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A new approach to fabrication of gradient WC–Co hardmetals
Int. Journal of Refractory Metals & Hard Materials 28 (2010) 228–237
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A new approach to fabrication of gradient WC–Co hardmetals I. Konyashin a,*, S. Hlawatschek a, B. Ries a, F. Lachmann a, A. Sologubenko b, T. Weirich b a b
Element Six Hard Materials, Element Six GmbH, Städeweg 12 – 24, 363151 Burghaun, Germany Central Facility for Electron Microscopy, Aachen University, Ahornstrasse 55, D-52074 Aachen, Germany
a r t i c l e
i n f o
Article history: Received 20 July 2009 Accepted 2 October 2009
Keywords: Hardmetal Gradient hardmetals Hardness Fracture toughness
a b s t r a c t WC–Co hardmetals with gradient structure comprising neither g-phase nor grain growth inhibitors were produced for the first time by regulating the WC re-crystallisation and carbon content in their near-surface layer and core. Hardmetals with low Co contents in the surface region were obtained by selective carburisation of the near-surface zone of green articles with the original low carbon content and their consequent liquid-phase sintering. The surface region of such gradient hardmetals has a hardness of up 150 Vickers units higher and fracture toughness significantly superior than those of the core. Gradient hardmetals with high Co contents in the surface region were obtained by selective decarburisation of the near-surface zone of green articles with the original high carbon content and their consequent liquidphase sintering. The new approach for fabrication of gradient WC–Co materials appears to be a unique tool for increasing both the hardmetal hardness and fracture toughness. Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction Fabrication of WC–Co hardmetals with gradient composition, microstructure and properties has been an issue of great interest in the hardmetal industry for a long time. If one can produce hardmetals with a near-surface region with lower Co contents than the average Co content, this region would have a high hardness leading to its better wear-resistance. It is also expected that the near-surface layer would possess higher fracture toughness than that of the core. The reason for that can be seen from Fig. 1, where a sample pressed from two WC–Co graded powders with different Co contents after consequent liquid-phase sintering is shown. The sample becomes bent after sintering, which is evidence for the presence of high residual compression stresses in the lower part of the sample comprising less Co. This occurs as a result of very different shrinkage rates during the hardmetal solidification after liquid-phase sintering in the two parts of the sample with various Co contests. As a result, the upper part of the sample with the high Co content shrinks significantly more intensively than the lower part. The presence of high residual compression stresses in the lower part should lead to its significantly higher fracture toughness than that of the hardmetal of the same composition and WC mean grain size produced conventionally. In the literature the two following major approaches to fabrication of gradient hardmetals are described. The first approach is based on the carburisation of fully sintered hardmetal articles with very low carbon contents originally com* Corresponding author. Tel.: +49 6652 82412; fax: +49 6652 82293. E-mail address: [email protected] (I. Konyashin). 0263-4368/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2009.10.003
prising g-phase [1,2]. In this case, there is no residual porosity in the articles and the carburisation occurs at temperatures above the eutectic melting point, so that the carbon diffusion proceeds through the liquid Co-based binder resulting in a Co drift from the surface toward the core. The microstructure of the gradient hardmetals obtained in such a way consists of (1) an upper layer with a low Co content comprising neither g-phase nor free carbon; (2) an intermediate layer with a high Co content comprising no gphase and (3) a core comprising much g-phase. The upper layer is characterised by both higher hardness and fracture toughness presumably as a result of high residual compression stresses [3]. The major disadvantage of the gradient hardmetals obtained by this approach is that the intermediate layer with the high Co content is very soft, so that if the hard upper surface layer is worn out during operation, the wear occurs within the intermediate layer very fast. It is very difficult to regulate the thickness of the hard upper layer, which is normally of nearly 1–2 mm in thickness, as it forms due to the carbon diffusion in the liquid binder with a very limited rate. The other major disadvantage of the gradient hardmetals obtained by this approach is the presence of the very brittle core comprising much g-phase, which cannot be eliminated. The only published work describing gradient hardmetals not comprising g-phase in the microstructure after the surface carburisation is Ref. [4]. However, fully sintered hardmetal articles without residual porosity are subjected to carburisation in this work, so that the thickness of the carburized layer with a decreased Co content is very insignificant (of the order of 100–150 lm). Another approach to fabrication of gradient WC–Co hardmetals is based on selective introducing WC grain growth inhibitors (mainly Cr and V) into the hardmetal near-surface layer as a result
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Residual Tensile Stress WC-15%Co WC-5%Co
Residual Compression Stress 10 mm Fig. 1. Appearance of a sample pressed from two graded WC–Co powders containing 15 wt.% Co and 5 wt.% Co after its sintering at 1420 °C for 75 min.
of applying the grain growth inhibitors or their precursors to the surface of green articles [5–7]. As a result of this technology, the near-surface region comprising much Cr or V becomes much finer than the core after liquid-phase sintering resulting in its high hardness and wear-resistance. Nevertheless, there might be significant difficulties related to applying the grain growth inhibitors or their precursors to the surface of hardmetal articles of complicated shape in the green state, as they are very brittle. Also, the presence of high concentrations of V and Cr in the near-surface layer can lead to its significantly decreased fracture toughness, as these grain growth inhibitors tend to segregate at the WC–Co interface [8,9]. In general, the following two major phenomena can be used for fabrication of gradient WC–Co materials: (1) If there are two adjacent layers of WC–Co with various carbon contents in the green state, cobalt drifts from the layer with the higher carbon content into the layer with the lower carbon content during liquid-phase sintering; (2) If there are two adjacent layers of WC–Co with various WC mean grain sizes in the green state, cobalt drifts from the layer with coarser microstructure into the layer with finer microstructure during liquid-phase sintering as a result of different capillary forces. These phenomena were described in detail in numerous publications with respect to bi-layer hardmetals pressed and sintered from various graded powders (see e.g. [10–12]) but, except for Ref. [4], never used for fabrication of gradient hardmetals comprising no g-phase. It is expected that if one can selectively regulate the WC mean grain size and/or carbon content in the hardmetal near-surface layer and core in the green state, it will be possible to create and control the Co drift during the final liquid-phase sintering thus obtaining gradient hardmetals. On the one hand, Co is thought to drift from the part with a higher carbon content to the part with a lower carbon contents together with the carbon diffusion. On the other hand, when taking into account that the WC re-crystallisation rate and consequently the WC mean grain size in WC–Co hardmetals can be regulated in a wide range by varying the carbon content [13], the Co drift is thought to occur from the part with a coarser WC grain size into the part with a finer grains size due to various capillary forces. Fig. 2 shows two hardmetal microstructures well illustrating the suppression of WC re-crystallization in WC–Co hardmetals with low carbon contents. The sample with a low carbon content produced by intensive milling of ultra-coarse WC powders shown in Fig. 2b has a significantly finer microstruc-
Fig. 2. Microstructure of hardmetals containing 6.5 wt.% Co produced from WC with the WC mean grain size of nearly 10 lm by milling in an attritor mill in hexane with the ball-to-powder ratio of 6:1 for 1 h and sintered at 1420 °C for 75 min: (a) hardmetal with ETC of 6.11% and (b) hardmetal with ETC of 5.99% (etching in the Murakhami reagent).
ture compared to the sample with a high carbon content shown in Fig. 2a. It can be seen that the fine-grain WC fraction formed as a result of milling completely dissolves and re-precipitates on the surface of large WC grains in the sample with the high carbon content. At the same time, the fine-grain WC fraction does not dissolve and re-crystallize in the sample with the low carbon content remaining in its microstructure. Recently we suggested a new approach to fabrication of WC–Co hardmetals with gradient structure [14] comprising the following major steps: (1) Fabrication of WC–Co mixtures with either low or high carbon contents at various milling conditions. The sintered samples of the mixtures with the low carbon content have no g-phase after conventional sintering, but are close to the border with the g-phase formation according to the WC–Co phase diagram with respect to the carbon content. The sintered samples of the mixtures with high carbon content have no free carbon after conventional sintering but are close to the border with the free carbon formation with respect to the carbon content according to the WC–Co
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Surface distance, mm Fig. 4. WC mean grain size, Co distribution and hardness of the alloy A sintered at 1460 °C for 60 min.
Fig. 3. Microstructures of the alloys A, B and C after their sintering at 1460 °C for 60 min (etching in the Murakhami reagent).
diagram. The milling conditions are adjusted in the following two ways. A very fine-grain WC fraction forms during milling, when the very intensive WC re-crystallisation in hardmetals is needed. The original WC grains are only very
gently de-agglomerated almost without crushing during milling when the intensive WC re-crystallisation during liquid-phase sintering is not desirable. (2) Pre-sintering of the green articles of the WC–Co mixtures in the solid state in a vacuum or neutral gas atmosphere at temperatures and time being sufficient to maintain open porosity within the surface region of the green articles. This allows one to obtain a certain level of gas permeability of just the near-surface layer with respect to the reactive gas phase and remain the core not reacted with the gas phase. (3) Heat-treatment of the pre-sintered green articles in the solid state in either carburising or de-carburising gas atmosphere to selectively carburise the near-surface region of the articles with the original low carbon content or selectively decarburise the near-surface surface region of the articles with the original high carbon content. The gas permeability and
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Surface distance, mm Fig. 6. WC mean grain size, Co distribution and hardness of the alloy A sintered at 1380 °C for 0 min.
Fig. 5. Microstructure of the alloy A sintered at 1460 °C for 60 min: (a) near-surface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
consequently the thickness of either carburised or de-carburised surface zone can be varied in a large scale by varying the parameters of the pre-sintering and heat-treatment.
(4) Final liquid-phase sintering of the articles in vacuum or vacuum + HIP to obtain their full density and allow carbon to diffuse and Co to drift in the liquid phase as well as the fine-grain fraction to re-crystallize to the needed extent. As a result of the liquid-phase sintering, Co drifts from the region with the local high carbon contents into the region with local low carbon contents. Co drifts also from the region with a coarser microstructure into the region with a finer microstructure and achieves its equilibrium concentration. As a result, a near-surface zone with a lower Co content forms on the surface of the samples with the originally low carbon content after their carburisation. A near-surface zone with a higher Co content forms on the surface of the samples with the originally high carbon content after their decarburisation.
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2. Experimental details
Fig. 7. Microstructure of the alloy A sintered at 1380 °C for 0 min: (a) near-surface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
The major objectives of the present work are to examine the influence of fabrication parameters on the microstructure and properties of the gradient WC–Co hardmetals comprising neither g-phase nor grain growth inhibitors obtained by use of the new approach.
Alloys of three types designated as A, B and C were made in the present work. Ultra-coarse grain WC (MAS3000-5000, H.C. Starck) and finegrained Co powder (EF, Umicore) were employed for the samples A and C. The mean grain size of the WC powder is found to be equal to 10.6 lm. The original ultra-coarse WC powder was milled in a ball mill in alcohol at the ball-to-powder ratio of 6:1 for 120 h. After subsequent drying the milled WC powder was analysed for carbon and oxygen contents and mixed with 10 wt.% Co and various amounts of carbon black or W metal with 2 wt.% paraffin wax in hexane for 1 h. Two alloys designated as alloys A and C with various carbon contents were made. The Equivalent Total Carbon (ETC) with respect to WC of the alloy A was equal to 5.98% and that of the alloy C was equal to 6.11%. The ETC values were adjusted to the sintering conditions in such a way that the samples A have the Specific Magnetic Saturation (SMS) of 76% and do not comprise gphase, and the samples C have SMS of 90% and do not comprise free carbon. Also an alloy with a low carbon content very close to that of the alloy A, designated as the alloy B, was produced by very gentle milling. The alloy B was produced on the basis of WC with an average grain size of about 2.5 lm (DS250, H.C. Starck), which was milled with 6 wt.% cobalt and fine W metal powder for one hour by means of a ball mill in a milling medium consisting of hexane with 2 wt.% paraffin wax, and using a powder-to-ball ratio of 1:2. A number of samples of each alloy were pre-sintered in a vacuum and sintered in a laboratory furnace (Vakkumanlagen Solms) in a graphite crucible of 200 mm in diameter and 500 mm in height. The samples A and B with the low carbon content were carburised in hydrogen–methane gas mixtures at 700–1100 °C and then sintered at different temperatures and times. The samples C with the high carbon content were de-carburised in hydrogen at 700 °C for 1 h and finally sintered at 1460 °C for 60 min. After the final liquid-phase sintering, metallurgical cross-sections were prepared and examined on an optical microscope and a SEM (a LeicaCambridge S360 instrument). The presence of g-phase and free carbon was established on the metallurgical cross-section with and without etching in the Murakhami reagent and finally confirmed by XRD. The measurement of WC mean grains size was carried out by the linear intercept method by use of the 4ai Analysis Grains Size software with the precision of ±5%. The Co contents
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in the samples after their carburisation or decarburisation and final liquid-phase sintering were measured by EDX on the SEM with the precision of roughly ±5%. Hardness measurements were carried out according to the DIN ISO 3878 at a load of 10 kg with the precision of roughly ±10 Vickers units. The fracture toughness was measured by the Palmquist method at a load of 100 kg with the precision of roughly ±5% after annealing of the cross-sections in a vacuum at 800 °C for 30 min. Transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM) of the hardmetal with ETC of 5.79% after conventional sintering in vacuum were carried out by use of the FEI Titan-T instrument.
Fig. 10. Microstructure of the alloy B sintered at 1420 °C for 60 min: (a) nearsurface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
3. Results and discussion Fig. 3 shows the microstructures of the alloys A, B and C after conventional sintering at 1460 °C for 1 h. It can be seen that the microstructure of the sample A is significantly finer than that of
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Fig. 11. WC–Co property map from Ref. [15] showing the points corresponding to the alloys A and B (blank circles).
the sample C, so tion formed due pressed at the employed in the of the samples A
that the WC re-crystallisation of fine-grain fracto the intensive milling of WC is strongly suplow carbon content. This phenomenon was selective carburisation of the near-surface zone and decarburisation of the samples C.
3.1. Alloy A with the original low carbon content after its carburisation and final sintering Fig. 4 shows the WC mean grain size, Co distribution and hardness, and Fig. 5 shows the microstructures of the alloy A sintered at 1460 °C for 60 min at various distances from the surface. It can be seen that the alloy sample contains neither g-phase nor free carbon. The XRD analysis confirms the absence of both g-phase and free carbon. The microstructure of the near-surface layer of the sample is coarser than that of the core, which is a result of intensive dissolution and re-crystallisation of the fine-grained fraction in the surface region with the high carbon content during liquidphase sintering. The WC re-crystallisation is noticeably suppressed in the core with the significantly lower carbon content remained after the selective surface carburisation. Fig. 4 clearly indicates the distinct Co gradient in the sample, so that the Co content decreases down to nearly 8% in the surface region and increases up to roughly 11.5% in the core. In the sample shown in Fig. 4 and Fig. 5 sintered at the relatively high temperature for the long time, which allows the intensive WC re-crystallisation in the surface region, the Co drift is thought to take place as a result of the difference in both the carbon content and capillary forces between the near-surface layer and the core. It can be seen that the hardness of the sample A is similar in both the surface region and core or even slightly decreases in the thin near-surface layer of 1 mm. This is presumably a result of the significantly coarser microstructure of the sample in the near-surface layer than in the core, even when the near-surface layer contains considerably less Co than the core. Fig. 6 shows the WC mean grain size, Co distribution and hardness, and Fig. 7 shows the microstructures of the alloy A sintered at 1380 °C for 0 min at various distances from the surface. Neither gphase nor free carbon is seen in the microstructure of the sample. The microstructure and WC mean grain size of the sample sintered at the relatively low temperatures for the very short time are similar in both the core and the surface region, because there was not enough time for dissolution and re-crystallisation of the fine-
grained WC fraction in the near-surface layer. The difference in the Co content between the near-surface zone and the region located at a distance of roughly 7 mm from the surface is nearly 3 wt.%. In this case, the Co drift is believed to be mainly a result of the difference in carbon contents between the surface region and the core, which are characterised by similar capillary forces. According to the results of Ref. [15] the diffusion coefficients of carbon and tungsten in the cobalt binder of WC–Co hardmetals are as follows:
Dc ¼ 8:72 106 expð149300=RTÞ Dw ¼ 7:0 105 expð282100=RTÞ From the data written above it is clear that the diffusion coefficient of carbon is much higher than that of tungsten, so that the carbon diffusion in the binder should occur very fast during sintering leading to the Co drift. It can be seen that the hardness dramatically increases from the core to the surface of the sample, so that the near-surface layer of the sample is harder by nearly 150 Vickers units compared to the core. Fig. 8 shows the curve indicating the distribution of the Palmquist fracture toughness in the samples A. The result on fracture toughness of the samples A provides evidence that their near-surface layer is significantly tougher than the core. The fracture toughness of the core is found to be equal to nearly 12 MPa m1/2, whereas that of the near-surface layer at the distance from the surface of 2 mm is equal to 16 MPa m1/2. The fracture toughness of the thin near-surface layer of roughly 1 mm is found to be very high and achieves the level of roughly 24 MPa m1/2. Such a high increase of fracture toughness is thought to be related to the presence of high compression stresses created in the surface layer as a result of the Co drift toward the core. 3.2. Alloy B with the original low carbon content after its carburisation and final sintering Fig. 9 shows the WC mean grain size, Co distribution and hardness, and Fig. 10 shows the microstructures of the alloy sample B sintered at 1420 °C for 60 min at various distances from the surface. Neither g-phase nor free carbon is present in the microstructure of the sample. The XRD analysis confirms the absence of both
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g-phase and free carbon. The microstructure and WC mean grain size of the sample are similar in both the core and surface region, because the alloy comprises little fine-grain fraction of WC, which is very active with respect to dissolution and re-crystallisation during liquid-phase sintering. The difference in the Co content between the near-surface zone and core is equal to roughly 2 wt.%. In this case, the Co drift is believed to be mainly a result of the difference in carbon contents between the surface region and the core. The near-surface layer is harder compared to the core by nearly 100 Vickers units due to the Co drift. The fracture toughness of the core is found to be equal to 13.1 MPa m1/2, which is typical for the hardmetals with the Co content and WC mean grain size of the core. The fracture toughness of the near-surface layer of 1–2 mm is found to be equal roughly to 16 MPa m1/2, which is noticeably higher than that of the core.
Fig. 13. Microstructure of the alloy C sintered at 1460 °C for 60 min: (a) nearsurface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
Fig. 11 shows the obtained combinations of hardness and fracture toughness in the near-surface region of the samples A and B in
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comparison with the typical combinations of hardness and fracture toughness for conventional WC–Co hardmetals with various Co contents and WC mean grain size according Roebuck et al. [16]. It is clearly seen that the near-surface region of the samples A and B are characterised by noticeably better combinations of hardness and fracture toughness compared to the conventional hardmetals, presumably as a result of high residual stresses forming in the near-surface region of the samples. Potential applications of the gradient hardmetals with the harder and tougher near-surface layer are obvious. They are mining and construction tools, wear-parts, tools for stamping, drawing etc. as well as cutting tools when the hardmetal is not subjected to re-grinding. For each application the design of the gradient WC–Co hardmetals with respect to the WC mean grain size, Co content and thickness of the near-surface layer should be optimised.
3.3. Alloy C with the original high carbon content after its decarburisation and final sintering Fig. 12 shows the WC mean grain size, Co distribution and hardness, and Fig. 13 shows the microstructures of the sample C sintered at 1460 °C for 60 min at various distances from the surface. It can be seen that, as in the case of the samples A and B, the microstructure of the samples C contains neither g-phase nor free carbon. The XRD analysis confirms the absence of both g-phase and free carbon. The microstructure of the near-surface layer of the sample after decarburisation is finer than that of the core, which is a result of the suppression of WC re-crystallisation in the surface region with the low carbon content after its selective decarburisation. In this case, after the relatively long sintering at the high temperature, the fine WC fraction easily dissolves in the liquid binder and re-crystallizes on large WC grains in the core region, in which the original high carbon contain is substantially remained, forming the relatively coarse microstructure. The distinct Co gradient in the sample C can be seen, so that the Co content increases up to roughly 14% in the surface region and decreases down to nearly 8% in the core. The Co drift is thought to take place due to the difference in both the carbon content and capillary forces between the near-surface layer and core. In spite of the presence of the significant Co gradient in the sample C, its hardness in both the near-surface layer and core is similar, presumably because the Co drift toward the surface is compensated by the difference in WC mean grain sizes between the core and the surface region. The fracture toughness of the samples C is found to be 12.2 MPa m½ in the near-surface layer and 14.0 MPa m½ in the core. Potential applications of the gradient WC–Co structures with higher Co contents in the surface region can be substrates for polycrystalline diamond (PCD) layers. PCD layers, or caps are typically bonded integrally to WC–Co substrates during the process of sintering the PCD. It can be useful in certain applications for the hardmetal substrate to be as stiff as possible in order to provide mechanical support for the PCD layer. Conventional hardmetal substrates presently employed for the PCD production typically contain 12–13% Co. For some applications the Co content in the hardmetal substrates for PCD should be as low as 6–8 wt.%. However, the practical use of hardmetals with low cobalt contents as substrates for PCD is limited by the fact that some of Co is required to migrate from the substrate into the PCD layer forming the socalled ‘‘Co depleted zone” adjacent to the PCD layer during the sintering process. The gradient hardmetals with relatively low Co contents (down to 8 or even 6 wt.% Co) as a whole, but with high Co contents in the surface region can potentially be of great importance as substrates for PCD.
Fig. 14. TEM and HRTEM images of hardmetal with ETC of 5.79%: (a) bright field TEM micrograph; the arrow shows the WC grain where the HRTEM images were taken; (b) and (c) HRTEM images of the WC–Co interface in various regions of the WC grain.
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3.4. HRTEM of WC–Co interfaces in hardmetals with the low carbon contents The high fracture toughness values of the gradient hardmetals mentioned above is a result of the absence of both g-phase and free carbon in their microstructure, which distinguishes them from the conventional gradient hardmetals comprising much g-phase in the core [1,2]. However, it is very important that no additional phases, which can reduce the hardmetal fracture toughness as in the case of the V and Cr-containing hardmetal [8,9], will form at the WC–Co interface in the constituent parts of the gradient hardmetals with low carbon contents. Fig. 14 shows TEM and HRTEM images of the hardmetal with the very low carbon content. The TEM image in Fig. 14a shows a rounded WC grain of nearly 0.5 lm surrounded by the Co-based binder; the HRTEM images were taken from different parts of this grain. The HRTEM images shown in Fig. 14b and c clearly indicate the absence of any additional phases at the WC–Co interface. As a result, the regulation of WC re-crystallisation and Co drifts by local decreasing the carbon content in hardmetals suggested in the present work allows one to increase the hardmetal hardness without loosing its fracture toughness. 4. Conclusions (1) A new approach to fabrication of gradient WC–Co hardmetals is described. Hardmetals with gradient structure and properties are produced by obtaining various carbon contents in the hardmetal near-surface layer and core. WC–Co hardmetals comprising neither g-phase nor grain growth inhibitors with various combinations of WC mean grain size, Co content, hardness and fracture toughness in their surface and core regions are obtained for the first time. (2) Hardmetals with low Co contents in the surface region are fabricated by carburisation of green articles with the original low carbon content and final liquid-phase sintering. The noticeable hardness gradient in these hardmetals can be achieved, so that their near-surface layer is harder by nearly 150 Vickers units and characterised by significantly superior fracture toughness compared to the core. Thus, it appears that the new approach to fabrication of gradient WC–Co hardmetals suggested in the present work is a unique tool for increasing both the hardmetal hardness and fracture toughness. (3) Hardmetals with high Co contents in the surface region are obtained by decarburisation of green articles with the original high carbon content and final liquid-phase sintering.
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Such hardmetals can potentially be useful as substrates for PCD comprising relatively low Co contents as a whole, but a significantly increased Co content in the near-surface layer. (4) HRTEM of hardmetals with the very low carbon content clearly indicates the absence of any additional phases at the WC–Co interface. As a result, the regulation of WC recrystallisation and Co drifts by local decreasing the carbon content in hardmetals suggested in the present work allows one to obtain the high hardmetal hardness without loosing its fracture toughness. References [1] Fischer U, Waldenstrom M, Hartzell T. Cemented carbide body with increased wear resistance. US Patent 5,856,626; 1999. [2] Fischer U, Hartzell E, Akerman J. Cemented carbide body used preferably for rock drilling and mineral cutting. US Patent 4,743, 515; 1988. [3] Li Z, Yuan-jie W, Xian-wnag Y, Sgu C, Xiangjun X. Crack propagation characteristic and toughness of functionally graded WC–Co cemented carbide. Int J Refract Met Hard Mater 2008;26:195–300. [4] Guo J, Fan P, Fang Z. A new method for making graded WC–Co by carburizing heat treatment of fully densified WC–Co. In: Sigl L, Rödhammer P, Wildner H, editors. Proceedings of the 17th international plansee seminar, vol. 2. Reutte; 2009. p. 50/1–6. [5] Collin M, Norgren S. Cemented carbide insert and method of making the same. European Patent EP1548136B1; 2008. [6] Glatzle J, Kosters R, Glatzle W. Hard metal component with a graduated structure and methods of producing the component. US Patent Application US2004/0009088; 2004. [7] Greenfield M. Sintered hard-alloy composites and tools manufactured with microstructural zones by grain refining US Patent 5,623,723; 1997. [8] Lay S, Hamar-Thibault S, Lackner A. Location of VC in VC Cr3C2 codoped WC– Co cermets by HRTEM and EELS. Int J Refract Met Hard Mater 2002;20:61–9. [9] Warbichel P, Hoferm F, Grogger W, Lackner A. EFTEM-EELS characterization of VC and Cr3C2 doped cemented carbides. In: Kneringer G, Rödhammer P, Wildner H, editors. Proceedings of the 15th international plansee seminar, vol. 2. Reutte; 2001. p. 65–74. [10] Liu Y, Wang H, Yang J, Huang B, Long Z. Formation mechanism of cobaltgradient structure in WC–Co hard alloy. J Mater Sci 2004;39:4397–9. [11] Eso O, Fang Z, Griffo A. Liquid phase sintering of functionally graded WC–Co composites. Int J Refract Met Hard Mater 2005;23:233–41. [12] Eso O, Fang P, Fang Z. A kinetic model for cobalt gradient formation during liquid phase sintering of functionally graded WC–Co. Int J Refract Met Hard Mater 2007;25:286–93. [13] Konyashin I, Hlawatschek S, Ries B, Lachmann F, Dorn F, Sologubenko A, et al. On the mechanism of WC coarsening in WC–Co hardmetals with various carbon contents. Int J Refract Met Hard Mater 2009;27:234–43. [14] Konyashin I, Hlawatschek S, Ries B, Lachmann F. Gradient WC–Co structures obtained by regulated WC re-crystallization without using grain growth inhibitors. In: Sigl L, Rödhammer P, Wildner H, editors. Proceedings of the 17th international plansee seminar, vol. 2. Reutte; 2009. p. 6/1–6/12. [15] Sun L, Jia C-C, Xian M. A research on the grain growth of WC–Co cemented carbide. Int J Refract Met Hard Mater 2007;25:121–4. [16] Roebuck B, Gee MG, Morrell R. Hardmetals – microstructural design testing and property maps. In: Kneringer G, Rödhammer P, Wildner H, editors. Proceedings of the 15th international plansee seminar, vol. 4. Reutte; 2001. p. 245–66.
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Int. Journal of Refractory Metals & Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
A new approach to fabrication of gradient WC–Co hardmetals I. Konyashin a,*, S. Hlawatschek a, B. Ries a, F. Lachmann a, A. Sologubenko b, T. Weirich b a b
Element Six Hard Materials, Element Six GmbH, Städeweg 12 – 24, 363151 Burghaun, Germany Central Facility for Electron Microscopy, Aachen University, Ahornstrasse 55, D-52074 Aachen, Germany
a r t i c l e
i n f o
Article history: Received 20 July 2009 Accepted 2 October 2009
Keywords: Hardmetal Gradient hardmetals Hardness Fracture toughness
a b s t r a c t WC–Co hardmetals with gradient structure comprising neither g-phase nor grain growth inhibitors were produced for the first time by regulating the WC re-crystallisation and carbon content in their near-surface layer and core. Hardmetals with low Co contents in the surface region were obtained by selective carburisation of the near-surface zone of green articles with the original low carbon content and their consequent liquid-phase sintering. The surface region of such gradient hardmetals has a hardness of up 150 Vickers units higher and fracture toughness significantly superior than those of the core. Gradient hardmetals with high Co contents in the surface region were obtained by selective decarburisation of the near-surface zone of green articles with the original high carbon content and their consequent liquidphase sintering. The new approach for fabrication of gradient WC–Co materials appears to be a unique tool for increasing both the hardmetal hardness and fracture toughness. Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction Fabrication of WC–Co hardmetals with gradient composition, microstructure and properties has been an issue of great interest in the hardmetal industry for a long time. If one can produce hardmetals with a near-surface region with lower Co contents than the average Co content, this region would have a high hardness leading to its better wear-resistance. It is also expected that the near-surface layer would possess higher fracture toughness than that of the core. The reason for that can be seen from Fig. 1, where a sample pressed from two WC–Co graded powders with different Co contents after consequent liquid-phase sintering is shown. The sample becomes bent after sintering, which is evidence for the presence of high residual compression stresses in the lower part of the sample comprising less Co. This occurs as a result of very different shrinkage rates during the hardmetal solidification after liquid-phase sintering in the two parts of the sample with various Co contests. As a result, the upper part of the sample with the high Co content shrinks significantly more intensively than the lower part. The presence of high residual compression stresses in the lower part should lead to its significantly higher fracture toughness than that of the hardmetal of the same composition and WC mean grain size produced conventionally. In the literature the two following major approaches to fabrication of gradient hardmetals are described. The first approach is based on the carburisation of fully sintered hardmetal articles with very low carbon contents originally com* Corresponding author. Tel.: +49 6652 82412; fax: +49 6652 82293. E-mail address: [email protected] (I. Konyashin). 0263-4368/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2009.10.003
prising g-phase [1,2]. In this case, there is no residual porosity in the articles and the carburisation occurs at temperatures above the eutectic melting point, so that the carbon diffusion proceeds through the liquid Co-based binder resulting in a Co drift from the surface toward the core. The microstructure of the gradient hardmetals obtained in such a way consists of (1) an upper layer with a low Co content comprising neither g-phase nor free carbon; (2) an intermediate layer with a high Co content comprising no gphase and (3) a core comprising much g-phase. The upper layer is characterised by both higher hardness and fracture toughness presumably as a result of high residual compression stresses [3]. The major disadvantage of the gradient hardmetals obtained by this approach is that the intermediate layer with the high Co content is very soft, so that if the hard upper surface layer is worn out during operation, the wear occurs within the intermediate layer very fast. It is very difficult to regulate the thickness of the hard upper layer, which is normally of nearly 1–2 mm in thickness, as it forms due to the carbon diffusion in the liquid binder with a very limited rate. The other major disadvantage of the gradient hardmetals obtained by this approach is the presence of the very brittle core comprising much g-phase, which cannot be eliminated. The only published work describing gradient hardmetals not comprising g-phase in the microstructure after the surface carburisation is Ref. [4]. However, fully sintered hardmetal articles without residual porosity are subjected to carburisation in this work, so that the thickness of the carburized layer with a decreased Co content is very insignificant (of the order of 100–150 lm). Another approach to fabrication of gradient WC–Co hardmetals is based on selective introducing WC grain growth inhibitors (mainly Cr and V) into the hardmetal near-surface layer as a result
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Residual Tensile Stress WC-15%Co WC-5%Co
Residual Compression Stress 10 mm Fig. 1. Appearance of a sample pressed from two graded WC–Co powders containing 15 wt.% Co and 5 wt.% Co after its sintering at 1420 °C for 75 min.
of applying the grain growth inhibitors or their precursors to the surface of green articles [5–7]. As a result of this technology, the near-surface region comprising much Cr or V becomes much finer than the core after liquid-phase sintering resulting in its high hardness and wear-resistance. Nevertheless, there might be significant difficulties related to applying the grain growth inhibitors or their precursors to the surface of hardmetal articles of complicated shape in the green state, as they are very brittle. Also, the presence of high concentrations of V and Cr in the near-surface layer can lead to its significantly decreased fracture toughness, as these grain growth inhibitors tend to segregate at the WC–Co interface [8,9]. In general, the following two major phenomena can be used for fabrication of gradient WC–Co materials: (1) If there are two adjacent layers of WC–Co with various carbon contents in the green state, cobalt drifts from the layer with the higher carbon content into the layer with the lower carbon content during liquid-phase sintering; (2) If there are two adjacent layers of WC–Co with various WC mean grain sizes in the green state, cobalt drifts from the layer with coarser microstructure into the layer with finer microstructure during liquid-phase sintering as a result of different capillary forces. These phenomena were described in detail in numerous publications with respect to bi-layer hardmetals pressed and sintered from various graded powders (see e.g. [10–12]) but, except for Ref. [4], never used for fabrication of gradient hardmetals comprising no g-phase. It is expected that if one can selectively regulate the WC mean grain size and/or carbon content in the hardmetal near-surface layer and core in the green state, it will be possible to create and control the Co drift during the final liquid-phase sintering thus obtaining gradient hardmetals. On the one hand, Co is thought to drift from the part with a higher carbon content to the part with a lower carbon contents together with the carbon diffusion. On the other hand, when taking into account that the WC re-crystallisation rate and consequently the WC mean grain size in WC–Co hardmetals can be regulated in a wide range by varying the carbon content [13], the Co drift is thought to occur from the part with a coarser WC grain size into the part with a finer grains size due to various capillary forces. Fig. 2 shows two hardmetal microstructures well illustrating the suppression of WC re-crystallization in WC–Co hardmetals with low carbon contents. The sample with a low carbon content produced by intensive milling of ultra-coarse WC powders shown in Fig. 2b has a significantly finer microstruc-
Fig. 2. Microstructure of hardmetals containing 6.5 wt.% Co produced from WC with the WC mean grain size of nearly 10 lm by milling in an attritor mill in hexane with the ball-to-powder ratio of 6:1 for 1 h and sintered at 1420 °C for 75 min: (a) hardmetal with ETC of 6.11% and (b) hardmetal with ETC of 5.99% (etching in the Murakhami reagent).
ture compared to the sample with a high carbon content shown in Fig. 2a. It can be seen that the fine-grain WC fraction formed as a result of milling completely dissolves and re-precipitates on the surface of large WC grains in the sample with the high carbon content. At the same time, the fine-grain WC fraction does not dissolve and re-crystallize in the sample with the low carbon content remaining in its microstructure. Recently we suggested a new approach to fabrication of WC–Co hardmetals with gradient structure [14] comprising the following major steps: (1) Fabrication of WC–Co mixtures with either low or high carbon contents at various milling conditions. The sintered samples of the mixtures with the low carbon content have no g-phase after conventional sintering, but are close to the border with the g-phase formation according to the WC–Co phase diagram with respect to the carbon content. The sintered samples of the mixtures with high carbon content have no free carbon after conventional sintering but are close to the border with the free carbon formation with respect to the carbon content according to the WC–Co
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Fig. 3. Microstructures of the alloys A, B and C after their sintering at 1460 °C for 60 min (etching in the Murakhami reagent).
diagram. The milling conditions are adjusted in the following two ways. A very fine-grain WC fraction forms during milling, when the very intensive WC re-crystallisation in hardmetals is needed. The original WC grains are only very
gently de-agglomerated almost without crushing during milling when the intensive WC re-crystallisation during liquid-phase sintering is not desirable. (2) Pre-sintering of the green articles of the WC–Co mixtures in the solid state in a vacuum or neutral gas atmosphere at temperatures and time being sufficient to maintain open porosity within the surface region of the green articles. This allows one to obtain a certain level of gas permeability of just the near-surface layer with respect to the reactive gas phase and remain the core not reacted with the gas phase. (3) Heat-treatment of the pre-sintered green articles in the solid state in either carburising or de-carburising gas atmosphere to selectively carburise the near-surface region of the articles with the original low carbon content or selectively decarburise the near-surface surface region of the articles with the original high carbon content. The gas permeability and
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Fig. 5. Microstructure of the alloy A sintered at 1460 °C for 60 min: (a) near-surface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
consequently the thickness of either carburised or de-carburised surface zone can be varied in a large scale by varying the parameters of the pre-sintering and heat-treatment.
(4) Final liquid-phase sintering of the articles in vacuum or vacuum + HIP to obtain their full density and allow carbon to diffuse and Co to drift in the liquid phase as well as the fine-grain fraction to re-crystallize to the needed extent. As a result of the liquid-phase sintering, Co drifts from the region with the local high carbon contents into the region with local low carbon contents. Co drifts also from the region with a coarser microstructure into the region with a finer microstructure and achieves its equilibrium concentration. As a result, a near-surface zone with a lower Co content forms on the surface of the samples with the originally low carbon content after their carburisation. A near-surface zone with a higher Co content forms on the surface of the samples with the originally high carbon content after their decarburisation.
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K1C, MPa m1/2
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2. Experimental details
Fig. 7. Microstructure of the alloy A sintered at 1380 °C for 0 min: (a) near-surface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
The major objectives of the present work are to examine the influence of fabrication parameters on the microstructure and properties of the gradient WC–Co hardmetals comprising neither g-phase nor grain growth inhibitors obtained by use of the new approach.
Alloys of three types designated as A, B and C were made in the present work. Ultra-coarse grain WC (MAS3000-5000, H.C. Starck) and finegrained Co powder (EF, Umicore) were employed for the samples A and C. The mean grain size of the WC powder is found to be equal to 10.6 lm. The original ultra-coarse WC powder was milled in a ball mill in alcohol at the ball-to-powder ratio of 6:1 for 120 h. After subsequent drying the milled WC powder was analysed for carbon and oxygen contents and mixed with 10 wt.% Co and various amounts of carbon black or W metal with 2 wt.% paraffin wax in hexane for 1 h. Two alloys designated as alloys A and C with various carbon contents were made. The Equivalent Total Carbon (ETC) with respect to WC of the alloy A was equal to 5.98% and that of the alloy C was equal to 6.11%. The ETC values were adjusted to the sintering conditions in such a way that the samples A have the Specific Magnetic Saturation (SMS) of 76% and do not comprise gphase, and the samples C have SMS of 90% and do not comprise free carbon. Also an alloy with a low carbon content very close to that of the alloy A, designated as the alloy B, was produced by very gentle milling. The alloy B was produced on the basis of WC with an average grain size of about 2.5 lm (DS250, H.C. Starck), which was milled with 6 wt.% cobalt and fine W metal powder for one hour by means of a ball mill in a milling medium consisting of hexane with 2 wt.% paraffin wax, and using a powder-to-ball ratio of 1:2. A number of samples of each alloy were pre-sintered in a vacuum and sintered in a laboratory furnace (Vakkumanlagen Solms) in a graphite crucible of 200 mm in diameter and 500 mm in height. The samples A and B with the low carbon content were carburised in hydrogen–methane gas mixtures at 700–1100 °C and then sintered at different temperatures and times. The samples C with the high carbon content were de-carburised in hydrogen at 700 °C for 1 h and finally sintered at 1460 °C for 60 min. After the final liquid-phase sintering, metallurgical cross-sections were prepared and examined on an optical microscope and a SEM (a LeicaCambridge S360 instrument). The presence of g-phase and free carbon was established on the metallurgical cross-section with and without etching in the Murakhami reagent and finally confirmed by XRD. The measurement of WC mean grains size was carried out by the linear intercept method by use of the 4ai Analysis Grains Size software with the precision of ±5%. The Co contents
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in the samples after their carburisation or decarburisation and final liquid-phase sintering were measured by EDX on the SEM with the precision of roughly ±5%. Hardness measurements were carried out according to the DIN ISO 3878 at a load of 10 kg with the precision of roughly ±10 Vickers units. The fracture toughness was measured by the Palmquist method at a load of 100 kg with the precision of roughly ±5% after annealing of the cross-sections in a vacuum at 800 °C for 30 min. Transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM) of the hardmetal with ETC of 5.79% after conventional sintering in vacuum were carried out by use of the FEI Titan-T instrument.
Fig. 10. Microstructure of the alloy B sintered at 1420 °C for 60 min: (a) nearsurface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
3. Results and discussion Fig. 3 shows the microstructures of the alloys A, B and C after conventional sintering at 1460 °C for 1 h. It can be seen that the microstructure of the sample A is significantly finer than that of
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Fig. 11. WC–Co property map from Ref. [15] showing the points corresponding to the alloys A and B (blank circles).
the sample C, so tion formed due pressed at the employed in the of the samples A
that the WC re-crystallisation of fine-grain fracto the intensive milling of WC is strongly suplow carbon content. This phenomenon was selective carburisation of the near-surface zone and decarburisation of the samples C.
3.1. Alloy A with the original low carbon content after its carburisation and final sintering Fig. 4 shows the WC mean grain size, Co distribution and hardness, and Fig. 5 shows the microstructures of the alloy A sintered at 1460 °C for 60 min at various distances from the surface. It can be seen that the alloy sample contains neither g-phase nor free carbon. The XRD analysis confirms the absence of both g-phase and free carbon. The microstructure of the near-surface layer of the sample is coarser than that of the core, which is a result of intensive dissolution and re-crystallisation of the fine-grained fraction in the surface region with the high carbon content during liquidphase sintering. The WC re-crystallisation is noticeably suppressed in the core with the significantly lower carbon content remained after the selective surface carburisation. Fig. 4 clearly indicates the distinct Co gradient in the sample, so that the Co content decreases down to nearly 8% in the surface region and increases up to roughly 11.5% in the core. In the sample shown in Fig. 4 and Fig. 5 sintered at the relatively high temperature for the long time, which allows the intensive WC re-crystallisation in the surface region, the Co drift is thought to take place as a result of the difference in both the carbon content and capillary forces between the near-surface layer and the core. It can be seen that the hardness of the sample A is similar in both the surface region and core or even slightly decreases in the thin near-surface layer of 1 mm. This is presumably a result of the significantly coarser microstructure of the sample in the near-surface layer than in the core, even when the near-surface layer contains considerably less Co than the core. Fig. 6 shows the WC mean grain size, Co distribution and hardness, and Fig. 7 shows the microstructures of the alloy A sintered at 1380 °C for 0 min at various distances from the surface. Neither gphase nor free carbon is seen in the microstructure of the sample. The microstructure and WC mean grain size of the sample sintered at the relatively low temperatures for the very short time are similar in both the core and the surface region, because there was not enough time for dissolution and re-crystallisation of the fine-
grained WC fraction in the near-surface layer. The difference in the Co content between the near-surface zone and the region located at a distance of roughly 7 mm from the surface is nearly 3 wt.%. In this case, the Co drift is believed to be mainly a result of the difference in carbon contents between the surface region and the core, which are characterised by similar capillary forces. According to the results of Ref. [15] the diffusion coefficients of carbon and tungsten in the cobalt binder of WC–Co hardmetals are as follows:
Dc ¼ 8:72 106 expð149300=RTÞ Dw ¼ 7:0 105 expð282100=RTÞ From the data written above it is clear that the diffusion coefficient of carbon is much higher than that of tungsten, so that the carbon diffusion in the binder should occur very fast during sintering leading to the Co drift. It can be seen that the hardness dramatically increases from the core to the surface of the sample, so that the near-surface layer of the sample is harder by nearly 150 Vickers units compared to the core. Fig. 8 shows the curve indicating the distribution of the Palmquist fracture toughness in the samples A. The result on fracture toughness of the samples A provides evidence that their near-surface layer is significantly tougher than the core. The fracture toughness of the core is found to be equal to nearly 12 MPa m1/2, whereas that of the near-surface layer at the distance from the surface of 2 mm is equal to 16 MPa m1/2. The fracture toughness of the thin near-surface layer of roughly 1 mm is found to be very high and achieves the level of roughly 24 MPa m1/2. Such a high increase of fracture toughness is thought to be related to the presence of high compression stresses created in the surface layer as a result of the Co drift toward the core. 3.2. Alloy B with the original low carbon content after its carburisation and final sintering Fig. 9 shows the WC mean grain size, Co distribution and hardness, and Fig. 10 shows the microstructures of the alloy sample B sintered at 1420 °C for 60 min at various distances from the surface. Neither g-phase nor free carbon is present in the microstructure of the sample. The XRD analysis confirms the absence of both
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g-phase and free carbon. The microstructure and WC mean grain size of the sample are similar in both the core and surface region, because the alloy comprises little fine-grain fraction of WC, which is very active with respect to dissolution and re-crystallisation during liquid-phase sintering. The difference in the Co content between the near-surface zone and core is equal to roughly 2 wt.%. In this case, the Co drift is believed to be mainly a result of the difference in carbon contents between the surface region and the core. The near-surface layer is harder compared to the core by nearly 100 Vickers units due to the Co drift. The fracture toughness of the core is found to be equal to 13.1 MPa m1/2, which is typical for the hardmetals with the Co content and WC mean grain size of the core. The fracture toughness of the near-surface layer of 1–2 mm is found to be equal roughly to 16 MPa m1/2, which is noticeably higher than that of the core.
Fig. 13. Microstructure of the alloy C sintered at 1460 °C for 60 min: (a) nearsurface layer, (b) intermediate layer and (c) core (etching in the Murakhami reagent).
Fig. 11 shows the obtained combinations of hardness and fracture toughness in the near-surface region of the samples A and B in
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comparison with the typical combinations of hardness and fracture toughness for conventional WC–Co hardmetals with various Co contents and WC mean grain size according Roebuck et al. [16]. It is clearly seen that the near-surface region of the samples A and B are characterised by noticeably better combinations of hardness and fracture toughness compared to the conventional hardmetals, presumably as a result of high residual stresses forming in the near-surface region of the samples. Potential applications of the gradient hardmetals with the harder and tougher near-surface layer are obvious. They are mining and construction tools, wear-parts, tools for stamping, drawing etc. as well as cutting tools when the hardmetal is not subjected to re-grinding. For each application the design of the gradient WC–Co hardmetals with respect to the WC mean grain size, Co content and thickness of the near-surface layer should be optimised.
3.3. Alloy C with the original high carbon content after its decarburisation and final sintering Fig. 12 shows the WC mean grain size, Co distribution and hardness, and Fig. 13 shows the microstructures of the sample C sintered at 1460 °C for 60 min at various distances from the surface. It can be seen that, as in the case of the samples A and B, the microstructure of the samples C contains neither g-phase nor free carbon. The XRD analysis confirms the absence of both g-phase and free carbon. The microstructure of the near-surface layer of the sample after decarburisation is finer than that of the core, which is a result of the suppression of WC re-crystallisation in the surface region with the low carbon content after its selective decarburisation. In this case, after the relatively long sintering at the high temperature, the fine WC fraction easily dissolves in the liquid binder and re-crystallizes on large WC grains in the core region, in which the original high carbon contain is substantially remained, forming the relatively coarse microstructure. The distinct Co gradient in the sample C can be seen, so that the Co content increases up to roughly 14% in the surface region and decreases down to nearly 8% in the core. The Co drift is thought to take place due to the difference in both the carbon content and capillary forces between the near-surface layer and core. In spite of the presence of the significant Co gradient in the sample C, its hardness in both the near-surface layer and core is similar, presumably because the Co drift toward the surface is compensated by the difference in WC mean grain sizes between the core and the surface region. The fracture toughness of the samples C is found to be 12.2 MPa m½ in the near-surface layer and 14.0 MPa m½ in the core. Potential applications of the gradient WC–Co structures with higher Co contents in the surface region can be substrates for polycrystalline diamond (PCD) layers. PCD layers, or caps are typically bonded integrally to WC–Co substrates during the process of sintering the PCD. It can be useful in certain applications for the hardmetal substrate to be as stiff as possible in order to provide mechanical support for the PCD layer. Conventional hardmetal substrates presently employed for the PCD production typically contain 12–13% Co. For some applications the Co content in the hardmetal substrates for PCD should be as low as 6–8 wt.%. However, the practical use of hardmetals with low cobalt contents as substrates for PCD is limited by the fact that some of Co is required to migrate from the substrate into the PCD layer forming the socalled ‘‘Co depleted zone” adjacent to the PCD layer during the sintering process. The gradient hardmetals with relatively low Co contents (down to 8 or even 6 wt.% Co) as a whole, but with high Co contents in the surface region can potentially be of great importance as substrates for PCD.
Fig. 14. TEM and HRTEM images of hardmetal with ETC of 5.79%: (a) bright field TEM micrograph; the arrow shows the WC grain where the HRTEM images were taken; (b) and (c) HRTEM images of the WC–Co interface in various regions of the WC grain.
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3.4. HRTEM of WC–Co interfaces in hardmetals with the low carbon contents The high fracture toughness values of the gradient hardmetals mentioned above is a result of the absence of both g-phase and free carbon in their microstructure, which distinguishes them from the conventional gradient hardmetals comprising much g-phase in the core [1,2]. However, it is very important that no additional phases, which can reduce the hardmetal fracture toughness as in the case of the V and Cr-containing hardmetal [8,9], will form at the WC–Co interface in the constituent parts of the gradient hardmetals with low carbon contents. Fig. 14 shows TEM and HRTEM images of the hardmetal with the very low carbon content. The TEM image in Fig. 14a shows a rounded WC grain of nearly 0.5 lm surrounded by the Co-based binder; the HRTEM images were taken from different parts of this grain. The HRTEM images shown in Fig. 14b and c clearly indicate the absence of any additional phases at the WC–Co interface. As a result, the regulation of WC re-crystallisation and Co drifts by local decreasing the carbon content in hardmetals suggested in the present work allows one to increase the hardmetal hardness without loosing its fracture toughness. 4. Conclusions (1) A new approach to fabrication of gradient WC–Co hardmetals is described. Hardmetals with gradient structure and properties are produced by obtaining various carbon contents in the hardmetal near-surface layer and core. WC–Co hardmetals comprising neither g-phase nor grain growth inhibitors with various combinations of WC mean grain size, Co content, hardness and fracture toughness in their surface and core regions are obtained for the first time. (2) Hardmetals with low Co contents in the surface region are fabricated by carburisation of green articles with the original low carbon content and final liquid-phase sintering. The noticeable hardness gradient in these hardmetals can be achieved, so that their near-surface layer is harder by nearly 150 Vickers units and characterised by significantly superior fracture toughness compared to the core. Thus, it appears that the new approach to fabrication of gradient WC–Co hardmetals suggested in the present work is a unique tool for increasing both the hardmetal hardness and fracture toughness. (3) Hardmetals with high Co contents in the surface region are obtained by decarburisation of green articles with the original high carbon content and final liquid-phase sintering.
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Such hardmetals can potentially be useful as substrates for PCD comprising relatively low Co contents as a whole, but a significantly increased Co content in the near-surface layer. (4) HRTEM of hardmetals with the very low carbon content clearly indicates the absence of any additional phases at the WC–Co interface. As a result, the regulation of WC recrystallisation and Co drifts by local decreasing the carbon content in hardmetals suggested in the present work allows one to obtain the high hardmetal hardness without loosing its fracture toughness. References [1] Fischer U, Waldenstrom M, Hartzell T. Cemented carbide body with increased wear resistance. US Patent 5,856,626; 1999. [2] Fischer U, Hartzell E, Akerman J. Cemented carbide body used preferably for rock drilling and mineral cutting. US Patent 4,743, 515; 1988. [3] Li Z, Yuan-jie W, Xian-wnag Y, Sgu C, Xiangjun X. Crack propagation characteristic and toughness of functionally graded WC–Co cemented carbide. Int J Refract Met Hard Mater 2008;26:195–300. [4] Guo J, Fan P, Fang Z. A new method for making graded WC–Co by carburizing heat treatment of fully densified WC–Co. In: Sigl L, Rödhammer P, Wildner H, editors. Proceedings of the 17th international plansee seminar, vol. 2. Reutte; 2009. p. 50/1–6. [5] Collin M, Norgren S. Cemented carbide insert and method of making the same. European Patent EP1548136B1; 2008. [6] Glatzle J, Kosters R, Glatzle W. Hard metal component with a graduated structure and methods of producing the component. US Patent Application US2004/0009088; 2004. [7] Greenfield M. Sintered hard-alloy composites and tools manufactured with microstructural zones by grain refining US Patent 5,623,723; 1997. [8] Lay S, Hamar-Thibault S, Lackner A. Location of VC in VC Cr3C2 codoped WC– Co cermets by HRTEM and EELS. Int J Refract Met Hard Mater 2002;20:61–9. [9] Warbichel P, Hoferm F, Grogger W, Lackner A. EFTEM-EELS characterization of VC and Cr3C2 doped cemented carbides. In: Kneringer G, Rödhammer P, Wildner H, editors. Proceedings of the 15th international plansee seminar, vol. 2. Reutte; 2001. p. 65–74. [10] Liu Y, Wang H, Yang J, Huang B, Long Z. Formation mechanism of cobaltgradient structure in WC–Co hard alloy. J Mater Sci 2004;39:4397–9. [11] Eso O, Fang Z, Griffo A. Liquid phase sintering of functionally graded WC–Co composites. Int J Refract Met Hard Mater 2005;23:233–41. [12] Eso O, Fang P, Fang Z. A kinetic model for cobalt gradient formation during liquid phase sintering of functionally graded WC–Co. Int J Refract Met Hard Mater 2007;25:286–93. [13] Konyashin I, Hlawatschek S, Ries B, Lachmann F, Dorn F, Sologubenko A, et al. On the mechanism of WC coarsening in WC–Co hardmetals with various carbon contents. Int J Refract Met Hard Mater 2009;27:234–43. [14] Konyashin I, Hlawatschek S, Ries B, Lachmann F. Gradient WC–Co structures obtained by regulated WC re-crystallization without using grain growth inhibitors. In: Sigl L, Rödhammer P, Wildner H, editors. Proceedings of the 17th international plansee seminar, vol. 2. Reutte; 2009. p. 6/1–6/12. [15] Sun L, Jia C-C, Xian M. A research on the grain growth of WC–Co cemented carbide. Int J Refract Met Hard Mater 2007;25:121–4. [16] Roebuck B, Gee MG, Morrell R. Hardmetals – microstructural design testing and property maps. In: Kneringer G, Rödhammer P, Wildner H, editors. Proceedings of the 15th international plansee seminar, vol. 4. Reutte; 2001. p. 245–66.