Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides

Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides

RMHM-04250; No of Pages 7 Int. Journal of Refractory Metals and Hard Materials xxx (2016) xxx–xxx Contents lists available at ScienceDirect Int. Jou...

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RMHM-04250; No of Pages 7 Int. Journal of Refractory Metals and Hard Materials xxx (2016) xxx–xxx

Contents lists available at ScienceDirect

Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM

Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides J. Weidow a,⁎, E. Halwax b, W.-D. Schubert b a b

Chalmers University of Technology, 41296 Göteborg, Sweden Vienna University of Technology, 1060 Vienna, Austria

a r t i c l e

i n f o

Article history: Received 30 March 2016 Received in revised form 13 May 2016 Accepted 22 May 2016 Available online xxxx Keywords: EBSD Electron backscatter diffraction Doping SEM Scanning electron microscopy

a b s t r a c t Powders of hexagonal (W,Ta)C were produced following a two-step carburization process with (W,Ta)2C powder as an intermediate product. XRD measurements indicate that higher temperatures during the first carburization step increase the fraction of Ta in the (W,Ta)2C structure where lower temperatures during the second carburization step seems to increase the fraction of Ta in the (W,Ta)C structure. Lower temperatures during the carburization steps increase the fraction of Σ2 WC/WC grain boundaries and cause the formation of what is interpreted as Σ4 boundaries. The (W,Ta)C powders can be successfully used to produce fine grained WC-Co based cemented carbides. The pre-alloying with Ta appears to have a softening effect on the material. © 2016 Elsevier Ltd. All rights reserved.

1. Introduction Cubic carbides such as TiC, NbC and TaC are frequently added to the powder mixture in the production of cemented carbides in order to improve the as-sintered material properties. Sintering of WC, Co and TaC powder mixtures causes Ta to partially substitute for W in the structure for reprecipitated parts of WC grains [1]. As an example, for a WC-TaC-Co material with 16.20 mass% Ta and 5.51 mass% Co sintered at 1430 °C, the Ta solubility in WC was found to be 0.959 ± 0.009 at.% [2]. Further increasing the amount of Ta in the WC structure is advantageous as it would provide an additional degree of freedom in the cemented carbide production process which possibly would result in improved material properties. One method to achieve a WC-Co based cemented carbide with increased Ta content in the WC structure would be to manipulate the WC powder so that it contains Ta from the very beginning, i.e. a (W,Ta)C powder. As metal diffusion is not expected to occur at normal sintering temperatures (1400–1450 °C), undissolved parts of WC grains should maintain the high concentration of Ta [3]. The phase diagrams for the W-Ta-C system show that the solubility of Ta in W2C is about 12 at.% at 1500 °C, and about 15% at 1750 °C, while the solubility of Ta in WC is known to be much smaller [4]. It has been shown that by performing a two-step carburization of a W + Ta powder mixture, with (W,Ta)2C as an intermediate product, it is possible to produce a (W,Ta)C powder with a Ta/(Ta + W) ratio of 0.086 [5]. This means ⁎ Corresponding author. E-mail address: [email protected] (J. Weidow).

that the concentration of Ta in the (W,Ta)C powder is over four times higher than what has been measured for Ta in WC grains of an as-sintered cemented carbide produced from pure WC powder. This two-step procedure follows the guidelines published in patent application [6]. It is the aim of this study to investigate how the powder production parameters can be optimized in order to increase the amount of Ta being dissolved in the WC structure. It is also the aim to investigate what effects the Ta pre-alloying has on the microstructure and material properties of fine-grained cemented carbides. 2. Experimental Two powder series were produced. In the first powder series, 0.84 μm FSSS (as-supplied) W powder was carburized together with Ta2O5 at 1450 °C resulting in a (W,Ta)2C powder (X1). This powder was carburized a second time at various temperatures leading to the formation of (W,Ta)C powders P1–P3. As a reference to these powders, a mixture of the same W powder and Ta2O5 powder was directly carburized in a one-step process to give (W,Ta)C (P4). The second series was similar to the first. Here a 0.83 μm FSSS (as-supplied) W powder was used and the first carburization step was instead performed at 1630 °C. This resulted in (W,Ta)2C powder X2 and (W,Ta)C powders P5–P7. The amount of Ta2O5 added was based on the aim to have a WC powder with Ta corresponding to 3.0 wt% TaC. The powders were analyzed with X-ray diffraction (XRD) using a PANalytical X'Pert PRO powder diffractometer (CuKα radiation, X'Celerator detector, Ni Kβ filter). Before analysis, the powders

http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012 0263-4368/© 2016 Elsevier Ltd. All rights reserved.

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012

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Table 1 (W,Ta)2C powder details, silicon used as an internal standard.

Powder

Carburization temperature (°C)

Lattice parameter a (nm)

Lattice parameter c (nm)

W2C X1 X2

– 1450 1630

0.51910 0.51931 0.51989

0.47279 0.47295 0.47357

10.6° but it was noted that the width of the peak is much smaller and a larger acceptance would therefore include more randomly oriented grain boundaries. The values for the fraction of Σ2 WC/WC grain boundaries will therefore appear smaller than using the Brandon criterion. However, the differences between different powders should be approximately the same. 3. Results

were mixed with Si powder with well-known lattice parameter. In the quantification calculation of different phases, Si was removed. The lattice parameters of the (W,Ta)C phase were determined from Rietveld refinements using the TOPAS software having the lattice parameter of Si fixed. For the modeling of (W,Ta)C, the pure WC phase was used in the quantification of the phase fractions of the powder mixture. Cemented carbides were produced from the seven different WC powders (P1–P7). 270 g WC, 30 g Co and 6 g paraffin were weighed in and loaded into a ball mill together will 1800 g cemented carbide milling balls (d = 9.1 mm). Cyclohexane was used as milling liquid and the milling was performed for 72 h (P1–P4) or 24 h (P5–P7). The cyclohexane was removed through distillation and the milled powder was sieved and pressed to cylinder-shaped green bodies, d = 10 mm, with 2000 kg load. The green bodies were sintered in a vacuum furnace where the temperature was increased with 10 °C/min to 1250 °C, after which a 20 min hold was made. The temperature was then increased with 3 °C/min to 1420 °C where a 1 h hold was made. The cemented carbides were named C1–C7 after the corresponding (W,Ta)C powders P1–P7. In addition, powders P5–P7 were also used to produce the cemented carbides C8–C10. The only difference between the sintering of the different powders was that for some powders, an initial step performed in hydrogen atmosphere was necessary in order to get as-sintered materials with similar magnetic saturation relative to Co. Scanning electron microscope (SEM) imaging of the powders and the as-sintered materials were carried out using a FEI Quanta 200 FEG SEM and a Leo Ultra 55 FEG SEM. The latter instrument was equipped with a HKL electron backscatter diffraction (EBSD) system using the HKL Channel 5 software. The powders were embedded in copper by melting of the powder mixtures and Cu granulates. The Cu embedments were polished with diamond spray with the smallest particle size being 1 μm. The outer surface layers were cleaned by a 2 h 4.0 kV Ar + ion sputtering in a Gatan precision polishing system, model 691. The incident angle was 2° and both guns were used during this procedure. The EBSD analysis was performed at 20 kV, in high-current mode and using a 60 μm aperture with the specimen being tilted by 70° and the camera inserted in the SEM as far as possible. The step size was set to 50 nm. Data was filtered using noise reduction level 4 and correction for wild spikes, and at least four pixels were needed to define a grain. In order to define a WC/WC grain boundary that was of the Σ2 type, i.e. a 90° rotation around the [1010] axis, a deviance of 3° was accepted. The Brandon criterion [7] would lead to an acceptance of a deviance of

The XRD analyses of what was aimed as W2C powders, X1 and X2, showed no presence of TaC phase. However, some WC was also found and in the case of X1 also a minor amount of W was present. The lattice parameters of the (W,Ta)2C powders, X1 and X2, as measured by XRD, are shown in Table 1. Both the a and c lattice parameters for the powders are larger than for pure W2C with powder X2 having substantially larger lattice parameters than X1. The results from the XRD measurements of powders P1–P7 are shown in Table 2. In addition to hexagonal (W,Ta)C, minor amounts of cubic (Ta,W)C were also found. The content of cubic (Ta,W)C in the powders becomes higher between P1 and P3 and between P5 and P7. The content of (Ta,W)C in P5–P7 is lower than in P1–P3. The highest content of (Ta,W)C is found in P4. Regarding the (W,Ta)C phase, both the a and c lattice parameters for all powders are greater than for pure WC. The lattice parameters are greater for powders P5–P7 than for powders P1–P3. In an earlier study, we have compared the lattice parameters of (W,Ta)C of a strongly doped powder with the composition as measured by atom probe tomography [5]. If the Ta content scales linearly with lattice parameter a, an increase by 0.1 pm would correspond to an increase of the Ta content by 2 at.%. For the c lattice parameter, the corresponding value is 2.1 at.% per 0.1 pm. The Ta content in the hexagonal (W,Ta)C phase, calculated as the average composition due to the changes in lattice parameters for the different powders, are shown in Table 2. From earlier experiments, we estimate the uncertainty in the lattice parameter to be in the range of ±0.00005–0.0001 nm. For the calculation of the uncertainty as shown in Table 2, we have chosen the lower limit of the uncertainty. SEM micrographs of powders P1–P7 are shown in Fig. 1. Both (W,Ta)C grain size and particle size increases from P1 to P3 and from P5 to P7. Data from EBSD analyses of powders P1–P7 are shown in Table 3. In general, the WC grain size increases and the aspect ratio as well as the fraction of Σ2 WC/WC grain boundaries decreases from P1 to P3 and from P5 to P7, respectively. In general, the second group of powders exhibits greater WC grain sizes, smaller aspect ratios and a lower fraction of Σ2 WC/WC than the first group of powders. Fig. 2 shows the WC/WC grain boundary distribution in powder P1 together with a distribution for a material with random grain boundaries. Apart from the peak at 90° corresponding to Σ2 WC/WC grain boundaries, there is also a peak at small angles and a smaller peak at 60°. The latter is observed in P1–P2 and P4–P6, barely for P3 and not at all for P7. As-sintered material properties of materials C1–C10 are shown in Table 4. The relative magnetic saturation is in the range 0.86–0.88 for

Table 2 (W,Ta)C powder details.

Powder

Originating (W,Ta)2C powder

Carburization temperature (°C)

Particle size (μm)

Total C content (wt%)

Cubic (Ta,W)C (at%)

Lattice parameter a (nm)

Lattice parameter c (nm)

Calculated Ta content in (W,Ta)C (at.%)

WC P1 P2 P3 P4 P5 P6 P7

– X1 X1 X1 – X2 X2 X2

– 1450 1630 1950 1950 1450 1630 1950

– 0.90 1.04 1.42 1.12 1.24 1.32 2.27

6.21 6.24 6.20 6.11 6.15 6.18 6.22

– 1.65 1.74 1.75 2.67 0.80 0.91 1.46

0.29063 0.29070 0.29072 0.29080 0.29076 0.29084 0.29084 0.29087

0.28375 0.28388 0.28383 0.28383 0.28384 0.28394 0.28387 0.28388

– 2.1 ± 0.7 1.7 ± 0.7 1.6 ± 0.7 2.2 ± 0.7 4.1 ± 0.7 3.4 ± 0.7 3.8 ± 0.7

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012

J. Weidow et al. / Int. Journal of Refractory Metals and Hard Materials xxx (2016) xxx–xxx

materials C1–C5 and C9–C10, somewhat higher for materials C6 and C7 and somewhat lower for material C8. It was also noted that the coercivity increased and the hardness decreased for materials C1 to C3, for C5 to C7 and for C8 to C10. SEM micrographs of cemented carbides C1–C7 are shown in Fig. 3. The WC grain size appears to increase from C1 to C3 and from C5 to C7. Data from EBSD analysis of as-sintered materials C1–C5 and C8–C10 are shown in Table 5. In general, the WC grain size increases and the fraction of Σ2 WC/WC grain decreases for materials C1 and C2 to C3 and from C5 and C8 to C9 and C10.

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4. Discussion The presence of both WC and W in powder X1 shows the difficulty in the production of a homogenous powder mixture. The absence of TaC is supposed to mean that all of the Ta should substitute for W in any of the phases being formed, i.e. (W,Ta)2C, (W,Ta)C and (W,Ta). The larger lattice parameters for the W2C phases strongly indicate that Ta has in fact substituted for W in this phase. From a comparison with literature values, the lattice parameters indicate that the change in lattice parameters corresponds to homogeneous (W,Ta)2C mixed crystals with 5 (X1)

Fig. 1. SEM backscatter detector micrographs of Cu embedments of powder P1–P7.

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012

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Table 3 WC measurements of the powders from EBSD (equivalent circle diameter) and fraction of Ʃ2 WC/WC grain boundaries. Powder

Mean grain size (μm)

Grain size std. dev. (μm)

Aspect ratio

Analyzed grains (#)

Fraction Ʃ2 boundaries

P1 P2 P3 P4 P5 P6 P7

0.363 ± 0.010 0.351 ± 0.007 0.440 ± 0.010 0.325 ± 0.005 0.404 ± 0.008 0.435 ± 0.011 0.743 ± 0.025

0.189 0.175 0.236 0.164 0.235 0.248 0.425

1.61 1.66 1.60 1.71 1.70 1.55 1.46

1516 1647 1342 2430 2025 1249 739

0.422 0.380 0.322 0.335 0.373 0.394 0.195

and 8–11 (X2) mole% TaC, respectively. The margin is due to the uncertainty in the lattice parameter determination and the uncertainty in the measurements in the literature [8]. The difference between the two powders can be attributed to a higher temperature in the carburization process enabling a larger fraction of Ta to replace W in the W2C structure. The relatively small change in lattice parameters compared to the uncertainty of the measurement makes XRD rather inappropriate to quantify the composition of the (W,Ta)C phase. The only significant difference is that the Ta content in the (W,Ta)C phase is greater for powders P5–P7 than powders P1–P4. Instead, the concentration of cubic (Ta,W)C phase is considered to be a better method as for a known total amount of Ta it shows how much Ta must be present in the hexagonal (W,Ta)C phase. It must be remembered that we do not know the exact composition of the cubic (Ta,W)C phase for the different carburization temperatures. For a comparison, the solubility of W in TaC is approximately 18 at.% at 1700 °C [8]. Being aware of these difficulties, the fraction of cubic (Ta,W)C for the different powders supports the result that powders P5–P7 have a larger fraction of Ta in the (W,Ta)C phase than powders P1–P3. Comparing with the results from the (W,Ta)2C powders, we believe that a higher temperature in the first step of the two-step carburization process is beneficial for the Ta content in the (W,Ta)2C structure and later also in the (W,Ta)C structure.

Table 4 As-sintered material properties.

Material

Originating (W,Ta)C powder

Sa

Hc (kA/m)b

Hardness HV30 (GPa)

C1 C2 C3 C4 C5 C6 C7 C8 C9 C10

P1 P2 P3 P4 P5 P6 P7 P5 P6 P7

0.873 0.864 0.864 0.880 0.864 0.924 0.926 0.832 0.858 0.871

18.29 18.01 16.24 16.93 17.21 16.33 12.25 17.46 16.97 13.71

15.42 (σ = 0.08) 15.41 (σ = 0.13) 15.02 (σ = 0.07) 15.23 (σ = 0.07) 14.72 (σ = 0.08) 14.62 (σ = 0.09) 14.03 (σ = 0.06) 14.92 (σ = 0.07) 14.83 (σ = 0.09) 14.31 (σ = 0.03)

a b

Magnetic saturation relative to Co (by weight). Coercivity.

As the fraction of cubic (Ta,W)C significantly increases from P5 to P7, we also believe that the Ta content in the hexagonal (W,Ta)C grains decreases from P5 to P7. This is in agreement with our previous study showing that a lower temperature in the second step of a two-step carburization process is beneficial for the Ta content in the (W,Ta)C structure [5]. Finally, the content of cubic (Ta,W)C for powder P4 is much greater than for the other powders. This together with the small lattice parameters suggests that the Ta content of (W,Ta)C in P4 is low, which means that a one-step carburization process gives less Ta dissolved in WC than a two-step process. We thus conclude that the two-step method we have used to produce a WC powder pre-alloyed with Ta is in fact working. The EBSD analyses of the powders showed the big importance of the temperature in the second carburization step when it comes to the grain size. They also showed that the larger this temperature is, the smaller is the aspect ratio, which means the more round the grains are. It is well known that Σ2 WC/WC grain boundaries are formed already in the WC production process [9]. The relatively low temperature in the carburization process means that the enthalpy dominates the expression for

Fig. 2. WC/WC grain boundary distribution in powder P1 together with a distribution for a material with random grain boundaries (R).

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012

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Fig. 3. SEM backscatter detector micrographs of cemented carbides C1–C7.

the Gibbs free energy and it is therefore easy to understand that there is a high probability for the formation of the special low energy Σ2 grain boundaries. That the fraction of Σ2 WC/WC grains decreased from P1 to P3 and from P5 to P7 is in agreement with our previous study for strongly pre-alloyed powers [5]. A higher temperature causes the enthalpy to be less dominating. The presence of a peak in the WC/WC grain boundary misorientation profile at 60° is believed to correspond to a Σ4 grain boundary. Lay and Loubradou conclude that a 60° rotation

around the [1210] axis results when two WC grains with different Σ2 relationships to a third grain are in contact with each other [10]. With the same argument as for the Σ2 grain boundary, the Σ4 grain boundary appears to a larger extent at lower carburization temperatures. The peak at low angles corresponds to low angle grain boundaries. Optical microscopy and SEM imaging (Fig. 3) showed that we successfully produced cemented carbides of the different powders P1– P7. As the carbon potential has a big effect not the least on the grain

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012

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Table 5 WC measurements of the as-sintered materials from EBSD (equivalent circle diameter) and fraction of Ʃ2 WC/WC grain boundaries. Material

Mean grain size (μm)

Grain size std. dev. (μm)

Aspect ratio

Analyzed grains (#)

Fraction Ʃ2 boundaries

C1 C2 C3 C4 C5 C8 C9 C10

0.534 ± 0.009 0.534 ± 0.008 0.616 ± 0.013 0.580 ± 0.013 0.646 ± 0.012 0.595 ± 0.012 0.688 ± 0.015 0.853 ± 0.022

0.285 0.283 0.346 0.299 0.346 0.321 0.358 0.459

1.59 1.55 1.61 1.54 1.57 1.55 1.57 1.51

2340 2871 1851 1355 2142 1880 1493 1043

0.189 0.189 0.183 0.182 0.179 0.167 0.156 0.100

growth during sintering, it is important to analyze materials with comparable magnetic saturation relative to Co. The initial step of the sintering performed in hydrogen atmosphere enabled this despite different amount of carbon in the powders. The comparable materials in this sense are C1–C5 and C9–C10. As the starting powders differed in grain size it is no surprise that the as-sintered materials also had different WC grain sizes. There is a good correlation between the coercivity and the grain size where a material with lower coercivity is coarser grained. The hardness for WC-Co cemented carbides depends on the fraction and composition of binder phase, the WC grain size, the volume fraction and grain size of possible cubic carbides and potentially also the composition of the WC phase. The fraction of binder phase of the materials was kept constant in this study. As the magnetic saturation relative to Co is almost the same for materials C1–C5 and C9–C10, the composition of the binder phase in these is assumed to be about the same. Also, material C8 has only a slightly lower relative magnetic saturation. Based on the XRD measurements of the powders, it is expected that the cemented carbides produced from powder P5 (C5, C8) have the smallest amount of TaC and that the cemented carbide produced from powder P4 (C4) the largest amount. The difference in the as-sintered materials is believed to be smaller than for the powders due to the lack of metal diffusion and continuous solution and reprecipitation process during sintering. TaC has a microhardness of 18.00 GPa compared with WC with a microhardness of 13.00 GPa in the (1010) direction and 22.00 GPa in the (0001) direction [11]. It is not therefore believed that the cubic TaC phase will have any significant positive effect on the hardness for these materials. In addition, optical microscopy and EBSD shows that the TaC grain size is larger than the WC grain size. It might therefore even be that the TaC phase makes the material softer. In order to investigate any effect on the hardness of

the Ta doping of the WC grains, the hardness was plotted as a function of the WC grain size (Fig. 4). The error bar of the hardness mean value is calculated with a 95% confidence interval (five indents per material). Two materials stand out, C5 and C8. Both are less hard than can be expected from their grain size. These are the materials we believe have the largest amount of Ta in the WC grains. We thus deduce from our results that doping of Ta in the WC crystals has a softening effect on the material. WC powder pre-alloyed with Ta has successfully been used to produce coarse grained WC-Co based cemented carbides with a confirmed high concentration of Ta remaining in the WC structure [12]. The Young's moduli of the doped crystals were significantly lower than for undoped WC, as measured by nano-indentation in agreement with the lowering of the elastic constants, from ab initio calculations [12]. Our results are therefore in agreement with this study. The fraction of Σ2 WC/WC grain boundaries in the as-sintered materials has decreased compared to what existed in the originating powders. This in agreement with the theory that these boundaries exists already in the powder and does not form during sintering [9]. It was noted that no peak existed at 60° in the misorientation profile of the WC/WC grain boundaries. This does not surprise as when the fraction Σ2 boundaries decrease due to the grain coarsening process, so do the Σ4 boundaries. 5. Conclusions Hexagonal (W,Ta)C powders were produced with (W,Ta)2C powder as an intermediate product following a two-step carburization process. XRD measurements indicate that a higher temperature during the first carburization step increases the fraction of Ta in the (W,Ta)2C structure. XRD measurements indicate that a lower temperature during the second carburization step increases the fraction of Ta in the (W,Ta)C structure. The pre-alloyed powders had a high fraction of Σ2 WC/WC grain boundaries and a small fraction of what has been interpreted as Σ4 WC/WC grain boundaries. The powders carburized at lower temperatures had a larger fraction of these low energy boundaries. The (W,Ta)C powders were successfully used to produce cemented carbides. The pre-alloying with Ta appeared to have a softening effect of the material. Acknowledgements Thanks for advice and helpful comments are due to Hans-Olof Andrén from Chalmers University of Technology, Susanne Norgren from Sandvik Mining, Ernesto Coronel, Marie Lundbäck and José García from AB Sandvik Coromant R&D, Jan Qvick from Seco Tools AB and Gabriele Kremser and Andreas Bock from Wolfram Bergbau und Hütten AG. References

Fig. 4. Hardness (HV30) as a function of WC grain size for cemented carbide materials.

[1] E. Rudy, B.F. Kieffer, E. Baroch, HfN coatings for cemented carbides and new hard-facing alloys on the basis (Mo, W)C-(Mo, W)2C, Planseeber. Pulvermetall. 26 (1978) 105–115. [2] J. Weidow, S. Johansson, G. Wahnström, H.-O. Andrén, Transition metal solubilities in WC in cemented carbide materials, J. Am. Ceram. Soc. 94 (2011) 605–611. [3] H.-O. Andrén, U. Rolander, P. Lindahl, Phase composition in cemented carbides and cermets, Int. J. Refract. Met. Hard Mater. 12 (1994) 107–113. [4] E. Rudy, Compendium of phase diagram data. Part V, Air Force Materials Laboratory, Ohio, 1969 (AFML-TR-65-2). [5] J. Weidow, E. Halwax, W.-D. Schubert, Analysis of WC with increased Ta doping, Int. J. Refract. Met. Hard Mater. 51 (2015) 56–60. [6] Wolfram Bergbau und Hütten AG, W.D. Schubert, Doped Hexagonal Tungsten Carbide and Method for Producing Same, 2012 WO 2012/145773 A1. [7] D.G. Brandon, The structure of high-angle grain boundaries, Acta Metall. 14 (1966) 1479–1484. [8] E. Rudy, E. Rudy, F. Benesovsky, Investigations of the Ta-W-C system, Monatsh Chem 93 (1962) 1176–1195. [9] J.-D. Kim, L. Kang S-J, J.-W. Lee, Formation of grain boundaries in liquid-phase-sintered WC–Co alloys, J. Am. Ceram. Soc. 88 (2005) 500–503. [10] S. Lay, M. Loubradou, Characteristics and origin of clusters in submicron WC-Co cermets, Philos. Mag. 83 (2003) 2669–2679. [11] H.E. Exner, Physical and chemical nature of cemented carbides, Int. Met. Rev. 4 (1979) 149–173. [12] J. Weidow, A. Blomqvist, J. Salomonsson, S. Norgren, Cemented carbides based on WC pre-alloyed with Cr or Ta, Int. J. Refract. Met. Hard Mater. 49 (2015) 36–41.

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012

J. Weidow et al. / Int. Journal of Refractory Metals and Hard Materials xxx (2016) xxx–xxx Jonathan Weidow was born in 1980 and studied physics at Chalmers University of Technology. After a post-doc at Vienna University of Technology he now has a position as university lecturer at Chalmers. His main interest is on characterization of the microstructure of metallic materials with a special focus on cemented carbides. Jonathan has more than 20 publications in refereed journals and was awarded with the 2015 Sawamura Award from the Iron and Steel Institute of Japan.

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Wolf-Dieter Schubert was born in 1951. He is professor at the Institute of Chemical Technologies and Analytics at Vienna University of Technology and head of a research group focused on the metallurgy of less common metals. His main expertise is on refractory metals and cemented carbides. He has more than 100 publications in refereed journals, several book contributions, co-author of the monography on TUNGSTEN: Properties, Chemistry, Technology of the Element, Alloys and Chemical Compounds (ISBN 0-306-45053-4). He is technical consultant of the International Tungsten Industry Association and has recently received the Skaupy-Award by the German Association of Powder Metallurgy.

Erich Halwax was born in 1951 and studied chemistry in the Vienna Technical University. He did his doctoral thesis with Horst Völlenkle on the crystal structure analyses of various alkali metal germanates. He has been an Assistant Professor since 1990 with main interests in powder diffraction, especially quantitative phase analysis.

Please cite this article as: J. Weidow, et al., Optimized production route of WC powder pre-alloyed with Ta for fine grained cemented carbides, Int J Refract Met Hard Mater (2016), http://dx.doi.org/10.1016/j.ijrmhm.2016.05.012