Pulse Electric Current Sintering of Cubic Boron Nitride/Tungsten

Materials Science & Engineering A 607 (2014) 490–497

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Materials Science & Engineering A

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Pulse electric current sintering of cubic boron nitride/tungsten carbide–cobalt (cBN/WC–Co) composites: Effect of cBN particle size and volume fraction on their microstructure and properties

Bo Wang a,b,n, Yi Qin a, Feng Jin b, Jian-Feng Yang a,nn, Kozo Ishizaki c a State Key Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xianning West Road No. 28, Xi'an 710049, Shaanxi Province, China b State Key Laboratory for Strength and Vibration of Mechanical Structures, School of Aerospace Engineering, Xi'an Jiaotong University, Xianning West Road No. 28, Xi'an 710049, Shaanxi Province, China c Department of Mechanical Engineering, Nagaoka University of Technology, Nagaoka 940-2188, Japan article info abstract

Article history: Cubic boron nitride/tungsten carbide–cobalt (cBN/WC–Co) composites were fabricated by pulse electric Received 15 February 2014 current sintering (PECS), using Ni–P as sintering additives to promote low temperature densification. The Received in revised form effect of cBN particle size and volume fraction on the densification, microstructure and mechanical 3 April 2014 properties of WC–Co composites was investigated. There was no phase transformation from cBN to hBN Accepted 4 April 2014 (hexagonal BN) with low-hardness due to low sintering temperature (1100–1200 1C) and short sintering Available online 13 April 2014 time. Smaller cBN particle led to lower sinter-ability of the composites due to its high specific surface Keywords: area. The 30 vol% 10–14 mm cBN/WC–Co composite (P14V30) exhibited high hardness (18.3 GPa, 1200 1C) Cubic boron nitride and high fracture toughness (15.6 MP m1/2,10001C). The high hardness resulted from the homo- Tungsten carbide geneously dispersed cBN particles, which had a good bonding with the WC matrix. Increased fracture Pulse electric current sintering toughness was mainly attributed to crack deflection or bridging and pullout of cBN grains. Hardness & Fracture toughness 2014 Elsevier B.V. All rights reserved.

1. Introduction alloys. The strength and the fracture toughness of the cemented carbides were improved by increasing the amount of the binding Cubic boron nitride (cBN) has the second highest hardness and phase from 5 wt% to 25 wt% cobalt, but the hardness decreased as thermal conductivity after diamond, and it was characterized the WC particle size increased [2,3]. by greater thermal stability and less reactivity with iron than diamond. The addition of cBN particles into WC–Co cemented carbide Owing to these excellent properties, polycrystalline cubic boron nitride could increase its hardness and wear resistance. In addition, has been widely used as strengthening particles for high-modulus superior fracture toughness could be obtained as a result of the cemented carbide for the application in super hard cutting tools, crack-deflection via cBN particles[4,5]. Thus, cBN dispersed WC–Co drilling tools, and wear resistant parts [1]. cemented carbide composites (CDCC) would be promising to Tungsten carbide–cobalt (WC–Co) cemented carbides have realize cutting tools with high hardness and fracture toughness. been widely used for fabricating cutting or heat- and wear- However, the addition of cBN particles into the cemented carbides resistant tools because of their superior mechanical properties. significantly reduced their sinter-ability due to strong covalent However, the tool life was short because WC has a lower wear bonds of cBN and low self-diffusion coefficient of B and N. More- resistance and hardness than other carbides. Usually, the mechan- over, the transformation from cBN to low hardness hexagonal ical properties of WC–Co cemented carbides were modified by boron nitride (hBN) with a graphite structure occurred at the changing the content of the Co binding phase or the grain size of eutectic point of WC/Co alloy (1320 1C) in regular liquid phase WC starting powder [2,3]. However, the hardness and toughness sintering, which would decrease the hardness of the composites. were inversely related for the conventional coarse-grained WC–Co Therefore, the production of CDCC composites were generally sintered by very costly techniques employing a high temperature and an ultra-high pressure (5–8GPa)[6–9] to achieve fully dense n Corresponding author at: State Key Laboratory for Mechanical Behavior of material without BN phase transformation. The application of Materials, Xi'an Jiaotong University, Xianning West Road No. 28, Xi'an 710049, sintering techniques at low-moderate pressure was limited due Shaanxi Province, China. Tel.: þ86 29 82667942 803; fax: þ86 29 82667942 804. nn to the low stability of the cBN at a high temperature. Corresponding author. E-mail addresses: [email protected] (B. Wang), Pulse electric current sintering (PECS) technique could be use [email protected] (J.-F. Yang). to heat specimens rapidly because the pulsed direct current used http://dx.doi.org/10.1016/j.msea.2014.04.029 0921-5093/& 2014 Elsevier B.V. All rights reserved. B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497 491 in this technique could pass though the conduct (graphite die system (Sumitomo Coal Mining Co. Ltd., Tokyo, Japan). The sinter- and punch rods) and powder mixture [10,11]. Consequently, the ing temperature was ranged from 950 1C to 1300 1C for 10 min densification time required for sintering could be shortened and with an axial pressure of 50 MPa. The sample was heated to 600 1C the phase transformation from cBN to hBN could be inhibited with a heating rate of 100 K/min, and then, it was heated to 700 1C [12,13]. On the other hand, the optimization of sintering tempera- for 10 min with a heating rate of 15 K/min. After that, it was heated ture was essential to produce cBN/WC–Co composite by PECS to 950–1300 1C for 10 min with a heating rate of 50 K/min (the technique under a moderate low pressure o100 MPa. heating rate was changed to 25 K/min during final 50 1Cto Ni–P has already been used as sintering additives in the nano- the objective temperature). Finally, it was cooled down with the structured WC–Ni–PorWC–Co–Ni–Pcoating[14,15] to improve furnace. The codes of P3V10 and P14V30 were used for the corrosion and wear resistance, and in the micro-grain sized WC–Co samples, where “P3” and “P14” represented the article size of cemented carbides to increase the sinter-ability and then decrease 2–3 and 10–14 μm, respectively, “V” represented the volume the sintering temperature [16]. In our previous research, near fully percent of cBN addition. dense cBN/WC–Co composites have been obtained at low tempera- The densities of the sintered compacts were determined by the ture 1100–1200 1C using Ni–P as sintering additive by PECS tech- Archimedes method. The hardness was measured by indentation test nology [17]. And the homogeneously dispersed cBN particles in the through Vickers method with load of 49 N for 15 s. The fracture

WC–Co composites promoted a slight increase of hardness, but a toughness KIC of the sintered cBN–WC/Co composites was deter- significant improvement of fracture toughness [17]. Beside the effect mined by measuring the crack length from the tip of the indentation of the cBN volume content [5,17], there was a clear dependence of made by Vicker's indentation load of 490 N. The fracture toughness sinter-ability and mechanical properties of cBN/WC–Co composites KIC was calculated according to Niihara's equations [18,19]: KIC ¼ 2=5 = pffiffiffi on the c-BN particle size [5]. This effect was more appreciable at low Hv c 1 2 3=5 0:035 1 Hv aϕ ,whereE is Young's modulus of pressure o100 MPa [5]. E a the cemented carbides, Hv is the hardness, a is the half length of the The purpose of this research was to fabricate cBN dispersed indent diagonal, c is the half-crack length from the center of the WC–Co cemented carbides by PECS technique at low sintering indent to tip of the crack, ϕ is a constraint factor (3). Young's temperature, using Ni–P as sintering additives. The influence of modulus of the cBN–WC/Co composites was calculated according to cBN particle size (2/3 and 10/14 mm grades) and cBN volume the mixture rule. Young's modulus of WC, Co, Ni–P, cBN are 696, 174, fraction (0–30 vol%) on the densification, microstructure and 205, and 909 GPa, respectively [20–25]. Young's modulus of cemen- mechanical properties of cBN/WC–Co composites was investi- ted carbides (WC–9.7 wt% Co–2.9 wt% Ni–P) was calculated as gated. The cBN reinforcing mechanism was also discussed. 590 GPa. And Young's modulus of the cBN–WC/Co composites was calculated at the same way. The cBN to hBN phase transformation was investigated by ™ 2. Experimental procedure X-ray diffraction (XRD, SHIMADZU XRD-6000 , Shimadzu Corp., Kyoto, Japan) and Fourier transform infrared spectroscope Tungsten carbide powder (WC, mean particle size: 0.2 mm, TIX (FT-IR8300, Shimadzu Corp., Kyoto, Japan). The microstructure of Holdings Co. Ltd., Tokyo, Japan) containing 9.7 wt% cobalt (Co, the polished and etched surface and the fracture surface of the mean particle size: 3 mm, TIX Holdings Co. Ltd., Tokyo, Japan) and cBN–WC/Co composites were observed by scanning electron 2.9 wt% nickel–phosphorus powder (Ni–P, 89 wt% Ni and 11 wt% P, microscopy (SEM, VE-7800, KEYENCE Corp., Osaka, Japan). The 5–20 mm, TIX Holdings Co. Ltd., Tokyo, Japan) as sintering additives average grain size of WC in the sintered composites was measured were used as starting materials. The powder mixture with differ- by the mean linear intercept method. ent WC–Co–Ni–P: cBN (2–3 mm and 10–14 mm, BN3600, LANDS Superabrasives, Co., NY, USA) volume ratios of 100:0, 90:10, 80:20, and 70:30 were prepared by wet milling in anhydrous alcohol for 3. Results and discussion 24 h in a plastic bottle. The morphology of the cBN particles with different particle size was shown in Fig.1. 3.1. Thermal stability of the cBN After milling, the isopropyl alcohol was removed and the powder mixture was sieved through mesh screen (25 μm) to It was well known that the thermal stability of cBN was highly break up the soft agglomerates. The powders were put into a reduced by the presence of cobalt, especially in the presence of 20 mm inner diameter graphite die directly and uniaxially pressed liquid [5,12,26]. The dissolution–precipitation process of boron under a pressure of 3 MPa. The pulse electric current sintering was nitride during liquid phase sintering would accelerate the hBN carried out under vacuum in the Dr. Sinter Model SPS-1050 T PECS formation. As a result, the hBN in the composites would

Fig. 1. SEM micrographs of different cubic boron nitride powder grades used in this work. (a) 2–3 μm and (b) 10–14 μm. 492 B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497

take place in the solid state at sintering temperature lower than 1200 1C, which could restrain the phase transformation from cBN to hBN. The 804 cm1 peak appeared clearly when the BN–WC/Co composites sintered at 1300 1C, conﬁrming that cBN to hBN phase transformation occurred at nearly the WC–Co carbide eutectic temperature (1320 1C). The hBN peak intensity of the composites with 10–14 mm cBN (P14V20) was slightly lower than that of P3V20 [17], indicating the lower thermal stability of cBN with smaller particle size. As expected, the thermal stability was highly reduced by the presence of Co liquid at 1300 1C. In this case, the boron nitride dissolution-precipitation dominant mechanism for the cBN to hBN phase transformation would be accelerated [5]. Fine cBN particles with high speciﬁc surface area would promote the dissolution–precipitation process. In other words, the cBN stability was proportional to the particle size.

Fig. 2. XRD patterns of starting powder of P3V30 and 30 vol% cBN–WC/Co composite at 1300 1C for 10 min under 50 MPa. 3.2. Densiﬁcation and sinter-ability

Fig. 4 demonstrates the relationship between the relative density of cBN/WC–Co composites sintered at different temperature and the cBN addition. As can be observed in the diagram, the addition of cBN particles to the WC/Co cemented carbides significantly impeded their densification. The particle size of the WC and Co starting powders was much less than the cBN powders. Therefore, the much finer WC starting powders yield a larger driving force for the densification compared to the coarser cBN powders. Moreover, the significant WC solubility in solid Co binder and the good wetting between Co and WC contribute to the higher solid state densification of WC–Co based hard metals when compared to the cBN–WC/Co based hard metals. This effect was more appreciable at low sintering temperature. Below 1100 1C, the relative density was considerably decreased with increasing the volume content of cBN. The density of 10–14 μm cBN/WC–Co composites sintered at 950 1C decreased to 97.4% of theoretical density (TD) with increasing the content of cBN to 30 vol% (P14V30 sample). In contrast, the density of P3V30 sample

Fig. 3. FT-IR spectra of the sintered cBN/WC–Co composite containing 20 vol% dramatically decreased to 85.8% of TD. The addition of larger 10–14 mm cBN at different temperatures for 10 min under 50 MPa. amounts of 2–3 cBN (30 vol%) to the WC–Co material led to a higher reduction of the shrinkage and densification of the compacts due to the higher specific area of cBN contact with WC–Co. considerably decrease the sinter-ability and the mechanical prop- The particle size of the cBN added to the basic WC/Co carbides erties of the WC–Co carbides. Fig. 2 show the XRD patterns of exhibited a significant influence on the sinter-ability of the starting powder of P3V30 and 30 vol% cBN–WC/Co composites sintered at 1300 1C. Besides WC and Co, only cBN phase was detected. The peak of hBN could not be observed for all samples. Moreover, the height of cBN peak was much lower than that expected peak intensities based on the volume fraction in the composites (30 vol%). However, it was not sure that there was no cBN to hBN phase transformation at 1300 1C because the mass- absorption coefficient of WC/Co was much stronger compared with those of cBN of hBN [3]. Thus the convertion from cBN to hBN at the interfaces in the cBN–WC/Co composites was difficultly identified and quantified. Fig. 3 shows the transmission spectra of the 20 vol% 10–14 mm cBN/WC–Co (P14V20) composites sintered at different temperature for 10 min under a pressure of 50 MPa. FT-IR peaks of cBN were at about 1245 and 1080 cm1, whereas hBN peaks appeared at about 1383 and 804 cm1 [27,28]. It could be observed that the cBN peak intensity gradually decreased with increasing the sintering temperature from 1100 1C to 1300 1C, and thoroughly disap- peared at 1500 1C. Below 1200 1C, no hBN peak was observed, Fig. 4. Relationship between the cBN addition and the relative density of cBN/WC– indicating that there was no phase transformation from cBN to Co composites, which are sintered at different temperatures for 10 min under hBN. It was important to emphasize the densification by PECS may 50 MPa. B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497 493

Fig. 5. SEM micrographs of the polished surface of cBN/WC–Co composites at 1000 1C: (a) P2V20, (b) P14V20, (c) P14V30; and at 1200 1C: (d) P2V20, (e) P14V20 and (f) P14V30.

composites. Smaller cBN particle led to lower sinter-ability of the phase transformation would create a volume increase, which WC–Co composites. reduced the relative density of the WC–Co materials and The density rapidly increased with increasing the sintering decreased the sinter-ability of the tungsten carbide even if it was temperature from 950 1C to 1200 1C. Near fully dense (497.7% 100% dense [29]. of TD) composites could be achieved for all 0–30 vol% cBN/WC–Co samples sintered at 1200 1C under 50 MPa for 10 min. And the 3.3. Microstructure characterization highest density (499% TD) could be obtained for 10–14 μm cBN/ WC–Co composites. At this optimum temperature of 1200 1C, the Fig. 5 shows the SEM micrographs of the polished surface of influence of cBN size and content was not noticeable, especially for cBN/WC–Co composites sintered at different temperature under the latter one. The adding of small amount of Ni–P allowed 50 MPa for 10 min. The cBN particles (in black contrast) were chemically-activated sintering at lower temperature of 1200 1C homogeneously distributed and had a good bonding with WC and it could effectively enhance the densification during solid state matrix due to Co binder. The morphology of the cBN particles in all sintering. Above 1200 1C, the density decreased, especially for the samples was similar. Additionally, no obvious pores were found in composites containing a large amount of cBN. The decrease of these specimens. However, when addition of cBN increasing to density from 1200 1C to 1300 1C was due to the phase transforma- 30 vol%, as shown in Fig. 5(f), some cBN particles were pulled out tion from cBN to hBN, which could be detected by FT-IR, or by the from the surface during the grinding or polishing operations for degradation of the bonding between the cBN particles and the metallographic preparation. For the samples sintered at low hard metal matrix at high temperature [5,17]. The cBN to hBN temperature of 1000 1C, WC–Co material exhibited low contrast 494 B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497 and a large amount of heterogeneously distributed Co (in gray 3.4. Mechanical properties contrast) could be distinguished (Fig. 5(b) and (c)). Large Co pools from the initial mixtures were still present, and spreading of the Fig. 7 shows the dependence of Vickers hardness of the cBN/ binder had only started on the neighboring WC particles, which WC–Co composites under 50 MPa for 10 min on the sintering indicated the whole densification took place at solid state during temperature. As shown in it, the addition of well-dispersed cBN sintering. In this condition (at low temperature of 1000 1C), the hard particles to WC–Co hard metal base materials resulted in homogeneous distribution of Co was more difficult. With increas- an increase of Vickers hardness. The variation of hardness of ing the sintering temperature to 1200 1C, large Co pools could cBN/WC–Co composites with the sintering temperature was quite not be observed from Fig. 5(e) and (f). Fig. 6 shows the high similar to the variation of density. As seen from Fig. 7(a), the magnification SEM micrographs of the polished surface of P14V20 hardness of P3V10 composites increased with the sintering tem- sintered at different temperature. The large Co pools could be perature until a maximum value (17.0 GPa, P3V10) appeared at observed clearly in Fig. 6(a). As the sintering temperature was 1200 1C. And then, the hardness decreased to 14.4 GPa with raised to 1200 1C, as seen in Fig. 6(b), Co binder between WC increasing the temperature to1300 1C due to the transformation grains considerably decreased and the grain size of WC was from hard cBN to soft hBN. The hardness of P14V10 samples was slightly increased in spite of the lack of Co liquid phase formed. lower than that of P3V10 with similar relative density. It may be This result indicates that the mass transport like surface diffusion due to the lower amount of cBN particles per unit volume on the or dissolution of the WC grains was very fast even during solid composites. However, the hardness of P14V30 samples was higher state sintering. On the other hand, the addition of cBN particles than that of P3V30 because of the lower relative density of P3V30. into WC–Co cemented carbides significantly reduced the Co The P14V30 sample with larger cBN particles sintered at 1200 1C pools due to few amount of Co diffuse to the interface between possessed the highest hardness of 18.3 GPa (Fig. 7(b)). cBN and WC grains. cBN particles exhibited smaller surface area Usually, the hardness of nanocrystalline cemented tungsten and smoother surface compared with the WC particles. As carbide was significant reduced due to grain growth [19,30]. Fig. 8 expected, cBN particles with larger particle size were more prone shows the SEM micrographs of fractured surface of WC–Co and to Co diffusing to the WC/cBN interfaces[17]. However, even 20 vol% cBN/WC–Co composites. As shown in Fig. 8, the added cBN though fully dense P14V20 sample (99% of TD) could be consoli- particles were distributed homogeneously and they had a good dated at 1200 1C by PECS process, Co binder between WC grains bonding with the WC matrix (Fig. 8(d)). Compared with the flat was observed. fracture surface of WC–Co, some cBN particles were pulled out For cBN/WC–Co composites, the densification mechanism was from the surface of cBN/WC–Co composites. The average grains the particle rearrangement, enhanced by the diffusion and viscous size of WC–Co was a little bit larger than that of 20 vol% cBN/WC– flow of the Co binder. In the present case, applied pressure of Co composites. For WC–Co materials, the rapid grain growth 50 MPa could collapse pore structures and induce Co plastic flow occurred by the coalescence of grains [31] or solution–reprecipita- and rapid bonding in powder compacts at elevated temperatures tion coarsening [32,33] during solid state sintering. The atomic (110 0 1C). At this stage, the WC particles, considered as perfectly mobility of atoms of the refractory phase to the Co binder phase rigid, could undergo substantial rearrangement due to the plastic was dominant for solid-state sintering. This grain growth was deformation of the cobalt phase. Beside the high-applied pressure, predominantly diffusion controlled. In this processing, the densi- the sintering temperatures selected in this work were in the range fication of solid-state sintering was mainly driven by the spreading of hot rolling of cobalt alloys 1000 1C [11], where large plastic of the metallic binder. Compared to cBN/WC–Co composites, thin deformation could be produced. Once WC particles formed a solid film of Co binder phase between WC grains in the WC–Co based skeleton in mechanical equilibrium, other mechanisms, as cobalt cemented carbides was facilitated at 1200 1C. Consequently, the creep flow or diffusion, needed to be taken into account to explain coalescence of grains and the grain growth by the diffusion of the removal of residual porosity. Both phenomena were possible carbon and tungsten atoms in cobalt took place more easily for the for cobalt at 1000–1200 1C(41/2 Co melting point). As a result, Co WC–Co based cemented carbides. The grain growth of WC could binder between WC grains considerably decreased with increasing be inhibited to some extent by the cBN particles. Furthermore, the the temperature to 1200 1C. Nevertheless, the cobalt distribution PECS process could also inhibit the WC grain growth. Compared was not perfect compared with that of WC–Co materials consoli- with the density, the slight grain growth only had a little influence dated by liquid phase sintering (Fig. 6(b)). on the hardness on micro-scale.

Fig. 6. High magniﬁcation SEM micrographs of the polished surface of P14V20 sample at (a) 1000 1C and (b) 1200 1C under 50 MPa for 10 min. B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497 495

Fig. 7. The dependence of the Vickers hardness on the sintering temperature of the cBN/WC–Co composites sintered at 50 MPa: (a) 10 vol%; and (b) 30 vol% cBN.

Fig. 8. SEM micrographs of the fracture surface of cBN/WC–Co composites sintered at 1200 1C: (a) WC–Co, (b) P2V20, and (c and d) P14V20.

The fracture toughness, calculated from the crack length induced different temperature on the particle size and volume content of by the Vicker's indentation under a load of 490 N is shown in Fig. 9. the cBN addition. For the 2–3 mm cBN/WC–Co composites sintered ThefracturetoughnessofcBN/WC–Co composites with lower density at 1100–1200 1C, the hardness increased and subsequently was obviously higher than that of WC–Co cemented carbides without decreased with increasing the cBN addition. For the 10–14 mm cBN particles, especially for the 10–14 mmcBN/WC–Co composites. cBN/WC–Co composites sintered at 1100–1200 1C, the hardness For the WC–Co and 2–3 mmcBN/WC–Co composites, the toughness increased with the content of cBN. As shown in Fig. 10(b), the increased slightly from 900 1Cto11001C and subsequently decreased fracture toughness of all samples increased with the content of with increasing the sintering temperature to 1300 1C. Highest tough- the cBN. ness of 14.9 MP m1/2 could be achieved for P14V30 samples sintered Fig. 11 shows the typical crack propagation of the cBN/WC–Co at 1000 1C. composites. The crack propagation path in WC–Co baseline hard Fig. 10 shows the dependence of (a) Vickers hardness and metals was much more straight compared with cBN/WC–Co (b) fracture toughness of the cBN/WC–Co composites sintered at composites. It could be clearly seen that the crack passed through 496 B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497

Fig. 9. The dependence of the fracture toughness on the sintering temperature of the cBN/WC–Co composites sintered at 50 MPa: (a) 10 vol% and (b) 30 vol% cBN.

Fig. 10. The dependence of (a) Vickers hardness and (b) fracture toughness on the cBN volume content of the cBN/WC–Co composites sintered at different temperatures.

Fig. 11. SEM micrographs of the crack propagation path of cBN/WC–Co composites: (a) WC, 1200 1C; (b) P3V20, 1200 1C; (c) P14V10, 1200 1C and (d) P14V20, 1000 1C. B. Wang et al. / Materials Science & Engineering A 607 (2014) 490–497 497 the boundaries of fine grains as well as the cBN particles oriented Specialized Research Fund for the Doctoral Program of Higher Educa- at low angles, resulting in the crack bridging and deflection by the tion of China (Grant nos. 20130201120003 and 20100201110036), and coarse cBN particles. The larger cBN particle size resulted in the by the Japan Society for Promotion of Science (Improvement of larger angle of crack deflection (comparing Fig. 11(b) with (d)). Research Environement for Young Researchers). These deflections were caused by the interaction of the growing crack front with the hard cBN particles. 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Betteridge, Cobalt and Its Alloys, Ellis Horwood Limited and Wiley, 1982. – the WC/Co carbides promoted an increase of hardness and [22] S.I. Cha, S.H. Hong, B.K. Kim, Mater. Sci. Eng. A 351 (2003) 31 38. [23] Y. Zhou, U. Erb, K.T. Aust, G. Palumbo, Scripta Mater. 48 (2003) 825–830. a significant improvement of fracture toughness. The 30 vol% [24] M.P. D’Evelyn, T. Taniguchi, Diam. Relat. Mater. 8 (1999) 1522–1526. 10–14 mm cBN/WC–Co composite (P14V30) sintered at 1200 1C [25] V.L. Solozhenko, S.N. Dub, N.V. Novikov, Diam. Relat. Mater. 10 (2001) showed excellent mechanical properties (Vickers hardness of 2228–2231. – 1/2 [26] K. Brookes, Met. Powder Rep. 62 (2007) 14 17. 18.3 GPa and fracture toughness KIC of 12.5 MP m ). The crack [27] G.H. Chen, X.W. Zhang, B. Wang, X. Song, B Cui, H Yan, Appl. Phys. Lett. 75 deflection or bridging and the pullout of cBN grains contributed to (1999) 10–12. high fracture toughness. [28] Z.P. Xia, Z.Q. Li, J. Alloy. Compd. 436 (2007) 170–173. [29] J.C. Garrett, I. Sigalas, M. Herrmann, E.J. Olivier, J.H. O’Connell, J. Eur. Ceram. Soc. 33 (2013) 2191–2198. [30] D.K. Shetty, I.G. Wright, P.N. Mincer, A.H. Clauer, J. Mater. Sci. 20 (1985) Acknowledgment 1873–1882. [31] X. Wang, Z.Z. Fang, H.Y. Sohn, Int. J. Refract. Met. Hard. Mater. 26 (2008) – This work was supported by the National Natural Science Founda- 232 241. [32] V. Bounhoure, J.M. Missiaen, S. Lay, E. Pauty, J. Am. Ceram. Soc. 92 (2009) tion of China (Grant nos. 51302208, 51272205, and 51072157), by the 1396–1402. China Postdoctoral Science Foundation (Grant no. 2012M521763), by [33] A.V. Laptev, Powder Metall. Met. Ceram. 46 (2007) 415–422.