Effect of Zrc Nano-Powder Addition on the Microstructure and Mechanical Properties of Binderless Tungsten Carbide Fabricated by Spark Plasma Sintering

Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407

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Int. Journal of Refractory Metals and Hard Materials

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Effect of ZrC nano-powder addition on the microstructure and mechanical properties of binderless tungsten carbide fabricated by spark plasma sintering

Xiaoyong Ren a, Zhijian Peng a,⁎,ChengbiaoWanga, Zhiqiang Fu a, Longhao Qi b, Hezhuo Miao b a Key Laboratory on Deep GeoDrilling Technology of the Ministry of Land and Resources, School of Engineering and Technology, China University of Geosciences, Beijing 100083, PR China b State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, PR China article info abstract

Article history: Binderless tungsten carbide with varied amounts of ZrC nano-powder (0–9 wt.%) were prepared by spark plasma Received 18 August 2014 sintering at 1600 °C under an applied pressure of 50 MPa. The effect of ZrC nano-powder addition on the densi- Received in revised form 23 October 2014 fication behavior, phase composition, microstructure and mechanical properties of the composites were investi- Accepted 26 October 2014 gated. With appropriate amount of ZrC nano-powder, the coarsening and abnormal growth of WC grains were Available online 29 October 2014 suppressed, resulting in more homogeneous microstructure of the materials. The Vickers hardness, flexural fi Keywords: strength and fracture toughness of the as-prepared composites rst increased and then decreased with increas- Binderless tungsten carbide ing addition fraction of ZrC nano-powder. When the added fraction of ZrC nano-powder was 1 wt.%, the hardness 1/2 ZrC nano-powder and fracture toughness reached their maximum values of about 2800 HV10 and 8.1 MPa·m , respectively. Spark plasma sintering When the added fraction of ZrC nano-powder was 3 wt.%, the flexural strength presented its maximum value Mechanical properties of about 1100 MPa. © 2014 Elsevier Ltd. All rights reserved.

1. Introduction it is almost impossible to sinter pure WC powder to dense materials with uniform grain size through conventional processes [11–13].Asa WC-based cemented carbides have been regarded as a crucial mate- result, such materials possess much lower flexural strength and fracture rial underpinning the world's tooling industries because of their wide toughness than WC–Co cemented carbides. applications in various fields, such as wood machining, metal cutting, With the development of field assisted sintering technique, especial- electronics industry, dentistry and rock drilling [1–3]. Normally, WC- ly spark plasma sintering (SPS), and the employment of nano-sized based cemented carbides consist of hard phase WC bonded together powder, the densification of binderless cemented carbides at relatively with metallic binders, in which cobalt has been the most commonly low sintering temperature has been possible [13–16]. It was stated used one because of its excellent wetting ability to WC and the out- briefly that pure WC powder could be consolidated by SPS at above standing mechanical properties of WC–Co cemented carbides [4,5]. 1900 °C in Omori's work [13]. But the WC grain was enlarged by However, WC–Co cemented carbides are sensitive to corrosive and/or 1 mm because of the long heating time (20–40 min). In Ref. [16],it oxidative environments because metal Co could be attacked easily in was reported that pure WC powder with a grain size in the range of such environments [6]. For example, oxygen-rich water at moderately 40–70 nm could be consolidated to near theoretical densities (99.1%) elevated temperatures is strong enough to erode the Co binder phase by SPS at 1750 °C without holding time under a uniaxial pressure of in cemented carbides [6]. In order to improve the corrosive and oxida- 126 MPa. But the decarburization of WC, which would form a W2C tive resistance of WC-based cemented carbides without consuming phase, was observed when fine initial WC powder was used. By adding their wear resistance, one effective route is to reduce the applied extra carbon, the formation of the W2C phase could be suppressed, but amount of the metallic binder even to nearly zero, in other words, to the abnormal growth of WC grains became serious [11,12,16]. In other fabricate binderless cemented carbides [6–10]. Indeed, such binderless words, although pure WC powder could be sintered to dense materials cemented carbides have shown higher corrosive and oxidative wear re- by SPS, the sintering temperature is still too high and the abnormal sistance than conventional WC–Co cemented carbides [6–10].However, growth of WC grains is usually serious. because of the high melting point of WC and absence of metallic binder, Recently, considerable studies have focused on the possibility of re- placing metallic binders by ceramic additives, so as to fabricate dense WC materials at a relatively low sintering temperature, thus improving ⁎ Corresponding author at: School of Engineering and Technology, China University of fl – Geosciences at Beijing, Beijing 100083, PR China. the materials' fracture toughness and exural strength [17 20].By E-mail address: [email protected] (Z. Peng). using TiC as an additive, nearly fully dense WC composites with a

http://dx.doi.org/10.1016/j.ijrmhm.2014.10.013 0263-4368/© 2014 Elsevier Ltd. All rights reserved. X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407 399 relative density up to 98.5% were obtained at 1400–1500 °C through Table 2 high-frequency induction heating sintering under a pressure of Basic parameters of ZrC nano-powder used in the present study. 60 MPa because of the formation of WC-TiC solid solution, but the Technique Color Purity (%) Average particle Specific surface Vickers hardness and fracture toughness of the samples decreased specification size (nm) area (m2/g) with increasing TiC content due to the existence of carbon in the grain ZrC nano-powder Black N99.0 20 11 boundaries [17]. In the work of Malek et al. [18], it was reported that the fracture toughness and flexural strength increased from 4.9 MPa·m1/2 and 570 MPa for pure WC to 6.1 MPa·m1/2 and consists of four stages: pre-heating, heating, transitional and holding

1536 MPa for WC–10 vol.% ZrO2 composites. By properly adding VC or stages, which is shown in Fig. 1. During the pre-heating stage, the tem- Cr3C2, the abnormal growth of WC grains could be suppressed and the perature reached 600 °C in 3 min from room temperature under a uni- fracture toughness of the resultant samples would be slightly improved axial pressure of 30 MPa. After this stage, the samples were heated up to [19]. After appropriately adding AlN nano-powder, the relative densi- 1550 °C in 9 min under a pressure of 50 MPa. However, between the ties, hardness and flexural strength of WC materials fabricated by SPS heating and holding stages, a short transitional stage was designed, dur- at 1600 °C under a pressure of 50 MPa increased although the fracture ing which the temperature was raised up from 1550 to 1600 °C in 1 min toughness would slightly decrease [20]. In summary, although the me- to ensure a smooth temperature transition into the holding stage. Dur- chanical strength of binderless WC-based cemented carbides might be ing the holding stage, the temperature was kept at 1600 °C under a improved by adding ceramic additives, such materials still possess sig- pressure of 50 MPa for 5 min. After sintering, the finally obtained spec- nificantly lower fracture toughness (about 5–6MPa·m1/2)thanCo- imens were disks of approximately 20 mm in diameter and 5 mm in containing cemented carbides (approximately 12 MPa·m1/2). height. Then the specimens were ground mechanically and cut into In literature, zirconium carbide (ZrC) is characteristic of high melting rectangular bars with a dimension of about 2 mm × 3 mm × 15 mm point (about 3420 °C), hardness (approximately 25.5 GPa), electrical by electrical discharge machining. After that, the bars were polished conductivity (78 × 10−6 Ω cm) and modulus of elasticity (about with diamond papers to mirror surface. 440 MPa) [21], being an excellent structural material itself. As an additive to WC–Co cemented carbides, ZrC powder can suppress the growth 2.2. Materials characterization of WC grains and improve the mechanical properties of the samples. However, there is no report on the addition effect of ZrC nano-powder The shrinkage displacement of the samples was measured by re- on binderless WC materials. cording the displacement of the lower punch of the SPS apparatus at In the present work, ultrafine binderless tungsten carbides with var- the heating, transitional and holding stages during sintering. The appar- ied amounts of ZrC nano-powder were fabricated by SPS at relatively ent density of the obtained samples was measured by Archimedes' low temperature (about 1600 °C) under a uniaxial pressure of method according to the international standard (ISO 18754). The theo- 50 MPa. The effect of ZrC nano-powder addition on the sintering behav- retical density for the designed samples with various compositions was ior, phase composition, microstructure and mechanical properties of calculated following the rule of mixtures, supposing that after sintering binderless WC cemented carbides was investigated. WC kept its state but ZrC reacted completely into ZrO2, and adopting the 3 theoretical densities of WC and ZrO2 as 15.7 and 5.8 g/cm ,respectively. 2. Experimental procedure The relative density of the obtained samples was defined as the percent- age of the apparent density to their corresponding theoretical density.

2.1. Sample preparation The sample hardness (HV10) was evaluated on a Vickers hardness tester (Model: VH-5, China) with a load of 98 N (10 kgf) and a dwell time of The raw WC powder used in this work was supplied by Zigong 15 s. The flexural strength of the samples was measured at room tem- Cemented Carbide Co. Ltd., Sichuan Province, China, and its basic pa- perature by a three-point bending method on an EHF-EG 50 kN tester rameters are presented in Table 1. The ZrC nano-powder was bought with a loading rate of 0.5 mm/min and a span length of 10 mm. The frac- 1/2 from Kaier Nanometer Energy & Technology Co. Ltd., Hefei City, China, ture toughness (KIC,MPa·m ) was calculated by measuring the crack and its basic parameters are listed in Table 2. The nominal compositions length near the indent resulting from the hardness test, according to of the designed samples with different amounts of ZrC nano-powder are the Shetty equation [22]: given in Table 3. During processing, the raw powders were mixed to- sffiffiffiffiffiffiffiffiffiffiffiffi gether according to the designed compositions in absolute alcohol, pffiffiffiffiffiffiffi P and then dispersed by ultrasonic vibration for 15 min to reduce the ag- K ¼ 0:0889 HV ð1Þ IC Σ4 l glomeration of ZrC nano-powder. After that, the slurries were milled for i¼1 i 44 h in a high energy attrition mill (Model: SY-1, China) with YG-6 cemented carbides balls (ISO: K20) as grinding media. During milling, where HV is the Vickers hardness (N/mm2) of the samples, P the load the ball-to-powder mass ratio was 10:1, and the rotational speed was (N) during indentation, and li the length of the ith crack emerging 200 rpm. After milling, the slurries were dried in a vacuum oven at from the hardness indent (μm). about 60 °C under a pressure of 0.01 MPa. After drying, the powder In addition, the phase identification was performed by X-ray chunks were crushed and sieved into fine powders. The resultant com- diffraction (XRD, Model: D8-Advance, Germany) with Cu Kα radiation posite powders were then sintered by a SPS apparatus (Model: SPS- (λ = 1.54178 Å) through a continuous scanning mode at a speed of 1050 T, Japan) under a vacuum of 6 Pa. For the applied SPS system, de- tails can be found in our previous work [20]. The sintering process Table 3 Nominal compositions of the designed samples.

Table 1 Samples No. WC (wt.%) ZrC (wt.%) Basic parameters of WC powder used in the present study. WC–0ZrC 100 0 Grade Grain size Speciﬁc Purity Chemical (%) WC–1ZrC 99 1 (nm) surface area (%) WC–3ZrC 97 3 Total Free Combined Oxygen (m2/g) WC–5ZrC 95 5 carbon carbon carbon WC–7ZrC 93 7 WC04 600 1.85 N99.5 6.22 0.069 6.15 0.23 WC–9ZrC 91 9 400 X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407

shrinkage displacement of the samples could reach its maximum value at no more than 1400 °C, demonstrating that it is possible to de- velop fairly dense binderless WC cemented carbides with ZrC nano- powder via SPS at 1400 °C under the applied uniaxial pressure of 50 MPa. This phenomenon would be generally attributed to the fine initial WC powder and unique joule heating of SPS. As is well known, the process of SPS is assumed to generate sparks between the sintered particles and clean the surface oxide of the particles to some extent, resulting in surface activation of the particles and overheating at the inter-particle boundaries, so an overshoot to the melting point can be expected on the particle surfaces, leading to localized melting [13–16]. In addition, as reported in Refs. [23,24], a small amount of contaminations (mainly Co) might also be brought into the composite powders during the high energy milling from the applied Co-rich cemented carbide milling balls, contributing to the sintering to a certain extent. After the sintering temperature raised higher than about 1400 °C, the further shrinkage of the samples was negligible, implying that the densification of the samples almost reached their maximum below 1400 °C. On the Fig. 1. Schematic representation of the temperature and pressure varied with the heating contrary, the thermal expansion of the samples and/or the punches time during SPS. were still existing, resulting in the shrinkage displacement decreased with the further increase in sintering temperature as shown in Fig. 2, 5°/min. The contact state of the WC and ZrC nano-particles in the raw which is consistent with the observation on the thermal expansion of powders was examined by transmission electron microscopy (TEM, the samples and/or the punches by Cha et al. [12]. Model: JEOL JEM-2100). An energy dispersive X-ray spectroscopy Fig. 2b shows the magnified view of the section A in Fig. 2a. It can be (EDX) analyzer attached to the TEM was used to measure the chemical seen that when the added fraction of ZrC nano-powder was lower than compositions of the particles. The microstructure examination on the 3 wt.%, the addition of ZrC nano-powder would not impose a significant polished and fresh fractured surfaces of the samples was carried out effect on the shrinkage behavior of the samples. With further increasing using a scanning electron microscope (SEM, Model: LEO-1530, addition of ZrC nano-powder (higher than 5 wt.%), the maximum Germany). And the elemental analysis of the samples was performed shrinkage displacement of the samples was obtained at about 1350 °C, by EDX spectroscopy attached to the SEM. which was about 50 °C lower than that of the sample without ZrC nano-powder. This phenomenon might be attributed to the high specif- 3. Results and discussion ic surface and lower sintering activation energy of the applied ZrC nano- powder. 3.1. Densification behavior The apparent and relative densities of the samples as a function of the addition fraction of ZrC nano-powder are presented in Fig. 3.Itcan From the heating to holding stages of SPS, the variations in the be seen that the apparent density decreased gradually with increasing shrinkage displacement and temperature with heating time for the as- addition fraction of ZrC nano-powder, which mainly resulted from the prepared WC-based cemented carbides with different amounts of ZrC much lower densities of the added ZrC (about 6.73 g/cm3)anditsreac- 3 nano-powder were recorded, which could reflect the densification be- tion product of ZrO2 during sintering (approximately 5.8 g/cm ) than havior of the samples under the applied sintering conditions. The results that of the matrix WC (about 15.7 g/cm3). The relative density, however, are illustrated in Fig. 2.Throughthisfigure, four stages can be visualized increased initially and then decreased with the increase in the addition in the shrinkage process: gentle shrinkage stage (at approximately 600- amount of ZrC nano-powder. When the added fraction of ZrC nano- 900 °C), rapid shrinkage stage (at about 900–1400 °C), expansion stage powder was about 3 wt.%, the relative density of the samples reached (at about 1400–1600 °C) and constant stage (at roughly 1600 °C). For its maximum of almost 100%, indicating that it is possible to improve all the samples, the rapid shrinkage started at about 900 °C and the the densification of binderless WC cemented carbides by adding the

Fig. 2. (a) Variations of the shrinkage displacement and temperature of the as-prepared WC-based cemented carbides with heating time from the heating to holding stages of SPS, and (b) the magniﬁed view of the section A in panel a. X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407 401

I WC WC WC WC I WC WC WC

(a) P (b)

I ZrC Nano-ZrC I

WC WC WC Fig. 3. Apparent and relative densities of the as-prepared WC-based cemented carbides as I a function of the added fraction of ZrC nano-powder. I WC appropriate amount of ZrC nano-powder. But it should be noted that the calculated relative density might be somewhat higher than its real P WC value, because for simplicity, its calculation did not consider the produc- (c) (d) 3 tion of W2C (17.1 g/cm ) from WC and the incomplete reaction of ZrC Nano-ZrC into ZrO2, both of which will be further discussed in the next section. Anyway, when the added fraction of ZrC nano-powder was lower than I 3wt.%,thesampledensification could be improved. And it might be at- WC WC I tributed to the follow reasons. Because the mean size of the applied WC particles (600 nm) is much larger than that of the used ZrC nano- WC WC powder (20 nm), it is reasonable to assume that the WC particles may WC I be surrounded by ZrC nano-particles after the mixing by ultrasonic vi- I ZrC bration dispersion and long-time ball-milling. In order to confirm this WC assumption, TEM micrographs on the ball-milled powders were taken Nano-ZrC to reveal the contact state of WC particles and ZrC nano-particles after P the applied mixing, and the corresponding schematic diagrams were (e) (f) drawn in Fig. 4. It can be seen that without ZrC nano-powder, the WC particles tended to gather together to form a large particle (see in WC WC particle ZrC nano-particle Fig. 4b). However, with 3 wt.% ZrC nano-powder added, most of the nano-particles tended to adhere to the submicron WC particles, as seen in Fig. 4d. The adherent nano-powder tends to suppress the Fig. 4. Schematic diagrams and corresponding TEM micrographs revealing the contact boundary diffusion of WC grains and to fill the inter-space produced state of WC and ZrC nano-particles during sintering: (a) and (b) sample without by particle rearrangement and plastic deformation of WC grains with ZrC nano-powder, (c) and (d) samples with 3 wt.% ZrC nano-powder, and (e) and the uniaxial pressure during SPS. However, nano-powders usually (f) samples with 9 wt.% ZrC nano-powder. have a tendency to form agglomerates due to their high specificsurface area. With further increasing addition of ZrC nano-powder into the sam- case of WC as seen from the W–C binary phase diagram [12,25], with ples, for instance, in the sample with 9 wt.% ZrC nano-powder as shown the help of the residual oxygen in the chamber and carried oxygen in Fig. 4f, the agglomeration of ZrC nano-powder became more and with the raw powder, the decarburization of WC would easily happen, more serious, resulting in decreased relative density of the samples thus resulting in the W C phase. The following reactions for the forma- after sintering. 2 tion of W2C may occur: (1) the direct decarburization of dissolved WC: 2WC → W2C + C; (2) the direct reaction between WC and O2: 3.2. Phase composition and microstructure 2WC + O2 → W2C+CO2; and (3) the indirect reaction between dissolved WC and ZrC and O2:2WC+ZrC+3O2 → W2C+ZrO2 +2CO2. Fig. 5 represents the XRD patterns of the initial WC and WC–9ZrC Moreover, with the increase in the added fraction of ZrC nano- composite powders used in this work, and the obtained binderless powder, the intensity of W2C phase increased at first and then de- WC-based cemented carbides, respectively. In the initial pure WC pow- creased. When the added fraction was higher than 3 wt.%, both W2C der, only WC phase (JCPDS PDF card: 65-8828) was detected. Further- phase and Co2W4C phase could not be identified in the detection limit more, in the initial WC–9ZrC composite powder, both WC and ZrC of XRD, but a new phase ZrO2 (JCPDS PDF card: 65-1025) was detected. phases (JCPDS PDF card: 65-8835) were detected. After sintering, how- Furthermore, in the detection limit of XRD, no ZrC phase was detected in ever, besides the WC phase, a small amount of W2C phase (JCPDS PDF the obtained WC–ZrC cemented carbides either. Because of the high card: 65-3896) and Co2W4C phase (JCPDS card: 06-0611) could be de- specific surface of ZrC nano-powder, with the increase in the added frac- tected in the samples without or with no more than 3 wt.% of ZrC tion of ZrC nano-powder, more and more oxygen would be brought into nano-powder. The Co2W4C phase should stem from the small amount the sintering system. According to the above-mentioned Reactions of contamination (mainly Co) during high energy milling from the Co- (2) and (3), more and more W2C phase would be formed, resulting in rich cemented carbide milling balls, the only Co source in the material the increased intensity of W2Catfirst. However, with further increasing system, and such phenomenon was also reported in Refs. [23,24]. Be- added fraction of ZrC nano-powder (more than 3 wt.%), more and more cause the stoichiometric range of carbon content is very narrow in the ZrC was dissolved in the contaminations (mainly Co), restraining the 402 X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407

and abnormal growth of WC grains could be observed very obviously in the sample without ZrC nano-powder as shown in Fig. 7a, and many prismatic WC grains existed in the sample. The measured longest intercept length of the prismatic WC grain was almost 2.5 μm, which was much longer than the diameter of the normal WC grains. The sharp edge of these prismatic grains could lead to tensile stress concentra- tions, resultantly promoting the initiation and propagation of cracks [26]. Under loading, the sharp edges were easily broken as shown in Fig. 8, thus reducing the fracture toughness of the samples. In addition, ﬁner microstructure can always result in better mechanical properties. Therefore, the abnormal growth of WC grains may also degrade the mechanical properties of the as-prepared cemented carbides. Generally, prismatic WC grains are formed through solution/re- precipitation during liquid sintering due to the difference in surface energy between different surface planes. However, because the sintering temperature (about 1600 °C) in this work was much lower than the melting point of WC (about 2785 °C), no liquid phase would form during sintering. So, solution/re-precipitation is not responsible for the formation of prismatic WC grains in the sample from pure WC powder. On the other hand, considering that the abnormal growth of WC grains was also observed in other authors' work during the fabrication of binderless WC materials by SPS [12,16], we suppose that this phenomenon may be attributed to the unique heating model of SPS. The generated sparks between the WC particles during SPS would cause overheating at the inter-particle boundaries, leading to localized melting of the particles. Moreover, some of the WC particles would agglomerate together with particle rearrangement and plastic deformation under the applied high uniaxial pressure. So, the grain boundary diffusion among the agglomerated WC particles would happen with the help of localized melting of the particles, facilitating the growth of WC grains. However, due to the difference in surface energy between the higher energy surface plane and lower surface energy one as reported in Ref. [26],thegrowth of WC grains along different surface planes was uneven (abnormal), thus resulting in the prismatic WC grains. In addition, because the liquid layer owing to localized melting was very thin and non-uniform in the sample, some pores between the agglomerated WC particles cannot be extruded out during the grain growth for the rapid sintering process of SPS. Therefore, some inner-pores existed in the grains as seen in – Fig. 5. XRD patterns of (a) the initial WC and WC 9ZrC composite powders used in this Fig. 8, which would cause and extend the cracks during the application work, and (b) the WC–χ wt.% ZrC cemented carbides obtained in this work. of cemented carbides. With an appropriate addition amount of ZrC nano-powder, the mi- dissolution of WC and most of the oxygen was reacted with crostructure of the as-prepared WC based cemented carbides can be im- the dissolved ZrC to form ZrO2 through the reaction: ZrC + 2O2 → proved, as seen in Fig. 7. For example, the average intercept length of the ZrO2 +CO2. On the other hand, because ZrC nano-powder can be oxi- WC grains was only about 700 nm in the sample with 1 wt.% ZrC nano- dized much easier than WC micron powder, less of WC would be oxi- powder, which was much shorter than that of the WC grains (about dized with more ZrC in the material system. Therefore, both W2C 1.3 μm) in the sample without ZrC nano-powder. The amount of pris- phase and Co2W4C phase could not be detected in the detection limit matic WC grains in the sample with ZrC nano-powder was also much of XRD but ZrO2 phase was detected. less than that in the sample from pure WC powder. The reduction of ab- In addition, as seen in Figs. 3 and 5, the higher detected intensity of normally grown grains might be attributed to the following: the added

W2C, the higher sintered density of the samples, suggesting that the for- ZrC nano-powder could adhere to WC grains during the preparation of mation of W2C would be related to the sintered density, which is similar WC–ZrC composite powder as indicated in Fig. 4d, and then during with the observations in Refs. [11,16]. sintering, the adherent ZrC nano-powder would share the load of the Fig. 6 shows typical SEM micrographs on the polished surfaces of the applied uniaxial pressure, relieve the plastic deformation of WC grains as-prepared binderless WC-based cemented carbides. It can be seen and impede the grain boundary diffusion during the localized melting that the amount and size of the micro-pores on the polished surface of of the WC grain surface. As a result, the abnormal growth of WC grains the samples initially decreased and then increased with increasing addi- was suppressed, and the microstructure became finer and more uni- tion amount of ZrC nano-powder, which is consistent with the variation form after an appropriate amount of ZrC nano-powder was added into trend of the relative density of the samples as seen in Fig. 3. In addition, the samples. However, because nano-powders always agglomerate eas- the amount of broken grains on the polished surface of the samples ily, when the added fraction of ZrC nano-powder was higher than were also decreased first and then increased under the same polishing 5 wt.%, the agglomeration of ZrC nano-powder became more and condition. When the added fraction of ZrC nano-powder was about 1– more serious, which can be seen from Fig. 7(d–f). In order to understand 3 wt.%, the amount of broken grains on the surface reached its mini- the microstructure of such samples, one typical SEM micrograph on the mum value, indicating a good grain boundary strength between WC fractured surface of the sample with 9 wt.% ZrC nano-powder is present- grains. ed in Fig. 9a. From this figure, two kinds of particles could be typically Fig. 7 displays typical SEM micrographs on the fractured surfaces of observed on the fractured surface of the sample: non-uniform distribut- the obtained binderless WC-based cemented carbides. The coalescence ed light particles and angular dark ones. In order to figure out their X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407 403

(a) (b)

(e) (f)

Fig. 6. Typical SEM micrographs on the polished surfaces of the as-prepared WC–χ wt.% ZrC cemented carbides, where χ is (a) 0, (b) 1, (c) 3, (d) 5, (e) 7 and (f) 9, respectively. compositions, EDX spectra were collected on different particles, in the sample microstructures in this work, the appropriate addition frac- which spectrums 1 and 3 were recorded on the dark particles, and spec- tion of ZrC nano-powder was supposed as 1–3wt.%. trums 2 and 4 were recorded on the lightones, respectively, as marked in Fig. 9a. The recorded spectra are shown in Fig. 9(b–e), and their cor- 3.3. Mechanical properties responding calculated constituents are listed in Table 4. The results indicated that the content of Zr atoms in the light particles was much higher Fig. 10 presents the Vickers hardness of the as-prepared WC–χ wt.% than that in the dark ones, demonstrating that the light particles were Zr ZrC cemented carbides as a function of the added fraction of ZrC nano- rich ones. On the contrary, the content of W atoms in the light particles powder. The Vickers hardness initially increased and then decreased was much lower than that in the dark ones, implying that the dark par- with increasing addition fraction of ZrC nano-powder. When the ticles were W rich ones. In combination with the feature of the initial added fraction of ZrC nano-powder was 1 wt.%, the hardness of the powders, it could be supposed that, the dark particles were of WC samples reached its maximum of about 2800 HV10, which is much grains, and the light particles were formed by the agglomeration of higher than those of the reported WC-based cemented carbide samples ZrC nano-powder. Thus, it can be safely concluded that with increasing with metal binders in literature [2], because the hardness of metal addition fraction of ZrC nano-powder, the amount of light particles in- binders are much lower than that of WC. Moreover, the increase in creased, that is, the agglomeration of ZrC became more and more seri- hardness of the samples with the addition fraction of ZrC nano- ous. When excessive ZrC nano-powder was added, it would reduce powder could be attributed to the enhancement in sample densiﬁcation the sample densiﬁcation, increase the number of defects (such as (relative density), reduction of WC grain size, and suppression of the ab- pores and the agglomerated ZrC) and thus degrade the mechanical normal growth of WC grains as mentioned in Sections 3.1 and 3.2.How- properties of the cemented carbides. And from the optimization on ever, with further increase in the added fraction of ZrC nano-powder, 404 X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407

(a) (b)

Abnormal grains

(e) (f)

Fig. 7. Typical SEM micrographs on the fractured surfaces of the as-prepared WC–χ wt.% ZrC cemented carbides, where χ is (a) 0, (b) 1, (c) 3, (d) 5, (e) 7 and (f) 9, respectively.

the agglomeration of ZrC or ZrO2 became more and more serious. The relative density of the sample decreased and the number of defects in- Inner-pores creased, resulting in a rapid decrease in the Vickers hardness of the as- prepared cemented carbides. The flexural strength of the as-prepared WC–χ wt.% ZrC cemented Micro-pores carbides as a function of the added fraction of ZrC nano-powder is illustrated in Fig. 11. It can be seen from this figure that the flexural strength Inner-pores does not change dramatically, which slightly increases first and then de- creases somewhat. It presented a maximum value of 1100 MPa when 3 wt.% ZrC was added into the sample. The increase in sample densification and reduction of the abnormal growth of WC grains as presented in Sections 3.1 and 3.2 may also be the main reasons responsible for the increase in flexural strength of samples. However, with further addition of Fragmentation of sharp edges ZrC nano-powder, the inhomogeneous dispersion of ZrC nano-powder became more and more serious and the number of defects in the sample increased, resulting in the decrease in flexural strength. The addition effect of ZrC nano-powder on the fracture toughness of – Fig. 8. Typical magnified SEM micrograph on the fractured surface of pure WC cemented the as-prepared WC ZrC cemented carbides is shown in Fig. 12.The carbides. fracture toughness of the samples first increased and then decreased X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407 405

(a)

Dark particles

Light particles

(b) (c)

(d) (e)

Fig. 9. (a) Typical SEM micrograph on the fractured surface of the sample WC–9 wt.% ZrC, and EDX spectra corresponding to the spots marked on the micrograph: (b) spectrum 1, (c) spectrum 2, (d) spectrum 3 and (e) spectrum 4. with increasing addition fraction of ZrC nano-powder, reaching its maximum value of about 8.1 MPa·m1/2 when 1 wt.% ZrC nano-powder was added. The increase of the fracture toughness at ﬁrst is mainly attributed to the improved sample densiﬁcation as discussed in Sections 3.1 and 3.2 and the more homogeneous microstructure of the WC–1 wt.% ZrC cemented carbides as shown in Fig. 7b after the addition of ZrC nano- powder. However, with further addition of ZrC nano-powder, the agglomeration of the added ZrC became more and more serious as seen in Figs. 7 and 9, the relative density of the samples decreased, and the

Table 4 The constituents of different spots on the fracture surface of the sample WC–9wt.%ZrC,as shown in Fig. 9.

Elements Spectrum 1 Spectrum 2 Spectrum 3 Spectrum 4

C (at.%) 36.92 21.42 44.57 17.36 Zr (at.%) 1.76 54.39 1.91 64.21 Fig. 10. Vickers hardness of the as-prepared WC–χ wt.% ZrC cemented carbides as a func- W (at.%) 61.32 24.19 53.52 18.43 tion of the added fraction of ZrC nano-powder. 406 X. Ren et al. / Int. Journal of Refractory Metals and Hard Materials 48 (2015) 398–407

of mechanical properties were observed when VC was added into binderless WC cemented carbides, and when the added fraction of VC was 1 wt.%, the hardness and fracture toughness of WC–VC cemented carbides reached their maximum values [27]. With increasing addition

fraction of Mo2C, the hardness of WC–Mo2C cemented carbides decreased and the fracture toughness increased [28].FromTable 5,itcan be seen that both the Vickers hardness and fracture toughness of the obtained binderless WC–1 wt.% ZrC cemented carbides were higher than those of the reported WC–20 at.% TiC, WC–1 wt.% VC and WC–1 wt.%

Mo2C cemented carbides. The enhanced hardness and fracture toughness of the as-prepared samples in this work may be resulted from the increased relative density, the reﬁnement of the microstructure and the suppression of the abnormal growth of WC grains. Therefore, it can be concluded that ZrC nano-powder was a very effective additive to improve the mechanical properties of binderless cemented carbides.

4. Conclusions

Fig. 11. Flexural strength of the as-prepared WC–χ wt.% ZrC cemented carbides as a func- Using spark plasma sintering, binderless WC cemented carbides tion of the added fraction of ZrC nano-powder. with varied addition fraction of ZrC nano-powder were prepared at 1600 °C under a pressure of 50 MPa. The rapid shrinkage of the as- prepared samples started at approximately 900 °C and ended at about 1400 °C. The relative density of the samples had a maximum of almost

100% when 1–3 wt.% ZrC nano-powder was added. W2C phases were detected in the samples without or with no more than 3 wt.% of ZrC nano-powder. In the sample without ZrC nano-powder, the coarsening and abnormal growth of WC grains were obviously serious, and by adding the appropriate amount of ZrC nano-powder, the grain growth was suppressed, resulting in more homogeneous microstructure. With further addition of ZrC nano-powder, the agglomeration of ZrC nano- powder would become more and more serious. The Vickers hardness, flexural strength and fracture toughness of the as-prepared samples first increased and then decreased with the increase in the addition fraction of ZrC nano-powder. When the added fraction of ZrC nano-powder was 1 wt.%, the hardness and fracture toughness of the cemented carbides reached their maximums of about 1/2 2800 HV10 and 8.1 MPa·m , respectively. When the added fraction of ZrC nano-powder was 3 wt.%, the flexural strength presented its maximum of about 1100 MPa.

Acknowledgments Fig. 12. Fracture toughness of the as-prepared WC–χ wt.% ZrC cemented carbides as a function of the added fraction of ZrC nano-powder. This work was supported by Grand Survey on Land and Nature Sources of China sponsored by China Geological Survey (grant no. number of defects in the samples increased, all of which would result in 1212010916026), the Excellent Adviser Foundation in China University a decrease in the fracture toughness of the samples. of Geosciences from the Fundamental Research Funds for the Central In order to further help the readers to understand the addition effect Universities, and National Key Technology R & D Program of China of ZrC nano-powder on the mechanical properties of the obtained (grant no. 2011BAB03B08). binderless WC cemented carbide, a comparison on Vickers hardness and fraction toughness between the obtained WC–1 wt.% ZrC samples References in this work and the binderless WC-based cemented carbides with ﬁ other carbides (TiC, VC and Mo C) reported in literature was presented [1] L.J. Prakash, Application of ne grained tungsten carbide based cemented carbides, 2 Int. J. Refract. Met. Hard Mater. 13 (1995) 257–264. in Table 5. It was reported that both the hardness and fracture tough- [2] X.Y. Ren, H.Z. Miao, Z.J. Peng, A review of cemented carbides for rock drilling: an old ness of binderless WC–TiC cemented carbides decreased with increas- but still tough challenge in geo-engineering, Int. J. Refract. Met. Hard Mater. 39 ing addition fraction of TiC, and when the added fraction of TiC was (2013) 61–77. [3] G. Gille, B. Szesny, K. Dreyer, H. Van Den Berg, J. Schmidt, T. Gestrich, G. Leitner, Sub- 20 at.%, the hardness and fracture toughness of such samples reached micron and ultraﬁne grained hardmetals for microdrills and metal cutting inserts, their maximum values as shown in Table 5 [17]. Similar variation trends Int. J. Refract. Met. Hard Mater. 20 (2002) 3–22.

Table 5 Comparison on Vickers hardness and fracture toughness of WC–1 wt.% ZrC cemented carbides obtained in this work with the reported values in literature.

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