Materials Science and Engineering A

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1 Materials Science and Engineering A 532 (2012) Contents lists available at SciVerse ScienceDirect Materials Science and Engineering A jo ur n al hom epage: Ultrafine WC Ni cemented carbides fabricated by spark plasma sintering Huiyong Rong a, Zhijian Peng a,, Xiaoyong Ren a, Ying Peng a, Chengbiao Wang a, Zhiqiang Fu a, Longhao Qi b, Hezhuo Miao b a School of Engineering and Technology, China University of Geosciences at Beijing, Beijing , PR China b State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing , PR China a r t i c l e i n f o Article history: Received 25 April 2011 Received in revised form 20 October 2011 Accepted 26 October 2011 Available online 11 November 2011 PACS: Mh Keywords: WC Ni cemented carbide Spark plasma sintering Mechanical properties Fracture mechanism a b s t r a c t With VC and TaC as WC grain growth inhibitors, ultrafine WC Ni cemented carbides with different fractions (6 10 wt%) of binder metal nickel were fabricated by utilizing high energy milling together with spark plasma sintering. In the obtained samples, only WC and Ni phases were detected in X-ray diffraction limit. The microstructure of the specimens was examined on fractural, polished, and polished/etched surfaces by scanning electron microscopy, and the results revealed that the average WC grain size of the WC Ni cemented carbides was about 330 nm, and there were lots of micro-pores in the samples. The relative density of the samples was all higher than 92%. But the measurement of hardness and flexural strength indicated that the existence of micro-pores had no significant influence on the performance of the obtained materials. On the basis of observation on the micro-fracture surface of the samples, it was found that fractures occurred along the binder metal, and the obtained ultrafine WC Ni cemented carbides showed a very short binder mean free path (about 22 nm), thus resulting in excellent performance in mechanical strength Elsevier B.V. All rights reserved. 1. Introduction WC Co cemented carbides have excellent properties, for instance, high hardness, high hot-hardness, high strength and toughness, and good wear resistance [1]; and thus they have been widely applied in many fields including cutting tools and geo-engineering equipment. However, because of the low corrosion and oxidation resistance of WC Co cemented carbides [2], and high price of the binder metal cobalt [3], their applications have been limited. Additionally, in order to further improve the performance (hardness, friction coefficient and so on) of WC Co cemented carbides, technological combination between WC Co cemented carbides and diamond films has been applied. However, in case of chemical vapor deposition of diamond films onto WC Co cemented carbides, the binder metal cobalt in cemented carbide surface layer is detrimental to the nucleation of diamond films from the gas phase and the formation of interfacial graphite, dramatically reducing the adhesion strength between diamond films and cemented carbide substrates [4,5]. Thus, engineers and scientists have endeavored for years to find new binder metals to replace cobalt in cemented carbides. To date, several metals have been proved to be possible substitutes for cobalt as binder metal phase in Corresponding author. Tel.: ; fax: addresses: pengzhijian@cugb.edu.cn, huiyong.rong@hotmail.com, pengzhijian@tsinghua.org.cn (Z. Peng). WC-based cemented carbides [6 8]. Among all those metals investigated, nickel is an exciting and promising candidate. It is not only because of its good wet ability to WC and relatively lower price than that of cobalt, but also due to the much better performance of WC Ni cemented carbides in oxidation resistance and corrosion [2] than that of WC Co cemented carbides. However, the mechanical properties (hardness and strength) of WC Ni cemented carbides are relatively inferior to those of WC Co cemented carbides [9]. The proposed ways to overcome the shortcomings of WC Ni cemented carbides and to improve their performance in previous works [2 16] include two aspects. One, according to Hall Petch formula, is to try to fabricate sub-micrometer or near-nanometer and even nanometer WC Ni cemented carbides in the way of inhibiting the growth of WC grains in cemented carbides during sintering [2 14]. Another is to directly increase the strength and/or hardness of cemented carbides by adding some materials with high strength and/or high hardness into the matrix of WC Ni cemented carbides [15,16]. In order to inhibit the growth of WC grains during sintering, two kinds of approaches have been proposed. One of them is to utilize some rapid sintering methods to shorten the duration of sintering so as to reduce the growth of WC grains [3,10,14]; and the results indicated that the hardness of the obtained WC Ni cemented carbides was significantly higher than those of both WC Co and WC Ni cemented carbides sintered by conventional sintering methods, but the fracture toughness was a little bit reduced. Another is, as in the fabrication of WC Co cemented /$ see front matter 2011 Elsevier B.V. All rights reserved. doi: /j.msea

2 544 H. Rong et al. / Materials Science and Engineering A 532 (2012) Table 1 Basic physical and chemical parameters of the used WC powder. Designation Distribution by turbid meter (%) Chemical (wt%) 0 1 m 1 2 m Total carbon Free carbon Combined carbon Oxygen GWC carbides, to choose and apply WC grain growth inhibitors into the WC Ni matrix [11 14]; and previous studies [12 14] proved that the most effective WC grain growth inhibitor in WC Ni cemented carbides was VC, followed by TaC, Cr 3 C 2, TiC and ZrC in sequence. Using solid solution technology, Correa et al. [15] effectively improved the hardness and strength of WC Ni cemented carbides by adding SiC powder into WC Ni cemented carbides. The Vickers hardness of their WC 10 wt% (Ni Si) cemented carbide was similar to that of conventional WC Co cemented carbide with the same content of binder metal, and the flexural strength was even higher that that of the later. Spark plasma sintering (SPS) is a kind of rapid sintering method; it enables the powder compact to be densified by Joule heating when the pulsed direct current goes through the powder sample. Because of the feature of rapid heating and cooling, short holding time and unique consolidation mechanism in SPS processes, the grain growth during sintering can be effectively inhibited [14,17]. Thus, combining with the above methods, the present work tailors the synthesis of WC Ni cemented carbides by using a rapid sintering method of SPS and applying VC and TaC as WC grain inhibitors to develop ultrafine WC Ni cemented carbides, trying to obtain a good combination of flexural strength and hardness for WC Ni cemented carbides. 2. Experimental procedures 2.1. Sample preparation The raw powders used in this study include WC, hydroxyl-nickel (Ni), VC and TaC, which were all commercially bought industrial reagents. Table 1 lists the basic physical and chemical parameters of the applied WC powder (the main recipe). The fraction of Ni in the cemented carbides was 6 10 wt%; and the fractions of VC and TaC were constantly 0.7 wt% and 0.3 wt%, respectively. The particle sizes of the Ni, VC, and TaC powders were about 0.7, 2 4, and m, respectively. The balls for attrition milling were made from cemented carbide YG6 (ISO: K20) and their diameters were about 5 mm. Owing to its face-centered cubic crystal structure of binder metal Ni, the Ni particle is susceptible to deformation and agglomeration during attrition milling. This would stimulate the formation of pores during sintering [13], which is detrimental for improving the relative density of the sintered samples. Therefore, in the present study, the raw powders were designedly mixed together at different milling stages and milled for disparate durations 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, the Ni powder was added into the slurry 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 mixtures were dried in a vacuum oven with a pressure of 0.01 atm at 35 C. The dried powder chunks were crashed into fine powder and sieved. After that, the resultant fine powder was loaded in a graphite die and sintered in a SPS oven (SPS-1050T, Japan). For sintering, the initial pressure applied on the graphite die was 30 MPa, the heating rate was 200 C/min, the sintering temperature was 1350 C, the duration of sintering was 6 min and the applied sintering pressure was 50 MPa, respectively. After sintering, the cooling rate was the same as that of heating. The sintered specimens were cylinders, and their dimensions were approximately 20 mm (diameter) 5 mm (height) Materials characterization X-ray diffraction (XRD, D/max2550HB+/PC, Cu K and = Å) was utilized to identify the phase composition of the as-prepared cemented carbides through continuous scanning mode with a speed of 5 /min. Three-point bending method was applied to measure the flexural strength of the samples on an AG-IC 20 kn Shimazu tester. The dimension of the testing bars was about 2 mm 3 mm 10 mm, and during the testing, the applied load rate was 0.5 mm/min. The microstructure of the specimens was examined on fractured (by LEO 1530 SEM), polished (by SSX-550 SEM), and polished/etched (by LEO 1530 SEM) surfaces, respectively. For the etching, the polished cemented carbide samples were soaked in a Murakami s reagent (1 g potassium ferricyanide, 2 g potassium hydroxide, and 30 g water) for about 2 min at room temperature. The apparent density of the samples was measured with Archimedes method according to international standard (ISO18754), and the relative density was the percentage of apparent density to theoretical density. The bulk Vickers hardness of the samples was evaluated by a LECO DM-400 hardness tester with a 1 kg load and 20 s dwell time. The WC grain size [18] and binder mean free path of the as-prepared samples were calculated by the linear intercept method from the SEM image with a Lince PC software, in which the binder mean free path was defined as the average thickness of the binder phase [19]. 3. Results and discussion 3.1. Phase composition and microstructure Fig. 1 presents the XRD patterns of the as-prepared WC Ni cemented carbide samples. It reveals that the main phases of the Fig. 1. XRD patterns of the as-prepared WC Ni cemented carbides samples with different fractions of binder metal Ni.

3 H. Rong et al. / Materials Science and Engineering A 532 (2012) Table 2 Properties of the as-prepared cemented carbides. Properties WC 6 wt% Ni WC 7 wt% Ni WC 8 wt% Ni WC 9 wt% Ni WC 10 wt% Ni Relative density (%) 92.3 ± ± ± ± ± 0.2 Vickers hardness (GPa) 24.0 ± ± ± ± ± 0.2 Flexural strength (MPa) 1609 ± ± ± ± ± 118 samples were WC and Ni, and there was no carbon phase detected in the samples, which would be helpful to improve the hardness and strength of the samples for the absence of low strength carbon carbon interface. The set of the as-prepared WC Ni cemented carbides with different contents of metal binder and their main attributes are listed in Table 2. From this table it can be seen that, when the mass fraction of the binder metal Ni was below 10%, the relative density of the samples had very little difference, fluctuating at about 92%, indicating that there might be some pores in the obtained materials. However, when the mass fraction of metal binder increased to 10%, the relative density had an obvious increase up to 97%, revealing that the number of the pores might decrease, and the samples became denser. Fig. 2 compares the relative densities of the asprepared WC Ni cemented carbides with those of some similar WC Co cemented carbides also prepared by rapid sintering techniques in previous works. The result shows that the relative density of the as-prepared WC Ni cemented carbides was lower than that of the previous work, when the applied binder fraction was the same. Fig. 3 shows the microstructure of the polished surface of typical WC Ni cemented carbides obtained in the present study. It is shown that there were lots of micro-pores in the samples, which was consistent with the supposition mentioned above and might be the reason why the obtained materials had low relative density. For the formation of the pores, the first reason would be that because the obtained cemented carbides were fabricated by powder metallurgy technology, liquid phase was formed during heating and the bond of WC grains relied on the formation of this liquid, i.e., the formation and spread of the liquid phase are the key factors which influence the microstructure of the as-prepared WC based cemented carbides. However, SPS is a rapid sintering method, and with the prompt heating rate (200 C/min) and short sintering duration (6 min), the formation and spread of liquid phase were not sufficient; furthermore, the migration and growth of WC grains were inhibited during fast sintering. And these factors would result in the inter-space between the grains which were not fully filled by metal binder, as can be seen from Fig. 4. The second reason would be that, although the binder metal Ni in the present study was mixed and milled at the second stage at a lower speed and for a shorter time during milling procedure, the particle deformation and agglomeration could still happen, which promoted the formation of pores during sintering [13]. The third one would be that although metal Ni has a good wet ability to WC during sintering, it is relatively lower than that of cobalt to WC, and this relatively lower ability would be harmful to the spread of the binder phase [13,14]. And the last one would be attributed to the formation of fine grain microstructure of WC due to the use of VC and TaC as WC grain inhibitor in the materials, which would be detrimental to the flow of liquid with aggravation of forming pores during sintering [13]. Thus, it could be reasonably concluded that the sufficient formation and spread of liquid phase during sintering were responsible Fig. 3. Typical micrograph of the polished surface of the as-prepared WC Ni cemented carbides with pores of far less than 1 m in diameter and distributed uniformly in the matrix. Fig. 2. Relative densities of the as-prepared WC Ni cemented carbides compared to the literature data: (a) WC 6 wt% Co fabricated by spark plasma sintering, (b) WC 8 wt% Co fabricated by pulsed current activated sintering, and (c) WC 10 wt% Co fabricated by spark plasma sintering. Fig. 4. Typical micrograph of the as-prepared WC Ni cemented carbides after etching in Murakami s reagent for about 2 min.

4 546 H. Rong et al. / Materials Science and Engineering A 532 (2012) for the obvious increase in the relative density of the samples as the mass fraction of binder metal Ni increased to 10 wt%, and it could be also reasonable to deduce that 10 wt% is the critical value for preparing dense samples under the designed experimental conditions in this study, i.e., when the fraction of Ni is 10 wt%, the formation and spread of liquid phase become sufficient enough, so the relative density of the samples increases obviously compared with the other samples in this study. On contrast, when the fraction of Ni is below 10 wt%, the increase of relative density of the samples is unobvious because the formation and spread of liquid phase is not sufficient enough as can be seen from the little change in the relative density when the binder fraction increased from 6 to 9 wt%. Fig. 4 shows the microstructure of a typical polished/etched surface of the as-prepared WC Ni cemented carbide samples. The average size of the WC grains calculated from the figure was about 330 nm, indicating that ultrafine WC Ni cemented carbides were successfully fabricated, which is helpful to improve the hardness and strength of WC cemented carbides according to Hall Petch formula. The fabrication of ultrafine WC Ni cemented carbides could be attributed to the short duration during spark plasma sintering and the use of VC and TaC as WC grain inhibitors. Typical micrograph of the fracture surface of the as-prepared WC Ni cemented carbides is presented in Fig. 5. It reveals that the fracture mode through binder phases was of predominance in the obtained materials, characterized by the deep dimples, the survival of binder phase, and the fracture mode along the WC WC interface. These results suggest that the fracture resistance of the obtained WC Ni cemented carbides during the measurement of mechanical properties was controlled by the intrinsic strength of the binder phase and the degree of squeletum formation or contiguity of the WC grains. From Fig. 4 it could be also observed that the binder mean free path in the obtained WC Ni cemented carbides was very short (22 nm). It would constrain the propagation of crack during fracture, thus resulting in high performance in hardness and strength [19,20]. In addition, the survival of binder phase indicated that the rupture of the binder layers occurred between WC particles, and this might be due to the rapid nucleation and propagation of microcracks in the WC WC interface of the squeletum for the high stress of WC particles suffered in the binder phase, while the rigid continuous squeletum of WC particles was submitted to plastic deformation; and this observation is well in line with the findings of Ref. [21]. Investigations on the fracture surfaces also showed that although there were lots of micro-pores in the material, they were not obviously found in the fracture surface, and this might indicate that the existence of micro-pores had no significant influence on the mechanical property of the obtained materials Mechanical properties The hardness and flexural strength of the as-prepared WC Ni cemented carbides presented different dependences on the content of binder metal Ni in the materials (see Table 2). The hardness decreased with the increasing mass fraction of binder metal Ni, but the flexural strength increased. The reason is that WC Ni cemented carbides were fabricated by powder metallurgy technology, and they contained WC ceramic grains (with relatively much high hardness but low strength and toughness) and metal binder phase (with relatively much low hardness but high strength and toughness). Thus, the mass fraction of binder metal in the materials would exert opposite influence on the hardness and flexural strength. Fig. 6 compares the hardness of the as-prepared WC Ni cemented carbides with those of the counterparts in previous works, which were also fabricated by rapid sintering techniques. The result shows that, with the same mass fraction of binder metal, the hardness of the as-prepared WC Ni cemented carbides was higher than that of the WC Co cemented carbides reported in the literature. The hardness of the WC 6 wt% Ni cemented carbide was as high as 24 GPa, and the hardness of the WC 10 wt% Ni cemented carbide was 18 GPa. This would be possibly because of the absence of carbon phases, the special fracture model (along the binder metal) of the materials and very short binder mean free path [19,20] detailedly discussed in the last section. And with more Ni was added, the sample hardness decreased due to the decrease in the content of hard phase WC. The data of the flexural strength of the as-prepared WC Ni cemented carbides cannot be able to compare with those in the other works for their different dimensions because in the present work the data were obtained using testing bars with size of 2 mm 3 mm 10 mm, whereas, the data in the literature were obtained using testing bars with size of 3 mm 4 mm 30 mm or 4 mm 6 mm 30 mm. But it is no doubt that the flexural strength of 1600 MPa when the binder fraction was 6 wt% and 2100 MPa when the binder fraction was 10 wt% were still relatively high values, and the decrease in flexural strength was only 500 MPa when the binder fraction increased from 6 to 10 wt%. The reason for the Fig. 5. Typical fracture surface of the as-prepared WC Ni cemented carbides in which the survival of binder could be clearly observed, revealing that the fracture mode of the hardmetals was a fracture along the binder metal. Fig. 6. Vickers hardnesses of the as-prepared WC Ni bulk cemented carbides compared to the literature data: (a) WC 6 wt% Co fabricated by spark plasma sintering, (b) WC 8 wt% Co fabricated by pulsed current activated sintering, and (c) WC 10 wt% Co fabricated by spark plasma sintering.

5 H. Rong et al. / Materials Science and Engineering A 532 (2012) high flexural strength would be also possibly because of the absence of carbon phase, and the along the binder metal fracture model in the materials. And with more Ni was added, the flexural strength of the samples increased due to the increase in the content of tough phase metal Ni. 4. Conclusions Ultrafine WC Ni cemented carbides were fabricated using spark plasma sintering and with the assistance of VC and TaC as WC grain inhibitors. The relative density of the obtained WC Ni cemented carbides was relatively a bit lower than that of WC Co cemented carbides with the same binder fraction in the literature, due to the existence of micro-pores. But the hardness of the obtained WC Ni cemented carbides was higher than those of WC Co cemented carbides fabricated by some rapid sintering methods in the literature, and at the same time the flexural strength still possessed very high value from 1600 to 2100 MPa when the mass fraction of the binder metal Ni increased from 6 to 10 wt%, due to the absence of carbon phase, very short binder mean free path, and the special fracture model along binder metals. Acknowledgements This work was supported by Grand Survey on Land and Nature Sources of China sponsored by China Geological Survey (Grant No ), Ph.D. Programs Foundation by Ministry of Education of China (Grant No ), the Cultivating Foundation for Young Scientists in China University of Geosciences at Beijing from the Fundamental Research Funds for the Central Universities (Grant No. 2011PY192) and State Key Laboratory of New Ceramic and Fine Processing in Tsinghua University (Grant No. KF0903). References [1] A.J. Gant, M.G. Gee, B. Roebuck, Wear 258 (2005) [2] S. Imasato, K. Tokumoto, T. Kitada, et al., Int. J. Refract. Met. Hard Mater. 13 (1995) [3] H.-C. Kim, I.-J. Shon, J.-K. Yoon, et al., Int. J. Refract. Met. Hard Mater. 24 (2006) [4] F.H. Sun, Z.M. Zhang, M. Chen, et al., Diam. Relat. 12 (2003) [5] C. Gil, G. Jan, L. Jörn, et al., Diam. Relat. Mater. 17 (2008) [6] B. Reichel, K. Wagner, D.S. Janisch, et al., Int. J. Refract. Met. Hard Mater. 28 (2010) [7] M. Aristizabal, N. Rodriguez, F. Ibarreta, et al., Int. J. Refract. Met. Hard Mater. 28 (2010) [8] E. Taheri-Nassaj, S.H. Mirhosseini, J. Mater. Process. Technol. 142 (2003) [9] E.A. Almond, B. Roebuck, Mater. Sci. Eng. A 105 (1988) [10] J.F. Zhao, T. Holland, Int. J. Refract. Met. Hard Mater. 27 (2009) [11] J. Zackrisson, B. Jansson, G.S. Uphadyaya, et al., Int. J. Refract. Met. Hard Mater. 16 (1998) [12] B. Wittmann, W.-D. Schubert, B. Lux, Int. J. Refract. Met. Hard Mater. 20 (2002) [13] Z.C. Jia, Powder Metall. Ind. 9 (1999) (in Chinese). [14] A. Mukhopadhyay, B. Basu, J. Mater. Sci. 46 (2011) [15] E.O. Correa, J.N. Santos, A.N. Klein, Int. J. Refract. Met. Hard Mater. 28 (2010) [16] K. Brookes, Met. Powder Rep. 62 (2007) [17] W.B. Liu, X.Y. Song, K. Wang, et al., Mater. Sci. Eng. A 499 (2009) [18] J.H. Han, D.Y. Kim, Acta Mater. 46 (1998) [19] K. Jia, T.E. Fischer, B. Gallois, Nanostruct. Mater. 10 (1998) [20] H.C. Lee, J. Gurland, Mater. Sci. Eng. 33 (1978) [21] G. Grathwohl, R. Warren, Mater. Sci. Eng. 14 (1974)