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
4.2.3 Density measurement 87
4.2.4 Hardness and fracture toughness 88
4.3 Effect of sintering temperature on microstructure and mechanical
properties of WC-10FeCrNi hardmetals
90
4.3.1 Phase identification and microstructure of sintered samples 90
4.3.2 Density measurement of sintered samples 93
4.3.3 Hardness and fracture toughness 95
4.4 Effect of sintering time on the microstructure and mechanical
properties of WC-10FeCrNi hardmetals
96
4.4.1 Phase identification and microstructure of sintered samples 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
4.5.2 Density of sintered samples 107
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
4.6.2 Density of sintered samples 113
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
4.7.2 Density of sintered samples 122
4.7.3 Hardness and fracture toughness 123
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
4.8.2 Density measurement 132
4.8.3 Hardness and fracture toughness 132
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
4.9.2 Density measurement 141
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
4.10.2 Density measurement 147
4.10.3 Hardness and fracture toughness 148
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
Melting point (oC) 3490
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
Melting point (oC) 3650
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.
4.3 Effect of sintering temperature on microstructure and mechanical
properties of WC-10FeCrNi hardmetals
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.
4.3.1 Phase identification and microstructure of sintered samples
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
4.3.3 Hardness and fracture toughness
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
properties of WC-10FeCrNi hardmetals
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.
4.4.1 Phase identification and microstructure of sintered samples
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
4.5.2 Density of sintered samples
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
4.6.2 Density of sintered samples
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
4.7.2 Density of sintered samples
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
4.7.3 Hardness and fracture toughness
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
4.8.2 Density measurement
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
4.8.3 Hardness and fracture toughness
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.
4.9 Effect of sintering temperature on microstructure and mechanical
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.
4.9.1 Phase identification and microstructure
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
4.9.2 Density measurement
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.
4.10.1 Phase identification and microstructure
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.
4.10.2 Density measurement
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
4.10.3 Hardness and fracture toughness
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
REFERENCES
Acchar, W., Zollfrank, C. and Greil, P. (2004). Microstructure and mechanical
properties of WC-Co reinforced with NbC. Materials Research, 7, p.445-450.
Adabavazeh, Z., Karimzadeh, F. and Enayati, M. H. (2012). Synthesis and structural
characterization of nanocrystalline (Ni,Fe)3Al intermetallic compound
prepared by mechanical alloying. Advanced Powder Technology, 23,
p.284-289.
Allibert, C. H. (2001). Sintering features of cemented carbides WC–Co processed
from fine powders. International Journal of Refractory Metals and Hard
Materials, 19, p.53-61.
Arató, P., Bartha, L., Porat, R., Berger, S. and Rosen, A. (1998). Solid or liquid phase
sintering of nanocrystalline WC/Co hardmetals. Nanostructured Materials, 10,
p.245-255.
Arenas, F., de Arenas, I. B., Ochoa, J. and Cho, S. A. (1999). Influence of VC on the
microstructure and mechanical properties of WC–Co sintered cemented
carbides. International Journal of Refractory Metals and Hard Materials, 17,
p.91-97.
Aristizabal, M., Ardila, L. C., Veiga, F., Arizmendi, M., Fernandez, J. and Sánchez, J.
M. (2012). Comparison of the friction and wear behaviour of WC–Ni–Co–Cr
and WC–Co hardmetals in contact with steel at high temperatures. Wear,
280–281, p.15-21.
Aristizabal, M., Rodriguez, N., Ibarreta, F., Martinez, R. and Sanchez, J. M. (2010).
Liquid phase sintering and oxidation resistance of WC–Ni–Co–Cr cemented
carbides. International Journal of Refractory Metals and Hard Materials, 28,
p.516-522.
154
Azcona, I., Ordóñez, A., Sánchez, J. M. and Castro, F. (2002). Hot isostatic pressing
of ultrafine tungsten carbide-cobalt hardmetals. Journal of Materials Science,
37, p.4189-4195.
Bai, Y., He, X., Zhu, C. and Chen, G. (2012). Microstructures, electrical, thermal,
and mechanical properties of bulk Ti2AlC synthesized by self-propagating
high-temperature combustion synthesis with pseudo hot isostatic pressing.
Journal of the American Ceramic Society, 95, p.358-364.
Barbatti, C., Garcia, J., Brito, P. and Pyzalla, A. R. (2009). Influence of WC
replacement by TiC and (Ta,Nb)C on the oxidation resistance of Co-based
cemented carbides. International Journal of Refractory Metals and Hard
Materials, 27, p.768-776.
Benjamin, J. (1970). Dispersion strengthened superalloys by mechanical alloying.
Metallurgical and Materials Transactions B, 1, p.2943-2951.
Benjamin, J. and Volin, T. (1974). The mechanism of mechanical alloying.
Metallurgical and Materials Transactions B, 5, p.1929-1934.
Bergström, M. (1977). The eta-carbides in the ternary system Fe-W-C at 1250 °C.
Materials Science and Engineering, 27, p.257-269.
Bhaumik, S. K., Upadhyaya, G. S. and Vaidya, M. L. (1996). Full density processing
of complex WC-based cemented carbides. Journal of Materials Processing
Technology, 58, p.45-52.
Breval, E., Cheng, J. P., Agrawal, D. K., Gigl, P., Dennis, M., Roy, R. and Papworth,
A. J. (2005). Comparison between microwave and conventional sintering of
WC/Co composites. Materials Science and Engineering: A, 391, p.285-295.
Carpinteri, A., Pugno, N. and Puzzi, S. (2009). Strength vs. toughness optimization
of microstructured composites. Chaos, Solitons & Fractals, 39,
155
p.1210-1223.
Cha, S. I., Hong, S. H., Ha, G. H. and Kim, B. K. (2001a). Mechanical properties of
WC–10Co cemented carbides sintered from nanocrystalline spray conversion
processed powders. International Journal of Refractory Metals and Hard
Materials, 19, p.397-403.
Cha, S. I., Hong, S. H., Ha, G. H. and Kim, B. K. (2001b). Microstructure and
mechanical properties of nanocrystalline WC-10Co cemented carbides.
Scripta Materialia, 44, p.1535-1539.
Chabretou, V., Allibert, C. H. and Missiaen, J. M. (2003). Quantitative analysis of the
effect of the binder phase composition on grain growth in WC-Co sintered
materials. Journal of Materials Science, 38, p.2581-2590.
Chabretou, V., Lavergne, O., Missiaen, J. and Allibert, C. (1999). Quantitative
evaluation of normal and abnormal grain growth of cemented carbides during
liquid phase sintering. Metals and Materials International, 5, p.205-210.
Chermant, J. L. and Osterstock, F. (1976). Fracture toughness and fracture of WC-Co
composites. Journal of Materials Science, 11, p.1939-1951.
Choi, K., Hwang, N. M. and Kim, D. Y. (2000). Effect of VC addition on
microstructural evolution of WC-Co alloy: mechanism of grain growth
inhibition. Powder Metallurgy, 43, p.168-172.
Cullity, B. D. and Stock, S. R. (2001). Elements of X-ray diffraction third edition,
Prentice-Hall, Inc. p. 385 - 433.
Da Silva, A. G. P., De Souza, C. P., Gomes, U. U., Medeiros, F. F. P., Ciaravino, C.
and Roubin, M. (2000). A low temperature synthesized NbC as grain growth
inhibitor for WC–Co composites. Materials Science and Engineering: A, 293,
p.242-246.
156
Daoush, W. M., Park, H. S., Lee, K. H., Moustafa, S. F. and Hong, S. H. (2009).
Effect of binder compositions on microstructure, hardness and magnetic
properties of (Ta,Nb)C–Co and (Ta,Nb)C–Ni cemented carbides.
International Journal of Refractory Metals and Hard Materials, 27,
p.669-675.
David, D. W., Morgan, W. K., Perkins, D. G. and Rubin, A. (1990). The Respiratory
Effects of Cobalt Arch. Intern. Med., 150, p.7.
De Boeck, M., Kirsch-Volders, M. and Lison, D. (2003). Cobalt and antimony:
genotoxicity and carcinogenicity. Mutation Research/Fundamental and
Molecular Mechanisms of Mutagenesis, 533, p.135-152.
de Macedo, H. R., da Silva, A. G. P. and de Melo, D. M. A. (2003). The spreading of
cobalt, nickel and iron on tungsten carbide and the first stage of hard metal
sintering. Materials Letters, 57, p.3924-3932.
Deng, X., Patterson, B. R., Chawla, K. K., Koopman, M. C., Fang, Z., Lockwood, G.
and Griffo, A. (2001). Mechanical properties of a hybrid cemented carbide
composite. International Journal of Refractory Metals and Hard Materials,
19, p.547-552.
Deorsola, F. A., Vallauri, D., Ortigoza Villalba, G. A. and Benedetti, B. D. (2010).
Densification of ultrafine WC–12Co cermets by pressure assisted fast electric
sintering. International Journal of Refractory Metals and Hard Materials, 28,
p.254-259.
El-Eskandarany, M. S., Mahday, A. A., Ahmed, H. A. and Amer, A. H. (2000).
Synthesis and characterizations of ball-milled nanocrystalline WC and
nanocomposite WC–Co powders and subsequent consolidations. Journal of
Alloys and Compounds, 312, p.315-325.
157
Eriksson, M., Radwan, M. and Shen, Z. Spark plasma sintering of WC, cemented
carbide and functional graded materials. International Journal of Refractory
Metals and Hard Materials.
Eso, O., Fang, Z. Z. and Griffo, A. (2007). Kinetics of cobalt gradient formation
during the liquid phase sintering of functionally graded WC–Co.
International Journal of Refractory Metals and Hard Materials, 25,
p.286-292.
Ettmayer, P. (1989). Hardmetals and Cermets. Annual Review of Materials Science,
19, p.145-164.
Fang, Z., Maheshwari, P., Wang, X., Sohn, H. Y., Griffo, A. and Riley, R. (2005). An
experimental study of the sintering of nanocrystalline WC–Co powders.
International Journal of Refractory Metals and Hard Materials, 23,
p.249-257.
Fernandes, C. M. 2008. Sputtering on the production of tungsten carbide based
composites. Ph.D, University of Aveiro.
Fernandes, C. M., Ferreira, V. M., Senos, A. M. R. and Vieira, M. T. (2003a).
Stainless steel coatings sputter-deposited on tungsten carbide powder
particles. Surface and Coatings Technology, 176, p.103-108.
Fernandes, C. M., Popovich, V., Matos, M., Senos, A. M. R. and Vieira, M. T.
(2009a). Carbide phases formed in WC–M (M = Fe/Ni/Cr) systems.
Ceramics International, 35, p.369-372.
Fernandes, C. M. and Senos, A. M. R. (2011). Cemented carbide phase diagrams: A
review. International Journal of Refractory Metals and Hard Materials, 29,
p.405-418.
Fernandes, C. M., Senos, A. M. R. and Vieira, M. T. (2003b). Sintering of tungsten
158
carbide particles sputter-deposited with stainless steel. International Journal
of Refractory Metals and Hard Materials, 21, p.147-154.
Fernandes, C. M., Senos, A. M. R. and Vieira, M. T. (2006). Particle surface
properties of stainless steel-coated tungsten carbide powders. Powder
Technology, 164, p.124-129.
Fernandes, C. M., Senos, A. M. R. and Vieira, M. T. (2007). Control of eta carbide
formation in tungsten carbide powders sputter-coated with (Fe/Ni/Cr).
International Journal of Refractory Metals and Hard Materials, 25,
p.310-317.
Fernandes, C. M., Senos, A. M. R., Vieira, M. T. and Fernandes, J. V. (2009b).
Composites from WC powders sputter-deposited with iron rich binders.
Ceramics International, 35, p.1617-1623.
Froschauer, L. and Fulrath, R. (1976). Direct observation of liquid-phase sintering in
the system tungsten carbide-cobalt. Journal of Materials Science, 11,
p.142-149.
Gee, M. G., Gant, A. and Roebuck, B. (2007). Wear mechanisms in abrasion and
erosion of WC/Co and related hardmetals. Wear, 263, p.137-148.
German, R. M., Smid, I., Campbell, L. G., Keane, J. and Toth, R. (2005). Liquid
phase sintering of tough coated hard particles. International Journal of
Refractory Metals and Hard Materials, 23, p.267-272.
Gille, G., Bredthauer, J., Gries, B., Mende, B. and Heinrich, W. (2000). Advanced
and new grades of WC and binder powder – their properties and application.
International Journal of Refractory Metals and Hard Materials, 18, p.87-102.
Gille, G., Szesny, B., Dreyer, K., van den Berg, H., Schmidt, J., Gestrich, T. and
Leitner, G. (2002). Submicron and ultrafine grained hardmetals for
159
microdrills and metal cutting inserts. International Journal of Refractory
Metals and Hard Materials, 20, p.3-22.
Gilman, P. S. and Benjamin, J. S. (1983). Mechanical alloying. Annual Review of
Materials Science, 13, p.279-300.
González, R., Echeberría, J., Sánchez, J. M. and Castro, F. (1995). WC-(Fe,Ni,C)
hardmetals with improved toughness through isothermal heat treatments.
Journal of Materials Science, 30, p.3435-3439.
Goodfellow. 2012. Stainless steel AISI304 [Online]. Available: www.goodfellow.com
[Accessed].
Guillermet, A. (1989a). Thermodynamic properties of the Co-W-C system.
Metallurgical and Materials Transactions A, 20, p.935-956.
Guillermet, A. F. (1989b). The Co-Fe-Ni-W-C phase diagram: a thermodynamic
description and calculated sections for (Co-Fe-Ni) bonded cemented WC
tools. Z Metallkunde, 80 (2), p.12.
Guo, Z., Xiong, J., Yang, M., Song, X. and Jiang, C. (2008). Effect of Mo2C on the
microstructure and properties of WC–TiC–Ni cemented carbide.
International Journal of Refractory Metals and Hard Materials, 26,
p.601-605.
Gustafson, P. (1988). A thermodynamic evaluation of the C-Cr-Fe-W system.
Metallurgical and Materials Transactions A, 19, p.2547-2554.
Habashi and Fathi (1997). Handbook of extractive metallurgy
Hanyaloglu, C., Aksakal, B. and Bolton, J. D. (2001). Production and indentation
analysis of WC/Fe–Mn as an alternative to cobalt-bonded hardmetals.
Materials Characterization, 47, p.315-322.
Hashe, N. G., Neethling, J. H., Berndt, P. R., Andrén, H.-O. and Norgren, S. (2007).
160
A comparison of the microstructures of WC–VC–TiC–Co and WC–VC–Co
cemented carbides. International Journal of Refractory Metals and Hard
Materials, 25, p.207-213.
Hewitt, S. A. and Kibble, K. A. (2009). Effects of ball milling time on the synthesis
and consolidation of nanostructured WC–Co composites. International
Journal of Refractory Metals and Hard Materials, 27, p.937-948.
Hewitt, S. A., Laoui, T. and Kibble, K. K. (2009). Effect of milling temperature on
the synthesis and consolidation of nanocomposite WC–10Co powders.
International Journal of Refractory Metals and Hard Materials, 27, p.66-73.
Huang, S. G., Li, L., Vanmeensel, K., Van der Biest, O. and Vleugels, J. (2007). VC,
Cr3C2 and NbC doped WC–Co cemented carbides prepared by pulsed
electric current sintering. International Journal of Refractory Metals and
Hard Materials, 25, p.417-422.
Huang, S. G., Liu, R. L., Li, L., Van der Biest, O. and Vleugels, J. (2008a). NbC as
grain growth inhibitor and carbide in WC–Co hardmetals. International
Journal of Refractory Metals and Hard Materials, 26, p.389-395.
Huang, S. G., Van der Biest, O. and Vleugels, J. (2008b). VC-doped WC–NbC–Co
hardmetals. Materials Science and Engineering: A, 488, p.420-427.
Huang, S. G., Vanmeensel, K., Li, L., Van der Biest, O. and Vleugels, J. (2008c).
Influence of starting powder on the microstructure of WC–Co hardmetals
obtained by spark plasma sintering. Materials Science and Engineering: A,
475, p.87-91.
Ishihara, K. N. and Shingu, P. H. (1990). A pseudo-HIP process applicable for
near-net-shape synthesis of intermetallic compounds. Journal of the Japan
Society of Powder and Powder Metallurgy, 37 (5), p.670-673.
161
Jia, C., Sun, L., Tang, H. and Qu, X. (2007). Hot pressing of nanometer WC–Co
powder. International Journal of Refractory Metals and Hard Materials, 25,
p.53-56.
Jia, K., Fischer, T. E. and Gallois, B. (1998). Microstructure, hardness and toughness
of nanostructured and conventional WC-Co composites. Nanostructured
Materials, 10, p.875-891.
Johansson, S. A. E. and Wahnström, G. (2010). Theory of ultrathin films at
metal–ceramic interfaces. Philosophical Magazine Letters, 90, p.599-609.
Jorn, L. B. (1985). Binder extrusion in sliding wear of WC-Co alloys. Wear, 105,
p.247-256.
Kang, K. Y., Roemer, J. G. and Ghosh, D. (2000). Microstructural characterization of
cemented carbide samples by image analysis techniques. Powder Technology,
108, p.130-136.
Kellner, F. J. J., Hildebrand, H. and Virtanen, S. (2009). Effect of WC grain size on
the corrosion behavior of WC–Co based hardmetals in alkaline solutions.
International Journal of Refractory Metals and Hard Materials, 27,
p.806-812.
Kim, B. K., Ha, G. H. and Lee, D. W. (1997). Sintering and microstructure of
nanophase WC/Co hardmetals. Journal of Materials Processing Technology,
63, p.317-321.
Kim, H.-C., Oh, D.-Y. and Shon, I.-J. (2004). Sintering of nanophase
WC–15vol.%Co hard metals by rapid sintering process. International Journal
of Refractory Metals and Hard Materials, 22, p.197-203.
Kim, H.-C., Shon, I.-J., Yoon, J.-K. and Doh, J.-M. (2007). Consolidation of ultra
fine WC and WC–Co hard materials by pulsed current activated sintering and
162
its mechanical properties. International Journal of Refractory Metals and
Hard Materials, 25, p.46-52.
Kim, S., Han, S.-H., Park, J.-K. and Kim, H.-E. (2003). Variation of WC grain shape
with carbon content in the WC–Co alloys during liquid-phase sintering.
Scripta Materialia, 48, p.635-639.
Kishino, J., Nomura, H., Shin, S. G., Matsubara, H. and Tanase, T. (2002).
Computational study on grain growth in cemented carbides. International
Journal of Refractory Metals and Hard Materials, 20, p.31-40.
Koutsospyros, A., Braida, W., Christodoulatos, C., Dermatas, D. and Strigul, N.
(2006). A review of tungsten: From environmental obscurity to scrutiny.
Journal of Hazardous Materials, 136, p.1-19.
Kurlov, A. and Gusev, A. (2006). Tungsten carbides and W-C phase diagram.
Inorganic Materials, 42, p.121-127.
Kurlov, A. S., Gusev, A. I. and Rempel, A. A. (2011). Vacuum sintering of WC–8
wt.% Co hardmetals from WC powders with different dispersity.
International Journal of Refractory Metals and Hard Materials, 29,
p.221-231.
Lay, S., Allibert, C. H., Christensen, M. and Wahnström, G. (2008). Morphology of
WC grains in WC–Co alloys. Materials Science and Engineering: A, 486,
p.253-261.
Lay, S., Hamar-Thibault, S. and Lackner, A. (2002). Location of VC in VC, Cr3C2
codoped WC-Co cermets by HREM and EELS. International Journal of
Refractory Metals and Hard Materials, 20, p.61-69.
Lay, S., Thibault, J. and Hamar-Thibault, S. (2003). Structure and role of the
interfacial layers in VC-rich WC-Co cermets. Philosophical Magazine, 83,
163
p.1175-1190.
Lee, H. C. and Gurland, J. (1978). Hardness and deformation of cemented tungsten
carbide. Materials Science and Engineering, 33, p.125-133.
Lee, H. R., Kim, D. J., Hwang, N. M. and Kim, D.-Y. (2003). Role of vanadium
carbide additive during sintering of WC–Co: Mechanism of grain growth
inhibition. Journal of the American Ceramic Society, 86, p.152-154.
Lee, W.-B., Kwon, B.-D. and Jung, S.-B. (2006). Effects of Cr3C2 on the
microstructure and mechanical properties of the brazed joints between
WC–Co and carbon steel. International Journal of Refractory Metals and
Hard Materials, 24, p.215-221.
Lin, C., Kny, E., Yuan, G. and Djuricic, B. (2004). Microstructure and properties of
ultrafine WC–0.6VC–10Co hardmetals densified by pressure-assisted critical
liquid phase sintering. Journal of Alloys and Compounds, 383, p.98-102.
Lison, D., Carbonnelle, P., Mollo, L., Lauwerys, R. and Fubini, B. (1995).
Physicochemical mechanism of the interaction between cobalt metal and
carbide particles to generate toxic activated oxygen species. Chemical
Research in Toxicology, 8, p.600-606.
Liu, S., Huang, Z.-L., Liu, G. and Yang, G.-B. (2006). Preparing nano-crystalline
rare earth doped WC/Co powder by high energy ball milling. International
Journal of Refractory Metals and Hard Materials, 24, p.461-464.
Mahboubi Soufiani, A., Karimzadeh, F. and Enayati, M. H. (2012). Formation
mechanism and characterization of nanostructured Ti6Al4V alloy prepared by
mechanical alloying. Materials & Design, 37, p.152-160.
Mahmoodan, M., Aliakbarzadeh, H. and Gholamipour, R. (2009). Microstructural
and mechanical characterization of high energy ball milled and sintered
164
WC–10wt%Co–xTaC nano powders. International Journal of Refractory
Metals and Hard Materials, 27, p.801-805.
Mannesson, K., Borgh, I., Borgenstam, A. and Ågren, J. (2011). Abnormal grain
growth in cemented carbides — Experiments and simulations. International
Journal of Refractory Metals and Hard Materials, 29, p.488-494.
Markström, A., Frisk, K. and Sundman, B. (2005). A revised thermodynamic
description of the Co-W-C system. Journal of Phase Equilibria and Diffusion,
26, p.152-160.
Menéndez, E., Sort, J., Concustell, A., Suriñach, S., Nogués, J. and Baró, M. D.
(2007). Microstructural evolution during solid-state sintering of ball-milled
nanocomposite WC–10 mass% Co powders. Nanotechnology, 18, p.185609.
Meredith, B. and Milner, D. R. (1976). Densification mechanisms in the tungsten
carbide-cobalt system. Powder Metallurgy, 19, p.38-45.
Metox. 2012. Niobium carbide (NbC) [Online]. Available:
http://en.metox.ru/niobium_carbide_nbc [Accessed].
Missiaen, J. M. and Roure, S. (1998). A general morphological approach of sintering
kinetics: application to WC–Co solid phase sintering. Acta Materialia, 46,
p.3985-3993.
Missiaen, J. M. and Roure, S. (2000). Automatic image analysis methods for the
determination of stereological parameters – application to the analysis of
densification during solid state sintering of WC–Co compacts. Journal of
Microscopy, 199, p.141-148.
Mondal, B., Das, P. K. and Singh, S. K. (2008). Advanced WC–Co cermet
composites with reinforcement of TiCN prepared by extended thermal plasma
route. Materials Science and Engineering: A, 498, p.59-64.
165
Morton, C. W., Wills, D. J. and Stjernberg, K. (2005). The temperature ranges for
maximum effectiveness of grain growth inhibitors in WC–Co alloys.
International Journal of Refractory Metals and Hard Materials, 23,
p.287-293.
Niels, H. (2004). Hall–Petch relation and boundary strengthening. Scripta Materialia,
51, p.801-806.
Niihara, K. (1983). A fracture mechanics analysis of indentation-induced Palmqvist
crack in ceramics. Journal of Materials Science Letters, 2, p.221-223.
Niihara, K., Morena, R. and Hasselman, D. P. H. (1982). Evaluation of KIC of brittle
solids by the indentation method with low crack-to-indent ratios. Journal of
Materials Science Letters, 1, p.13-16.
Norley, J. 2011. The role of natural graphite in electronics cooling [Online].
Available:
http://www.electronics-cooling.com/2001/08/the-role-of-natural-graphite-in-e
lectronics-cooling/ [Accessed 2011].
Park, K., Hong, J. and Hwang, S. (1997). Effect of Cu addition on consolidating
Ti5Si3 by the elemental powder-metallurgical method. Metallurgical and
Materials Transactions A, 28, p.223-228.
Park, Y., Hwang, N. and Yoon, D. (1996). Abnormal growth of faceted (WC) grains
in a (Co) liquid matrix. Metallurgical and Materials Transactions A, 27,
p.2809-2819.
Penrice, T. (1987). Alternative binders for hard metals. Journal of Materials Shaping
Technology, 5, p.35-39.
Peters, C. T. (1979). The relationship between Palmqvist indentation toughness and
bulk fracture toughness for some WC-Co cemented carbides. Journal of
166
Materials Science, 14, p.1619-1623.
Petersson, A. (2004). Sintering shrinkage of WC–Co and WC–(Ti,W)C–Co materials
with different carbon contents. International Journal of Refractory Metals
and Hard Materials, 22, p.211-217.
Pollock, C. and Stadelmaier, H. (1970). The eta carbides in the Fe−W−C and
Co−W−C systems. Metallurgical and Materials Transactions B, 1,
p.767-770.
Ponton, C. B. and Rawlings, R. D. (1989). Vickers indentation fracture toughness test
Part 1 Review of literature and formulation of standardised indentation
toughness equations. Materials Science and Technology, 5, p.865-872.
Raghavan, V. (2007). C-Co-Fe-Ni-W (Carbon-Cobalt-Iron-Nickel-Tungsten).
Journal of Phase Equilibria and Diffusion, 28, p.284-285.
Ramnath, V. and Jayaraman, N. (1987). Quantitative phase analysis by X-ray
diffraction in the Co-W-C system. Journal of Materials Science Letters, 6,
p.1414-1418.
Richter, V. and Ruthendorf, M. v. (1999). On hardness and toughness of ultrafine and
nanocrystalline hard materials. International Journal of Refractory Metals
and Hard Materials, 17, p.141-152.
Roebuck, B. (2006). Extrapolating hardness-structure property maps in WC/Co
hardmetals. International Journal of Refractory Metals and Hard Materials,
24, p.101-108.
Sánchez, J. M., Ordóñez, A. and González, R. (2005). HIP after sintering of ultrafine
WC–Co hardmetals. International Journal of Refractory Metals and Hard
Materials, 23, p.193-198.
Schroter, K. 1925. Hard-metal alloy and the process of making same. Germany
167
patent application. 11 Aug, 1925.
Schubert, W. D., Lassner, E. and Böhlke, W. 2010. Cemented carbides - a success
story [Online]. Available:
http://www.itia.info/assets/files/Newsletter_2010_06.pdf [Accessed June
2010].
Schubert, W. D., Neumeister, H., Kinger, G. and Lux, B. (1998). Hardness to
toughness relationship of fine-grained WC-Co hardmetals. International
Journal of Refractory Metals and Hard Materials, 16, p.133-142.
Seo, O., Kang, S. and Lavernia, E. J. (2003). Growth inhibition of nano WC particles
in WC-Co alloys during liquid-phase sintering. Materials transactions - JIM,
44, p.2339-2345.
Shamah, A. M., Ibrahim, S. and Hanna, F. F. (2011). Formation of nano
quasicrystalline and crystalline phases by mechanical alloying. Journal of
Alloys and Compounds, 509, p.2198-2202.
Shatov, A. V., Ponomarev, S. S. and Firstov, S. A. (2008). Fracture of WC–Ni
cemented carbides with different shape of WC crystals. International Journal
of Refractory Metals and Hard Materials, 26, p.68-76.
Shetty, D. K., Wright, I. G., Mincer, P. N. and Clauer, A. H. (1985). Indentation
fracture of WC-Co cermets. Journal of Materials Science, 20, p.1873-1882.
Shon, I.-J., Jeong, I.-K., Ko, I.-Y., Doh, J.-M. and Woo, K.-D. (2009). Sintering
behavior and mechanical properties of WC–10Co, WC–10Ni and WC–10Fe
hard materials produced by high-frequency induction heated sintering.
Ceramics International, 35, p.339-344.
Silva, A. G. P. d., Schubert, W. D. and Lux, B. (2001). The role of the binder phase in
the WC-Co sintering. Materials Research, 4, p.59-62.
168
Sivaprahasam, D., Chandrasekar, S. B. and Sundaresan, R. (2007). Microstructure
and mechanical properties of nanocrystalline WC–12Co consolidated by
spark plasma sintering. International Journal of Refractory Metals and Hard
Materials, 25, p.144-152.
Soares, E., Malheiros, L. F., Sacramento, J., Valente, M. A. and Oliveira, F. J. (2012).
Solid and Liquid Phase Sintering of Submicrometer Carbides with Different
Cobalt Contents. Journal of the American Ceramic Society, p.n/a-n/a.
Soleimanpour, A. M., Abachi, P. and Simchi, A. (2012). Microstructure and
mechanical properties of WC–10Co cemented carbide containing VC or (Ta,
Nb)C and fracture toughness evaluation using different models. International
Journal of Refractory Metals and Hard Materials, 31, p.141-146.
Sommer, M., Schubert, W.-D., Zobetz, E. and Warbichler, P. (2002). On the
formation of very large WC crystals during sintering of ultrafine WC–Co
alloys. International Journal of Refractory Metals and Hard Materials, 20,
p.41-50.
Spriggs, G. E. (1995). A history of fine grained hardmetal. International Journal of
Refractory Metals and Hard Materials, 13, p.241-255.
Srivatsan, T. S., Woods, R., Petraroli, M. and Sudarshan, T. S. (2002). An
investigation of the influence of powder particle size on microstructure and
hardness of bulk samples of tungsten carbide. Powder Technology, 122,
p.54-60.
Stremchemical. 2012. Tungsten-carbide (WC) [Online]. Available:
http://www.strem.com [Accessed].
Suetin, D. V., Shein, I. R. and Ivanovskii, A. L. (2009). Structural, electronic and
magnetic properties of η carbides (Fe3W3C, Fe6W6C, Co3W3C and Co6W6C)
169
from first principles calculations. Physica B: Condensed Matter, 404,
p.3544-3549.
Sun, L., Jia, C.-C. and Xian, M. (2007). A research on the grain growth of WC–Co
cemented carbide. International Journal of Refractory Metals and Hard
Materials, 25, p.121-124.
Suryanarayana, C. (2001). Mechanical alloying and milling. Progress in Materials
Science, 46, p.1-184.
Sutthiruangwong, S. and Mori, G. (2005). Influence of refractory metal carbide
addition on corrosion properties of cemented carbides. Materials and
Manufacturing Processes, 20, p.47-56.
Torres, C. S. and Shaeffer, L. (2010). Effect of high energy milling on the
microstruture and properties of WC-Ni composite. Materials Research, 13(3),
p.293-298.
Tracey, V. A. (1992). Nickel in hardmetals. International Journal of Refractory
Metals and Hard Materials, 11, p.137-149.
Tsuchida, T., Suzuki, K. and Naganuma, H. (2001). Low-temperature formation of
ternary carbide Fe3M3C (M=Mo, W) assisted by mechanical activation. Solid
State Ionics, 141–142, p.623-631.
Uhrenius, B., Pastor, H. and Pauty, E. (1997). On the composition of
Fe-Ni-Co-WC-based cemented carbides. International Journal of Refractory
Metals and Hard Materials, 15, p.139-149.
Upadhyaya, A., Sarathy, D. and Wagner, G. (2001). Advances in alloy design aspects
of cemented carbides. Materials & Design, 22, p.511-517.
Upadhyaya, G. S. 1998. Cemented Tungsten Carbides Production- Properties- and
Testing (Materials Science and Process Technology), Noyes Publications.
170
Upadhyaya, G. S. (2001). Materials science of cemented carbides — an overview.
Materials & Design, 22, p.483-489.
Upadhyaya, G. S. 2010. Sintered metallic and ceramic materials: Preparation,
properties and applications, Kanpur, John wiley & Sons, LTD.
Upadhyaya, G. S. and Bhaumik, S. K. (1988). Sintering of submicron
WC-10wt.%Co hard metals containing nickel and iron. Materials Science and
Engineering: A, 105–106, Part 1, p.249-256.
Vaßen, R. and Stöver, D. (1999). Processing and properties of nanophase ceramics.
Journal of Materials Processing Technology, 92–93, p.77-84.
Viswanadham, R. and Lindquist, P. (1987). Transformation-toughening in cemented
carbides: Part I. Binder composition control. Metallurgical and Materials
Transactions A, 18, p.2163-2173.
Voitovich, V. B., Sverdel, V. V., Voitovich, R. F. and Golovko, E. I. (1996). Oxidation
of WC-Co, WC-Ni and WC-Co-Ni hard metals in the temperature range
500–800 °C. International Journal of Refractory Metals and Hard Materials,
14, p.289-295.
Wang, B., Matsumaru, K., Yang, J., Fu, Z. and Ishizaki, K. (2012). The effect of cBN
additions on densification, microstructure and properties of WC–Co
composites by pulse electric current sintering. Journal of the American
Ceramic Society, p.n/a-n/a.
Wang, X., Fang, Z. Z. and Sohn, H. Y. (2008). Grain growth during the early stage of
sintering of nanosized WC–Co powder. International Journal of Refractory
Metals and Hard Materials, 26, p.232-241.
Wang, Y., Heusch, M., Lay, S. and Allibert, C. H. (2002). Microstructure evolution in
the cemented carbides WC–Co I. Effect of the C/W ratio on the morphology
171
and defects of the WC grains. physica status solidi (a), 193, p.271-283.
Webelements. 2012. Carbon: allotropes [Online]. Available:
http://www.webelements.com/carbon/allotropes.html [Accessed].
Weidow, J., Norgren, S. and Andrén, H.-O. (2009a). Effect of V, Cr and Mn additions
on the microstructure of WC–Co. International Journal of Refractory Metals
and Hard Materials, 27, p.817-822.
Weidow, J., Zackrisson, J., Jansson, B. and Andrén, H.-O. (2009b). Characterisation
of WC-Co with cubic carbide additions. International Journal of Refractory
Metals and Hard Materials, 27, p.244-248.
Wittmann, B., Schubert, W.-D. and Lux, B. (2002). WC grain growth and grain
growth inhibition in nickel and iron binder hardmetals. International Journal
of Refractory Metals and Hard Materials, 20, p.51-60.
Xi, X., Pi, X., Nie, Z., Song, S., Xu, X. and Zuo, T. (2009). Synthesis and
characterization of ultrafine WC–Co by freeze-drying and spark plasma
sintering. International Journal of Refractory Metals and Hard Materials, 27,
p.101-104.
Xiao, D.-h., He, Y.-h., Luo, W.-h. and Song, M. (2009). Effect of VC and NbC
additions on microstructure and properties of ultrafine WC-10Co cemented
carbides. Transactions of Nonferrous Metals Society of China, 19,
p.1520-1525.
Xinghong, Z., Chuncheng, Z., Wei, Q., Xiaodong, H. and Kvanin, V. L. (2002).
Self-propagating high temperature combustion synthesis of TiC/TiB2
ceramic–matrix composites. Composites Science and Technology, 62,
p.2037-2041.
Xu, Q., Zhang, X., Han, J., He, X. and Kvanin, V. L. (2003). Combustion synthesis
172
and densification of titanium diboride–copper matrix composite. Materials
Letters, 57, p.4439-4444.
Xueming, M. A., Gang, J. I., Ling, Z. and Yuanda, D. (1998). Structure and
properties of bulk nano-structured WC–CO alloy by mechanical alloying.
Journal of Alloys and Compounds, 264, p.267-270.
Yao, Z., Stiglich, J. J. and Sudarshan, T. S. (1998). WC-Co enjoys proud history and
bright future. Metal powder report, J-53A, p.5.
Yao, Z., Stiglich, J. J. and Sudarshan, T. S. (1999). Nano-grained tungsten
carbide-cobalt (WC/Co). Materials Modification p.1-27.
Zawrah, M. F. (2007). Synthesis and characterization of WC–Co nanocomposites by
novel chemical method. Ceramics International, 33, p.155-161.
Zhang, D., Massalski, T. and Paruchuri, M. (1994). Formation of metastable and
equilibrium phases during mechanical alloying of Al and Mg powders.
Metallurgical and Materials Transactions A, 25, p.73-79.
Zhang, F., Shen, J. and Sun, J. (2004a). The effect of phosphorus additions on
densification, grain growth and properties of nanocrystalline WC–Co
composites. Journal of Alloys and Compounds, 385, p.96-103.
Zhang, F., Shen, J. and Sun, J. (2004b). Processing and properties of carbon
nanotubes-nano-WC-Co composites. Materials Science and Engineering: A,
381, p.86-91.
Zhang, L., Chen, S., Wang, Y.-j., Yu, X.-w. and Xiong, X.-j. (2008). Tungsten carbide
platelet-containing cemented carbide with yttrium containing dispersed phase.
Transactions of Nonferrous Metals Society of China, 18, p.104-108.
Zhang, L. and Sun, B. (1996). A new hardmetal for mining with Ni-Co binder.
International Journal of Refractory Metals and Hard Materials, 14,
173
p.245-248.
Zhang, W., Liu, Y., Liu, B., Li, H.-z. and Tang, B. (2010). Deformability and
microstructure transformation of PM TiAl alloy prepared by pseudo-HIP
technology. Transactions of Nonferrous Metals Society of China, 20,
p.547-552.
Zhao, S., Song, X., Wang, M., Liu, G. and Zhang, J. (2009). Preparation of ultrafine
WC-Co cermets by combining pretreatment and consolidation with spark
plasma sintering. Rare Metals, 28, p.391-395.
174
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