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Microstructure and mechanical properties of ultrafine WC–Ni–VC–TaC–cBN cemented carbides fabricated by spark plasma sintering
Int. Journal of Refractory Metals and Hard Materials 29 (2011) 733–738
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Int. Journal of Refractory Metals and Hard Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / I J R M H M
Microstructure and mechanical properties of ultrafine WC–Ni–VC–TaC–cBN cemented carbides fabricated by spark plasma sintering Huiyong Rong a, Zhijian Peng a,⁎, Xiaoyong Ren a, Chengbiao Wang a, Zhiqiang Fu a, Longhao Qi b, Hezhuo Miao b a b
School of Engineering and Technology, China University of Geosciences, Beijing 100083, PR China State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, PR China
a r t i c l e
i n f o
Article history: Received 28 November 2010 Accepted 16 June 2011 Keywords: Spark plasma sintering Cubic BN WC–Ni cemented carbides Mechanical properties
a b s t r a c t Ultrafine WC–Ni–VC–TaC cemented carbides with different amounts of cubic boron nitride (cBN) were fabricated by spark plasma sintering, and the microstructure and mechanical properties of the as-prepared cermets were investigated. Scanning electron microscopy observations showed that the size of WC grains in the cermet samples was 0.2–0.4 μm. After the addition of cBN, the samples were still quite dense with the highest relative density of almost 98% when the addition fraction of cBN was 50 vol.%, although some micropores might exist in the samples. X-ray diffraction results indicated that no phase transformation of cBN was detected. The relative density and hardness of the cemented carbides increased with the addition fraction of cBN, but their strength decreased. When the fraction of cBN increased from 0 up to 50 vol.%, the hardness of the samples increased from 2100 to 3200 HV, but the flexural strength decreased from 1950 to 1250 MPa. © 2011 Elsevier Ltd. All rights reserved.
1. Introduction Tungsten carbide cermets utilizing cobalt as the metal binder phase have excellent properties, such as high hardness, high hothardness, high toughness and strength, and good wear resistance [1]; thus they have been widely applied in many fields including cutting tools and drilling equipment. However, because metal cobalt is of high price [2], and the corrosion and oxidation resistance of cemented carbides with cobalt as the binder phase is low [3], their applications have been limited. Additionally, in the case of the chemical vapor deposition of diamond films onto WC–Co cemented carbides, the cobalt in the cemented carbide surface layer is detrimental to the diamond nucleation from the gas phase, and promotes the formation of interfacial graphite which dramatically reduces the adhesion strength between diamond films and substrates [4,5]. Thus scientists and engineers have endeavored for years to find new metals to replace cobalt in cemented carbides. To date, several metals have been found to be possible substitutes for cobalt as metal binder phase in WC cemented carbides [6–8]. Among all those metals investigated, nickel is an exciting and promising candidate as a binder phase metal, not only because of its relatively lower price than cobalt, but also due to its much better performance in corrosion and oxidation resistance [3]
⁎ Corresponding author at: School of Engineering and Technology, China University of Geosciences at Beijing, Beijing 100083, PR China. Tel.: + 86 10 82320255; fax: + 86 10 82322624. E-mail addresses: [email protected] (Z. Peng), [email protected] (H. Rong). 0263-4368/$ – see front matter © 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2011.06.004
than those of cobalt. However, the mechanical properties (hardness and strength) of WC–Ni cemented carbides are relatively inferior to those of WC–Co cemented carbides [9]. To overcome the shortcoming of WC–Ni cemented carbides, two ways have been proposed in previous work [2–20] to improve their performance. One is to try to decrease the size of WC grains in the cermets during sintering, i.e. fabricating sub-micrometer or nearnanometer and even nanometer WC–Ni cemented carbides [2–19]. Another is to add some materials with high hardness/strength into the matrix of WC–Ni cemented carbides, increasing directly the hardness/strength of the cermets [20]. In order to fabricate ultrafine WC–Ni cemented carbides, some new rapid sintering methods to reduce the growth of WC grains were proposed [2,10], and the results showed that the hardness of the specimens was significantly higher than those of WC–Co and WC–Ni cemented carbides sintered with conventional sintering methods, and their fracture toughnesses were not reduced much. The other way of inhibiting the growth of WC grains during sintering is that, as in WC–Co cemented carbides [11], WC grain growth inhibitors must be chosen and utilized [12–19]. Vanadium carbide, VC, has been proven the most effective WC grain growth inhibitor in WC–Ni cemented carbides [13,14], followed in sequence by TaC, Cr3C2, TiC and ZrC. Correa et al. [20] presented a different but effective way to improve the hardness and strength of WC–Ni cemented carbides by adding SiC powder into WC–Ni cemented carbides. The obtained WC–10 wt.% Ni–Si cemented carbide presented similar Vickers hardness to its conventional WC–Co cemented carbide counterpart, and their flexural strength was improved.
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The research works to date [2–20] encourage the combination of different methods to improve the hardness and strength of WC–Ni cemented carbides by properly utilizing the rapid speed sintering method, adding WC grain growth inhibitors, and using harder/ tougher materials. Cubic boron nitride (cBN) is the secondary hardest material (only inferior to diamond) in the world [21], has good chemical inertness with respect to ferrous alloys at high temperatures up to 1000 °C, and outperforms diamond, especially in many harsh service conditions [22]. Moreover, it has good oxidation resistance and thermal stability [22], which is very important during sintering at high temperature and in application. Thus, it has potential if it could be sintered in WC–Ni cemented carbides, and its much lower price than diamond is also favorable. So, the aim of this work is to fabricate ultrafine WC–Ni–cBN cemented carbides by spark plasma sintering (SPS) and by utilizing VC and TaC as WC grain inhibitors, and to investigate the effect of cBN addition on the microstructure and mechanical properties of WC–Ni–VC–TaC–cBN cemented carbides. 2. Experimental procedure 2.1. Sample preparation The raw material powders used in the present study include tungsten carbide (WC), hydroxyl-nickel (Ni), vanadium carbide (VC), tantalum carbide (TaC) and cubic boron nitride (cBN) powders, which were all commercially bought. The basic physical and chemical parameters of the applied WC powder (the main recipe) are given in Table 1. The particle sizes of the Ni, VC, TaC and cBN powders were about 0.7, 2–4, 1–1.5 and 20–40 μm, respectively, and they were all of industrial reagent grade. In order to avoid liquid phase attack to cBN particles [23], and improve the pullout strength of cBN particles from the sintered bodies of cemented carbide matrix, the cBN particles were coated with about 500-nm-thickness nickel metal thin film using a physical vapor deposition method before use. The balls for attrition milling were made from cemented carbide YG6 (ISO: K20) and their diameters were about 5 mm. Table 2 lists the nominal compositions of the designed samples. Because Ni has a face-centered cubic crystal structure, it is susceptible to particle deformation and agglomeration during attrition milling, which encourages the formation of pores during sintering [13], and is thus detrimental for improving the relative density of the sintered samples. Additionally, cBN particles are easily broken during milling, especially in long-time high-energy milling. Therefore, in this study, the raw powders were designedly mixed together at different milling stages and milled for disparate times in absolute alcohol with a high energy attrition mill (Model: SY-1, China). The powder mixture of WC, VC and TaC was milled for 44 h in the first stage at a speed of 300 rpm; after that, Ni and cBN powders were added into the mixture and milled for another 4 h at a speed of 80 rpm. For the milling, the mass ratio of ball-to-powder was 10:1 and that of powder-to-absolute alcohol was 3:1. After milling, the powder slurries were dried in a vacuum oven at 35 °C. The dried powder mixture was sintered by SPS (SPS-1050 T, Japan). For sintering, the initial pressure applied on the graphite die was 30 MPa, the heating rate was 200 °C/min, and the holding time was 6 min at the highest temperature (1350 °C) with a pressure of 50 MPa. After sintering, the cooling rate was the same as that of heating. The sintered specimens were cylinders, and Table 1 Basic physical and chemical parameters of WC powder used in the present study. Grade
Distribution by turbid meter (%)
Chemical (%)
0–1 μ 1–2 μ
Total carbon Free carbon Combined carbon Oxygen
GWC002 97.5
2.5
6.11
0.04
6.07
0.31
Table 2 Nominal compositions of the designed samples. Sample no.
Sp Sp Sp Sp Sp Sp
1 2 3 4 5 6
WC
Ni
VC
TaC
cBN
(wt.%)
(wt.%)
(wt.%)
(wt.%)
(vol.%)
(wt.%)
91.0 87.9 84.8 81.7 78.6 75.5
8.0 8.0 8.0 8.0 8.0 8.0
0.7 0.7 0.7 0.7 0.7 0.7
0.3 0.3 0.3 0.3 0.3 0.3
0.0 10.0 20.0 30.0 40.0 50.0
0.0 3.1 6.2 9.3 12.4 15.5
their dimensions were approximately 20 mm in diameter and 5 mm in height (Φ20 mm × 5 mm). 2.2. Material characterization Before measurement, the top and bottom surfaces of each sintered specimen were ground with a diamond wheel with particle size of 60 μm into a dimension of Φ20 mm × 3 mm. The samples were examined to identify their phase composition by X-ray diffraction (XRD, D/max2550HB+/PC, Cu Kα and λ = 1.5418 Å) through a continuous scanning mode with a speed of 5°/min. After XRD examination, each specimen was cut into testing bars of 2 mm × 3 mm × 10 mm, and then the bars were embedded in resin, ground, and polished to a mirror finish with seven-grade diamond pastes. The particle sizes of the abrasives in the sequentially applied pastes were 40, 28, 14, 7, 3.5, 1 and the last 0.5 μm, respectively. The polished bars were used to measure flexural strength with threepoint bending experiment on an AG-IC 20 kN Shimazu tester. During the testing, the applied load rate was 0.5 mm/min. After that, the apparent density of the samples was measured with Archimedes method according to international standard (ISO18754) with the resultant broken bars, and the relative density was defined as the percentage of the apparent density to their corresponding theoretical density. Then the microstructure of the specimens was examined by scanning electron microscope (SEM) on fresh fractured (with LEO 1530 SEM), polished (with SSX-550 SEM), and polished/etched (with LEO 1530 SEM) surfaces, respectively. The etching of samples was carried out in a Murakami's reagent (1 g potassium ferricyanide, 2 g potassium hydroxide, and 30 g water) for about 2 min at room temperature. With the polished specimen, bulk Vickers hardness (HV) of the samples was measured by a LECO DM-400 hardness tester with an indenting load of 1 kgf and a dwell time of 20 s. The binder mean free path, which was defined as the average thickness of the binder phase [24], and WC grain size [25] of the as-prepared samples were calculated by the linear intercept method from the obtained SEM images with a Lince PC software. 3. Results and discussion 3.1. Composition and microstructure The phase compositions of the sintered samples were derived by XRD. Typical XRD patterns of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples are shown in Fig. 1. WC, Ni and carbon phases were detected in all the sintered samples. Cubic BN phases were detected in the samples with the addition fraction of cBN more than 20 vol.%. The XRD results also revealed that no cBN phase transformation happened during sintering, meaning that the excellent properties of cBN can be potentially kept. Thus sintering by SPS should promote the increase in hardness of WC–Ni–VC–TaC cemented carbides. Typical SEM images are shown in Figs. 2–4. From the images on the polished/etched surface of typical sample as shown in Fig. 2, the mean size of WC grains was calculated about 0.312 ± 0.095 μm, indicating
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735
a
Micropores
Fig. 1. XRD patterns of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples, showing the volume% of cBN.
that ultrafine WC–Ni cemented carbides were fabricated, which is helpful to obtain materials of high hardness and strength. The ultrafine WC–Ni cemented carbides could be attributed to the use of VC and TaC as WC grain inhibitors, and the short reaction time during SPS. From Fig. 2 it could be also conducted that the binder mean free path in the obtained WC–Ni cemented carbides was very short (only about 22 nm) and the sample was of a much fine dispersion, which would result in high performance in hardness for the obtained sample [24,26]. Typical micrographs on the fracture surface of the as-prepared WC–Ni cemented carbides are presented in Fig. 3. It reveals that the fracture mode could be the typical brittle fracture in ceramic materials, because the characteristics of cleavage and dimple could be obviously observed on the fracture surface of the samples in the figure, and both intergranular fracture and transgranular fracture of WC grains and transgranular fracture of cBN particles occur in the samples. This result implies that the fracture resistance of the obtained WC–Ni cemented carbides during flexural test stemmed from the intrinsic strengths of the binder phase, WC grains and cBN particles. Therefore, with the increase of cBN addition fraction, the mechanical strength of the obtained WC–Ni cemented carbides would reasonably decrease because of the low strength of cBN particles. However, as can be seen from Fig. 3a and b, the dimples on the
b
cBN
cBN
c
Matrix
cBN
Fig. 3. Typical secondary electron SEM images on the fracture surface of the as-prepared WC–Ni cemented carbide samples: (a) without cBN, (b) with cBN, and (c) showing strong bonding force between the cBN particles and cemented carbide matrix. From this figure, the fracture mode could be observed as brittle fracture.
WC grains Binder
Micropores
Fig. 2. Typical secondary electron SEM image of the WC–Ni–VC–TaC–cBN cemented carbide samples after etching. Etching: in Murakami's reagent for about 2 min.
fracture surface of the WC–Ni–cBN cemented carbide sample are deeper than those in WC–Ni sample without cBN, and so the decrease in mechanical strength of the WC–Ni–cBN could be compensated to a certain extent after the addition of cBN particles into WC–Ni cemented carbides. The microstructure of the as-prepared WC–Ni–10 vol.% cBN was also compared with those of samples sintered by other rapid sintering techniques in previously reported works. The results are listed in Table 3. From this table it could be concluded that the relative density of the as-prepared WC–Ni-10 vol.% cBN is much lower than their counterparts of WC–Co cemented carbides in literature, indicating that there might be some pores in the sample. This is consistent with the observations from Figs. 2 to 4. The formation of the micropores
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a
b Micropores
Matrix Micropores
10 µm
cBN
10 µm
c
d Matrix
cBN
cBN
Matrix
10 µm
10 µm
Fig. 4. Typical secondary electron SEM images on the polished surface of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples with different addition fractions of cBN: (a) 0 vol.%, (b) 10 vol.%, (c) 30 vol.% and (d) 50 vol.%.
Ref. [27] [27] [28] This work
WC-based cemented carbides
Sintering technique
Relative density (%)
Grain size (nm)
HV (GPa)
8Co 8Co 8Co 8Ni-10 vol.% cBN
HFIHS PCAS PCAS SPS
99.2 99.2 98.4 91.9
410 440 380 320
19.5 19.2 18.8 22.6
13.5 100
Apparent density Relative density
13.0
Apparent density (g/cm3)
Table 3 Comparison on physical properties between the obtained WC–Ni-10 vol.% cBN sample in this study and the ones sintered by high frequency induction heated sintering (HFIFS) and pulsed current activated sintering (PCAS) reported in previous work.
However, with increasing cBN, the size and number of micropores decreased, and when the addition faction of cBN reached 50 vol.%, the pores almost disappeared (as can be seen from Fig. 4), meaning that as the fraction of cBN increased, the densification of the samples would increase, and this was confirmed by Fig. 5. As the addition fraction of cBN increased, although the apparent density of the obtained samples decreased due to the low density of the additive cBN, the relative density of the obtained samples still increased, indicating that the samples become denser with increased cBN. The maximum relative
98
12.5
96
12.0
94
11.5
92
11.0
90
10.5
88
10.0
86
9.5
84
9.0
0
10
20
30
40
50
Relative density (%)
can be explained as follows. Firstly, the adopted binder metal Ni has a relatively low wet ability in the cermets during sintering in comparison with Co, which would be harmful for the spread of the binder phase during sintering [13]. Secondly, although Ni was mixed and milled at the second stage at a lower speed and for a shorter time, particle deformation and agglomeration could still happen, which promoted the formation of pores during sintering [13]. Thirdly, since the as-prepared cemented carbides were fabricated by powder metallurgy, liquid phase was formed during heating and the bond of particles relied on the formation of this liquid. Thus, the formation and flow of the liquid phase are key factors which influence the microstructure of WC-based cemented carbides. However, because the formed liquid was not sufficient at the relatively low sintering temperature and the flow of liquid was not enough during the rapid sintering by SPS, pores would inevitably be formed. Moreover, the use of VC and TaC as WC grain inhibitors in WC–Ni–VC–TaC–cBN cemented carbides would cause a fine microstructure, and this fine microstructure would also be against the flow of liquid with aggravation of forming pores during sintering [13].
82
Adding fraction of cBN (vol%) Fig. 5. Apparent density and relative density of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples as a function of the addition fraction of cBN.
H. Rong et al. / Int. Journal of Refractory Metals and Hard Materials 29 (2011) 733–738
density of about 98% was reached when 50 vol.% cBN was added in this study. The reason for this phenomenon is possibly that the capillary forces are proportional to the surface tension divided by the particle size, and a micrometer-sized particle would result in an internal pressure equivalent to 10 atmospheres of external pressure during sintering [23]. Therefore, for the added micrometer-sized cBN powder, the internal pressure would increase obviously and keep increasing with more cBN added. The increased internal pressures would impel the micropores to shrink. Thus the micropores gradually disappeared in the samples and the structure became denser after the addition of cBN. This could also explain why the relative density of the samples presented a sharp increase when 10 to 20 vol.% cBN was added compared to the WC–Ni cemented carbide without cBN. And the density increase of the as-prepared WC–Ni cemented carbides would promote the mechanical properties, such as hardness and flexural strength. Moreover, it is worthwhile mentioning that the WC grain size of WC–Ni–cBN samples might be smaller than that of WC–Ni sample without cBN as can be seen from Fig. 3a and b, and this might be the result of the above-mentioned pressure, under which the growth of WC grain during sintering would be inhibited. Moreover, transgranular fracture of cBN and the tear of the cemented carbide matrix could be clearly observed from Fig. 3c, indicating that the binding force of the interface between the cBN particle and cemented carbide matrix was very high, and the pullout strength of cBN from the matrix was also much improved. Otherwise, the debonding of cBN from the materials during fracture would happen at the interface between the cBN particle and cemented carbide matrix.
The Vickers hardness of the as-prepared cemented carbides as a function of the addition fraction of cBN is shown in Fig. 6. From this figure, it can be seen that the Vickers hardness keeps increasing with increasing addition fraction of cBN. In order to compare the hardness of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide with their WC–Co cemented carbide counterparts prepared by other rapid sintering technique, the hardness of WC–Ni–VC–TaC-10 vol.% cBN was also measured by a VH-5 hardness tester (Everone) with a comparable indenting load of 30 kgf and a dwell time of 15 s with literatures as recommended by ISO 3878, and the result is listed in Table 3. From the table, it can be concluded that the hardness of the as-prepared WC–Ni–VC–TaC-10 vol.% cBN cemented carbide is comparably higher than their counterparts of WC–Co cemented carbides sintered by other rapid sintering technique [27,28], due to its
2000
1800
3000
1700
2800
1600
2600
1500 2400 1400
Vickers Hardness
Flexural Strength (MPa)
3200
Flexural Strength Vickers Hardness
2200
1300
2000
1200 0
10
finer WC grains, the addition of much harder material of cBN and very short binder mean free path [24]. The flexural strength of the cemented carbides as a function of the addition fraction of cBN is also shown in Fig. 6. It can be concluded from this figure that, with increasing addition fraction of cBN, the flexural strength of the cemented carbides decreased to about 1250 MPa with the addition fraction of cBN more than 40 vol.%. It is easily understood that the hardness of the cemented carbides increases with increasing addition fraction of cBN. One factor for the hardness increase is attributed to the strong binding force on the interface between the Ni-coated cBN particle and cemented carbide matrix (as can be seen from Fig. 3c). Another factor is due to no phase transformation of cBN during sintering (as can be seen from Fig. 1). In addition, the increase in density of the samples after cBN was added would be surely helpful to improve the hardness, and this explained why the hardness of the samples dramatically increased after 10 to 20 vol.% of cBN particles was added. And with more cBN added, the hardness of the cemented carbides increased, while the densification of the samples also increased. The reason for the decrease of flexural strength is attributed to the transgranular fracture mode of cBN particles and the weak strength of cBN. Thus the more cBN added, the lower flexural strength obtained. However, the improvement of relative density with increasing addition fraction of cBN, and the deeper dimples in WC–Ni–cBN would be helpful to improve the flexural strength and compensate the decrease of flexural strength to a certain extent. Thus, the flexural strength of samples only decreased from 1950 to 1250 MPa with the addition fraction of cBN increased from 0 to 50 vol.%. 4. Conclusions
3.2. Mechanical properties
1900
737
20
30
40
50
Adding fraction of cBN (vol%) Fig. 6. Vickers hardness and flexural strength of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples as a function of the addition fraction of cBN.
Ultrafine WC–Ni–cBN cemented carbides were successfully fabricated by spark plasma sintering using VC and TaC as WC grain inhibitors. During sintering no cBN phase transformation happened. The Ni-coated cBN particles had a strong binding force with the cemented carbide matrix in the sintered body. The relative density of WC–Ni–VC–TaC–cBN cemented carbides increased with the addition fraction of cBN, and when the addition fraction of cBN reached 50 vol.%, the structure of WC–Ni–VC–TaC–cBN cemented carbide was almost fully dense. The flexural strength of WC–Ni–VC–TaC–cBN cemented carbides decreased from 1950 to 1250 MPa with the addition fraction of cBN increased from 0 to 50 vol.%, while the hardness of the samples increased from 2400 to 3150 HV. Acknowledgments The work was supported by Grand Survey on Land and Nature Sources of China sponsored by China Geological Survey (grant no. 1212010916026), Ph.D. Programs Foundation by Ministry of Education of China (grant no. 20100022110002), State Key Laboratory of New Ceramic and Fine Processing in Tsinghua University (grant no. KF0903), the Cultivating Foundation for Young Scientists in China University of Geosciences at Beijing from the Fundamental Research Funds for the Central Universities (Grant No. 2011PY0192) and National Lab on Scientific Drilling in China University of Geosciences at Beijing (grant no. NLSD200801). References [1] Gant AJ, Gee MG, Roebuck B. Rotating wheel abrasion of WC/Co hardmetals. Wear 2005;258:178–88. [2] Kim HC, Shon IJ, Yoon JK, Doh JM, Munir ZA. Rapid sintering of ultrafine WC–Ni cermets. Int J Refract Met Hard Mater 2006;24:427–31. [3] Imasato S, Tokumoto K, Kitada T, Sakaguchi S. Properties of ultrafine grain binderless cemented carbide RCCFN. Int J Refract Met Hard Mater 1995;13: 305–12.
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Int. Journal of Refractory Metals and Hard Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / I J R M H M
Microstructure and mechanical properties of ultrafine WC–Ni–VC–TaC–cBN cemented carbides fabricated by spark plasma sintering Huiyong Rong a, Zhijian Peng a,⁎, Xiaoyong Ren a, Chengbiao Wang a, Zhiqiang Fu a, Longhao Qi b, Hezhuo Miao b a b
School of Engineering and Technology, China University of Geosciences, Beijing 100083, PR China State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, PR China
a r t i c l e
i n f o
Article history: Received 28 November 2010 Accepted 16 June 2011 Keywords: Spark plasma sintering Cubic BN WC–Ni cemented carbides Mechanical properties
a b s t r a c t Ultrafine WC–Ni–VC–TaC cemented carbides with different amounts of cubic boron nitride (cBN) were fabricated by spark plasma sintering, and the microstructure and mechanical properties of the as-prepared cermets were investigated. Scanning electron microscopy observations showed that the size of WC grains in the cermet samples was 0.2–0.4 μm. After the addition of cBN, the samples were still quite dense with the highest relative density of almost 98% when the addition fraction of cBN was 50 vol.%, although some micropores might exist in the samples. X-ray diffraction results indicated that no phase transformation of cBN was detected. The relative density and hardness of the cemented carbides increased with the addition fraction of cBN, but their strength decreased. When the fraction of cBN increased from 0 up to 50 vol.%, the hardness of the samples increased from 2100 to 3200 HV, but the flexural strength decreased from 1950 to 1250 MPa. © 2011 Elsevier Ltd. All rights reserved.
1. Introduction Tungsten carbide cermets utilizing cobalt as the metal binder phase have excellent properties, such as high hardness, high hothardness, high toughness and strength, and good wear resistance [1]; thus they have been widely applied in many fields including cutting tools and drilling equipment. However, because metal cobalt is of high price [2], and the corrosion and oxidation resistance of cemented carbides with cobalt as the binder phase is low [3], their applications have been limited. Additionally, in the case of the chemical vapor deposition of diamond films onto WC–Co cemented carbides, the cobalt in the cemented carbide surface layer is detrimental to the diamond nucleation from the gas phase, and promotes the formation of interfacial graphite which dramatically reduces the adhesion strength between diamond films and substrates [4,5]. Thus scientists and engineers have endeavored for years to find new metals to replace cobalt in cemented carbides. To date, several metals have been found to be possible substitutes for cobalt as metal binder phase in WC cemented carbides [6–8]. Among all those metals investigated, nickel is an exciting and promising candidate as a binder phase metal, not only because of its relatively lower price than cobalt, but also due to its much better performance in corrosion and oxidation resistance [3]
⁎ Corresponding author at: School of Engineering and Technology, China University of Geosciences at Beijing, Beijing 100083, PR China. Tel.: + 86 10 82320255; fax: + 86 10 82322624. E-mail addresses: [email protected] (Z. Peng), [email protected] (H. Rong). 0263-4368/$ – see front matter © 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2011.06.004
than those of cobalt. However, the mechanical properties (hardness and strength) of WC–Ni cemented carbides are relatively inferior to those of WC–Co cemented carbides [9]. To overcome the shortcoming of WC–Ni cemented carbides, two ways have been proposed in previous work [2–20] to improve their performance. One is to try to decrease the size of WC grains in the cermets during sintering, i.e. fabricating sub-micrometer or nearnanometer and even nanometer WC–Ni cemented carbides [2–19]. Another is to add some materials with high hardness/strength into the matrix of WC–Ni cemented carbides, increasing directly the hardness/strength of the cermets [20]. In order to fabricate ultrafine WC–Ni cemented carbides, some new rapid sintering methods to reduce the growth of WC grains were proposed [2,10], and the results showed that the hardness of the specimens was significantly higher than those of WC–Co and WC–Ni cemented carbides sintered with conventional sintering methods, and their fracture toughnesses were not reduced much. The other way of inhibiting the growth of WC grains during sintering is that, as in WC–Co cemented carbides [11], WC grain growth inhibitors must be chosen and utilized [12–19]. Vanadium carbide, VC, has been proven the most effective WC grain growth inhibitor in WC–Ni cemented carbides [13,14], followed in sequence by TaC, Cr3C2, TiC and ZrC. Correa et al. [20] presented a different but effective way to improve the hardness and strength of WC–Ni cemented carbides by adding SiC powder into WC–Ni cemented carbides. The obtained WC–10 wt.% Ni–Si cemented carbide presented similar Vickers hardness to its conventional WC–Co cemented carbide counterpart, and their flexural strength was improved.
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The research works to date [2–20] encourage the combination of different methods to improve the hardness and strength of WC–Ni cemented carbides by properly utilizing the rapid speed sintering method, adding WC grain growth inhibitors, and using harder/ tougher materials. Cubic boron nitride (cBN) is the secondary hardest material (only inferior to diamond) in the world [21], has good chemical inertness with respect to ferrous alloys at high temperatures up to 1000 °C, and outperforms diamond, especially in many harsh service conditions [22]. Moreover, it has good oxidation resistance and thermal stability [22], which is very important during sintering at high temperature and in application. Thus, it has potential if it could be sintered in WC–Ni cemented carbides, and its much lower price than diamond is also favorable. So, the aim of this work is to fabricate ultrafine WC–Ni–cBN cemented carbides by spark plasma sintering (SPS) and by utilizing VC and TaC as WC grain inhibitors, and to investigate the effect of cBN addition on the microstructure and mechanical properties of WC–Ni–VC–TaC–cBN cemented carbides. 2. Experimental procedure 2.1. Sample preparation The raw material powders used in the present study include tungsten carbide (WC), hydroxyl-nickel (Ni), vanadium carbide (VC), tantalum carbide (TaC) and cubic boron nitride (cBN) powders, which were all commercially bought. The basic physical and chemical parameters of the applied WC powder (the main recipe) are given in Table 1. The particle sizes of the Ni, VC, TaC and cBN powders were about 0.7, 2–4, 1–1.5 and 20–40 μm, respectively, and they were all of industrial reagent grade. In order to avoid liquid phase attack to cBN particles [23], and improve the pullout strength of cBN particles from the sintered bodies of cemented carbide matrix, the cBN particles were coated with about 500-nm-thickness nickel metal thin film using a physical vapor deposition method before use. The balls for attrition milling were made from cemented carbide YG6 (ISO: K20) and their diameters were about 5 mm. Table 2 lists the nominal compositions of the designed samples. Because Ni has a face-centered cubic crystal structure, it is susceptible to particle deformation and agglomeration during attrition milling, which encourages the formation of pores during sintering [13], and is thus detrimental for improving the relative density of the sintered samples. Additionally, cBN particles are easily broken during milling, especially in long-time high-energy milling. Therefore, in this study, the raw powders were designedly mixed together at different milling stages and milled for disparate times in absolute alcohol with a high energy attrition mill (Model: SY-1, China). The powder mixture of WC, VC and TaC was milled for 44 h in the first stage at a speed of 300 rpm; after that, Ni and cBN powders were added into the mixture and milled for another 4 h at a speed of 80 rpm. For the milling, the mass ratio of ball-to-powder was 10:1 and that of powder-to-absolute alcohol was 3:1. After milling, the powder slurries were dried in a vacuum oven at 35 °C. The dried powder mixture was sintered by SPS (SPS-1050 T, Japan). For sintering, the initial pressure applied on the graphite die was 30 MPa, the heating rate was 200 °C/min, and the holding time was 6 min at the highest temperature (1350 °C) with a pressure of 50 MPa. After sintering, the cooling rate was the same as that of heating. The sintered specimens were cylinders, and Table 1 Basic physical and chemical parameters of WC powder used in the present study. Grade
Distribution by turbid meter (%)
Chemical (%)
0–1 μ 1–2 μ
Total carbon Free carbon Combined carbon Oxygen
GWC002 97.5
2.5
6.11
0.04
6.07
0.31
Table 2 Nominal compositions of the designed samples. Sample no.
Sp Sp Sp Sp Sp Sp
1 2 3 4 5 6
WC
Ni
VC
TaC
cBN
(wt.%)
(wt.%)
(wt.%)
(wt.%)
(vol.%)
(wt.%)
91.0 87.9 84.8 81.7 78.6 75.5
8.0 8.0 8.0 8.0 8.0 8.0
0.7 0.7 0.7 0.7 0.7 0.7
0.3 0.3 0.3 0.3 0.3 0.3
0.0 10.0 20.0 30.0 40.0 50.0
0.0 3.1 6.2 9.3 12.4 15.5
their dimensions were approximately 20 mm in diameter and 5 mm in height (Φ20 mm × 5 mm). 2.2. Material characterization Before measurement, the top and bottom surfaces of each sintered specimen were ground with a diamond wheel with particle size of 60 μm into a dimension of Φ20 mm × 3 mm. The samples were examined to identify their phase composition by X-ray diffraction (XRD, D/max2550HB+/PC, Cu Kα and λ = 1.5418 Å) through a continuous scanning mode with a speed of 5°/min. After XRD examination, each specimen was cut into testing bars of 2 mm × 3 mm × 10 mm, and then the bars were embedded in resin, ground, and polished to a mirror finish with seven-grade diamond pastes. The particle sizes of the abrasives in the sequentially applied pastes were 40, 28, 14, 7, 3.5, 1 and the last 0.5 μm, respectively. The polished bars were used to measure flexural strength with threepoint bending experiment on an AG-IC 20 kN Shimazu tester. During the testing, the applied load rate was 0.5 mm/min. After that, the apparent density of the samples was measured with Archimedes method according to international standard (ISO18754) with the resultant broken bars, and the relative density was defined as the percentage of the apparent density to their corresponding theoretical density. Then the microstructure of the specimens was examined by scanning electron microscope (SEM) on fresh fractured (with LEO 1530 SEM), polished (with SSX-550 SEM), and polished/etched (with LEO 1530 SEM) surfaces, respectively. The etching of samples was carried out in a Murakami's reagent (1 g potassium ferricyanide, 2 g potassium hydroxide, and 30 g water) for about 2 min at room temperature. With the polished specimen, bulk Vickers hardness (HV) of the samples was measured by a LECO DM-400 hardness tester with an indenting load of 1 kgf and a dwell time of 20 s. The binder mean free path, which was defined as the average thickness of the binder phase [24], and WC grain size [25] of the as-prepared samples were calculated by the linear intercept method from the obtained SEM images with a Lince PC software. 3. Results and discussion 3.1. Composition and microstructure The phase compositions of the sintered samples were derived by XRD. Typical XRD patterns of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples are shown in Fig. 1. WC, Ni and carbon phases were detected in all the sintered samples. Cubic BN phases were detected in the samples with the addition fraction of cBN more than 20 vol.%. The XRD results also revealed that no cBN phase transformation happened during sintering, meaning that the excellent properties of cBN can be potentially kept. Thus sintering by SPS should promote the increase in hardness of WC–Ni–VC–TaC cemented carbides. Typical SEM images are shown in Figs. 2–4. From the images on the polished/etched surface of typical sample as shown in Fig. 2, the mean size of WC grains was calculated about 0.312 ± 0.095 μm, indicating
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735
a
Micropores
Fig. 1. XRD patterns of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples, showing the volume% of cBN.
that ultrafine WC–Ni cemented carbides were fabricated, which is helpful to obtain materials of high hardness and strength. The ultrafine WC–Ni cemented carbides could be attributed to the use of VC and TaC as WC grain inhibitors, and the short reaction time during SPS. From Fig. 2 it could be also conducted that the binder mean free path in the obtained WC–Ni cemented carbides was very short (only about 22 nm) and the sample was of a much fine dispersion, which would result in high performance in hardness for the obtained sample [24,26]. Typical micrographs on the fracture surface of the as-prepared WC–Ni cemented carbides are presented in Fig. 3. It reveals that the fracture mode could be the typical brittle fracture in ceramic materials, because the characteristics of cleavage and dimple could be obviously observed on the fracture surface of the samples in the figure, and both intergranular fracture and transgranular fracture of WC grains and transgranular fracture of cBN particles occur in the samples. This result implies that the fracture resistance of the obtained WC–Ni cemented carbides during flexural test stemmed from the intrinsic strengths of the binder phase, WC grains and cBN particles. Therefore, with the increase of cBN addition fraction, the mechanical strength of the obtained WC–Ni cemented carbides would reasonably decrease because of the low strength of cBN particles. However, as can be seen from Fig. 3a and b, the dimples on the
b
cBN
cBN
c
Matrix
cBN
Fig. 3. Typical secondary electron SEM images on the fracture surface of the as-prepared WC–Ni cemented carbide samples: (a) without cBN, (b) with cBN, and (c) showing strong bonding force between the cBN particles and cemented carbide matrix. From this figure, the fracture mode could be observed as brittle fracture.
WC grains Binder
Micropores
Fig. 2. Typical secondary electron SEM image of the WC–Ni–VC–TaC–cBN cemented carbide samples after etching. Etching: in Murakami's reagent for about 2 min.
fracture surface of the WC–Ni–cBN cemented carbide sample are deeper than those in WC–Ni sample without cBN, and so the decrease in mechanical strength of the WC–Ni–cBN could be compensated to a certain extent after the addition of cBN particles into WC–Ni cemented carbides. The microstructure of the as-prepared WC–Ni–10 vol.% cBN was also compared with those of samples sintered by other rapid sintering techniques in previously reported works. The results are listed in Table 3. From this table it could be concluded that the relative density of the as-prepared WC–Ni-10 vol.% cBN is much lower than their counterparts of WC–Co cemented carbides in literature, indicating that there might be some pores in the sample. This is consistent with the observations from Figs. 2 to 4. The formation of the micropores
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a
b Micropores
Matrix Micropores
10 µm
cBN
10 µm
c
d Matrix
cBN
cBN
Matrix
10 µm
10 µm
Fig. 4. Typical secondary electron SEM images on the polished surface of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples with different addition fractions of cBN: (a) 0 vol.%, (b) 10 vol.%, (c) 30 vol.% and (d) 50 vol.%.
Ref. [27] [27] [28] This work
WC-based cemented carbides
Sintering technique
Relative density (%)
Grain size (nm)
HV (GPa)
8Co 8Co 8Co 8Ni-10 vol.% cBN
HFIHS PCAS PCAS SPS
99.2 99.2 98.4 91.9
410 440 380 320
19.5 19.2 18.8 22.6
13.5 100
Apparent density Relative density
13.0
Apparent density (g/cm3)
Table 3 Comparison on physical properties between the obtained WC–Ni-10 vol.% cBN sample in this study and the ones sintered by high frequency induction heated sintering (HFIFS) and pulsed current activated sintering (PCAS) reported in previous work.
However, with increasing cBN, the size and number of micropores decreased, and when the addition faction of cBN reached 50 vol.%, the pores almost disappeared (as can be seen from Fig. 4), meaning that as the fraction of cBN increased, the densification of the samples would increase, and this was confirmed by Fig. 5. As the addition fraction of cBN increased, although the apparent density of the obtained samples decreased due to the low density of the additive cBN, the relative density of the obtained samples still increased, indicating that the samples become denser with increased cBN. The maximum relative
98
12.5
96
12.0
94
11.5
92
11.0
90
10.5
88
10.0
86
9.5
84
9.0
0
10
20
30
40
50
Relative density (%)
can be explained as follows. Firstly, the adopted binder metal Ni has a relatively low wet ability in the cermets during sintering in comparison with Co, which would be harmful for the spread of the binder phase during sintering [13]. Secondly, although Ni was mixed and milled at the second stage at a lower speed and for a shorter time, particle deformation and agglomeration could still happen, which promoted the formation of pores during sintering [13]. Thirdly, since the as-prepared cemented carbides were fabricated by powder metallurgy, liquid phase was formed during heating and the bond of particles relied on the formation of this liquid. Thus, the formation and flow of the liquid phase are key factors which influence the microstructure of WC-based cemented carbides. However, because the formed liquid was not sufficient at the relatively low sintering temperature and the flow of liquid was not enough during the rapid sintering by SPS, pores would inevitably be formed. Moreover, the use of VC and TaC as WC grain inhibitors in WC–Ni–VC–TaC–cBN cemented carbides would cause a fine microstructure, and this fine microstructure would also be against the flow of liquid with aggravation of forming pores during sintering [13].
82
Adding fraction of cBN (vol%) Fig. 5. Apparent density and relative density of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples as a function of the addition fraction of cBN.
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density of about 98% was reached when 50 vol.% cBN was added in this study. The reason for this phenomenon is possibly that the capillary forces are proportional to the surface tension divided by the particle size, and a micrometer-sized particle would result in an internal pressure equivalent to 10 atmospheres of external pressure during sintering [23]. Therefore, for the added micrometer-sized cBN powder, the internal pressure would increase obviously and keep increasing with more cBN added. The increased internal pressures would impel the micropores to shrink. Thus the micropores gradually disappeared in the samples and the structure became denser after the addition of cBN. This could also explain why the relative density of the samples presented a sharp increase when 10 to 20 vol.% cBN was added compared to the WC–Ni cemented carbide without cBN. And the density increase of the as-prepared WC–Ni cemented carbides would promote the mechanical properties, such as hardness and flexural strength. Moreover, it is worthwhile mentioning that the WC grain size of WC–Ni–cBN samples might be smaller than that of WC–Ni sample without cBN as can be seen from Fig. 3a and b, and this might be the result of the above-mentioned pressure, under which the growth of WC grain during sintering would be inhibited. Moreover, transgranular fracture of cBN and the tear of the cemented carbide matrix could be clearly observed from Fig. 3c, indicating that the binding force of the interface between the cBN particle and cemented carbide matrix was very high, and the pullout strength of cBN from the matrix was also much improved. Otherwise, the debonding of cBN from the materials during fracture would happen at the interface between the cBN particle and cemented carbide matrix.
The Vickers hardness of the as-prepared cemented carbides as a function of the addition fraction of cBN is shown in Fig. 6. From this figure, it can be seen that the Vickers hardness keeps increasing with increasing addition fraction of cBN. In order to compare the hardness of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide with their WC–Co cemented carbide counterparts prepared by other rapid sintering technique, the hardness of WC–Ni–VC–TaC-10 vol.% cBN was also measured by a VH-5 hardness tester (Everone) with a comparable indenting load of 30 kgf and a dwell time of 15 s with literatures as recommended by ISO 3878, and the result is listed in Table 3. From the table, it can be concluded that the hardness of the as-prepared WC–Ni–VC–TaC-10 vol.% cBN cemented carbide is comparably higher than their counterparts of WC–Co cemented carbides sintered by other rapid sintering technique [27,28], due to its
2000
1800
3000
1700
2800
1600
2600
1500 2400 1400
Vickers Hardness
Flexural Strength (MPa)
3200
Flexural Strength Vickers Hardness
2200
1300
2000
1200 0
10
finer WC grains, the addition of much harder material of cBN and very short binder mean free path [24]. The flexural strength of the cemented carbides as a function of the addition fraction of cBN is also shown in Fig. 6. It can be concluded from this figure that, with increasing addition fraction of cBN, the flexural strength of the cemented carbides decreased to about 1250 MPa with the addition fraction of cBN more than 40 vol.%. It is easily understood that the hardness of the cemented carbides increases with increasing addition fraction of cBN. One factor for the hardness increase is attributed to the strong binding force on the interface between the Ni-coated cBN particle and cemented carbide matrix (as can be seen from Fig. 3c). Another factor is due to no phase transformation of cBN during sintering (as can be seen from Fig. 1). In addition, the increase in density of the samples after cBN was added would be surely helpful to improve the hardness, and this explained why the hardness of the samples dramatically increased after 10 to 20 vol.% of cBN particles was added. And with more cBN added, the hardness of the cemented carbides increased, while the densification of the samples also increased. The reason for the decrease of flexural strength is attributed to the transgranular fracture mode of cBN particles and the weak strength of cBN. Thus the more cBN added, the lower flexural strength obtained. However, the improvement of relative density with increasing addition fraction of cBN, and the deeper dimples in WC–Ni–cBN would be helpful to improve the flexural strength and compensate the decrease of flexural strength to a certain extent. Thus, the flexural strength of samples only decreased from 1950 to 1250 MPa with the addition fraction of cBN increased from 0 to 50 vol.%. 4. Conclusions
3.2. Mechanical properties
1900
737
20
30
40
50
Adding fraction of cBN (vol%) Fig. 6. Vickers hardness and flexural strength of the as-prepared WC–Ni–VC–TaC–cBN cemented carbide samples as a function of the addition fraction of cBN.
Ultrafine WC–Ni–cBN cemented carbides were successfully fabricated by spark plasma sintering using VC and TaC as WC grain inhibitors. During sintering no cBN phase transformation happened. The Ni-coated cBN particles had a strong binding force with the cemented carbide matrix in the sintered body. The relative density of WC–Ni–VC–TaC–cBN cemented carbides increased with the addition fraction of cBN, and when the addition fraction of cBN reached 50 vol.%, the structure of WC–Ni–VC–TaC–cBN cemented carbide was almost fully dense. The flexural strength of WC–Ni–VC–TaC–cBN cemented carbides decreased from 1950 to 1250 MPa with the addition fraction of cBN increased from 0 to 50 vol.%, while the hardness of the samples increased from 2400 to 3150 HV. Acknowledgments The work was supported by Grand Survey on Land and Nature Sources of China sponsored by China Geological Survey (grant no. 1212010916026), Ph.D. Programs Foundation by Ministry of Education of China (grant no. 20100022110002), State Key Laboratory of New Ceramic and Fine Processing in Tsinghua University (grant no. KF0903), the Cultivating Foundation for Young Scientists in China University of Geosciences at Beijing from the Fundamental Research Funds for the Central Universities (Grant No. 2011PY0192) and National Lab on Scientific Drilling in China University of Geosciences at Beijing (grant no. NLSD200801). References [1] Gant AJ, Gee MG, Roebuck B. Rotating wheel abrasion of WC/Co hardmetals. Wear 2005;258:178–88. [2] Kim HC, Shon IJ, Yoon JK, Doh JM, Munir ZA. Rapid sintering of ultrafine WC–Ni cermets. Int J Refract Met Hard Mater 2006;24:427–31. [3] Imasato S, Tokumoto K, Kitada T, Sakaguchi S. Properties of ultrafine grain binderless cemented carbide RCCFN. Int J Refract Met Hard Mater 1995;13: 305–12.
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