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Effect of C content on the microstructure evolution during early solid state sintering of WC–Co alloys
Int. Journal of Refractory Metals and Hard Materials 44 (2014) 27–34
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Effect of C content on the microstructure evolution during early solid state sintering of WC–Co alloys V. Bounhoure a, S. Lay a,⁎, F. Charlot b, A. Antoni-Zdziobek a, E. Pauty c, J.M. Missiaen a a b c
SIMAP, Grenoble INP-CNRS-UJF, BP 75, 38402 Saint-Martin d'Hères, France CMTC, Grenoble INP, 38402 Saint Martin d'Hères, France SANDVIK Hard Materials, 51 avenue Rhin et Danube, 38100 Grenoble, France
a r t i c l e
i n f o
Article history: Received 30 July 2013 Accepted 20 December 2013 Available online 29 December 2013 Keywords: WC–Co alloys Cobalt spreading Solid state sintering WC/Co interfaces
a b s t r a c t The densification of WC–10Co (wt %) alloys is studied in relation with interface formation, below the melting temperature of the Co base binder phase. Two alloys containing an excess of carbon or tungsten are investigated. The evolution of the pore and Co phase size distribution is quantified at several temperatures using image analysis. While shrinkage and Co spreading start earlier in the W rich alloy, the reduction of porosity and Co spreading is more uniform in the C rich alloy. Preferential orientation relationships are frequently found at WC basal facets for WC/Co interfaces especially in the W rich alloy. A WC1-x film is also sometimes observed at randomly oriented WC/Co interfaces whatever the C content. The formation of low energy interfaces may explain the early spreading of the binder phase and the onset of densification in the solid state. © 2014 Elsevier Ltd. All rights reserved.
1. Introduction A large part of hard tool materials are cemented carbides based on WC–Co because they combine a high hardness and sufficient toughness. These materials are produced by powder metallurgy and consist of a skeleton of WC grains embedded in a Co base binder. The suitable binder exists for a tight range of C and W contents. According to the literature that is briefly summarised in [1,2], densification initially proceeds by WC particle rearrangement, induced by spreading of the binder phase between WC grains, first in the solid state then when the binder phase is liquid. Solid state rearrangement requires the previous spreading of Co on the WC particle surfaces [3], a process governed by the energy of the WC/Co interfaces. The carbon potential has an influence on the shrinkage behaviour in the solid state [4,5]. In particular, the shrinkage starts earlier in W rich alloys and it was proposed to be related to the magnetic properties of the binder that would influence the interface energy [6]. The energy also depends on the interface atomic structure. While few experimental data on the structure of WC/Co interfaces are available for alloys sintered by conventional methods with a liquid phase [7–9] or using a field assisted sintering technique [10,11], a scarce attention was paid to the interfaces formed in the solid state [6]. The present paper investigates the effect of carbon content on the microstructure of WC–10Co (wt%) alloys on heating in the solid state. Attention is focused on the effect of composition on the early spreading of the binder on WC grains and on the onset of solid state shrinkage. Shrinkage is followed by dilatometry. The porosity evolution and Co ⁎ Corresponding author. E-mail address: [email protected] (S. Lay). 0263-4368/$ – see front matter © 2014 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijrmhm.2013.12.012
spreading are quantified by image analysis. The WC/Co interfaces are characterised and their structure is observed at the atom scale by transmission electron microscopy. The WC/Co interface characteristics formed in the solid state are compared with interfacial structures obtained after liquid phase sintering. 2. Experimental The two-phase field (WC + Co binder) suitable for the cemented carbides is tight and its limits change with temperature [12]. Therefore, to study the C potential effect, two compositions located in the threephase fields (WC + Co binder + Cg) and (WC + Co binder + η) at 1400 °C were selected where Cg refer to the graphite and η to the mixed (W,Co)6C carbide [13]. The studied alloys called WC–Co,C and WC–Co,W, correspond to two values of carbon potential which are fixed along the process due to the presence of Cg or η phase. The alloys were prepared from WC and Co powders both with a mean size of 0.9 μm. W powder with a mean size of 0.7 μm or Cg were added in order to provide the required C potential in the alloys (Table 1). The powder mixtures were attritor milled for 5 h with acetone and polyethylene glycol (2 wt.%) as an organic binder. For the dilatometry analysis, cylinders (8 mm in diameter, about 3 g in mass) were obtained by uniaxial pressing at 200 MPa. The green density was 53% for both mixtures. Samples were sintered in a SETARAM™ Setsys Evolution dilatometer. A slow heating at 1 °C/min under H2/He atmosphere was first applied up to 400 °C for debinding, followed by heating up to 1400 °C at 3 °C/min or 1 °C/min under Ar atmosphere where samples were held for 30 min and then cooled down. Interrupted experiments were performed on heating for microstructural characterisation. The constitution of the
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Table 1 Compositions of the initial WC–Co,C and WC–Co,W powder mixtures. Content (at%)
Co
W
C
WC–Co,C WC–Co,W
15.1 15.1
41.6 43.3
43.3 41.6
alloys was studied by X-ray diffraction using Cu-Kα radiation. For scanning electron microscopy (SEM) observations, the specimens were embedded in a fluid resin under vacuum then grinded using a standard procedure. For transmission electron microscopy (TEM) and high resolution transmission electron microscopy (HRTEM) investigations, thin foils were prepared by mechanical grinding and Ar ion milling. The JEOL 3010 and 2010 microscopes were used for the observations. 3. Results 3.1. Microstructure and constitution evolution In order to characterize the microstructure evolution of the alloys during solid state shrinkage, interrupted experiments were scheduled from the dilatometric plots. Fig. 1 shows the linear shrinkage of the alloys at a heating rate of 3 °C/min. No significant difference was observed with a heating rate of 1 °C/min. Shrinkage starts above 800 °C and strongly accelerates in the 1050 °C–1100 °C temperature range. A significant difference is observed at the onset of the shrinkage that starts earlier in WC–Co,W. A special attention will be paid in the following to explain this effect. Shrinkage rate then becomes naturally larger at higher temperature for the WC–Co,C composition as the remaining driving force is larger. Above 1200 °C, the differences between the alloys are partly explained from the melting temperature lower for the C rich than W rich binder. The sharp peak of the shrinkage rate plots corresponds to the beginning of liquid phase sintering that occurs at 1293 °C for WC–Co,C and 1369 °C for WC–Co,W. These values agree with the solidus temperatures determined with Thermocalc™ software which are 1298 and 1368 °C, respectively. In the alloys, a maximum of shrinkage rate is observed during solid state sintering. It may be attributed to the end of primary particle rearrangement, when a dense packing of WC particles has formed [6]. Microstructural characterisations were conducted at 1100 °C, 1220 °C, 1240 °C and 1260 °C. 1100 °C corresponds to the beginning of sintering, before the first peak of solid state densification for both alloys. 1220 °C is after the first peak of densification for WC–Co,W and 1260 °C is close to the maximum of the
densification rate for WC–Co,C. Temperature of 1240 °C corresponds to an intermediate stage between the two maxima of the densification rate. The constitution of the powder mixtures was checked by X-ray diffraction (Fig. 2). In the starting powder, a peak of tungsten is observed in the mixture of the W rich alloy, as expected. In the WC–Co,C powder mixture, only the peaks associated with WC and Co are found as the carbon is not detected due to the small atomic scattering factor of C atoms. Moreover, only the cubic crystal lattice for Co is detected in the powder. The microstructure of the powder mixtures was observed by SEM (Fig. 3). The starting compacts show the presence of Co pools as large as 2 μm. A large amount of WC grains smaller than 0.4 μm is present. Alloys heated at 3 °C/min were used for the SEM characterizations. After heating at 1100 °C, large Co pools from initial powder mixtures are still present, but they form agglomerates with WC grains. The Co binder phase has started to spread between WC grains. Some grains in contact with Co are facetted. Co spreading is more important in WC–Co,W. Moreover, in this alloy, the amount of very small grains has decreased. They have likely been dissolved in the binder. The dissolution of WC affects the lattice parameter of Co that increases in both alloys (Fig. 2). The η' phase ((Co,W)12C) is detected by X ray diffraction in WC–Co,W and observed by SEM in agreement with thermodynamic data [13]. At 1220 °C, the densification level is higher. A continuous skeleton of WC grains has formed due to further Co spreading and rearrangement of WC grains under capillary forces. The remaining Co pools are smaller in the W rich alloy. The η' phase is still present. At 1260 °C, the microstructure of the two alloys is rather similar. The Co binder has spread further between WC grains has still increased. There seems to be a slightly higher amount of Co pools in WC–Co,W than at 1220 °C. From XRD analysis, the η' phase has disappeared and the equilibrium η phase is present. In order to get accurate information on the microstructure evolution of the alloys during solid state sintering, image analysis was performed from SEM pictures. Due to the high level of porosity in the alloys, it was difficult to polish the samples, so the validity of the method was first checked using the porosity parameter. The amounts of porosity deduced from image analysis and from the measured density of the samples were compared (Table 2). A good agreement was found between the two methods at each temperature. Image analysis was then used to quantify the evolution of the porosity in the alloys. Fig. 4 shows the intercept length distribution of the pores in WC–Co,C and WC–Co,W specimens for the experiments interrupted at 1100 °C, 1220 °C and 1260 °C. The porosity reduction is uniform in WC–Co,C on heating. The pore volume fraction starts to decrease at a lower temperature for
9
Tc(WC-Co,W)
6
Tc(WC-Co,C)
0.1
WC
-0.1
0
-0.2
-3
-0.3
-6
-0.4
-9
-0.5
-12
WC-Co,C
-15
WC-Co,W
-0.6
WC
η
Co WC η
η'
η η
η'
η'
intensity
Shrinkage (%)
3
Shrinkage rate (%/min)
0
WC-Co,W 1260°C WC-Co,W 1220°C
η' η
WC-Co,W 1100°C
W
WC-Co,W powder
Co WC-Co,C 1260°C WC-Co,C 1220°C
-0.7
-18 -21 600
700
800
900
WC-Co,C 1100°C
-0.8 1000 1100 1200 1300 1400
Temperature (°C) Fig. 1. Shrinkage and shrinkage rate of the studied alloys during the thermal treatment (heating rate of 3 °C/min). The Curie temperature of the binder is indicated for each alloy by an arrow and the melting of the binder by a star. Dotted vertical lines indicate the temperatures (1100, 1220, 1240 and 1260 °C) at which experiments were interrupted for microstructural characterisations.
WC-Co,C powder 30
35
40
45
50
2θ (°) Fig. 2. Evolution of the X-ray diagram of the studied alloys versus the temperature. The arrows indicate the Co peaks. The shift of the Co peaks is related to the dissolution of WC in the binder.
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Fig. 3. SEM images of the WC–Co,C and WC–Co,W powder compacts and alloys. WC appears in bright, Co in grey and pores in black.
the WC–Co,W alloy, but large pores about 2 μm in size are difficult to eliminate and are still present at 1260 °C. Spreading of the binder was also studied by image analysis from SEM images. The volume fraction of the binder in the solid phase was
Table 2 Volume fraction of porosity and binder (%vol) deduced from density measurements (DM) or calculated by Thermocalc™ software (TC) compared with the values measured by image analysis on SEM images (IA).
Pores
WC–Co,C WC–Co,W
Binder
WC–Co,C WC–Co,W
DM IA DM IA TC IA TC IA
1100 °C
1220 °C
1260 °C
42.8 47.0 39.6 38.5 16.7 26.8 17.1 14.4
29.6 26.4 21.5 23.3 17.2 12.5 17.6 8.1
16.2 14.5 13.5 12.4 17.4 9.1 17.7 10.7
first calculated as a function of the temperature using Thermocalc™ software by taking into account the solubility of C and W atoms in the binder. The calculated values were compared with the volume fractions measured by image analysis to evaluate the accuracy of the method (Table 2). At 1100 °C, the fraction of binder is overestimated in WC– Co,C, likely due to the removal of WC grains during the grinding stage in this highly porous sample. For 1220 °C and 1260 °C, the volume of binder is underestimated in both alloys, as the thinnest Co areas between WC grains cannot be detected (Fig. 3). These latter are incorporated in the WC phase during image processing. Since the volume fraction of intercepts in the analysis is referred to the total volume of solid, spreading is then characterized by the decrease of the large intercept fraction which is not affected by the missing thin films. Fig. 5 shows the intercept length distribution of Co in the specimens after heating. The evolution of the distributions features the spreading of the binder when the temperature increases. The fraction of large intercepts decreases between 1220 °C and 1260 °C in WC–Co,C and between 1100 °C and 1220 °C in WC–Co,W. The increase observed at 1260 °C
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a
0.7
1100°C
WC-Co,C Volume fraction of pores
in WC–Co,W is interpreted as the decomposition of η' that releases some Co. Variation of intercept distributions in the binder phase at low temperature and comparison of the C rich and W rich alloys at 1220 °C (Fig. 5c) points out the earlier spreading of the binder in WC–Co,W.
+ 1220°C
0.6
1260°C
0.5 0.4
3.2. WC grain shape and orientation relationships of WC/Co interfaces
0.3
As one of the driving force for sintering is the decrease of interface energy, the structure of the WC/Co interfaces created in the first stages of sintering is an important parameter to characterise. As discussed above, this parameter will control the onset of spreading on the WC particle surface and hence densification at the solid state. The starting mixture of WC and Co powder as well as alloys heated up to 1100 °C and 1240 °C were studied by TEM. A heating rate of 1 °C/min was used. Before sintering, the WC powders are partly delimited by (0001) basal and (10–10) prismatic facets (Fig. 6). The grain surfaces at the local scale look either rounded or flat. After sintering at 1100 °C WC grains are much more facetted in both alloys. A lot of grains show basal and prismatic facets especially when they are in contact with the binder (Fig. 7). The transition between basal and prismatic facets is rather angular whatever the carbon content. From modelling and experiments on the kinetics of diffusion, it can be assumed that the binder composition is at the equilibrium in the alloys [14]. For example, at 1100 °C, the composition of the binder is close to Co–3.1C–2.4W (at%) in WC–Co,C and to Co–0.9C–9.0W (at%) in WC–Co,W from the calculation with Thermocalc™ software. Although the composition of the binder is different in the alloys, the grain shape is similar in agreement with the observations of [15]. The grain morphology is rather similar after heating at 1240 °C. These observations indicate that low energy WC/Co interfaces form early on heating the powder mixtures before the melting of the binder. A striking feature of the WC/Co interfaces observed after heating in the solid state is the frequency of special orientation relationships.
1100°C 1220°C
0.2
1260°C 0.1 0 0
2
4
6
8
10
Intercept length (µm)
b
0.7
1100°C
Volume fraction of pores
WC-Co,W
+ 1220°C
0.6
1260°C
0.5 0.4 0.3 0.2
1260°C
1220°C
1100°C
0.1 0 0
2
4
6
8
10
Intercept length (µm) Fig. 4. Intercept length distribution of pores in the alloys during the thermal treatment. The volume fraction of pores corresponds to the fraction of pores with reference to the total volume of solid (WC + Co). a) WC–Co,C. b) WC–Co,W. Arrows indicate the maximum intercept length at each temperature.
b
0.5
+ 1220°C
WC-Co,C
1260°C
0.4
Volume fraction of Co
Volume fraction of Co
a
0.3 0.2
1260°C
0.1
1220°C
0
0.5
1100°C
WC-Co,W
+ 1220°C
0.4
1260°C
0.3 0.2
1220°C
0.1
1100°C 1260°C
0 0
1
2
3
4
0
1
c
Volume fraction of Co
Intercept length (µm)
2
3
4
Intercept length (µm)
0.2
WC-Co,C
1220 C 0.15
WC-Co,W
0.1
WC-Co,W
WC-Co,C
0.05 0 0
1
2
3
4
Intercept length (µm) Fig. 5. Intercept length distribution of the binder in the alloys during the thermal treatment. The volume fraction of Co corresponds to the fraction of Co with reference to the total volume of solid (WC + Co). a) Intercept length distributions in WC–Co,C. b) in WC–Co,W. c) Comparison of the intercept length distributions in the alloys sintered at 1220 °C. Arrows indicate the maximum intercept length.
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Fig. 6. TEM images of the powder before sintering. The trace of the (0001) basal facet (B) or of the (10–10) prismatic facets (P) is drawn on the grains. a) WC grain viewed along b0001N showing some prismatic facets. b) WC grain viewed along b1–210N partly delimited by basal and prismatic facets. c) Polycrystalline powder consisting of two grains with a rather rounded shape. The insets correspond to the diffraction patterns of the grains.
Fig. 8 is an example of Co pools where the (111) plane of the Co lattice aligns parallel to the basal facet of the WC grain. The particular OR1 orientation ((0001)WC // (111)Co with [2–1–10]WC // [−110]Co) is often found either for the whole Co pool or for only a small Co area in contact with the WC grain. This OR1 relationship was already noticed for Co precipitates entrapped in WC grains [9] or for Co small areas mainly lying between WC grains after liquid phase sintering [7]. The orientation relationship called OR3 in [7] was also occasionally observed ((0001) WC // (001) Co with [2–1–10] WC // [110] Co). In other cases, the (001)Co or (111)Co planes were found parallel to the (0001) WC plane. It is to notice that no special orientation was recorded for the prismatic facet of WC. The WC/Co orientation relationship was determined for 25 WC basal facets in each alloy sintered at 1100 °C and 1240 °C (Fig. 9). The results indicate the large amount of special orientation relationships occurring in the alloys sintered in the solid state. This amount is higher in WC–Co,W alloy even at 1100 °C. 3.3. Atomic structure of the WC/Co interfaces In order to get a better insight on the interface atomic structure, the basal and prismatic WC/Co facets were observed at the atom scale in the alloys heated at 1100 °C and 1260 °C. In each specimen, several interfaces with a random orientation were chosen. It must be noticed that TEM foils were hardly thinned as the specimens were porous. The large thickness of the foils was often an obstacle for the observations
and the number of studied interfaces was about three in each sample. The observations depict various results according to the temperature and the composition. The investigations on the basal interface indicate that the last (0001) atomic plane in contact with Co is shifted from the regular hexagonal WC lattice position in WC–Co,W at 1100 °C and WC–Co,C at 1260 °C. In the other alloys, no change of the atom position is detected at the interface (Fig. 10). When the atoms are shifted, the two last planes of WC coincide with the cubic WC1-x lattice with the same orientation as for (Cr,W)C1-x films [16]. A value of 0.416 nm was found to fit with the lattice parameter of the cubic layer. This value corresponds to the smallest lattice parameter determined for the high temperature WC1-x cubic phase [17]. Elastic distortions are expected in the film in order to match with the lattice parameter of the hexagonal WC crystal as the lattice misfit is about 1.2% on the basal facet of WC grains. Such a thin cubic film is also detected on the prismatic facets of WC grains in WC–Co,C alloy sintered at 1260 °C. No WC1-x film is visible on the HRTEM images at WC/Co interfaces when the Co pool is oriented according to OR1. It can be due to the difficulty to detect the film as it is very thin and has a projected crystal lattice close to the one of the binder. It can also be related to the low energy of the interface with the OR1 relationship that would not favour the film formation. Fig. 11 is an example of a WC/Co interface with the OR1 orientation relationship in the WC–Co,W alloy heated at 1260 °C. A nearly periodic array of dislocations is observed at the interface with a mean spacing of 1.7 nm. The Burgers vector of the dislocations is the same
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Fig. 7. TEM images of WC grains viewed along b2–1–10N direction after heating (a) at 1100 °C in WC–Co,C alloy, (b) at 1100 °C in WC–Co,W alloy, (c) at 1240 °C in WC–Co,C alloy, (d) at 1240 °C in WC–Co,W alloy. The basal and prismatic facets that are parallel to the electron beam direction are denoted by B and P, respectively.
as the one found after liquid phase sintering, i.e. b = ½b−101NCo parallel to the interface [7]. After cooling, the lattice parameter of the binder is equal to 0.358 nm in WC–Co,W. At the interface, there is a parametric misfit between atomic planes in Co and WC, with 8d(−1–12)Co = 7d(01–10)WC = 1.7 nm where d(−1–12)Co and d(01–10)WC are the spacing between (−1–12) and (01–10) planes in Co and WC, respectively. The dislocations exactly compensate for the parametric misfit between Co and WC. This observation indicates that the interface structure is at the equilibrium. 4. Discussion The onset of sintering is delayed to higher temperature for the C rich alloys and the shrinkage rate is consequently larger at high temperature, as the remaining driving force is larger. Liquid phase sintering occurs at a lower temperature for the C rich alloys, according to thermodynamic calculations, and solid state sintering may also be enhanced due to a
higher diffusivity in the binder near the solidus temperature. Sintering at higher temperature for the C rich alloys results in a more uniform evolution of the microstructure. Also the initial presence of W particles in WC–Co,W locally induces the dissolution of WC and the formation of η' phase that decomposes on heating. These transformations favour local particle rearrangements resulting in the formation of dense agglomerates surrounded by large pores in this alloy. The earlier sintering observed in W rich alloys is explained by an earlier spreading of the binder phase in these alloys. Solubility differences in W rich or C rich alloys or η(η') phase formation could play a role. But the fraction of WC that is consumed for dissolution or for η(η') phase formation has been shown to be very small and similar in W rich or C rich alloys [6]. A justification for the earlier spreading in WC– Co,W may be the earlier ferromagnetic–paramagnetic transformation in this alloy. From thermodynamic data, the binder composition depends on the carbon activity and the amount of W in Co is much larger on the W rich side. The magnetic properties of the binder are related to
Fig. 8. TEM images of Co areas in contact with the basal facet (B) of a WC grain. In WC–Co,C after heating at 1240 °C: (a) bright field image, (b) dark field image using the (111)Co reflexion, (c) diffraction pattern of the WC and Co1 grains. In WC–Co,W after heating at 1100 °C (d) bright field image, (e) dark field image using the (111)Co reflexion, (f) diffraction pattern of the WC and Co grains displaying the OR1 orientation [7].
Fraction of special interfaces
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1
1240°C
1100°C 0.75
OR OR 0.5
OR 0.25
OR1
OR OR1 OR1
OR1
0 C-rich
W-rich
C-rich
W-rich
Fig. 9. Fraction of WC/Co interfaces at basal facets showing a special orientation relationship: OR1 corresponds to the relation (0001)WC // (111)Co with [2–1–10]WC // [−110]Co and OR is used when the interface plane is (0001)WC // (111)Co or (0001)WC // (001)Co.
the composition and an increase in W content results in a decrease of the Curie temperature [18]. In TiC–Co alloys, ab-initio calculations predict that the TiC/Co interface energy is smaller above the Curie temperature [19]. By analogy, the interface energy in WC–Co alloys is expected to decrease above 1051 °C in WC–Co,C and 933 °C in WC–Co,W, so spreading of the binder should be easier above these temperatures [6]. The TEM observations point out the frequency of preferential orientation relationships of the binder at basal facets of WC grains after solid state sintering. Such relationships were already identified in alloys after liquid phase sintering [7] and were shown to correspond to a low interface energy [20]. It was also proposed that coherent WC/Co interfaces could arise in alloys sintered by spark plasma sintering [10] while it is
Fig. 11. HRTEM image of a WC/Co interface with the OR1 orientation in WC–Co,W alloy after heating at 1260 °C.
difficult to conclude if they are formed in the liquid or solid state as the temperature distribution in the alloy is hardly controlled [21]. The present study indicates that the composition of the binder influences the occurrence of coherent interfaces after solid state sintering. They are twice as numerous in WC–Co,W. In this alloy, a higher amount of W atoms is dissolved in the binder. This induces higher elastic stresses and may facilitate the recrystallisation of stress-free cobalt grains with low interface energy with WC [6]. At the atomic scale, it is observed that the OR1 interfaces contain a periodic array of misfit dislocations which indicates that they are close to the equilibrium. Moreover, this
Fig. 10. HRTEM images of the WC/Co interface for a WC basal facet after heating (a) at 1100 °C in WC–Co,C alloy, (b) at 1100 °C in WC–Co,W alloy, (c) at 1260 °C in WC–Co,C alloy, (d) at 1260 °C in WC–Co,W alloy.
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OR1 orientation corresponds to the most stable WC/Co interface according to ab initio calculations [20,22]. The magnetic and the recrystallization effects could add their contributions to explain the earlier spreading of the binder in WC–Co,W. Effect of composition on atomic mobility or a dynamic coupling between dissolution and diffusion/spreading cannot be excluded but they are not so well documented. A cubic WC1-x film is observed at many WC/Co interfaces in solid state sintered alloys. It is questionable if the film is stable or if it forms on cooling as the solubility of WC in the binder decreases. According to simulations [14], it is expected to form on cooling even for the alloys sintered at 1100 °C. However, the presence of a film is characteristic of a lower energy for the WC/Co interface in the solid state. Therefore the formation of a cubic film at the interface may also facilitate Co spreading, but the presence of this film was not correlated to the C activity. 5. Conclusion The present experimental study emphasises the effect of the carbon activity on the microstructure evolution during solid state sintering of WC–Co alloys and gives a better insight on the WC/Co interface formation: • Spreading of the binder on WC grains and shrinkage starts earlier in W rich alloys during solid state sintering of WC–Co alloys. This may be explained by a lower interface energy in the paramagnetic state as shown by ab initio calculations in the literature. Indeed the transition from the ferromagnetic to the paramagnetic state occurs earlier in W rich alloys. In addition, many WC/Co interfaces adopt special orientation relationships at WC basal facets to ensure a good matching of Co lattice with WC grains, especially in W rich alloys. A diffusion-induced recrystallization of the binder at WC/Co interfaces may then also enhance the formation of low energy interfaces and spreading of the binder phase, as discussed in a previous paper [6]. It would be interesting to study the effect of the orientation relationship on the fracture resistance of the interface. If the effect was confirmed, solid state sintering could be considered to improve mechanical properties of the alloys. • Shrinkage is naturally enhanced at higher temperature in C rich alloys, due to the higher remaining driving force and reduction of porosity is more uniform in these alloys. Local effects related to W dissolution in W rich alloys may also explain the more heterogeneous porosity in these samples. • WC grains adopt a faceted shape on heating in the solid state. A thin cubic film is observed at some WC/Co facets whatever the
carbon activity. This may also facilitate spreading of the binder on WC grains. Acknowledgments The authors thank SANDVIK Hard Materials for technical and financial support. References [1] Allibert CH. Sintering features of cemented carbides WC–Co processed from fine powders. Int J Refract Met Hard Mater 2001;19:53–61. [2] Missiaen JM. Solid-state spreading and sintering of multiphase materials. Mater Sci Eng A 2008;475:2–11. [3] Haglund S, Ågren J, Uhrenius B. Solid-state sintering of cemented carbides. An experimental study. Z Metallkd 1998;89:316–22. [4] Delanoë A, Missiaen JM, Allibert CH, Lay S, Pauty E. Effect of the C potential and of Cr doping on the densification of alloys WC–Co. In: Kneringer G, Rodhammer P, Wilner H, editors. Proceedings of the 16th International Plansee Seminar, vol. 2; 2005. p. 642–52. [5] Petersson A. Sintering shrinkage of WC–Co and WC–(Ti, W)C–Co materials with different carbon contents. Int J Refract Met Hard Mater 2004;22:211–7. [6] Bounhoure V, Missiaen J-M, Lay S, Pauty E. Discussion of nonconventional effects in solid-state sintering of cemented carbides. J Am Ceram Soc 2009;92:1396–402. [7] Bounhoure V, Lay S, Loubradou M, Missiaen JM. Special WC/Co orientation relationships at basal facets of WC grains in WC–Co alloys. J Mater Sci 2008;43:892–9. [8] 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. [9] Mohan K, Strutt PR. Microstructure of spray converted nanostructured tungsten carbide–cobalt composite. Mater Sci Eng A 1996;209:237–42. [10] Song X, Gao Y, Liu X, Wei C, Wang H, Xu W. Effect of interfacial characteristics on toughness of nanocrystalline cemented carbides. Acta Mater 2013;61:2154–62. [11] Wei CB, Song XY, Fu J, Liu XM, Gao Y, Wang HB, et al. Microstructure and properties of ultrafine cemented carbides — differences in spark plasma sintering and sinterHIP. Mater Sci Eng A 2012;552:427–33. [12] Johansson T, Uhrenius B. Phase equilibriums, isothermal reactions, and a thermodynamic study in the cobalt–tungsten–carbon system at 1150 °C. Met Sci 1978;12:83–94. [13] Akesson L. An experimental and thermodynamic study of the Co–W–C system in the temperature range 1473–1698 K: University of Stockholm (Sweden); 1982. [14] Haglund S, Ågren J. W content in Co binder during sintering of WC–Co. Acta Mater 1998;46:2801–7. [15] Delanoë A, Lay S. Evolution of the WC grain shape in WC–Co alloys during sintering: effect of C content. Int J Refract Met Hard Mater 2009;27:140–8. [16] Delanoë A, Bacia M, Pauty E, Lay S, Allibert CH. Cr-rich layer at the WC/Co interface in Cr-doped WC–Co cermets: segregation or metastable carbide? J Cryst Growth 2004;270:219–27. [17] Nagender Naidu SV, Rama Rao P. Phase diagrams of binary tungsten alloys. Calcutta: Indian Institute of Metals; 1991. [18] Guillermet AF. Composition dependence of the Curie temperature of Co–W–C alloys predicted from thermodynamic data. Z Metallkd 1989;80:549–55. [19] Dudiy SV. Effects of Co magnetism on Co/TiC(0 0 1) interface adhesion. A 1stprinciples study. Surf Sci 2002;497:171–82. [20] Christensen M, Wahnström G, Lay S, Allibert CH. Morphology of WC grains in WC– Co alloys: theoretical determination of grain shape. Acta Mater 2007;55:1515–21. [21] Liu X, Song X, Zhang J, Zhao S. Temperature distribution and neck formation of WC– Co combined particles during spark plasma sintering. Mater Sci Eng A 2008;488:1–7. [22] Johansson S. A computational study of interface structures and energetics in cemented carbides and steels. Sweden: Chalmers University of Technology; 2010.
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Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Effect of C content on the microstructure evolution during early solid state sintering of WC–Co alloys V. Bounhoure a, S. Lay a,⁎, F. Charlot b, A. Antoni-Zdziobek a, E. Pauty c, J.M. Missiaen a a b c
SIMAP, Grenoble INP-CNRS-UJF, BP 75, 38402 Saint-Martin d'Hères, France CMTC, Grenoble INP, 38402 Saint Martin d'Hères, France SANDVIK Hard Materials, 51 avenue Rhin et Danube, 38100 Grenoble, France
a r t i c l e
i n f o
Article history: Received 30 July 2013 Accepted 20 December 2013 Available online 29 December 2013 Keywords: WC–Co alloys Cobalt spreading Solid state sintering WC/Co interfaces
a b s t r a c t The densification of WC–10Co (wt %) alloys is studied in relation with interface formation, below the melting temperature of the Co base binder phase. Two alloys containing an excess of carbon or tungsten are investigated. The evolution of the pore and Co phase size distribution is quantified at several temperatures using image analysis. While shrinkage and Co spreading start earlier in the W rich alloy, the reduction of porosity and Co spreading is more uniform in the C rich alloy. Preferential orientation relationships are frequently found at WC basal facets for WC/Co interfaces especially in the W rich alloy. A WC1-x film is also sometimes observed at randomly oriented WC/Co interfaces whatever the C content. The formation of low energy interfaces may explain the early spreading of the binder phase and the onset of densification in the solid state. © 2014 Elsevier Ltd. All rights reserved.
1. Introduction A large part of hard tool materials are cemented carbides based on WC–Co because they combine a high hardness and sufficient toughness. These materials are produced by powder metallurgy and consist of a skeleton of WC grains embedded in a Co base binder. The suitable binder exists for a tight range of C and W contents. According to the literature that is briefly summarised in [1,2], densification initially proceeds by WC particle rearrangement, induced by spreading of the binder phase between WC grains, first in the solid state then when the binder phase is liquid. Solid state rearrangement requires the previous spreading of Co on the WC particle surfaces [3], a process governed by the energy of the WC/Co interfaces. The carbon potential has an influence on the shrinkage behaviour in the solid state [4,5]. In particular, the shrinkage starts earlier in W rich alloys and it was proposed to be related to the magnetic properties of the binder that would influence the interface energy [6]. The energy also depends on the interface atomic structure. While few experimental data on the structure of WC/Co interfaces are available for alloys sintered by conventional methods with a liquid phase [7–9] or using a field assisted sintering technique [10,11], a scarce attention was paid to the interfaces formed in the solid state [6]. The present paper investigates the effect of carbon content on the microstructure of WC–10Co (wt%) alloys on heating in the solid state. Attention is focused on the effect of composition on the early spreading of the binder on WC grains and on the onset of solid state shrinkage. Shrinkage is followed by dilatometry. The porosity evolution and Co ⁎ Corresponding author. E-mail address: [email protected] (S. Lay). 0263-4368/$ – see front matter © 2014 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijrmhm.2013.12.012
spreading are quantified by image analysis. The WC/Co interfaces are characterised and their structure is observed at the atom scale by transmission electron microscopy. The WC/Co interface characteristics formed in the solid state are compared with interfacial structures obtained after liquid phase sintering. 2. Experimental The two-phase field (WC + Co binder) suitable for the cemented carbides is tight and its limits change with temperature [12]. Therefore, to study the C potential effect, two compositions located in the threephase fields (WC + Co binder + Cg) and (WC + Co binder + η) at 1400 °C were selected where Cg refer to the graphite and η to the mixed (W,Co)6C carbide [13]. The studied alloys called WC–Co,C and WC–Co,W, correspond to two values of carbon potential which are fixed along the process due to the presence of Cg or η phase. The alloys were prepared from WC and Co powders both with a mean size of 0.9 μm. W powder with a mean size of 0.7 μm or Cg were added in order to provide the required C potential in the alloys (Table 1). The powder mixtures were attritor milled for 5 h with acetone and polyethylene glycol (2 wt.%) as an organic binder. For the dilatometry analysis, cylinders (8 mm in diameter, about 3 g in mass) were obtained by uniaxial pressing at 200 MPa. The green density was 53% for both mixtures. Samples were sintered in a SETARAM™ Setsys Evolution dilatometer. A slow heating at 1 °C/min under H2/He atmosphere was first applied up to 400 °C for debinding, followed by heating up to 1400 °C at 3 °C/min or 1 °C/min under Ar atmosphere where samples were held for 30 min and then cooled down. Interrupted experiments were performed on heating for microstructural characterisation. The constitution of the
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Table 1 Compositions of the initial WC–Co,C and WC–Co,W powder mixtures. Content (at%)
Co
W
C
WC–Co,C WC–Co,W
15.1 15.1
41.6 43.3
43.3 41.6
alloys was studied by X-ray diffraction using Cu-Kα radiation. For scanning electron microscopy (SEM) observations, the specimens were embedded in a fluid resin under vacuum then grinded using a standard procedure. For transmission electron microscopy (TEM) and high resolution transmission electron microscopy (HRTEM) investigations, thin foils were prepared by mechanical grinding and Ar ion milling. The JEOL 3010 and 2010 microscopes were used for the observations. 3. Results 3.1. Microstructure and constitution evolution In order to characterize the microstructure evolution of the alloys during solid state shrinkage, interrupted experiments were scheduled from the dilatometric plots. Fig. 1 shows the linear shrinkage of the alloys at a heating rate of 3 °C/min. No significant difference was observed with a heating rate of 1 °C/min. Shrinkage starts above 800 °C and strongly accelerates in the 1050 °C–1100 °C temperature range. A significant difference is observed at the onset of the shrinkage that starts earlier in WC–Co,W. A special attention will be paid in the following to explain this effect. Shrinkage rate then becomes naturally larger at higher temperature for the WC–Co,C composition as the remaining driving force is larger. Above 1200 °C, the differences between the alloys are partly explained from the melting temperature lower for the C rich than W rich binder. The sharp peak of the shrinkage rate plots corresponds to the beginning of liquid phase sintering that occurs at 1293 °C for WC–Co,C and 1369 °C for WC–Co,W. These values agree with the solidus temperatures determined with Thermocalc™ software which are 1298 and 1368 °C, respectively. In the alloys, a maximum of shrinkage rate is observed during solid state sintering. It may be attributed to the end of primary particle rearrangement, when a dense packing of WC particles has formed [6]. Microstructural characterisations were conducted at 1100 °C, 1220 °C, 1240 °C and 1260 °C. 1100 °C corresponds to the beginning of sintering, before the first peak of solid state densification for both alloys. 1220 °C is after the first peak of densification for WC–Co,W and 1260 °C is close to the maximum of the
densification rate for WC–Co,C. Temperature of 1240 °C corresponds to an intermediate stage between the two maxima of the densification rate. The constitution of the powder mixtures was checked by X-ray diffraction (Fig. 2). In the starting powder, a peak of tungsten is observed in the mixture of the W rich alloy, as expected. In the WC–Co,C powder mixture, only the peaks associated with WC and Co are found as the carbon is not detected due to the small atomic scattering factor of C atoms. Moreover, only the cubic crystal lattice for Co is detected in the powder. The microstructure of the powder mixtures was observed by SEM (Fig. 3). The starting compacts show the presence of Co pools as large as 2 μm. A large amount of WC grains smaller than 0.4 μm is present. Alloys heated at 3 °C/min were used for the SEM characterizations. After heating at 1100 °C, large Co pools from initial powder mixtures are still present, but they form agglomerates with WC grains. The Co binder phase has started to spread between WC grains. Some grains in contact with Co are facetted. Co spreading is more important in WC–Co,W. Moreover, in this alloy, the amount of very small grains has decreased. They have likely been dissolved in the binder. The dissolution of WC affects the lattice parameter of Co that increases in both alloys (Fig. 2). The η' phase ((Co,W)12C) is detected by X ray diffraction in WC–Co,W and observed by SEM in agreement with thermodynamic data [13]. At 1220 °C, the densification level is higher. A continuous skeleton of WC grains has formed due to further Co spreading and rearrangement of WC grains under capillary forces. The remaining Co pools are smaller in the W rich alloy. The η' phase is still present. At 1260 °C, the microstructure of the two alloys is rather similar. The Co binder has spread further between WC grains has still increased. There seems to be a slightly higher amount of Co pools in WC–Co,W than at 1220 °C. From XRD analysis, the η' phase has disappeared and the equilibrium η phase is present. In order to get accurate information on the microstructure evolution of the alloys during solid state sintering, image analysis was performed from SEM pictures. Due to the high level of porosity in the alloys, it was difficult to polish the samples, so the validity of the method was first checked using the porosity parameter. The amounts of porosity deduced from image analysis and from the measured density of the samples were compared (Table 2). A good agreement was found between the two methods at each temperature. Image analysis was then used to quantify the evolution of the porosity in the alloys. Fig. 4 shows the intercept length distribution of the pores in WC–Co,C and WC–Co,W specimens for the experiments interrupted at 1100 °C, 1220 °C and 1260 °C. The porosity reduction is uniform in WC–Co,C on heating. The pore volume fraction starts to decrease at a lower temperature for
9
Tc(WC-Co,W)
6
Tc(WC-Co,C)
0.1
WC
-0.1
0
-0.2
-3
-0.3
-6
-0.4
-9
-0.5
-12
WC-Co,C
-15
WC-Co,W
-0.6
WC
η
Co WC η
η'
η η
η'
η'
intensity
Shrinkage (%)
3
Shrinkage rate (%/min)
0
WC-Co,W 1260°C WC-Co,W 1220°C
η' η
WC-Co,W 1100°C
W
WC-Co,W powder
Co WC-Co,C 1260°C WC-Co,C 1220°C
-0.7
-18 -21 600
700
800
900
WC-Co,C 1100°C
-0.8 1000 1100 1200 1300 1400
Temperature (°C) Fig. 1. Shrinkage and shrinkage rate of the studied alloys during the thermal treatment (heating rate of 3 °C/min). The Curie temperature of the binder is indicated for each alloy by an arrow and the melting of the binder by a star. Dotted vertical lines indicate the temperatures (1100, 1220, 1240 and 1260 °C) at which experiments were interrupted for microstructural characterisations.
WC-Co,C powder 30
35
40
45
50
2θ (°) Fig. 2. Evolution of the X-ray diagram of the studied alloys versus the temperature. The arrows indicate the Co peaks. The shift of the Co peaks is related to the dissolution of WC in the binder.
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Fig. 3. SEM images of the WC–Co,C and WC–Co,W powder compacts and alloys. WC appears in bright, Co in grey and pores in black.
the WC–Co,W alloy, but large pores about 2 μm in size are difficult to eliminate and are still present at 1260 °C. Spreading of the binder was also studied by image analysis from SEM images. The volume fraction of the binder in the solid phase was
Table 2 Volume fraction of porosity and binder (%vol) deduced from density measurements (DM) or calculated by Thermocalc™ software (TC) compared with the values measured by image analysis on SEM images (IA).
Pores
WC–Co,C WC–Co,W
Binder
WC–Co,C WC–Co,W
DM IA DM IA TC IA TC IA
1100 °C
1220 °C
1260 °C
42.8 47.0 39.6 38.5 16.7 26.8 17.1 14.4
29.6 26.4 21.5 23.3 17.2 12.5 17.6 8.1
16.2 14.5 13.5 12.4 17.4 9.1 17.7 10.7
first calculated as a function of the temperature using Thermocalc™ software by taking into account the solubility of C and W atoms in the binder. The calculated values were compared with the volume fractions measured by image analysis to evaluate the accuracy of the method (Table 2). At 1100 °C, the fraction of binder is overestimated in WC– Co,C, likely due to the removal of WC grains during the grinding stage in this highly porous sample. For 1220 °C and 1260 °C, the volume of binder is underestimated in both alloys, as the thinnest Co areas between WC grains cannot be detected (Fig. 3). These latter are incorporated in the WC phase during image processing. Since the volume fraction of intercepts in the analysis is referred to the total volume of solid, spreading is then characterized by the decrease of the large intercept fraction which is not affected by the missing thin films. Fig. 5 shows the intercept length distribution of Co in the specimens after heating. The evolution of the distributions features the spreading of the binder when the temperature increases. The fraction of large intercepts decreases between 1220 °C and 1260 °C in WC–Co,C and between 1100 °C and 1220 °C in WC–Co,W. The increase observed at 1260 °C
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a
0.7
1100°C
WC-Co,C Volume fraction of pores
in WC–Co,W is interpreted as the decomposition of η' that releases some Co. Variation of intercept distributions in the binder phase at low temperature and comparison of the C rich and W rich alloys at 1220 °C (Fig. 5c) points out the earlier spreading of the binder in WC–Co,W.
+ 1220°C
0.6
1260°C
0.5 0.4
3.2. WC grain shape and orientation relationships of WC/Co interfaces
0.3
As one of the driving force for sintering is the decrease of interface energy, the structure of the WC/Co interfaces created in the first stages of sintering is an important parameter to characterise. As discussed above, this parameter will control the onset of spreading on the WC particle surface and hence densification at the solid state. The starting mixture of WC and Co powder as well as alloys heated up to 1100 °C and 1240 °C were studied by TEM. A heating rate of 1 °C/min was used. Before sintering, the WC powders are partly delimited by (0001) basal and (10–10) prismatic facets (Fig. 6). The grain surfaces at the local scale look either rounded or flat. After sintering at 1100 °C WC grains are much more facetted in both alloys. A lot of grains show basal and prismatic facets especially when they are in contact with the binder (Fig. 7). The transition between basal and prismatic facets is rather angular whatever the carbon content. From modelling and experiments on the kinetics of diffusion, it can be assumed that the binder composition is at the equilibrium in the alloys [14]. For example, at 1100 °C, the composition of the binder is close to Co–3.1C–2.4W (at%) in WC–Co,C and to Co–0.9C–9.0W (at%) in WC–Co,W from the calculation with Thermocalc™ software. Although the composition of the binder is different in the alloys, the grain shape is similar in agreement with the observations of [15]. The grain morphology is rather similar after heating at 1240 °C. These observations indicate that low energy WC/Co interfaces form early on heating the powder mixtures before the melting of the binder. A striking feature of the WC/Co interfaces observed after heating in the solid state is the frequency of special orientation relationships.
1100°C 1220°C
0.2
1260°C 0.1 0 0
2
4
6
8
10
Intercept length (µm)
b
0.7
1100°C
Volume fraction of pores
WC-Co,W
+ 1220°C
0.6
1260°C
0.5 0.4 0.3 0.2
1260°C
1220°C
1100°C
0.1 0 0
2
4
6
8
10
Intercept length (µm) Fig. 4. Intercept length distribution of pores in the alloys during the thermal treatment. The volume fraction of pores corresponds to the fraction of pores with reference to the total volume of solid (WC + Co). a) WC–Co,C. b) WC–Co,W. Arrows indicate the maximum intercept length at each temperature.
b
0.5
+ 1220°C
WC-Co,C
1260°C
0.4
Volume fraction of Co
Volume fraction of Co
a
0.3 0.2
1260°C
0.1
1220°C
0
0.5
1100°C
WC-Co,W
+ 1220°C
0.4
1260°C
0.3 0.2
1220°C
0.1
1100°C 1260°C
0 0
1
2
3
4
0
1
c
Volume fraction of Co
Intercept length (µm)
2
3
4
Intercept length (µm)
0.2
WC-Co,C
1220 C 0.15
WC-Co,W
0.1
WC-Co,W
WC-Co,C
0.05 0 0
1
2
3
4
Intercept length (µm) Fig. 5. Intercept length distribution of the binder in the alloys during the thermal treatment. The volume fraction of Co corresponds to the fraction of Co with reference to the total volume of solid (WC + Co). a) Intercept length distributions in WC–Co,C. b) in WC–Co,W. c) Comparison of the intercept length distributions in the alloys sintered at 1220 °C. Arrows indicate the maximum intercept length.
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Fig. 6. TEM images of the powder before sintering. The trace of the (0001) basal facet (B) or of the (10–10) prismatic facets (P) is drawn on the grains. a) WC grain viewed along b0001N showing some prismatic facets. b) WC grain viewed along b1–210N partly delimited by basal and prismatic facets. c) Polycrystalline powder consisting of two grains with a rather rounded shape. The insets correspond to the diffraction patterns of the grains.
Fig. 8 is an example of Co pools where the (111) plane of the Co lattice aligns parallel to the basal facet of the WC grain. The particular OR1 orientation ((0001)WC // (111)Co with [2–1–10]WC // [−110]Co) is often found either for the whole Co pool or for only a small Co area in contact with the WC grain. This OR1 relationship was already noticed for Co precipitates entrapped in WC grains [9] or for Co small areas mainly lying between WC grains after liquid phase sintering [7]. The orientation relationship called OR3 in [7] was also occasionally observed ((0001) WC // (001) Co with [2–1–10] WC // [110] Co). In other cases, the (001)Co or (111)Co planes were found parallel to the (0001) WC plane. It is to notice that no special orientation was recorded for the prismatic facet of WC. The WC/Co orientation relationship was determined for 25 WC basal facets in each alloy sintered at 1100 °C and 1240 °C (Fig. 9). The results indicate the large amount of special orientation relationships occurring in the alloys sintered in the solid state. This amount is higher in WC–Co,W alloy even at 1100 °C. 3.3. Atomic structure of the WC/Co interfaces In order to get a better insight on the interface atomic structure, the basal and prismatic WC/Co facets were observed at the atom scale in the alloys heated at 1100 °C and 1260 °C. In each specimen, several interfaces with a random orientation were chosen. It must be noticed that TEM foils were hardly thinned as the specimens were porous. The large thickness of the foils was often an obstacle for the observations
and the number of studied interfaces was about three in each sample. The observations depict various results according to the temperature and the composition. The investigations on the basal interface indicate that the last (0001) atomic plane in contact with Co is shifted from the regular hexagonal WC lattice position in WC–Co,W at 1100 °C and WC–Co,C at 1260 °C. In the other alloys, no change of the atom position is detected at the interface (Fig. 10). When the atoms are shifted, the two last planes of WC coincide with the cubic WC1-x lattice with the same orientation as for (Cr,W)C1-x films [16]. A value of 0.416 nm was found to fit with the lattice parameter of the cubic layer. This value corresponds to the smallest lattice parameter determined for the high temperature WC1-x cubic phase [17]. Elastic distortions are expected in the film in order to match with the lattice parameter of the hexagonal WC crystal as the lattice misfit is about 1.2% on the basal facet of WC grains. Such a thin cubic film is also detected on the prismatic facets of WC grains in WC–Co,C alloy sintered at 1260 °C. No WC1-x film is visible on the HRTEM images at WC/Co interfaces when the Co pool is oriented according to OR1. It can be due to the difficulty to detect the film as it is very thin and has a projected crystal lattice close to the one of the binder. It can also be related to the low energy of the interface with the OR1 relationship that would not favour the film formation. Fig. 11 is an example of a WC/Co interface with the OR1 orientation relationship in the WC–Co,W alloy heated at 1260 °C. A nearly periodic array of dislocations is observed at the interface with a mean spacing of 1.7 nm. The Burgers vector of the dislocations is the same
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Fig. 7. TEM images of WC grains viewed along b2–1–10N direction after heating (a) at 1100 °C in WC–Co,C alloy, (b) at 1100 °C in WC–Co,W alloy, (c) at 1240 °C in WC–Co,C alloy, (d) at 1240 °C in WC–Co,W alloy. The basal and prismatic facets that are parallel to the electron beam direction are denoted by B and P, respectively.
as the one found after liquid phase sintering, i.e. b = ½b−101NCo parallel to the interface [7]. After cooling, the lattice parameter of the binder is equal to 0.358 nm in WC–Co,W. At the interface, there is a parametric misfit between atomic planes in Co and WC, with 8d(−1–12)Co = 7d(01–10)WC = 1.7 nm where d(−1–12)Co and d(01–10)WC are the spacing between (−1–12) and (01–10) planes in Co and WC, respectively. The dislocations exactly compensate for the parametric misfit between Co and WC. This observation indicates that the interface structure is at the equilibrium. 4. Discussion The onset of sintering is delayed to higher temperature for the C rich alloys and the shrinkage rate is consequently larger at high temperature, as the remaining driving force is larger. Liquid phase sintering occurs at a lower temperature for the C rich alloys, according to thermodynamic calculations, and solid state sintering may also be enhanced due to a
higher diffusivity in the binder near the solidus temperature. Sintering at higher temperature for the C rich alloys results in a more uniform evolution of the microstructure. Also the initial presence of W particles in WC–Co,W locally induces the dissolution of WC and the formation of η' phase that decomposes on heating. These transformations favour local particle rearrangements resulting in the formation of dense agglomerates surrounded by large pores in this alloy. The earlier sintering observed in W rich alloys is explained by an earlier spreading of the binder phase in these alloys. Solubility differences in W rich or C rich alloys or η(η') phase formation could play a role. But the fraction of WC that is consumed for dissolution or for η(η') phase formation has been shown to be very small and similar in W rich or C rich alloys [6]. A justification for the earlier spreading in WC– Co,W may be the earlier ferromagnetic–paramagnetic transformation in this alloy. From thermodynamic data, the binder composition depends on the carbon activity and the amount of W in Co is much larger on the W rich side. The magnetic properties of the binder are related to
Fig. 8. TEM images of Co areas in contact with the basal facet (B) of a WC grain. In WC–Co,C after heating at 1240 °C: (a) bright field image, (b) dark field image using the (111)Co reflexion, (c) diffraction pattern of the WC and Co1 grains. In WC–Co,W after heating at 1100 °C (d) bright field image, (e) dark field image using the (111)Co reflexion, (f) diffraction pattern of the WC and Co grains displaying the OR1 orientation [7].
Fraction of special interfaces
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1
1240°C
1100°C 0.75
OR OR 0.5
OR 0.25
OR1
OR OR1 OR1
OR1
0 C-rich
W-rich
C-rich
W-rich
Fig. 9. Fraction of WC/Co interfaces at basal facets showing a special orientation relationship: OR1 corresponds to the relation (0001)WC // (111)Co with [2–1–10]WC // [−110]Co and OR is used when the interface plane is (0001)WC // (111)Co or (0001)WC // (001)Co.
the composition and an increase in W content results in a decrease of the Curie temperature [18]. In TiC–Co alloys, ab-initio calculations predict that the TiC/Co interface energy is smaller above the Curie temperature [19]. By analogy, the interface energy in WC–Co alloys is expected to decrease above 1051 °C in WC–Co,C and 933 °C in WC–Co,W, so spreading of the binder should be easier above these temperatures [6]. The TEM observations point out the frequency of preferential orientation relationships of the binder at basal facets of WC grains after solid state sintering. Such relationships were already identified in alloys after liquid phase sintering [7] and were shown to correspond to a low interface energy [20]. It was also proposed that coherent WC/Co interfaces could arise in alloys sintered by spark plasma sintering [10] while it is
Fig. 11. HRTEM image of a WC/Co interface with the OR1 orientation in WC–Co,W alloy after heating at 1260 °C.
difficult to conclude if they are formed in the liquid or solid state as the temperature distribution in the alloy is hardly controlled [21]. The present study indicates that the composition of the binder influences the occurrence of coherent interfaces after solid state sintering. They are twice as numerous in WC–Co,W. In this alloy, a higher amount of W atoms is dissolved in the binder. This induces higher elastic stresses and may facilitate the recrystallisation of stress-free cobalt grains with low interface energy with WC [6]. At the atomic scale, it is observed that the OR1 interfaces contain a periodic array of misfit dislocations which indicates that they are close to the equilibrium. Moreover, this
Fig. 10. HRTEM images of the WC/Co interface for a WC basal facet after heating (a) at 1100 °C in WC–Co,C alloy, (b) at 1100 °C in WC–Co,W alloy, (c) at 1260 °C in WC–Co,C alloy, (d) at 1260 °C in WC–Co,W alloy.
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OR1 orientation corresponds to the most stable WC/Co interface according to ab initio calculations [20,22]. The magnetic and the recrystallization effects could add their contributions to explain the earlier spreading of the binder in WC–Co,W. Effect of composition on atomic mobility or a dynamic coupling between dissolution and diffusion/spreading cannot be excluded but they are not so well documented. A cubic WC1-x film is observed at many WC/Co interfaces in solid state sintered alloys. It is questionable if the film is stable or if it forms on cooling as the solubility of WC in the binder decreases. According to simulations [14], it is expected to form on cooling even for the alloys sintered at 1100 °C. However, the presence of a film is characteristic of a lower energy for the WC/Co interface in the solid state. Therefore the formation of a cubic film at the interface may also facilitate Co spreading, but the presence of this film was not correlated to the C activity. 5. Conclusion The present experimental study emphasises the effect of the carbon activity on the microstructure evolution during solid state sintering of WC–Co alloys and gives a better insight on the WC/Co interface formation: • Spreading of the binder on WC grains and shrinkage starts earlier in W rich alloys during solid state sintering of WC–Co alloys. This may be explained by a lower interface energy in the paramagnetic state as shown by ab initio calculations in the literature. Indeed the transition from the ferromagnetic to the paramagnetic state occurs earlier in W rich alloys. In addition, many WC/Co interfaces adopt special orientation relationships at WC basal facets to ensure a good matching of Co lattice with WC grains, especially in W rich alloys. A diffusion-induced recrystallization of the binder at WC/Co interfaces may then also enhance the formation of low energy interfaces and spreading of the binder phase, as discussed in a previous paper [6]. It would be interesting to study the effect of the orientation relationship on the fracture resistance of the interface. If the effect was confirmed, solid state sintering could be considered to improve mechanical properties of the alloys. • Shrinkage is naturally enhanced at higher temperature in C rich alloys, due to the higher remaining driving force and reduction of porosity is more uniform in these alloys. Local effects related to W dissolution in W rich alloys may also explain the more heterogeneous porosity in these samples. • WC grains adopt a faceted shape on heating in the solid state. A thin cubic film is observed at some WC/Co facets whatever the
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