-10FeCrNi-2Cgr-1NbC

sintered by PHIP at 1300oC for 45 min including 15 min pressing at 20 MPa and the

crack lengths at the corners of the indentation (L1-L4)

73

CHAPTER 4

RESULTS AND DISCUSSION

4.1 Raw material analysis

4.1.1 Tungsten carbide (WC) powder

Fig. 4.1 shows the XRD pattern of WC powders. The ICDD of PDF number

of WC selected in XRD database is 03-065-4539. The result of XRD analysis shows

that WC has hexagonal, -WC, with space group P6m2 and lattice parameters; a =

0.29 nm, c = 0.283 nm. Three strongest peaks appear at 31.582, 35.724, 48.406 of 2

degree corresponding to Miller index (001), (100) and (101), respectively. Fig. 4.2a

shows the FESEM image of WC powders. WC grains have faceted shape and

agglomerated with particle size less than 1 m. The C content is about 5.79 wt.% in

WC powders as seen in EDX result of WC powder, Fig. 4.2b.

Figure 4.1 XRD patterns of WC raw powders

74

Figure 4.2 a) FESEM image and b) EDX result of WC raw powder

4.1.2 FeCrNi powder

Fig. 4.3 presents XRD pattern of FeCrNi powders. ICDD of PDF number of

FeCrNi number selected in XRD database is 00-021-0619. FeCrNi has an

fcc-structure (-FeCrNi) with space group Fm-3m and lattice parameter: a = 0.36 nm.

Three mains peaks appear at 43.467, 50.685 and 74.763 of 2 degree corresponding

to Miller index (111), (200) and (220), respectively.

75

Figure 4.3 XRD patterns of FeCrNi powders

Fig. 4.4a shows FESEM image of FeCrNi powders. FeCrNi powders have

rounded shape with average grain size of 50 m. Composition of FeCrNi powders is

shown in Fig. 4.4b, EDX result of FeCrNi powder. EDX result shows that FeCrNi

powders contain about 64.48 wt.% Fe, 19.64 wt.% Cr, 10.45 wt.% Ni, 1.11wt.% Mn

and 1.45 wt.% Si. Besides that the containing of high content of C, 0.93 wt.% may be

due to the effect of C tape on FESEM holder. The presence of O2, 1.93 wt.% may be

resulted from the oxide during the exposal of powder the atmosphere.

76

Figure 4.4 a) FESEM image and b) EDX result of FeCrNi powders

4.1.3 Graphite (Cgr) powders

Fig. 4.5 presents XRD pattern of Cgr powder using in this work. ICDD of

PDF number selected in XRD database is 01-075-2078. The two main peaks appear

at 26.456 and 54.574 of 2 degree corresponding to Miller index (002) and (004),

respectively. Cgr has a hexagonal structure, space group P 63/mmc and lattice

parameters; a = 0.246 nm and c = 0.671 nm. Fig. 4.6a shows the FESEM image of

graphite powders. Cgr powders have plate shape with average particle size is about 10

m. EDX results of Cgr is shown in Fig. 4.6b. The result shows that graphite is pure

with 100 % C element.

77

Figure 4.5 XRD patterns of Cgr powders

Figure 4.6 a) FESEM image and b) EDX result of Cgr powders

78

4.1.4 Niobium carbide (NbC) powders

Fig. 4.7 shows the XRD pattern of NbC powders using in this work. The

ICDD of PDF number of NbC selected in XRD database is 01-089-3830. XRD

shows that NbC has cubic structure, space group cF8 (diamond structure). Three

main peaks appear at 34.768, 40.362 and 58.382 of 2 degree corresponding to

Miller index (111), (200) and (220), respectively.

Figure 4.7 XRD patterns of NbC powders

Fig. 4.8a shows the FESEM of NbC powders with particle size is about 3

m while Fig.4.8b shows the EDX result of NbC powders. The result shows that the

C content containing in NbC is about 9.51 wt.%.

79

Figure 4.8 a) SEM image and b) EDX result of NbC powders

4.2 Effect of milling time on microstructure and mechanical properties of

WC-10FeCrNi hardmetals

In order to study the effect of milling, WC-10FeCrNi hardmetal powders

were fabricated by mixing in the planetary ball milling machine at different milling

time; 5, 10, 15 and 20 h. The milling process was set up with 200 rpm and under Ar

gas with the ball-to-powder ratio of 10:1. After milling, the powder was cold

compacted at 300 MPa in a 10 mm cylindrical pellet and sintering at 1300oC under

vacuum for 1 h.

80

4.2.1 Phase identification and microstructure of as-milled powder

The phase analysis results of as-milled powders are shown in Fig. 4.9. There

are two phases; main phase is -WC with hexagonal structure and the minor phase is

-FeCrNi with fcc-structure (main peaks at 43.467, 50.685 and 74.763 degree of 2)

but only peak at 43.467o of 2 of -FeCrNi can be observed because other peaks

were overlapped by WC peaks. No other phases were detected indicating that phase

transformation did not occur during milling. Increasing the milling time led to a

decrease in the intensity of both WC and -FeCrNi peaks. And those peaks were

broaden due to the refinement of crystallite size and an increase in atomic level

strains resulted from intensive deformation from high energy of ball milling.

Figure 4.9. XRD patterns of WC-10FeCrNi powders at different milling time

The results of crystallite size calculation (Fig. 4.10) show the effect of

milling time on the crystallite size of WC (DWC). DWC has a tendency to decrease

with the increase of milling time. A very strong decrease of DWC at 5 h and 10 h

81

milling was observed (from about 27.5 to 10.3 nm). The decrease of DWC was slow

from 10 to 20 h milling. In the first stage of milling, the ductile particles, FeCrNi,

underwent deformation, their morphology changing from equiaxed to flattened, Fig.

4.11 (5 h and 10 h milling ), while brittle particle WC underwent fragmentation.

Then, when the ductile FeCrNi particles started to weld, the brittle WC particles

came into between two or more ductile particles at the instant of the ball collision. As

a result, fragmented WC particles would be placed in the interfacial boundaries of the

welded FeCrNi particles, and the result was the formation of a composite of

WC-FeCrNi. As welding is predominant mechanism in the process, the particles

changed their morphology by piling up laminar particles. These phenomena,

deformation, welding and dispersion of brittle phase in ductile phase, hardened the

material and increased the facture process. The fracture and welding of particles were

repeated continuously and, consequently, FeNiCr and WC were refined, leading to

the dramatically decreasing of WC crystalline size. The ability of particles to accept

further plastic deformation decreased with increasing milling time (15 and 20 h). At

long milling time, welding and fracture mechanisms then reach equilibrium,

promoting the formation of composite particles with randomly orientated interfacial

boundaries. The agglomeration of particles increased when the particle size of

fragmented WC decreased, thereby increases the fracture resistances. The increase of

fracture resistance and greater cohesion between the particles, with decreasing

particle size, caused agglomeration (Hewitt et al., 2009) and therefore, the ability to

continue fracturing of WC particles was inhibited, leading to the slow decrease of

WC crystallite size (DWC) at higher milling time.

82

Figure 4.10 WC crystallite size at different milling time

To investigate the morphology and distribution of mixture powders, the

as-milled powders were mixed with resin and hardener, then polished using SiC

paper and diamond paste, and observed under FESEM with back-scattered detector.

Fig. 4.11 shows the FESEM images of powders at different milling times. The results

reveal an agreement with crystalline size results (Fig. 4.10). Increasing the milling

time leads to further grain refining and better distribution of components. After 5 and

10 h milling, FeCrNi particles are still coarse (grey particles in Fig. 4.11) and the

particle size in the range of 20 - 40 µm. After 15 and 20 h milling, FeCrNi powders

are finer, 5 -10 µm in length with 10 h milling, and not clear after 20 h milling, and

the distribution of mixture powders is also better in these conditions. FeCrNi

particles were not only reduced in the size but also changed in the morphology. The

changing of FeCrNi from round shape (initial powders) to bar shape (flattened)

shows the plastic deformation of soft FeCrNi particles.

83

Figure 4.11. Back scattered electron FESEM images of WC-10FeCrNi at various

milling time: a) 5 h, b) 10 h, c) 15 h and d) 20 h

4.2.2 Phase identification and microstructure of sintered samples

The XRD results of sintered samples as shown in Fig. 4.12 show that WC

still dominated in the microstructure of samples. However, -FeCrNi transformed to

-FeCrNi and -phase (Fe3W3C) was detected in the microstructure of sintered

samples. The formation of -phase shows a similar trend of sintering of WC with Co

or Fe-rich binders. In WC-Co system, -phase Co6W6C is formed at about 800oC

while Co3W3C is formed at about 1000oC by phase reaction between Co6W6C and

carbon (Kurlov et al., 2011). In WC-Fe rich binder, -phase has been reported to be

formed at low temperature of about 750oC, and its amount gradually increases until

84

1100oC of WC-5.7 AISI304 stainless steel (Fernandes et al., 2007) and about 800oC

during sintering of mixture powder of W-Fe-C (Tsuchida et al., 2001).

The formation of -phase was generally attributed to the C deficit caused by

the decarburization during sintering that lowers the C content in the chemical

stoichiometric of phase diagram as shown in Fig. 2.5. Decarburization might be

related to a oxygen pick-up phenomena during ball milling or the residual air in the

furnace (Sommer et al., 2002). This oxygen presented at the surfaces of WC and Co

particles and then be released later during sintering at the temperature ranging from

500-1000oC as CO and CO2 in sintering of WC-Co hardmetals. Consequently,

C-deficient regions are generated. When these areas contacted with metal binder,

-phase is formed at the interfaces (Sommer et al., 2002). The more contacts of WC

and binder phase lead to more -phase formation and reduce the binder amount

(Fernandes et al., 2007) and it might be resulted in the transformation of -FeCrNi to

-FeCrNi as seen in XRD results (Fig. 4.12).

Figure 4.12 a) XRD patterns of sintered samples at 1300oC and b) magnification at

40-50 of 2 degree

85

In this work, increasing the milling time led to an increase of -phase

formation and consequently results in a decrease of -FeCrNi (Fig. 4.12 b). The

longer milling time led to the finer and the more homogeneity of WC and FeNiCr

particle distribution and hence, more contact areas between the WC and the binder

phase are formed. Consequently, shorter diffusion distance was available to transport

the species for -phase formation (Fernandes et al., 2007). Once the -phase is

formed, naturally -FeCrNi is reduced. This finding shows similar trend with

composites from WC powders sputter-deposited with iron rich binders (Fernandes et

al., 2007; Fernandes et al., 2009b).

The -phase fraction (wt.%) formed during sintering (calculated based on

XRD results by Rietveld refinement method) vs milling time is shown in Fig.4.13.

Figure 4.13 -phase fraction vs. milling time of sintered samples at 1300oC

The fraction of -phase increases with the milling time from 3.65 wt.% (5 h

milling) to 7.3 wt.% (20 h milling). However, it is can be seen that the formation

speed of -phase slow down at high milling time (not significant difference between

86

15 h and 20 h milling). This show that the fraction of -phase may obtain the critical

value or the formation of -phase was slow down by the dissolution of this phase in

liquid binder phase (Fernandes et al., 2007).

Microstructure plays a very important role in the mechanical properties of

sintered samples. Homogeneity of WC grain size and distribution in the binder phase

is a requirement for WC-based hardmetals in all applications. Back scattered electron

FESEM images (Fig. 4.14) show the effect of milling time on the microstructure of

sintered samples. The longer milling time led to a less porosity in microstructure of

sintered samples. The arrows in the images show the pores in the microstructure of

sintered samples. As shown in FESEM images of milled powders (Fig. 4.11), the

microstructure of 5 h and 10 h milled samples indicated several coarse FeCrNi

particles. These milling times were not sufficient to produce a homogeneous

structure. During sintering, there is a possibly that FeCrNi molten filled the

boundaries between WC grains and hence leaving the pores. The longer of milling

time (15 and 20 h) led to a finer particle of WC, FeCrNi and better distribution

between WC and FeCrNi particles (15 h and 20 h), therefore, promoting

pre-compaction and densification of sintered samples. The pores in the

microstructure can be eliminated by increasing sintering time or sintering

temperature to achieve full density; however, this may cause significant grain growth

of WC.

87

Figure 4.14 Back scattered electron FESEM images of sintered samples at different

milling time; arrows show the pores in the microstructures

4.2.3 Density measurement

Density is one of the important parameters to evaluate densification of

materials during heat treatment process. Relative densities of pre-compacted (green

density) and sintered samples were plotted in Fig. 4.15. The green densities of

samples are nearly similar, about 63 % of theoretical density. After sintering, the

relative density of samples increases up to more than 90 % (92.9 % for sintered

sample with 20 h milling). The results show an agreement with the microstructure of

samples shown in Fig. 4.14. An increasing nearly 30% of relative density was

obtained after sintering. A previous study has reported that a low densification of

WC coated AISI304 stainless steel at temperature of 1150oC is due to the solid state

sintering while for higher temperature than 1150oC, an abrupt increase in the

88

densification was observed which may associate to the formation of a liquid phase

(Fernandes et al., 2007). The FESEM images and the results of high relative density

show that the samples were completely sintered by liquid phase at 1300oC in 1 h.

The different of relative density of sintered samples shows the effect of milling time.

The increase of milling time led to finer and better distribution of WC and FeCrNi

particles and improved of densification.

Figure 4.15 Relative densities of pre-compacted and sintered samples at 1300oC

4.2.4 Hardness and fracture toughness

Figure 4.16 presents the Vickers hardness of sintered samples at various

milling time. The hardness increased with milling time. This was resulted from

higher density and better microstructure of sintered samples with further milling time.

At 5 h milling, the average value of hardness is 1240 kg/mm2 and it increases up to

1510 kg/mm2 at 20 h milling. There is a strong increase of hardness from 10 h to 15

h milling of sintered samples, from 1288 to 1448 kg/mm2. This caused by the better

89

distribution of WC and FeCrNi during milling and the less porosity of sintered

samples with high milling time (15 and 20 h milling).

Figure 4.16 Vickers hardness and fracture toughness vs. milling time

In general, the formation of -phase is unwanted phase during sintering of

WC-based hardmetals because this brittle phase has a negative effect on fracture

toughness of materials (Uhrenius et al., 1997; Allibert, 2001; Fernandes et al., 2007).

Although the formation of -phase in this work was enhanced with prolonged

milling, the fracture toughness of sintered samples increased with the milling time,

from about 7.2 to 8.25 MPa.m1/2 (Fig. 4.16). The more homogeneity of WC grains

and binder phase at longer milling may be attributed to the higher of the fracture

toughness. Besides that, the appearance of high porosity in the microstructure of

lower milling time samples is also the reason for lower the hardness and toughness

values. At 20 h milling, the increase of hardness is predominant than the increase of

fracture toughness compared to 15 h milled sample. This may caused by the effect of

more -phase formation in 20 h milled sample.

90

Although the lower -phase formation during sintering achieved at short

milling time of milled powder, the sintered samples contained high porosity and low

mechanical properties. More -phase formed in sintered samples with high milling

time, however the microstructure contained less porosity and mechanical properties

of sintered samples presents an optimum values at 15 – 20 h milling, higher hardness

and fracture toughness. So the optimal milling time for mixed powder should be

chosen from 15 to 20 h milling. In the continuous works, 15 h milling was selected

as the milled condition.



In order to investigate the effect of sintering temperature, WC-10FeCrNi

powders were milled for 15 h under Ar gas and subjected to sinter at different

temperature; 1250, 1300 and 1350oC for 1 h in vacuum furnace, respectively.


The results of phase analysis and microstructure as well as crystallite size of

as-milled powder can be seen in Fig.s 4.9 - 4.11. The XRD results of sintered

samples are shown in Fig. 4.17a. Besides WC and FeCrNi phases, -phase was

identified in all sintering temperature indicating that -phase was formed at

temperature lower than 1250oC. Besides that, the peaks of -phase were observed to

increase with sintering temperature. Moreover, the decrease of binder phase was

detected as seen in Fig. 4.17b. The reduction of binder phase was caused by its

91

reaction with the carbon-deficient WC grains to form -phase. As the temperature

increases, the spreading speed of liquid phase among of WC grains was enhanced by

the capillary force which leads to rearrangement of WC grain (Bhaumik et al., 1996,

Silva et al., 2001) and hence, more reaction between WC and binder phase occurred.

Consequently, more -phase was formed accompanied with the decrease of binder

phase (Fernandes et al., 2007). Insignificant increase of -phase from 1300 to

1350oC compared to the range of 1250 to 1300oC was observed. Quantification

analysis of -phase are presented in Table 4.1. At 1250oC, 4.8 wt.% of -phase was

formed. The fraction of -phase increased to 7.1wt.% at 1300oC and 7.4 wt.% at

1350oC. As the liquid binder phase formation, some dissolution of the -phase in

binder phase occurs (Fernandes et al., 2007). This dissolution may retard the

formation of -phase, thus the increase of -phase fraction is not significant at

1350oC in comparison with 1300oC.

Figure 4.17 a) XRD patterns of sintered WC-10FeCrNi at different temperatures and

b) magnification at 35-55 of 2

92

Table 4.1 -phase fraction formed vs. sintered temperatures

Sintered temperature 1250oC 1300oC 1350oC

-phase (wt.%) 4.8 7.1 7.4

The microstructure of polished sintered samples were observed by FESEM

under back scattered mode is shown in Fig. 4.18. As the liquid phase formed, liquid

binder spread out to fill in the interfaces among WC grains and completed

densification as the liquid phase fully filled to produce bulk material. As seen in Fig

4.18a, at 1250oC, there is a presence of rich binder phase areas (grey colors pointed

by arrows) where the binder phase is not completely spread out into WC interface

(white colors) and poor binder areas (circle) leaving pores and poor contacts between

WC grains and binder phase. This led to a lower density and an inhomogeneous

microstructure of sintered sample. The equi-axial or round shape of WC can be seen

indicating that the grain growth rate was small at this temperature. At 1300oC, the

binder seems to be having a better distribution in the microstructure of material (Fig.

4.18b). As the binder phase spread out at WC interfaces, the grain growth of WC also

can be observed in comparison to 1250oC sintered sample. Grain growth by an

Ostwald ripening may occur corresponding to solution-precipitation mechanism of

matter transport (Xiao et al., 2009). Grain growth is more pronounced in sample

sintered at 1350oC (Fig. 4.18c) with faceted platelet shape of WC and with higher

densification. Fig. 4.18d shows the EDX of binder phase (grey colour in 4.18a ).

There is a presence of WC in the binder phase. This is due to the dissolution of WC

in binder phase as a similar trend with WC-Co hardmetals (Sutthiruangwong and

Mori, 2005; Kellner et al., 2009; Fernandes and Senos, 2011) or it may be due to the

93

dissolution of -phase in the binder as the liquid binder phase formed (Fernandes et

al., 2007).

Figure 4.18 FESEM back scattered electron images of sintered WC-10FeCrNi at

different temperatures:a) 1250oC, b)1300oC, d) 1350oC for 1 h and EDX result of

point X in Fig. 4.18a

4.3.2 Density measurement of sintered samples

The results of density and relative density of sintered samples are shown in

Fig. 4.19. The densification of WC-FeCrNi hardmetals increases with the sintering

temperature. At 1250oC, the lowest density (87.6 % relative density) was obtained

resulted from the incompletely densification of liquid phase and high residual pores

in the microstructure. Increasing sintered temperature led to a lower viscosity of

94

liquid phase. This means that the liquid phase is easy to diffuse in a matrix and fill in

the spaces among WC grains, and hence increases the densification.

Figure 4.19 Relative density of WC-FeCrNi at different sintering temperature

At 1300oC, as the liquid phase completed to spread on WC interfaces, the

relative density increased about 5 % compared to 1250oC. At higher temperature,

1350oC, the density continued to increase but in a slow rate. Only 0.5 % of relative

density higher than that at 1300oC was obtained. This shows that the changing in

density is small at 1300oC and 1350oC. The different between two sintered

temperatures is from the grain size and shape of WC (Fig. 4.18). The increase of

sintered density with sintered temperature is in agreement with WC-Co system

(Arató et al., 1998; El-Eskandarany et al., 2000).

95


The Vickers hardness and fracture toughness of sintered samples are shown

in Fig. 4.20. Vickers hardness, HV30, shows an increase trend with the sintering

temperature. When the liquid phase binder not completely filled on the WC grain

matrix at 1250oC, there is a presence of several binder rich areas and pores which

lower hardness of sample. At higher temperature, the higher densifications are

obtained as well as the rearrangement of WC grains in the microstructure. More

homogeneous microstructure and distribution of binder phase are observed. The

hardness of sintered sample, therefore, increased. At 1350oC, the liquid phase

completely distributed among WC grains, however, the densification increases

slowly. The slight increasing of hardness for sintered sample at 1350oC compared to

that at 1300oC is observed. At higher sintering temperature, the more densification is

attained but the more grain growth of WC was also observed as well. This is the

reason for the slowly increase of hardness.

In contrast, the fracture toughness of sintered samples decreased with the

sintering temperature. This shows the effect of -phase on the fracture strength of

hardmetals. As shown in XRD results and -phase fraction calculation, the formation

of -phase was more pronounced at high temperature, therefore reducing the fracture

toughness despite of the better microstructure at higher sintering temperature.

Besides that more grain growth of WC and high hardness for higher sintering

temperature may result in the decrease of fracture toughness (Chermant and

Osterstock, 1976; Schubert et al., 1998; Vaßen and Stöver, 1999).

96

Figure 4.20 Vicker hardness and fracture toughness of sintered samples vs. sintering

temperatures

4.4 Effect of sintering time on the microstructure and mechanical


WC-10FeCrNi was subjected to sintering with different soaking time in

order to investigate the effect of sintering time on microstructure and mechanical

properties of sintered sample. Sintering was done at 1350oC for 15, 30, 45 and 60

min in vacuum furnace, respectively.


Fig. 4.21a shows the effect of sintering time on phase composition of

sintered samples. After sintering, WC was a dominant phase in the microstructure.

97

Stainless steel binder presents as -FeCrNi. The decrease of -FeCrNi peak was

observed with the increase of sintering time. This resulted from the formation of

-phase. -phase was formed as the deficiency of C during sintering. Increase the

sintering time enhances the movement of liquid binder phase filling into the spaces

among WC particle, consequently more contacts between binder phase and WC

particles and hence, more -phase formed. The increase of -phase peaks and the

decrease of -FeCrNi peak are shown more clearly in Fig. 4.21b. Quantification of

-phase fraction calculated based on the Rietveld refinement method is seen in Fig.

4.22. At 15 min sintering, 4.5 wt.% of -phase was formed. The formation of

-phase was then more pronounced with the increase of sintering time. At 60 min

sintering, nearly 7.4 wt.% of -phase was formed leading to the significant reduction

of -FeCrNi binder phase.

Figure 4.21 a) XRD patterns of sintered samples at different sintering time; 15, 30,

45 and 60 min, and b) magnification at 35-55 degree of 2

98

Figure 4.22 -phase fraction various with sintering time

Fig. 4.23 shows the back scattered electron images of sintered samples at

different sintering time. It is clear to see the effect of sintering time on the

microstructure of samples sintered at 1350oC. The white grains are WC, the grey

colours are FeCrNi binder and the dark colours are pores in the microstructure of

samples. The size of binder phase areas reduced with the sintering time as the liquid

phase filled in the spaces among WC grains by capillary force and matter transfers.

At the 15 and 30 min of sintering, the liquid binder phase formed and started to

spread over the WC grains, however, this is not sufficient to complete the

densification. Further densification was achieved after longer sintering time. At 45

min sintering, the binder seems completing to fill in the WC grain boundary, however,

high amount of pores is observed in the microstructures. At 60 min, the

microstructure gained the highest densification of all samples. The distribution of

99

WC grains, FeCrNi binders and pores are getting better. The results of EDX analysis

(Fig. 4.23 X, Y) at the points X, Y in Fig. 4.23 show that there is dissolution of WC

in binder phase as the liquid binder phase formed (Sutthiruangwong and Mori, 2005;

Kellner et al., 2009; Fernandes and Senos, 2011) or due to the dissolution of -phase

in binder phase in liquid phase (Fernandes et al., 2007). Besides that it may be due to

the agglomerates of WC and binder phase during milling process. The binder phase

is not observed clearly at longer sintering time. The existence of binder phase in the

microstructure is confirmed by EDX results of areas (Z and W) in the microstructure

of samples (Fig. 4.23 Z and W). The FESEM image also shows the effect of sintering

time on the grain growth of WC. Increasing the sintering time caused the larger WC

grains. The grain growth of WC is resulted from the migration of the fine grains in to

larger grains during sintering (Xiao et al., 2009). The longer sintering time, the more

grain growth occurs. The observation had been confirmed by previous study (Sun et

al., 2007) in which the grain size will grow rapidly along with the keeping time at the

same sintering temperature in WC-Co system and the slow grain growth rate was

observed at low sintering temperature. Moreover, WC has a tendency to change from

the rounded particles of WC to the facetted shapes for longer sintering, which show

the grain growth process of WC increases with sintering time.

100

Figure 4.23 Back scattered electron images of sintered WC-10FeCrNi at different sintering time: a, 15min; b) 30min; c) 45min and d) 60min, and EDX analysis of X, Y,

Z and W points in SEM images

101

4.4.2 Density of sintered samples

Fig. 4.24 shows the densities of sintered samples varied with the sintering

time. The density of sintered samples increases with the soaking time. At the

sintering temperature, liquid phase formed and spread in the WC boundaries and fill

in the pores. This activates the densification process rapidly. The consolidation

process completes with the sufficient amount of liquid binder phase and sintering

time. The increase of density with sintering time presents a similar trend in

comparison with WC-Co system. In WC-Co systems (WC-3Co, WC-10Co,

WC-18Co), rapid densification was observed within 30 sec., and the rate of

densification was then markedly decreased as sintered at 1340oC (Kim et al., 1997).

The full density of sintered samples obtained depends on the sintering time,

temperature, grain size of raw materials, content of binder phase.

Figure 4.24 Effect of sintering time on the density of sintered samples

In this study, the relative density of WC-10FeCrNi increases from 84% at 15

102

min sintering and reaches the value of 93 % theoretical density after 60 min sintering.

To achieve higher value of density, longer sintering time and higher sintering

temperature are the requirements. However, these requirements also lead to

pronounce grain growth of WC which may detrimental to the properties of

hardmetals.

4.4.3 Mechanical properties of sintered samples

Vickers hardness and fracture toughness of samples varied with sintering

time as seen in Fig. 4.25. The hardness of sintered samples increases with sintering

time. The longer sintering time resulted in higher densification which enhances the

bonding between WC and binder phase. When the liquid phase formed, it starts to

move and fill in the spaces between WC grains and enforce the coherence among

WC grains and binder phase, therefore strengthened the sintered sample. At 15 min

of sintering, time is not sufficient for liquid phase formation and completely filling to

enforce the bonding between components as seen in FESEM image (Fig. 23a). Thus,

the hardness of sample after 15 min sintering obtained is the lowest value. When the

sintering time increased, the densification was more pronounced which means that

the coherence among WC grains also increased, and hence, the hardness increased

with sintering time. However, as the sintering time increases, the grain growth of WC

also occurs which cause a reduction of hardness according to Hall-Petch relation

(Niels, 2004).

103

Figure 4.25 Vickers hardness and fracture toughness of samples at various sintering

time

Although higher density was obtained at longer sintering time, the fracture

toughness of sintered samples showed a tendency to decrease. At 15 min sintering,

the fracture toughness is about 9.5 MPa.m1/2 and the value reduced to 8.1 MPa.m1/2

at 60 min sintering. The decrease of fracture toughness resulted from the increase of

-phase formation (González et al., 1995; Fernandes et al., 2007) and the increasing

trend of hardness of sintered samples (Chermant and Osterstock, 1976; Schubert et

al., 1998; Vaßen and Stöver, 1999).

4.5 Role of binder phase composition

In order to study the role of binder content in microstructure and mechanical

properties of WC-FeCrNi hardmetals, WC was mixed with 8, 10, 12 and 15 wt.%

FeCrNi binder for 15 h ball milled and subjected to sintering at 1350oC for 1 h,

104

respectively. The composition of mixed powders and code of samples are presented

in Table.4.2

Table 4.2 Composition of samples

Sample code S1 S2 S3 S4

WC (wt.%) 92 90 88 85

FeCrNi (wt.%) 8 10 12 15

4.5.1 The role of binder in phase and microstructure of sintered samples

Fig. 4.26 presents XRD results of samples after sintering at 1350oC for 1h

under vacuum. The formation of -phase was observed in all samples and it

increases with initial binder content. Besides that an increase of binder phase in the

sintered samples was also observed with the increase of initial binder content. The

amount of -phase calculated is shown in Fig. 4.27. The amount of -phase is 8.15

wt.% formed at 15 wt.% of starting binder phase. This result shows a similar trend

with a previous study (Fernandes et al., 2007). In their research, -phase fraction was

observed to increase after sintering at 1325oC for 3h in either normal mixing powder

or sputtered coating. A linear relation between -phase and binder content with a

positive slope of ~1.7 was found in sintered samples with coated powder and the

slope of ~0.8 was observed in samples with normal mixing (Fernandes et al., 2007).

The increase binder content led to more contacts between WC and binder phase and

hence, more -phase is formed.

105

Figure 4.26 XRD patterns of samples various with binder content sintered at 1350oC

Figure 4.27 -phase fraction vs. initial binder contents

Microstructure of sintered samples is shown in Fig. 4.28. It can be seen that

the increase in initial binder content led to a decrease in pore amount in the sintered

106

samples. At low binder content, 8 wt.%, there are several pores (dark colours) in the

sample. At higher content of binder, the pores decreased in the amount and the shape.

At 8 wt.% of binder (S1), the liquid binder phase seems not to be sufficient to

completely filling in the spaces between WC grains and leaving the pores in the

microstructure. The liquid phase was attributed to minimize the total surface energy,

which is the main driving force for densification in all stages of sintering (Daoush et

al., 2009). Thus, at high content of binder phase (15 wt.%, S4) more densification of

sample was obtained since low porosity was observed.

Figure 4.28 SEM back-scatter electron images of samples sintered at 1350oC with

different binder contents

107


As discussed previously, the higher content of binder phase, the higher

densification was obtained. Fig. 4.29 shows the effect of binder content on the

density and relative density of sintered samples. Increasing the binder content led to

a decrease in density, as explained by the lower density of FeCrNi in comparison

with WC. However, the relative density increased with the binder content; this shows

that the higher liquid binder phase enhanced sintering process which resulting in a

higher relative density.

Figure 4.29 Density and relative density of samples vs. initial binder contents

4.5.3 Hardness and fracture toughness of sintered samples

Binder phase plays an important role in mechanical properties of WC-based

hardmetals. Fig. 4.30 presents Vickers hardness and fracture toughness of sintered

samples at different FeCrNi binder contents. It is apparent that the hardness of

overall hardmetals decreased as the fraction of binder phase increased. The binder

108

phase is softer than WC, therefore, increasing the binder phase results in the decrease

of hardness. Vickers hardness is about 1515 kg/mm2 at 8 wt.% binder phase and

decreased to the value of 1340 kg/mm2 at 15 wt.% binder phase. These results show

a similar trend with WC-Co system (Ettmayer, 1989; Deng et al., 2001; Daoush et al.,

2009). Although there is an increase of -phase fraction with the increase of binder

content, the fracture toughness, KIC, of sintered samples also increased. In phase

analysis (Fig. 4.26), the binder peak () increased with the binder content. Increasing

the binder composition led to an increase of fracture toughness. At 8 wt.% (S1) of

starting binder phase, the value of fracture toughness is about 7.8 MPam1/2, and this

value increases up to 9.4 MPam1/2 at 15 wt.% of binder composition (S4). This trend

agrees with those of WC-Co systems in previous studies (Kang et al., 2000; Deng et

al., 2001). The results of Vickers hardness and fracture toughness also shows that the

hardness decreases leading to an increase of fracture toughness (Vaßen and Stöver,

1999).

Figure 4.30 Vickers hardness and fracture toughness of sintered samples with

different binder content

109

4.6 Effect of Cgr addition on microstructure and mechanical properties of

WC-10FeCrNi hardmetals

In order to investigate the ability of Cgr to eliminate the formation of -phase

during sintering; 0, 1, 1.5, 2, 2.5 and 3 wt.% Cgr was added in WC-10FeCrNi before

milling 1h under Ar gas, respectively. The samples were sintered at 1350oC for 1 h.

The compositions of mixed powders and codes of samples are shown in Table.4.3.

Table 4.3 Composition of samples with Cgr addition

Sample wt.% WC wt.% FeCrNi wt.% Cgr

HC0 90.0 10 0.0

HC1 89.0 10 1.0

HC2 88.5 10 1.5

HC3 88.0 10 2.0

HC4 87.5 10 2.5

HC5 87.0 10 3.0

4.6.1 Phase identification and microstructure

XRD results of as-milled WC-10FeCrNi mixture powders containing

different amounts of Cgr are shown in Fig. 4.31. It can be seen that the composites

have only two phases; WC and -stainless steel. In the XRD spectra, only the peak of

(111) which is the Miller index of -stainless steel is clear with low intensity. Other

peaks, (200) and (220) of stainless steel are overloaded by WC peaks which are

supposed to be observed at 50.686 and 74.673 degree of 2-theta scale, respectively.

Peaks of Cgr were not observed in XRD spectra even at 3 wt.% of Cgr addition

110

suggesting that Cgr may transfer to amorphous state during milling process or due to

the limitation of XRD detector. No phase reaction among WC, FeCrNi and Cgr phase

was detected in all compositions by XRD results.

Fig. 4.31 XRD patterns of mixed powders with Cgr addition

The phase evolution of hardmetals can be seen in XRD patterns of samples

after sintering at 1350oC for 1 h in a vacuum furnace, as shown in Fig. 4.32. The

formation of -phase (Fe3W3C) in the sample without Cgr addition, HC0, has been

discussed previously.

111

Figure 4.32 XRD patterns of sintered samples vs. Cgr addition

As mentioned previously, the presence of this brittle phase should be avoided

because of the tendency to decrease the mechanical properties of hardmetals. To

eliminate the occurrence of -phase, Cgr powder was added in milling process to

recover the C losing during sintering cycle. The effect of Cgr content was indicated in

the XRD results (Fig. 4.32). The intensity of -phase decreases with the addition of

Cgr contents and this can be explained by the similar Eq. 4.1 in WC-Co systems (Eso

et al., 2007; Menéndez et al., 2007).

Fe3W3C + 2C 3Fe + 3WC (4.1)

The reaction between graphite and -phase depends on many factors such as

Cgr content, diffusion of graphite to the surface of -phase, higher temperature and

longer sintering time (Eso et al., 2007). The decrease of -phase with Cgr added have

been reported in Fe-rich binder hardmetals (Fernandes et al., 2007). In present work,

the peaks of -phase were not detected starting from 1.5 wt.% of Cgr. Decreasing of

-phase also leads to an increase of binder phase. From the XRD results (Fig. 4.32b),

112

a significant increase of -FeCrNi was observed after adding 1 wt.% Cgr. However,

at higher Cgr amounts, a tendency of -FeCrNi formation was also obtained and more

pronounced than -FeCrNi. When 2.5 wt.% Cgr was added (HC4), the trace of Cgr

peak was detected and at 3 wt.% the peak of Cgr became clear. Therefore, addition of

1.5 to 2 wt.% of Cgr can compensate the C lost during sintering and hence, the total C

content was in the range of two-phase region of hardmetals.

Back scattered electron FESEM images in Fig. 4.33, show the respective

microstructures of samples with different content of Cgr addition. Slightly finer WC

grain size can be observed in the sample without Cgr addition, HC0, in comparison

with other samples. This might be attributed to the formation of -phase in the

microstructure. As mentioned previously, the C-deficient regions were generated at

interfaces of WC grain and resulted in the formation -phase layers surrounding WC

grains during sintering. And hence, these -phase may inhibit the grain growth of

WC (Menéndez et al., 2007). As the formation of -phase was reduced with the Cgr

addition; the grain growth of WC increased and generated coarser WC grains. When

the -phase was significantly eliminated, the microstructures were not very different

in WC grains up to 2.5 wt.% added Cgr. However, the difference may come from the

more stable of -FeCrNi phase at higher Cgr addition as seen in XRD results (Fig.

4.32). Several coarse pores were observed in sample HC5 (3 wt.% Cgr). This may be

due to the appearance of free Cgr in the microstructure. The excess of Cgr addition

(2.5 – 3 wt.%) in composition may produce agglomerates of residual Cgr during the

sintering process. These agglomerates of Cgr may prevent the solidification of

microstructure and produced the pores or these pores may be the location of free Cgr

during sintering and they might be segregated during the polishing process.

113

Figure 4.33 FESEM backscattered electron images of samples after sintering at

1350oC in vacuum furnace various with Cgr contents: 0 wt.% HC0, 1 wt.% HC1, 1.5

wt.% HC2, 2 wt.% HC3, 2.5wt.% HC4, 3 wt.% HC5


Density of pre-compacted samples was calculated based on their weight and

dimensions. After sintering, density of WC-FeCrNi was measured by Archimedes’

principle. Relative densities were calculated by taking theoretical density of

114

WC+10FeCrNi as the reference density. The influence of Cgr addition on the relative

density of sintered samples is pointed out in Fig. 4.34. The relative densities of

pre-compacted samples decreased with the increase of added Cgr from 62.8 % (HC0,

0 wt.% graphite) to 58.1 % (HC5, 3 wt.% Cgr) which resulted from the lower density

of Cgr in comparison with WC and FeCrNi. However, the relative densities of

sintered samples increased with Cgr content up to 2 wt.% from 93 % in HC0 to

94.5 % in HC3 (2 wt.% graphite). At higher Cgr contents, the relative densities

reduced to 89.8 % in HC5 (3 wt.% Cgr). The increase of relative densities up to 2

wt.% was resulted from the decrease of -phase with Cgr content. The theoretical

density of Fe3W3C (-phase) is 14.464 g/cm3 as calculated in a previous study

(Suetin et al., 2009). This value is lower than theoretical density of WC (15.63

g/cm3). The increase in Cgr content led to a decrease in -phase and hence, enhanced

the density of sintered samples. The reduction in relative densities at higher than 2

wt.% Cgr decrease was attributed to the appearance of residual free Cgr in the

microstructure of these samples.

Figure 4.34 Relative densities of sintered samples vs. Cgr contents

115

4.6.3 Vickers hardness and fracture toughness

Variation of Vickers hardness and fracture toughness of WC-10%FeCrNi with

the content of Cgr addition is highlighted in Fig. 4.35. The results reveals that the

hardness of hardmetals was also increased up to 2 wt.% of Cgr from 1473 kg/mm2

(HC0, 0 wt.% Cgr) to 1625 kg/mm2 in HC3 (2 wt.% Cgr). At higher amount of Cgr,

Vickers hardness decreased to 1287 kg/mm2 in HC5 (3 wt.% Cgr). The fracture

toughness, KIC, also exhibited a tendency to increase with Cgr content up to 2 wt.%

Cgr from 8.1 MPa.m1/2 of HC0 to 10.2 MPa.m1/2 of HC3. Although -phase was not

detected in HC2 but the value of hardness and fracture toughness are slightly lower

than those of HC3. This is caused by -FeCrNi that was more dominating in HC3

compared to HC2. The increase of both Vickers hardness and fracture toughness may

attribute to the reduction of brittle -phase formation in microstructure and the

increase of relative density of sintered samples (Uhrenius et al., 1997; Fernandes et

al., 2007; Fernandes et al., 2009b). At higher Cgr content, HC4 (2.5 wt.%) and HC5

(3 wt.%) the fracture toughness starts to decrease to 9.4 MPam1/2. The reduction of

both Vickers hardness and fracture toughness in HC4 and HC5 might be resulted

from the appearance of free Cgr and pores in their microstructure.

116

Figure 4.35 Vickers hardness, HV30, and fracture toughness of sintered samples with

Cgr addition

4.6.4 Summary

WC-10AISI304 hard metals were produced by powder technique. -phase

(Fe3W3C) was observed in the samples after sintering at 1350oC. Formation of this

phase can be controlled by addition of Cgr. Increasing the Cgr content has enabled to

reduce the amount of -phase (Fe3W3C) and this phase was completely disappeared

at about 1.5 - 2 wt.% graphite. Free Cgr was detected in microstructure of sintered

samples starting at 2.5 wt.% graphite addition. The relative density of sintered

samples increased with the decrease of -phase but reduced with the appearance of

free Cgr.

Addition of Cgr improved both hardness and fracture toughness which was

resulted from the decrease of -phase. The highest value of Vickers hardness and

117

fracture toughness were obtained at about 1625 kg/mm2 and 10.2 MPa.m1/2 in

samples with 2 wt.% added Cgr. At higher than 2 wt.% Cgr, both hardness and

fracture toughness decreased with the occurrence of free Cgr.

4.7 Role of NbC addition on vacuum sintered WC-10FeCrNi-2Cgr hardmetals

In order to investigate the role of NbC as a WC grain growth inhibitor to

WC-FeCrNi hardmetals, different contents of NbC; 1, 1.5, 2 and 5 wt.% were added

in WC-10FeCrNi powders before milling for15 h under Ar atmosphere. The samples

were then sintered at 1300oC for 1 h in vacuum furnace. 2 wt.% of Cgr was also

added in all composition to eliminate the formation of -phase during sintering as the

results of previous section. The composition of mixed powders and code of samples

are presented in Table x

Table 4.4 Composition of sample with NbC addition

Sample code WC

(wt.%)

FeCrNi

(wt.%)

Cgr

(wt.%)

NbC

(wt.%)

HN1 88 10 2 0

HN2 87 10 2 1

HN3 86.5 10 2 1.5

HN4 86 10 2 2

HN5 83 10 2 5

4.7.1 Phase identification and microstructures

Fig. 4.36a shows the XRD patterns of the as-milled powders. The main phase

was WC. FeCrNi appeared in two types: -FeCrNi (main peaks at 43.467, 50.686 and

74.673 degree of 2) and -FeCrNi (main peaks at 44.653, 64.973 and 82.299 degree

118

of 2). All peaks except the peaks at 43.467 and 44.653 are overlapped with WC

peaks. With a higher NbC content (5 wt.%), FeCrNi had a tendency to be -FeCrNi.

As seen in Fig. 4.36b, the peaks of NbC became stronger as the content of NbC

increased. Peaks of NbC at sample HN2 (1 wt.% NbC) was not observed. It might

caused by the sensitivity of XRD detector. Peaks of Cgr were not observed in all

compositions suggesting that Cgr may transfer to amorphous state during milling

process as mentioned previously.

Figure 4.36 a) XRD patterns of milled samples various with NbC contents and b)

magnification at 40-50 of 2

XRD results of sintered samples are presented in Fig. 4.37. WC peaks were still

the main peaks dominating in the XRD patterns. After sintering, FeCrNi presented

mainly in type of -FeCrNi. Peaks of free Cgr or -phase was not detected in all

compositions suggesting additional of 2 wt.% graphite could suppress the C lost

during sintering. The intensity of NbC peaks increased with the NbC content and

became significant at 5 wt.%. The peaks of NbC in samples with 1 and 1.5 wt.%

NbC could not be observed; this may have happened due to the XRD detector low

sensitivity towards low contents of NbC.

119

Figure 4.37 XRD patterns of vacuum sintered samples with vs. NbC contents

The influences of NbC on the microstructure of as-sintered samples are shown in

FESEM images in Fig. 4.38. Increasing the NbC content led to a decrease in WC

grain size. At 2 wt.% NbC addition, the hardmetal achieved the optimal WC grain

size with more uniform WC grain distribution. At higher content of NbC (5

wt.%NbC), fine WC grains were achieved, however, there was an appearance of

coarse (Nb,W)C grains. The formation of (Nb,W)C is confirmed by EDX result in

Fig. 4.39.

120

Figure 4.38 FESEM back scattered electron images of sintered samples with different

NbC contents

A previous study (Xiao et al., 2009) has suggested that the mechanism of

WC grain growth during sintering is resulted from an Ostwald-ripening process due

to dissolution of the smaller WC grains in the binder, followed by their

re-precipitation onto larger grains, thereby reducing the interface area of the system.

The mechanism of inhibition of grain growth is assumed to be either an alteration of

121

the interface energies or an interference of the growth inhibitor with the interfacial

dissolution-nucleation-reprecipitation steps (Wittmann et al., 2002). Grain growth of

WC is inhibited by controlling one or more of the following processes: face-specific

adsorption, face-orientated deposition and blocking of active growth center of

crystals, including a change in edge energy (Wittmann et al., 2002; Guo et al., 2008).

We proposed that the addition of grain growth inhibitors (NbC in the present study)

might change the interface energy and interfere with the dissolution and

re-precipitation stages. In another study (Huang et al., 2008c), they indicated that up

to 10 wt.% WC can be dissolved into the cubic NbC phase when hot pressing for 10

min at 1300oC. And even higher solubility was observed in liquid phase sintering for

1 h at 1360oC (Huang et al., 2008c). This high solubility of WC in NbC also

inhibited the re-precipitation process of WC, thereby restricting the grain growth of

WC during sintering.

The dissolution of WC in NbC is confirmed by EDX analysis on the dark grain

in the 5 wt.% added NbC sample, point x of Fig. 4.38 (HN5) as shown in Fig. 4.39.

The result shows that up to 24.69 wt.% of WC was dissolved in NbC grain.

Therefore, this caused the formation of coarse (Nb,W)C grains in the microstructure

of the sample. Small amount of Fe, Cr, Ni elements dissolved in NbC also can be

observed by EDX result. The EDX result of this research is in agreement with

previous studies (Huang et al., 2007; Huang et al., 2008a; Huang et al., 2008c). A

small amount of Fe, Cr, Ni was also observed to be dissolved in NbC. Calculated

WC grain size by line intercept method is presented in Table 4.5. The grain size of

0.7 m was calculated in sample without NbC. The WC grain size was then

decreased due to the inhibition of NbC. In case of 2 wt.% NbC, the finest grain

shows the best effect of NbC to inhibit WC grain growth. Average grain size of 5.27

122

m of WC was calculated. At higher NbC content (5 wt.%), the grain size of WC

increased slightly to 5.33 m. This shows that NbC may have no further grain

growth inhibition for WC at high contents.

Figure 4.39 EDX result of point X in Fig. 4.30 (HN5)

Table 4.5 WC grain size of vacuum sintered samples with NbC addition

NbCWt.% 0.0 1.0 1.5 2.0 5.0

WC grain size (m) 0.690 0.650 0.584 0.527 0.533


The influence of NbC addition on the density of sintered samples is shown in

Fig. 4.40. The results of density measurement show that the density of sintered

samples decreased with the increase of NbC content. This is due to the lower

theoretical density of NbC compared to WC. The content of FeCrNi is a constant in

all compositions, thus increasing NbC means that the content of WC decreases.

123

Figure 4.40 Density of sintered sample vs. NbC content


The results of Vickers hardness and fracture toughness of sintered samples

are shown in Fig. 4.41. The hardness of samples increased with increasing NbC

content due to a decrease in WC grain size. The highest hardness, HV30 (1660

kg/mm2), was achieved with 2 wt.% NbC addition in comparison to HV30

(1500kg/mm2) of the sample without NbC addition. The effect of grain size on the

hardness of materials has been explained by Hall-petch relation as the strength of

materials depends on the grain size (Niels, 2004). Consequently, a decrease in grain

size leads to an increase in strength. A slightly decrease of Vickers hardness after 5

wt.% NbC addition in the present work could be caused by the presence of coarse

(Nb,W)C grains in the microstructure. Fracture toughness of sintered samples

decreases with the increase of NbC addition. The decrease of fracture toughness was

resulted from the increase of hardness of materials (Chermant and Osterstock, 1976;

124

Schubert et al., 1998; Vaßen and Stöver, 1999). KIC reached the value of 9.6

MPa.m1/2 in the sample without the addition of NbC and decreases to 8.3 MPa.m1/2 in

the sample with 5 wt.% NbC. Although the highest Vickers hardness was obtained

with 2 wt.% NbC, the fracture toughness reached the value of about 8.4 MPa.m1/2.

Figure 4.41 Vickers hardness of as-sintered samples with NbC addition

4.7.4 Summary

WC-FeCrNi-Cgr hardmetals were fabricated by sintering at 1300oC in a vacuum

furnace. The effect of NbC as the grain growth inhibitor to WC-FeCrNi-Cgr

hardmetals was investigated based on the evolution of microstructure and hardness of

samples. The optimum NbC content for fine microstructure and uniform WC grain

distribution obtained at 2 wt.%. The highest hardness was achieved at HV30 of 1660

kg/mm2 with the addition of 2 wt.% NbC, however; the fracture toughness of this

sample gained a moderate value of 8.4 MPa.m1/2. (Nb,W)C phase appeared at higher

content of NbC addition (5 wt.% NbC) which was responsible for the decreased in

hardness.

125

4.8 The role of NbC addition as WC grain growth inhibitor for PHIP

sintered WC-10FeCrNi-2Cgr hardmetals

In this section, different content of NbC was added to inhibit WC grain

growth during sintering; however, the sintering process was assisted by pressure

under PHIP process as mentioned in Chapter 3, Section 3.2.3. The as-milled powders

were obtained by the same milling process and powder compositions as in Section

4.7. PHIP process was done at 1300oC under vacuum. For comparison, a sample with

a composition of WC+10FeCrNi with 2 wt.% Cgr added was similarly sintered but

without applied pressure and was coded as HS6. Composition of mixed powders is

presented in Table 4.6.

Table 4.6 Composition of sample with NbC addition sintered by PHIP

Sample code WC

(wt.%)

FeCrNi

(wt.%)

Cgr

(wt.%)

NbC

(wt.%)

HS1 88 10 2 0

HS2 87 10 2 1

HS3 86.5 10 2 1.5

HS4 86 10 2 2

HS5 83 10 2 5

HS6 88 10 2 0

4.8.1 Phase identification and microstucture

Fig. 4.42 shows the XRD patterns of samples HS1 and HS6. HS1 was

sintered by PHIP process at 1300oC while HS6 was sintered at 1300oC in vacuum for

45 min without applied pressure. Only phases of WC and -FeCrNi were observed in

HS6. No free Cgr or -phase was detected. In HS1, in addition to WC and -FeCrNi,

126

-phase (Fe3W3C) also appeared in the microstructure. The -FeCrNi peak of HS1

was considerably lower than that of HS6. The low -FeCrNi peak of HS1 was due to

significant formation of -phase due to the reaction between FeCrNi and WC

particles (Fernandes et al., 2007; Fernandes et al., 2009b). The formation of -phase

was attributed to the C deficit caused by the decarburization during sintering which

was discussed in previous sections.

Figure 4.42 XRD patterns of samples HS1 and HS6 after sintering

As the decarburization started to occur, C deficient regions were generated on

the surface of WC grains; hence, the formation of -phase took place as the binder

phase is in close contact with these C deficient regions. The application of pressure at

high temperatures during PHIP generates more contact areas between WC grains and

the binder phase compared to pressureless sintering in vacuum. Consequently,

-phase was easily formed as a result of the shorter distance needed for the elements

to diffuse (Fernandes et al., 2007). No -phase peaks were detected in HS6 which

127

suggests that 2 wt.% Cgr addition can eliminate the -phase formation by sintering in

vacuum without applied pressure.

FESEM back scattered electron micrographs of as-sintered HS1 and HS6

are shown in Fig. 4.43a and b, respectively. Higher porosity was observable in the

HS6 samples (Fig. 4.43b). The applied pressure caused higher density in HS1

microstructure in comparison with HS6 samples, and it also produced higher

contiguity between the WC grains in the microstructure of HS1, as seen in Fig. 4.36a.

This observation suggests more contact among WC grains as well as between WC

and FeCrNi grains, therefore, more opportunity for the coalescence of WC and the

formation of -phase (Fernandes et al., 2007; Aristizabal et al., 2010). Although there

is no meaningful difference in the grain sizes between the products of the two

methods, less porosity was observed in Fig. 4.36a. Elimination of pores was due to

improve densification during PHIP which indirectly reduced diffusion distance for

-phase formation; this supports the argument that PHIP accelerated -phase

formation.

Figure 4.43 FESEM back-scattered electron images of as-sintered samples: (a) HS1

and (b) HS6

128

As seen in the XRD patterns of as-sintered samples fabricated by PHIP at

1300oC and 20MPa, Fig. 4.44, the evolution of phases varies with NbC content.

While WC continued to be the dominant phase, the intensity of -phase peaks

decreased with increasing NbC content. Furthermore, the height of FeNiCr peaks

also increased. FeCrNi and NbC peaks were displayed more clearly in HS4 and HS5

samples. The peaks of NbC in samples with 1 and 1.5 wt.% NbC could not be

observed that is similar to the results observed in the samples sintered by vacuum

alone as seen in Fig. 4.37, Section 4.7. The -phase peaks disappeared at about 2

wt.% NbC. This finding is similar to that found in a previous study (Xiao et al.,

2009); they showed that the addition of NbC reduced the formation of -phase and,

at the same time, worked as grain growth inhibitor during sintering of WC-Co. They

suggested that the addition of NbC may decrease the concentration of oxygen on the

surface of WC and Co particles, thereby reducing the volume fraction of Co3W3C

phases. Besides that, the formation of (Nb,W)C grains as seen in FESEM images

(Fig. 4.45) and EDX results (Fig. 4.46) may be attributed to a C source in preventing

-phase formation. Table 4.7 shows the effects of NbC content on the -phase

amount sintered by PHIP. At 0 wt.% NbC, 7.3 wt.% -phase is observed. The amount

of -phase decreases with the increase of NbC content and was not detectable in

samples with 2 and 5 wt.% NbC.

129

Figure 4.44 XRD patterns of as-sintered samples with various amounts of NbC; 0

wt.% NbC (HS1), 1 wt.% NbC (HS2), 1.5 wt.% NbC (HS3), 2 wt.% NbC (HS4) and

5 wt.% NbC (HS5)

Table 4.7 Amount of -phase of PHIP-sintered samples various with NbC contents

Sample code HS1 HS2 HS3 HS4 HS5

(wt.%) 7.3 6.5 3.1 - -

Fig. 4.45 presents FESEM back scattered electron images showing the effect

of different contents of NbC on the microstructure of samples which indicate that the

grain size of WC decreased with increasing NbC content. All the NbC-doped

samples exhibited a finer microstructure than that of the undoped sample. A

homogenous grain size distribution was achieved with 2 wt.% NbC (HS4). With the

addition of 5 wt.% NbC (HS5), WC grains still remained in the fine microstructure.

However, the presence of (Nb,W)C grains are observed in the WC matrix. They were

130

confirmed by EDX results (Fig. 4.46). The presence of large (Nb,W)C grains might

be attributed to the agglomeration of NbC during milling, which leads to a

re-crystallization of NbC during sintering and a dissolution of WC in the NbC grains

as discussed in Section 4.7. The mechanism of grain growth inhibitor of NbC by

changing the interface energy and interfere with the dissolution and re-precipitation

stages of WC was proposed as in Section 4.7.

Figure 4.45 FESEM back-scattered images of as-sintered samples with: (a) 1 wt.%

NbC (HS2), (b) 1.5 wt.% NbC (HS3), (c) 2 wt.% NbC (HS4) and (d) 5 wt.% NbC

(HS5)

131

Figure 4.46 EDX analysis of a) point X in Fig. 4.45c and b) point Y in Fig. 4.45d

EDX analysis on point X in Fig. 4.45c and point Y in Fig. 4.45d are shown

in Fig. 4.46. The analysis indicates that WC dissolved in NbC grains during sintering.

At point X of the HS4 sample, up to 28.49 wt.% of WC was dissolved in NbC (2

wt.% added NbC) while at point Y of sample HS5, 22.94 wt.% of WC was dissolved

in NbC (5 wt.% added NbC). As similar to samples sintered in the vacuum furnace,

the dissolution of WC in NbC caused the formation of (Nb,W)C carbide in the

microstructure and the increase in additional NbC content led to the formation of

(Nb,W)C grains, perhaps accounting for the presence of coarse (Nb,W)C grains in

the HS5 sample. The calculation of WC grain size based on AGI method is shown in

Table 4.8. The results show a similar trend that was observed in samples sintered in

vacuum furnace. About 0.86 m of WC grains obtained in sample without NbC

addition, the size of WC was then decreased with the addition of NbC up to 2 wt.%.

The smallest WC grain size was obtained about 0.56 . The effect of NbC on the

inhibition to WC grain growth was then decreased at higher NbC content which

resulted in an increase of WC grain size, about 0.64 m was calculated in sample

with 5 wt.% NbC.

132

Table 4.8 WC grain size of PHIP sintered samples with NbC addition

NbC (wt.%) 0.0 1.0 1.5 2.0 5.0

WC grain size (m) 0.86

0.78

0.72

0.55

0.64


Table 4.3 shows the relative density of samples sintered by PHIP sintering

and samples sintered in vacuum without pressure. An increase of nearly 4.5 % in

HS1 compared to HS6 indicates that the higher densification was generated by PHIP.

Increasing NbC content led to a slight decrease of relative density. This must be due

to the lower density of NbC compared to WC.

Table 4.9 Relative densities of sintered samples

Sample

code

NbC addition

(wt.%)

Relative density (%)

Vacuum sintering PHIP sintering

HS1 0.0 96.85

HS2 1.0 96.51

HS3 1.5 96.30

HS4 2.0 96.14

HS5 5.0 95.34

HS6 0.0 92.4


Fig. 4.47 presents the Vickers hardness (HV30) and fracture toughness of

samples with various contents of added NbC. It can be seen that the hardness of

as-sintered samples improved as NbC content increased up to 2 wt.%, leading to

133

decreased WC grain size and, therefore, increased hardness. This shows a similar

trend to the hardness of samples sintered in vacuum alone, Section 4.7.3. The

increase of hardness with the reduction of WC grain size is explained by the

Hall-Petch relation, the strength of a material depends on its grain size. Consequently,

a decrease in grain size leads to an increase in strength (Niels, 2004). The highest

hardness, about 1820 kg/mm2, was exhibited by sample HS4 with 2 wt.% added NbC.

The lower hardness in the HS5 (5 wt.% NbC) compared to HS4 (2 wt.% NbC) was

due to the presence of large (Nb,W)C grains in WC matrix.

Figure 4.47 Hardness and fracture toughness of PHIP-sintered samples vs

with NbC addition

The fracture toughness of samples decreased with increasing NbC content.

The highest value of KIC was achieved at about 10 MPa.m1/2 in HS1 with 0 wt.%

NbC. Although the highest hardness was achieved at 2 wt.% added NbC (HS4), the

134

value of KIC for this sample fell to 7.7 MPa.m1/2. The slightly higher KIC in the HS4

sample relative to that of the HS3 sample (1.5 wt.% added NbC) which might be

due to HS4’s more uniform microstructure. With 5 wt.% added NbC, fracture

toughness was slightly increased while the hardness decreased, perhaps attributable

to the presence of large NbC grains in the microstructure. The increase of hardness

leading to a decrease of fracture toughness is a general trend that has been reported

elsewhere Chermant and Osterstock, 1976; Schubert et al., 1998; Vaßen and Stöver,

1999.

In comparison with other works, several data from selected references was

compared with mechanical properties of samples produced in this work. Fig.4.48a

and Fig.4.48b shows the hardness and fracture toughness of samples with different

NbC content sintered in vacuum furnace and by PHIP. The content of NbC vs.

hardness and fracture toughness is only for this work. Data of hardness and fracture

toughness from several reports was also plotted for comparison. Fig. 4.48 shows that

PHIPed samples had higher Vickers hardness than those in sample sintered in

vacuum without pressure. As a result, fracture toughness of samples sintered in

vacuum was also higher than those sintered by PHIP at all composition containing

NbC.

135

Figure 4.48 (a)-Hardness and (b)-fracture toughness of this work compared with

references (Jia et al., 1998; Richter and Ruthendorf, 1999; Kim et al., 2007; Mondal

et al., 2008; Fernandes et al., 2009b)

Fig. 4.48a and b also show the results of Vickers hardness and fracture

toughness of other works in literature in comparison with the results of this work.

Ultrafine WC-Co powders was consolidated by pulsed current activated sintering

(Kim et al., 2007). They reported that Vickers hardness of sintered samples was

136

obtained in the range of 1375 – 2480 (kg/mm2) while fracture toughness was

obtained from 12.2 to 6.6 MPa.m1/2. High Vickers hardness was obtained by

producing ultrafine WC grains (~ 400 nm), however, low fracture toughness was

achieved about 6 - 7 MPa.m1/2 (Richter and Ruthendorf, 1999). In another study, high

Vickers hardness (1900-2300 kg/mm2) with moderate fracture toughness (8 – 8.5

MPa.m1/2) of WC-Co hardmetals were produced from nano WC-Co powders (Jia et

al., 1998). High fracture toughness (12.5 -18.8 MPa.m1/2) and moderate hardness

(1560 – 1700 kg/mm2) were obtained in WC-5Co added TiCN as a WC grain growth

inhibitor (Mondal et al., 2008). High hardness (1840 kg/mm2) and fracture toughness

(9.7 MPa.m1/2) of WC-Fe4.3Ni0.6Cr1 hardmetals were achieved by sintering WC

sputter coated iron rich binder (Fernandes et al., 2009b).

It is clear to see that the values of hardness and fracture toughness of sample

with 1- 5 wt.% NbC addition sintered by both vacuum alone and PHIP are in the

ranges provided by other references.

4.8.5 Summary

WC-10FeCrNi hardmetals were fabricated using PHIP sintering at 1300oC

and 20 MPa pressure. PHIP sintering improved densification of WC-10FeCrNi hard

metal, but this method also produced a greater amount of -phase. -phase was not

observed in the sample sintered in vacuum without pressure suggesting that the

addition of 2 wt.% Cgr can be used to avoid -phase formation in WC-10FeCrNi

sintered in vacuum alone. Hardness and fracture toughness were evaluated for each

137

WC-FeCrNi specimen sintered by PHIP. This study showed that 2 wt.% addition of

NbC was the suitable amount of NbC to be used in the production of WC-FeCrNi

hardmetals with fine microstructure. This finding was supported by FESEM

observations. The experimental data showed that Vickers hardness increases with

increasing addition of NbC up to 2 wt.%, albeit accompanied by a slight decrease in

fracture toughness, while the addition of 5 wt.% NbC caused the formation of large

(Nb,W)C grains in microstructure which reduced the Vickers hardness. Besides that,

the results show NbC’s ability to reduce -phase formation during PHIP sintering.

Besides that comparison with other results, WC-10FeCrNi-2Cgr with NbC addition

sintered by either vacuum sintering or PHIP sintering in this work has a potential to

produce hardmetals for cutting tool application.


properties of samples sintered by PHIP

In this section, WC-10FeCrNi-2 Cgr -1NbC powders were subjected to sintering

using PHIP at different temperatures in order to study the effect of PHIP sintering

temperature on the microstructure and mechanical properties of sintered samples.

Considering PHIP as a useful technique to produce high densification, PHIP was also

carried out at lower temperatures than those in vacuum sintering for comparison.

PHIP were done at 1150, 1200, 1250 and 1300oC at pressure of 20 MPa, respectively.


Fig. 4.49 shows the XRD results of PHIP-sintered samples at different

temperature. As discussed in Section 4.6.1, -phase formation was completely

138

eliminated by addition of 2 wt.% Cgr in the case of sintering at 1350oC in vacuum.

However, the XRD results of samples sintered by PHIP shows that -phase

formation increased with the temperature even though with the presence of 2 wt.%

Cgr in the composition. -phase was not observed only at the sintering temperature of

1150oC. As discussion in previous section, -phase fraction increases with

temperature and can be eliminated by addition of Cgr. The assistance of pressure

produced more contacts between FeCrNi binder and C deficit areas of WC and

therefore, improved the formation of -phase. The addition of 2 wt.% Cgr e can only

compensate the C loss during PHIPed at 1150oC and hence, -phase was not

observed in this sample. At higher temperature, the liquid phase is more flexible to

contact with WC by the capillary force (Bhaumik et al., 1996; Silva et al., 2001);

consequently, -phase formation in our samples was more pronounced and became

significant at 1300oC. The more flexible of liquid phase also leads to some

dissolution of -phase in liquid phase (Fernandes et al., 2007) and improve the

diffusion of graphite to the deficit C areas. However, the formation of -phase rate

may higher than the dissolution of -phase in liquid and the diffusion rate of Cgr, thus

the presence of -phase was observed. The -phase fraction calculated is shown in

Table 4.4. The reduction of binder phase in sintered samples was also observed as the

result of -phase formation.

139

Figure 4.49 XRD patterns of WC/10FeCrNi-1NbC sintered by PHIP at different

temperature

Table 4.4 -phase fraction formed at different temperature sintered by PHIP

Temperature (oC) 1150 1200 1250 1300

-phase fraction (%) - ~1 4.4 6.5

Fig. 4.50 shows the FESEM back-scattered images of PHIPed-samples at

different temperatures. It is seen that the higher porosity of sample at lower sintered

temperature. Besides that the incomplete liquid formation of binder phase was

observed at low temperature, 1150 and 1200oC although the liquid phase has been

reported to be formed at about 1150oC (Fernandes et al., 2007). This is a similar

trend to that of sample sintered in vacuum alone. However, at 1250oC, the binder

phase is not observed clearly as shown in sample sintered by vacuum without

pressure. This may due to the assistance of applied pressure during PHIP. The

applied force may enhance the capillary force and hence, accelerating the diffusion

of liquid binder between WC grains.

140

Figure 4.50 FESEM back-scattered images of PHIPed-samples at: (a) 1150oC, (b)

1200oC, (c) 1250oC and (d) 1300oC

FESEM images also show that the higher sintered temperature led to more

grain growth of WC. At 1300oC, the less porosity was observed and coarser WC

grain size is obtained. The grain morphology of WC also changed from round shape

at low temperature to faceted platelet shape at higher temperature (1300oC). As the

binder phase was completely liquid phase at 1250 and 1300oC, the microstructure of

samples shows a better densification (Fig. 4.50 c and d). FESEM images show the

coarser WC grains at higher temperature, however, WC grain size did not indicate

significant difference between 1250oC and 1300oC. This is due to the role of NbC as

a grain growth inhibitor.

141


The density of sintered samples was measured and shown in Fig. 4.50. As

similar result to samples sintered by vacuum without pressure, the density of PHIPed

samples also increases with temperature. Besides that, the result also shows a drastic

increase value at high temperature. At low temperature, 1150oC and 1200oC,

although there was an assistance of applied pressure, the densities of sintered

samples were lower than 90% of relative density. This may be caused by the

incompletely liquid formation of binder phase as seen in FESEM images (Fig. 4.50).

As the binder phase has completely become liquid, the density quickly increased

higher than 90 % of relative density (94.2 % at 1250oC and 96.5 % at 1300oC). This

show the efficiency of applied pressure as the liquid phase formed.

Figure 4.51 Density of PHIPed samples at different temperature

142

4.9.3 Hardness and fracture toughness of sintered samples

Vickers hardness and fracture toughness of PHIPed samples are shown in Fig.

4.52. Vickers hardness of samples increased with sintered temperature. As the

temperature increased, the densification was promoted with higher density thus the

bonding between WC and FeCrNi and among WC grains were enhanced. The grain

growth of WC was also occurred with the increase of temperature. As the grain

growth happened, hardness decreased with the larger grain size (Niels, 2004),

however; grain growth was lowered by the addition of NbC. A slightly bigger in

grain size of WC was observed at 1300oC compared to that at 1250oC. However;

with higher densification, hardness of sample sintered at 1300oC (about 1640

kg/mm2) was higher than at 1250oC (about 1570 kg/mm2).

Figure 4.52 (a) Vickers hardness and fracture toughness of

WC-10FeCrNi-2graphite-1NbC sintered by HIP

143

The fracture toughness results show a different trend in comparison with Vickers

hardness. KIC values decreased with the increase of sintering temperature. However,

there is an irregularity at 1200oC. Fracture toughness of sample sintered at 1200oC

(about 10.4 MPa.m1/2) was higher than that at 1150oC (about 10.26 MPa.m1/2) despite

of the formation of -phase at 1200oC. This may be attributed to the higher

densification sample even though small amount of -phase was formed. The

decrease of fracture toughness at higher temperature (1250 and 1300oC) was resulted

from the increase of hardness (Chermant and Osterstock, 1976; Schubert et al., 1998;

Vaßen and Stöver, 1999) and the significant formation of -phase presented in

samples at high temperature that reduced fracture toughness as well.

4.9.4 Summary

The results of this section show that PHIP process produced higher sintered

density compared to samples sintered in vacuum without assisted pressure at the

same sintering temperature. 2 wt.% graphite can only compensate the C loss during

sintering WC-10FeCrNi-2Cgr-1NbC at low temperature (1150oC). At higher

temperature, more contacts between WC and binder phase were generated by applied

pressure in PHIP process and hence, forming more -phase. The addition of NbC can

inhibit WC grain growth, and hence samples obtained higher Vickers hardness.

Although PHIP caused more -phase formation, the addition of 1 wt.% NbC and 2

wt.% Cgr bring to a higher fracture toughness. Increasing sintering temperature led to

a higher hardness but also lowered the fracture toughness of sintered samples.

144

4.10 Effect of pressure on the microstructure and mechanical properties of

PHIP sintered samples

WC-10FeCrNi composition of samples used in this part contained 2 wt.%

graphite and 1 wt.% NbC (WC-10FeCrNi-2Cgr -1NbC). To study the effect of

pressure in PHIP process on the microstructure and mechanical properties of sintered

sample, PHIP was carried out at 1300oC with different applied pressure: 0, 5, 15, 20

and 25 MPa, respectively.


The phase evolution of samples after sintering is presented in Fig. 4.53. -phase

was not observed in sample sintering for 45 min without the applied pressure. The

addition of 2 wt.% Cgr and 1 wt.% NbC enabled to eliminate the formation of

-phase sintered in vacuum alone. -phase formed with the presence of pressure

which was discussed and defined in the previous section 4.8. Calculated -phase

fraction is shown in Fig. 4.54. At 5 MPa, -phase was detected about 2.8 % and

became significant at the pressure of 15 MPa as about 6.2 %. The -phase fraction

increased with the increase of the pressure. However, -phase formation rate was

slow suggesting that the formation of -phase may reach the critical value for this

composition of WC/10FeCrNi hardmetal. It is also observed that the binder phase

became not clear at the sample pressed at 25 MPa.

145

Figure 4.53 XRD patterns of PHIPed samples vs. applied pressure

Figure 4.54-phase fraction vs. applied pressure in PHIP process

FESEM images of sintered samples are shown in Fig. 4.55. WC grain size

and shape of WC are similar in all samples. This shows that the pressure did not give

a significant effect on the size and the shape of WC in the range of pressure selected

in this work. However, high porosity was observed in the sample without pressure.

146

Figure 4.55 FESEM back scattered electron images of samples vs. pressure in PHIP

Increasing the pressure led to a decrease of pores in the microstructure of

sintered samples. At 5 MPa, the pores was not very different compared to sample

without pressure. A significant reduction of pores occurred with pressure of 15 MPa.

At higher pressure than 15 MPa, the microstructure did not change significantly.

Small amount of pores was still remaining in the microstructure. Thus the full

147

density was not obtained in this range of pressure. To get full density, the pressure

needs to be increased or longer sintering time and higher temperature should be

improved. The higher densification of sample with pressure is confirmed by density

measurement as seen in Fig. 4.56.


Density and relative density of sintered samples with different pressure are

shown in Fig.4.56. The density of sample increased with the pressure. The results

also show that a fast increase of density was obtained at pressure of 15 MPa since an

increase of nearly 4.3% relative density was obtained at this pressure compared to

that of sample without pressure (91.9%). With pressure higher than 15 MPa, 15 and

20 MPa, the density continued to increase with slow rate, ~ 4.6 and ~ 4.8 %

increasing of relative density were obtained, respectively. The results of density

show appropriateness with FESEM images as seen in Fig. 4.55.

Figure 4.56 Density vs. applied pressure of samples sintered by PHIP

148


Vickers hardness and fracture toughness of sintered sample are presented in

Fig. 4.57. Vickers hardness of samples increased with the pressure. The maximum

Vickers hardness value was obtained at the pressure of 25 MPa (about 1660 kg/mm2).

All though -phase formation increased with the applied pressure, the increase of

hardness was resulted from the higher density with less porosity of samples. The

higher densification of sintered sample led to a higher strengthened bonding among

WC grains and WC grains with binder phase, and consequently increased hardness of

sintered samples. Therefore, using PHIP for sintering enhanced the hardness of

samples by the improving densification of WC-FeCrNi-2Cgr -1NbC hardmetal.

Figure 4.57 Hardness and fracture toughness vs. pressure of PHIPed samples

Both of Vickers hardness and fracture toughness increased with pressure up

to 15 MPa (from 9.35 to 9.6 MPa.m1/2); this may be caused by the higher density of

149

sintered samples. Previous study (Sánchez et al., 2005) also shows that the fracture

strength of WC-Co improves due to the removal of residual porosity which promote

crack propagation resulting in the reduction of fracture strength. At higher pressure,

the density of samples continued to increase, however; the fracture toughness

decreased. This may be due to the more formation of -phase at higher applied

pressure and the increase of hardness of sintered samples.

4.10.4 Summary

WC/10FeCrNi added with 2 wt.% Cgr and 1 wt.% NbC hardmetals were

sintered by PHIP at different pressure. The results show that the densification of

sintered samples was improved by the applied pressure. Increasing the pressure led to

higher density of samples with less porosity. No significance in grain size and shape

of WC was observed upon the pressure application. However, more -phase was

formed at high pressure which resulted in more contacts between WC and binder

phases. Both Vickers hardness and fracture toughnes of sintered samples increased

with the pressure up to 15 MPa. At higher applied pressure, hardness of samples

increased, however; fracture toughness indicated a tendency to decrease. This may

attributed to the increase of hardness and the more formation of -phase.

150

CHAPTER 5

CONCLUSIONS AND RECOMMENDATIONS

5.1 Conclusion

WC-AISI304 hardmetals were successfully produced by sintering in both

vacuum furnace and PHIP. During sintering, -phase (Fe3W3C) was formed in the

microstructure of sintered samples. The formation of -phase was found to be

increased with milling time of mixed powder, sintering time, sintering temperature,

binder content and applied pressure (PHIP sintering). In this work, -phase was

eliminated by the addition of graphite powder prior to mixing in planetary ball

milling. The added 2 wt.% Cgr completely compensated the C loss during sintering

for WC-10AISI304 up to 1350oC in vacuum, thus no -phase was observed in the

microstructure of sintered samples. When the formation of -phase was eliminated,

both hardness and fracture toughness of WC-10AISI304 increased up to 1620

kg/mm2 and 10.2 MPa.m1/2, respectively.

Grain growth of WC was observed to increase with sintering time and sintering

temperature. The applied pressure did not cause significant effect in the WC grain

growth (in the range of pressure (PHIP sintering) by this work). WC grain growth of

WC-10FeCrNi-2Cgr can be inhibited by the addition of NbC up to 2 wt.% after

sintering at 1300oC by both vacuum sintering and PHIP sintering. The addition of

NbC more than 2 wt.% caused the formation of several coarse (Nb,W)C grains which

was resulted from the dissolusion of WC in NbC during sintering. As the WC grain

growth was inhibited, the hardness of sintered samples are highly increased,

151

however; their fracture toughness are decreased with the increase of hardness.

Besides that NbC was found to have the ability to reduce the formation of -phase

during sintering.

PHIP showed a high potential to sinter samples with higher density compared

to vacuum sintering. In general, the PHIPed samples are having about 4-5 % higher

relative density compared to samples produced by vacuum sintering. Consequently,

higher hardness was obtained for samples sintered by PHIP. However, more contacts

among WC and binder phase generated by applied pressure in PHIP led to a more

formation of -phase which resulted in the decrease of fracture toughness

In this work, WC-10AISI304-2Cgr-NbC hardmetals that was sintered at 1300oC

after planetary ball milling 15 h under Ar gas gave the Vickers hardness in the range

of 1600 to 1660 kg/mm2 and about 8.7 to 8.3 MPa.m1/2 for fracture toughness, KIC,

by vacuum sintering. However, the same samples produced via PHIP sintering gave

the Vickers hardness from 1640 to 1820 kg/mm2 and fracture toughness, KIC, from

10 to 7.3 MPa.m1/2. These range of hardness and fracture toughness are in the

intermediate range compared to other systems provided by literatures. This shows

that WC-10AISI304 hardmetals produced by this work can be applied to produce

cutting inserts for some specific application such as wood cutting or metal cutting.

And hence, AISI304 could be proposed to replace Co binder in order to produce

cutting inserts.

152

5.2 Recommendations

Other advanced sintering methods such as hot isostatic pressing (HIP) or spark

plasma sintering (SPS), etc. are suggested to get full densification so that mechanical

properties (hardness and fracture toughness or wear resistance) of the similar

hardmetals product could be compared with products of this work

Research on the corrosion resistance is suggested to investigate the corrosion

behaviour of WC-AISI304 hardmetals in corrosive environment.

Studies on the effect of other carbides such as VC or TaC grain growth

inhibitor are also possible in order to find out the optimum composition for better

mechanical properties.

153

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APPENDIX

LIST OF PUBLICATIONS & CONFERENCE

1 T.B. Trung, H. Zuhailawati,A.A. Zainal and N. K. Ishihara, Grain

growth, phase evolution and properties of NbC carbide-doped

WC-10AISI304 hardmetals produced by pseudo hot isostatic

pressing, Journal of Alloys and Compounds 552 (2013) pp 20–25

2 T.B. Trung, H. Zuhailawati,A.A. Zainal and N. K. Ishihara, Role

of NbC grain growth inhibitor on microstructure and hardness of

WC-stainless steel hard metals, Advanced Materials Research

Vol. 620 (2013) pp 424-428

3 T.B. Trung, H. Zuhailawati,A.A. Zainal and N. K. Ishihara,

Microstructure evolution of WC/10 wt.%AISI304-stainless steel

hardmetal fabricated by powder metallurgy, 3rd RC-NMR

Regional conference interdisciplinary on nature resources and

materials engineering, 25th – 26th October 2010, Bayview hotel,

Langkawi, Malaysia

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