Effect of Cr addition on solid state sintering of WC–Co alloys

Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

Contents lists available at ScienceDirect

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

Effect of Cr addition on solid state sintering of WC–Co alloys V. Bounhoure a, S. Lay a,⁎, S. Coindeau a,b, S. Norgren c, E. Pauty d, J.M. Missiaen a a

CNRS & Université de Grenoble-Alpes, SIMAP, F-38000 Grenoble, France Université de Grenoble-Alpes, CMTC, F-38000 Grenoble, France c Sandvik Mining R&D Rock Tools, SE-126 80 Stockholm, Sweden d Sandvik Hard Materials, F-38100 Grenoble, France b

a r t i c l e

i n f o

Article history: Received 14 January 2015 Received in revised form 23 April 2015 Accepted 2 May 2015 Available online 2 May 2015 Keywords: WC–Co alloys Cr addition Solid-state sintering Binder spreading Cr oxides

a b s t r a c t The effect of Cr addition and C content on densification and microstructural evolution of WC–Co alloys in the solid state was studied. Two alloys containing Cr and an excess of carbon or tungsten were investigated. They were compared with Cr free alloys. The intercept length distribution of the pore and binder phase was measured at several temperatures using image analysis. Cr addition delays the shrinkage whatever the C potential whilst better binder spreading and more effective rearrangement occur at lower temperature in the W rich alloy. When the temperature increases, the densification is accelerated in the C rich alloy. It is found that Cr addition has no significant effect on the solubility of W in the binder. The influence of Cr on the shrinkage behaviour is discussed putting emphasis on the grain boundary energies and on Cr oxides. © 2015 Elsevier Ltd. All rights reserved.

1. Introduction WC–Co materials are mainly used to manufacture drilling, construction or mining tools. They are particularly adapted to achieve cutting tools owing to their high hardness combined with sufficient ductility. These alloys are prepared by the powder metallurgy route. Limiting grain growth during the sintering process is necessary to get finegrained alloys and improve mechanical properties [1]. The optimized microstructure results from a well-controlled carbon potential during sintering as well as the use of grain growth inhibitors [2,3]. Small additions of carbides like VC, TaC, NbC or Cr3C2 have proved to reduce grain growth [4]. It was also established that these additives influence the shrinkage behaviour. The major part of shrinkage occurs in the solid state for submicron sized cemented carbides as summarized in [5]. Shrinkage starts with the wetting of WC grains by cobalt and is followed by the spreading of the binder into the porosity under the action of capillary forces, leading to the rearrangement of WC grains [6–11]. The spreading/rearrangement process proceeds by steps and sequential pore filling occurs [12–14]. An excess of W shifts the onset of densification at lower temperature by promoting the spreading of the binder onto WC grains. It was proposed that this effect is related to a reduction of the interface energy owing to a lower Curie temperature of the W rich binder [15,16]. Besides, addition of VC or Cr3C2 delays the shrinkage in the solid state compared with undoped alloys [17–21]. Whilst thermal analyses emphasize the reduction of chromium oxides in a temperature

⁎ Corresponding author.

http://dx.doi.org/10.1016/j.ijrmhm.2015.05.002 0263-4368/© 2015 Elsevier Ltd. All rights reserved.

range higher than those of Co and WC oxides, the exact effect of Cr addition on the retardation of densification at low temperature is still a question. Since carbon content also influences solid state shrinkage, this study aims to investigate the combined effect of Cr3C2 addition and carbon potential on solid state sintering of WC–Co alloys by studying the microstructure evolution of Cr doped alloys. The influence of Cr doping on the composition of the binder and Curie temperature is considered. 2. Experimental procedure Four mixtures were prepared from WC and Co powders both with a FSSS average particle size of 0.9 μm (Table 1). Cr3C2 powder with an average particle size of 3.1 μm was added in two powder mixtures to get the ratio Cr/(Cr + Co) equal to 5 at.%, that is below the solubility limit of Cr in the liquid binder [22]. W powder with an average size of 0.7 μm or graphite was added in order to fix the C potential in the alloys. As a result, the alloys are located in the three-phase fields (WC + Co (Cr) binder + Cg) and (WC + Co (Cr) binder + η) at 1400 °C where Cg refers to the graphite and η to the mixed (W,Co)6C or (W,Cr,Co)6C carbide. The studied alloys are called WC–Co,C and WC–Co,W, WC– Co,Cr,C and WC–Co,Cr,W. The powder mixtures were attritor milled for 5 h with acetone and polyethylene glycol (2 wt.%) as an organic binder. Thermogravimetric analyses (TGA) and dilatometry experiments were performed with a SETARAM™ Setsys 16/18 device and SETARAM™ Setsys Evolution dilatometer, respectively. Cylinders (8 mm in diameter, about 3 g in mass) obtained by uniaxial pressing at 200 MPa were used for these experiments. The green density was

22

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

Table 1 Composition of the initial powder mixtures. Content (at.%)

Co

W

C

Cr

WC–Co,C WC–Co,Cr,C WC–Co,W WC–Co,Cr,W

15.1 14.3 15.1 14.3

41.6 41.6 43.3 43.3

43.3 43.3 41.6 41.6

– 0.8 – 0.8

0

0

Mass loss (%)

-0.5 -1

W oxydes reduction

Co oxydes reduction

Cr oxydes reduction

-0.01

PEG elimination

-1.5

-0.02

WC-Co,Cr,C

-2

-2.5

Mass loss

-0.03

Mass loss rate (%/min)

Rate

WC-Co,Cr,W

-3 0

200

400

600

800

1000

1200

-0.04 1400

by heating at 3 °C/min up to 1400 °C under Ar atmosphere where samples were held for 30 min and then cooled down at a rate of 30 °C/min. Interrupted sintering experiments were then performed in the dilatometer at 1100, 1220 and 1260 °C with a cooling rate of 30 °C/min. It will be seen that each temperature corresponds to a specific step of the solid state sintering according to the dilatometric curves. The samples were cut, embedded under vacuum in a fluid resin to fill the porosity and polished. Micrographs were taken using a field electron gun scanning electron microscope (FEG-SEM) equipped with an X-ray energy dispersive spectroscopy (EDS) device. The observations were performed in the centre of the sample to avoid the decarburisation layer (≈ 100 μm) appearing after sintering under the surface of samples. The constitution of the alloys was studied by X-ray diffraction (XRD) using Cu-Kα radiation. Microanalyses of the binder phase were carried out by means of a transmission electron microscope (TEM) equipped with a X-ray EDS device and using a beam size of 20 nm. The amount of Co, Cr and W atoms was determined in between ten and twenty Co pools in each alloy. The analyses are not quantitative as no standard was used to determine with accuracy the Cliff Lorimer factors of the EDS system [23]. They were only used to compare the studied alloys. 3. Results

Temperature (°C)

between 51.9 and 53.3% for the mixtures. Differential thermal analyses (DTA) were conducted to determine the melting temperature of the binder using 0.4 g of the powder mixtures. The same thermal cycle was applied for all experiments. A slow heating at 1 °C/min under H2/ He atmosphere was first applied up to 400 °C for debinding, followed

6

-0,2

Shrinkage (%)

3 0

-0,6

-3

-0,8

-6

-1

-9 -12 -15

WC-Co,C WC-Co,Cr,C

-1,4 -1,6

-21

-1,8 1400

900

1000 1100 1200 Temperature (°C)

1300

9

-0,4

1220°C

0

-0,6

-3

-0,8

-6 -12

1260°C

WC-Co,Cr,C

-1 -1,2

WC-Co,Cr,W

-15

-1,4

-18

-1,6

-21 800

0

6

-0,2

3

-0,4

0

-0,6

-3

-0,8

-6

-1

-9 -12 -15

WC-Co,W

-1,2

WC-Co,Cr,W

-1,4 -1,6

-18 -21 800

900

1000 1100 1200 Temperature (°C)

1300

-1,8 1400

-0,2

1100°C

3

-9

9

0

6

Shrinkage (%)

-1,2

-18

800

c

-0,4

b

Shrinkage (%)

0

Shrinkage rate (%/min)

9

900

1000 1100 1200 Température (°C)

1300

Shrinkage rate (%/min)

a

According to the literature, the onset of solid state spreading of cobalt on WC grains and start of shrinkage is dependent on the reduction of the oxides at the surface of the powder particles [24,25]. The reduction temperature of the oxides present in the powder mixtures was determined by TGA experiments in order to interpret the evolution of the microstructures. The main processes leading to a weight loss are indexed on the typical TGA curve of Fig. 1. The weight loss was principally due to the debinding of the samples that represents almost 80% of the total weight loss. The main other contributions result from the

Shrinkage rate (%/min)

Fig. 1. Thermogravimetric plot of the WC–Co,Cr,C and WC–Co,Cr,W alloys.

-1,8 1400

Fig. 2. Dilatometric plots of the undoped and Cr doped WC–Co alloys. Triangles denote peaks on the shrinkage rate occurring in the solid state whilst a star is associated with the presence of a liquid phase. (a) Effect of Cr addition in the C rich alloy. (b) Effect of Cr addition in the W-rich alloy. (c) Comparison of the Cr doped alloys. Three temperatures used to study the microstructures are indicated on the graph.

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

WC–Co,C WC–Co,Cr,C WC–Co,W WC–Co,Cr,W

Tonset (°)

TDil (°)

TDTA (°)

TCalc (°)

924 1040 850 935

1293 ± 1 1277 ± 1 1368 ± 2 1323 ± 1

1299 ± 5 1257 ± 20 1369 ± 5 1326 ± 5

1298 1223 1368 1272

WC-Co,C

ΔT

Table 2 Temperatures for the onset of solid state shrinkage (Tonset) measured by dilatometry and for liquid formation determined from different methods: TDil is measured on the dilatometric plots, at the onset of the shrinkage rate peak corresponding to the liquid phase shrinkage and TDTA is measured on the DTA plot when the curve starts to deviate from a base-line. The solidus temperature, TCalc, is calculated from Thermocalc™ software [22,29].

WC-Co,W WC-Co,Cr,C

WC-Co,Cr,W

1200

reduction of cobalt, tungsten and also chromium oxides in the doped samples. The reduction is observed at around 200 °C for cobalt oxides, in the 400–700 °C and 800–1100 °C temperature range for tungsten oxides and chromium oxides, respectively, in agreement with literature results [25–28]. From the TGA plots of WC–Co,C and WC–Co,W samples, it is observed that the carbon content has nearly no effect on the reduction temperatures of cobalt and WC oxides [15]. On the other hand, the reduction of Cr oxides in WC–Co,Cr,C sample is shifted to a temperature range about 100 °C lower than in the WC–Co, Cr,W sample. Note that the continuous mass loss occurring at higher temperature is likely related to the decarburisation of the alloy due to the presence of oxygen traces in the Ar gas. The combined effect of Cr addition and C content on the densification behaviour was studied using dilatometry (Fig. 2). A Cr addition delays the densification whatever the carbon potential. The onset of solid state shrinkage is shifted by about 100 °C compared to the undoped alloys (Table 2). Although C addition promotes the reduction of Cr oxides at a lower temperature, shrinkage starts earlier in the W rich alloy. When the temperature increases, one or two peaks are observed for the shrinkage rate in the solid state for all alloys. The presence of Cr lowers the temperature of liquid formation and reduces the solid state sintering stage (Table 2). These results are in agreement with previous experimental studies [17,20]. The effect of composition on the temperature of liquid formation is also consistent with thermodynamic calculations [22,29] but the experimental temperature is higher than predicted for Cr doped alloys. As also noticed with TGA experiments, decarburisation occurs during the thermal treatment and the composition of the alloys slightly evolves towards a lower C content as the temperature increases. However, the value of the solidus temperature should not be drastically modified for the studied alloys (Fig. 3). DTA measurements also indicate a higher solidus temperature than the calculations (Fig. 4).

Fig. 3. Phase diagram of the (W,C,Co,Cr) system calculated for a Co content of 15.1 at.% and Cr content of 0.8 at.% [22]. The composition of WC–Co,Cr,C and WC–Co,Cr,W before the experiment (solid lines) is indicated on the graph. The dotted line shows the measured composition of the WC–Co,Cr,C alloy after the dilatometry test.

23

1250

1300

1350

1400

Temperature (°C) Fig. 4. DTA plots of the undoped and Cr doped WC–Co alloys. The temperature of liquid formation is indicated by the vertical line.

The constitution and the microstructure evolution of the alloys along the thermal cycle in the solid state were investigated using interrupted experiments (Fig. 2c). 1100 °C corresponds to the beginning of sintering for the alloys and is located before the first peak of densification of all alloys. For the Cr doped alloys, 1220 °C is at the onset of the first peak of densification whilst 1260 °C corresponds to the maximum shrinkage rate for WC–Co,Cr,W and is located after the maximum for WC–Co,Cr,C. The powder mixtures were characterized by XRD and show similar features. The Cr3C2 phase was not detected. In the sintered samples, only Co and WC are found up to 1260 °C in WC–Co,Cr,C alloy. The additional phases M6C (η) or M12C (η′) with M = (W,Co(,Cr)) are present in the W rich alloys. The mixed carbide η is found in WC–Co,Cr,W above 1100 °C. In WC–Co,W alloy, η′ is present at 1100 °C and 1220 °C and η at 1260 °C (Fig. 5) [15]. Characterization of the microstructure is also important to study the densification of the alloys [6–11]. The effect of Cr addition on the microstructure evolution may be qualitatively deduced by examination of the SEM images. The powder mixtures were first examined. In the Cr doped alloys, Cr3C2 grains with a size of 1–2 μm are observed in the material after milling (Fig. 6). The WC grain size is smaller than 1 μm and numerous small WC grains in the range size of 200–400 nm are present. On the other hand, most Co pools have a size close to 1 μm. In the samples heated at 1100 °C, Co spreading is delayed in Cr doped alloys especially in the C rich sample (Fig. 7). In WC–Co,Cr,W, locally densified WC–Co clusters are present but their size is smaller than in WC–Co,W. At 1220 °C, large Co pools are still present in WC–Co,Cr,C whilst spreading is accelerated between 1220 °C and 1260 °C in this alloy as the temperature is close to the solidus temperature for this composition. Shrinkage at 1260 °C results in the formation of large pores in WC–Co,Cr,C compared to WC–Co,C. The microstructure of the W rich

Fig. 5. Evolution of the X-ray diagram of the Cr doped 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.

24

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

Fig. 6. SEM images of the WC–Co,Cr,W powder mixture after compaction. (a) Back scattered electron image, (b) EDS map showing the distribution of Cr3C2 grains.

samples is rather similar at 1220 °C and 1260 °C but larger pores form between 1100 °C and 1220 °C in WC–Co,Cr,W. Image analysis was performed from SEM pictures in order to get a quantitative analysis of the porosity evolution and Co spreading during solid state sintering. The validity of the method was first analysed by checking the accuracy of porosity and Co volume fractions estimated by image analysis. The porosity measured by image analysis was compared to the value determined from the density of the samples. The measured Co volume fraction was compared to the value

1220°C

1260°C

W,Cr

W

C,Cr

C

1100°C

deduced from the initial composition. A good agreement was found for the two parameters except for the C rich specimens sintered at 1100 °C. The mismatch was due to the difficulty to get safe polished samples in these highly porous alloys. The intercept length distributions in the pore phase and in the binder phase were determined in the other specimens. The volume fraction of Co is underestimated at higher temperature in all alloys. This effect is likely due to the formation of thin Co pools and films that cannot be detected on SEM images and are incorporated in the WC phase during image processing.

Fig. 7. SEM images of the microstructure in the alloys heated at 1100 °C, 1220 °C and 1260 °C.

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

b

0,7

1100°C

WC-Co,W

0,6

Volume fraction of pores

Volume fraction of pores

a

WC-Co,Cr,W

0,5 0,4 0,3

WC-Co,Cr,W

0,2

WC-Co,W

0,1 0

0,3

WC-Co,Cr,W 0,2

WC-Co,Cr,W

0,1

WC-Co,W

2

4

6

8

10

0

2

4

Intercept length (µm)

WC-Co,W

0,15

WC-Co,Cr,W

0,12 0,09 0,06

WC-Co,W WC-Co,Cr,W

0,03

d

0,7

Volume fraction of pores

Volume fraction of pores

0,18

1260°C

6

8

10

12

Intercept length (µm)

0,6

0

1100°C 1220°C 1260°C

WC-Co,Cr,W

0,5 0,4 0,3

1220°C

0,2 0,1

1100°C

1260°C

0 0

2

4

6

8

10

0

2

4

Intercept length (µm)

f 0,3

Volume fraction of pores

1220°C

WC-Co,C WC-Co,Cr,C

0,2

0,1

WC-Co,C

6

8

10

12

Intercept length (µm)

e Volume fraction of pores

WC-Co,W

1220°C

0 0

c

25

WC-Co,Cr,C

0

0,18

1260°C

WC-Co,C

0,15

WC-Co,Cr,C 0,12 0,09

WC-Co,Cr,C 0,06

WC-Co,C 0,03 0

0

5

10

15

20

0

Intercept length (µm)

2

4

6

8

10

Intercept length (µm)

Fig. 8. Evolution of the intercept length size distribution of pores in the alloys during solid state sintering. (a–d) Comparison of the W rich alloys versus the temperature. (e–f) Comparison of the C rich alloys. Arrows indicate the maximum intercept length for each alloy.

The volume fractions are then referred to the total volume of solid (WC + Co + η or η′), so that Co spreading is characterized by a decrease of the large intercept fraction which is not affected by the missing films. The smallest intercept lengths in the binder phase will not be considered in the interpretation. Fig. 8 shows the effect of Cr addition on the intercept length distribution in the pore phase. A larger volume fraction of pores and a wider distribution are observed at each temperature in the Cr doped alloy, indicating a retardation of the densification. The reduction of porosity is not uniform in the W rich alloy as the fraction of large pores increases with the temperature (Fig. 8d). The elimination of pores between 1220 °C and 1260 °C is higher in WC–Co,Cr,C than in WC–Co,Cr,W owing to the lower melting temperature of the binder in the C rich alloy that accelerates the densification process. The effect of Cr addition on the intercept length distribution of the binder phase was also studied. In the W rich alloys, the larger intercept values recorded in the Cr doped alloy at 1100 °C and 1220 °C point out the delay of binder spreading in this alloy (Fig. 9). On the other hand, at 1260 °C, the spreading of the binder is better achieved in the Cr doped alloy. The comparison of the graphs (Fig. 9d) as a function of the temperature indicates a uniform thickness reduction of the binder in WC– Co,Cr,W.

Regarding the C rich alloys, the binder is slightly coarser at 1220 °C in the Cr doped alloy and more or less similar at 1260 °C in both alloys (Fig. 9e–f). Comparing Cr doped alloys shows that spreading is better achieved at 1220 °C in W rich alloy compared to C rich alloy and similar at 1260 °C (Fig. 9g–h). 4. Discussion The dilatometry and microstructural investigations reveal that Cr addition influences the solid state densification of WC–Co alloys whatever the C content. It is also observed that solid state shrinkage is delayed both in W and C rich alloy. Solid-state shrinkage starts with the wetting of WC grains by the Co binder, and is followed by the spreading of the binder into the porosity. This leads to the rearrangement of WC grains. The onset of shrinkage depends on the ability of the binder to wet WC grains. Several tracks may be explored to understand the effect of Cr addition on Co wetting. Formation of η or η′ carbides is observed in WC–Co,W and WC– Co,Cr,W (Fig. 5) [15]. However, in WC–Co,W, the fraction of WC that is consumed for phase formation remains very low [16]. So the effect on the onset of sintering is probably limited.

26

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

b

0,5

1100°C

WC-Co,W

0,4

Volume fraction of Co

Volume fraction of Co

a

WC-Co,Cr,W

0,3 0,2

WC-Co,Cr,W WC-Co,W

0,1

0,5

1220°C

WC-Co,Cr,W

0,3 0,2

WC-Co,W

0,1

0 1

2

3

4

5

6

0

Intercept length (µm)

d Volume fraction of Co

Volume fraction of Co

WC-Co,Cr,W

0,3 0,2

WC-Co,W WC-Co,Cr,W

0 1

2

3

0,3 0,2

1100°C 1220°C 1260°C

0,1

4

0

f

0,3

WC-Co,Cr,C WC-Co,C

0,1

6

0,5

Volume fraction of Co

WC-Co,Cr,C

0,2

4

1260°C

WC-Co,C

0,4

2

Intercept length (µm)

0,5

1220°C Volume fraction of Co

4

1100°C 1220°C 1260°C

WC-Co,Cr,W 0,4

Intercept lenght (µm)

WC-Co,C

0,4

WC-Co,Cr,C

0,3

WC-Co,C

0,2

WC-Co,Cr,C

0,1 0

0 0

2

4

6

0

8

h 1220°C

Volume fraction of Co

0,5 0,4

WC-Co,Cr,C

WC-Co,Cr,W 0,3 0,2

1

2

3

Intercept length (µm)

Intercept length (µm)

Volume fraction of Co

3

0

0

g

2

0,5

WC-Co,W

0,4

0,1

1

Intercept length (µm)

0,5

1260°C

e

WC-Co,Cr,W

0 0

c

WC-Co,W

0,4

WC-Co,Cr,C WC-Co,Cr,W

0,1

0

0,5

1260°C

0,4

WC-Co,Cr,C WC-Co,Cr,W

0,3

WC-Co,Cr,C

0,2

WC-Co,Cr,W

0,1

0 0

2

4

6

8

Intercept length (µm)

0

1

2

3

Intercept length (µm)

Fig. 9. Evolution of the intercept length distribution of the binder in the alloys during solid state sintering. (a–d) Comparison of the W rich alloys versus the temperature. (e–f) Comparison of the C rich alloys. (g–h) Comparison of the Cr doped alloys. Arrows indicate the maximum intercept length for each alloy.

Table 3 Mean composition of the binder measured by TEM/EDS (at.%) in the alloys heated at 1220 °C.

W Cr

WC–Co,C

WC–Co,Cr,C

WC–Co, W

WC–Co,Cr,W

1.5 ± 0.4 –

2.1 ± 0.7 3.9 ± 0.2

9.2 ± 1.3 –

8.0 ± 1.2 3.1 ± 0.5

The solubility of W in the binder influences the spreading/rearrangement process at the solid state. Indeed the grain dissolution leads to a more important free volume and may improve the WC grain rearrangement. A high solubility also improves the shape accommodation processes and so the reorganization of WC grains. By analysing the Co binder after extraction it was found that an addition of Cr3C2 lowers

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

Fig. 10. TEM image of a cobalt region surrounded by WC grains in WC–Co,Cr,W alloy sintered at 1220 °C. EDS analyses were carried out in the binder (1–2) and at the WC/Co interface (3). The results are shown in Table 4.

the W content of the binder in alloys sintered at 1200 °C [20]. This result is in contradiction with the analyses of diffusion couples showing no significant influence of Cr on the WC solubility at 1200 °C [30]. To further understand the effect of Cr, the composition of the binder was investigated in the alloys heated at 1220 °C by TEM/EDS (Table 3). Considering the confidence intervals, there is no effect of Cr on the solubility of WC in the binder at 1220 °C. The results in the Cr doped alloys are in good agreement with Thermo-Calc™ calculations for a similar alloy sintered at 1200 °C [31]. The influence of Cr addition on binder spreading could also be related to a different WC/Co interface structure and energy in the Cr doped alloys. It is now established that Cr mainly dissolved in Co also has a small solubility in WC [32,33]. The presence of Cr atoms in WC and Co lattices results in a modified atomic bonding through the interface. Moreover, a segregation of Cr at interfaces occurs at the solid state, as checked in this work in the W rich alloy (Fig. 10, Table 4) and already shown in a C rich alloy [31]. This may be related to the formation of a stable thin (W,Cr)C cubic films at interfaces in the doped alloys as was observed in undoped alloys [34,15]. It should lead to decrease the interface energy and improve spreading which is not observed in Cr rich alloys. Another effect of Cr addition on interface energy could be a change in Curie temperature of the binder. In Cr free alloys, the earlier spreading of Co on WC grains in the W rich alloy was ascribed to the earlier ferromagnetic–paramagnetic transition of the binder during the heating of the specimens that is prone to decrease the interface energy [16]. In WC–Co alloys, W and C atoms dissolved in the cobalt affect the magnetic properties of the binder as well as doping elements [18,35,36]. The Curie temperature of the binder in the Cr doped alloys is expected to be lower than those found in the undoped alloys i.e. 1051 °C and 933 °C for WC– Co,Cr,C and WC–Co,Cr,W, respectively. These temperatures are lower or in the same range as the reduction temperature of the Cr oxides. The positive effect of Cr on Curie temperature is probably not observed because of the presence of Cr oxides. The examination of possible effects of Cr addition emphasizes the very likely role of surface Cr oxides in delaying Co spreading at the beginning of densification. Surface oxides interfere with wetting in metal–metal systems. While studying the spreading mechanism, it was proved that there is a correlation between the end of the deoxidization of WC surfaces

Table 4 EDS analyses (in at.%) at positions shown in Fig. 10 in WC–Co,Cr,W alloy sintered at 1220 °C. The higher Cr/Co amount at position 3 indicates a Cr enrichment of the WC/Co interface.

1 2 3

W

Co

Cr

Cr/Co

8.6 8.1 64.3

88.1 88.6 33.8

3.3 3.4 1.9

3.7 3.8 5.6

27

and the beginning of shrinkage of WC–Co alloys [24]. Our results show that shrinkage begins at around 800 °C for the undoped alloys and at around 1000–1100 °C for the doped alloys. These temperatures coincide with reduction temperatures of Cr oxides. However a lower reduction temperature was found in C rich alloy whilst spreading is delayed in this alloy. The results indicate that there is a combined effect of Cr oxides and C potential on the onset of spreading. At higher temperature the shrinkage rate increases for Cr doped alloys which is mainly due to a higher densification driving force induced by a lower density. Also Cr addition slightly inhibits grain growth at 1260 °C (Fig. 7) what could enhance sintering for these alloys. An additional effect of Cr addition is to magnify differential shrinkage during solid state sintering. The formation of large pores due differential shrinkage has often been reported in solid state sintering of WC–Co materials [13]. It is due to the difficulty to deform and to rearrange WC–Co previously densified clusters [12]. One could expect the evolution of the microstructure to be more favourable when sintering begins at higher temperatures as the phase distribution would be rather homogeneous. The strengthening of WC grain boundaries could be responsible for a higher rigidity of the WC skeleton in these alloys due to the modification of grain boundary energies. TEM/EDS or atom probe analyses have shown the segregation of Cr in addition to Co to WC grain boundaries after solid or liquid phase sintering [31,37,38]. Moreover, atomistic calculations predict that WC grain boundaries are strengthened by segregants, especially by V and Cr [39]. These latter would enhance the grain boundary resistance to cobalt infiltration. Experimentally, the stabilization of grain boundaries by Cr segregation is in agreement with the larger contiguity and higher fraction of Σ2 grain boundaries measured in a Cr doped alloy compared to a reference material [40]. The addition of Cr could increase the rigidity of the WC skeleton which would delay the shrinkage. 5. Conclusion The influence of Cr addition on the densification and microstructure evolution of WC–Co alloys during solid state sintering has been studied. Alloys with two different levels of carbon were used for the investigations to thoroughly discern the effect of Cr. The onset of densification is delayed whatever the C potential. The delaying effect due to Cr addition is explained by the presence of stable Cr oxides which would inhibit the wetting of WC grains by the binder. Although Cr oxides are reduced at a lower temperature in the C rich alloy, a better wetting of WC grains and a more effective rearrangement occurs in the W rich alloy. It is likely the effect of the higher solubility of WC in the binder of this alloy as no reduction in solubility of W atoms in the solid state was assessed with Cr addition in agreement with previous investigations [30]. At higher temperature the shrinkage rate increases for Cr doped alloys which is mainly due to a higher densification driving force induced by a lower density. The formation of large pores due to differential shrinkage is enhanced for these compositions as densification proceeds. As Cr addition delays the shrinkage, the heterogeneity could be related to the difficulty to deform the WC skeleton built at low temperature, especially as Cr could increase the stability of the WC grain boundaries and strengthen the WC skeleton. Acknowledgments The authors thank SANDVIK Hard Materials for technical and financial support. References [1] G. Gille, J. Bredthauer, B. Gries, B. Mende, W. Heinrich, Advanced and new grades of WC and binder powder — their properties and application, Int. J. Refract. Met. Hard Mater. 18 (2000) 87–102. [2] A. Bock, W.D. Schubert, B. Lux, Inhibition of grain growth on submicron cemented carbides, Powder Metall. Int. 24 (1992) 20–26.

28

V. Bounhoure et al. / Int. Journal of Refractory Metals and Hard Materials 52 (2015) 21–28

[3] W.D. Schubert, A. Bock, B. Lux, General aspects and limits of conventional ultrafine WC powder manufacture and hard metal production, Int. J. Refract. Met. Hard Mater. 13 (1995) 281–296. [4] K. Hayashi, Y. Fuke, H. Suzuki, Effects of addition carbides on the grain size of WC–Co alloy, J. Jpn. Soc. Powder Powder Metall. 19 (1972) 67–71. [5] C.H. Allibert, Sintering features of cemented carbides WC–Co processed from fine powders, Int. J. Refract. Met. Hard Mater. 19 (2001) 53–61. [6] G. Gille, B. Szesny, K. Dreyer, H. van dB, J. Schmidt, T. Gestrich, et al., Submicron and ultrafine grained hardmetals for microdrills and metal cutting inserts, Int. J. Refract. Met. Hard Mater. 20 (2002) 3–22. [7] B. Meredith, D.R. Milner, Densification mechanisms in the tungsten carbide–cobalt system, Powder Metall. 1 (1976) 38–45. [8] J.M. Missiaen, S. Roure, A general morphological approach of sintering kinetics: application to WC–Co solid phase sintering, Acta Mater. 46 (1998) 3985–3993. [9] R.J. Nelson, D.R. Milner, Densification processes in the tungsten carbide–cobalt system, Powder Metall. 15 (1972) 346–363. [10] W.D. Schubert, A. Bock, B. Lux, Aspects of ultrafine hardmetal sintering, Adv. Powder Metall. Part. Mater. 3 (1999) 10/23–10/36. [11] S. Haglund, J. Agren, B. Uhrenius, Solid-state sintering of cemented carbides. An experimental study, Z. Metallkd. 89 (1998) 316–322. [12] J.M. Missiaen, Solid-state spreading and sintering of multiphase materials, Mater. Sci. Eng. A 475 (2008) 2–11. [13] A. Petersson, J. Agren, Rearrangement and pore size evolution during WC–Co sintering below the eutectic temperature, Acta Mater. 53 (2005) 1673–1683. [14] S. Roure, J.M. Missiaen, Macroscopic and microstructural evolution of WC–Co compacts during solid state sintering, Euro'PM96 European Conference on Advances in Hard Materials Production, EPMA, Stockholm 1996, pp. 77–84. [15] V. Bounhoure, S. Lay, F. Charlot, A. Antoni-Zdziobek, E. Pauty, J.M. Missiaen, Effect of C content on the microstructure evolution during early solid state sintering of WC– Co alloys, Int. J. Refract. Met. Hard Mater. 44 (2014) 27–34. [16] V. Bounhoure, J.-M. Missiaen, S. Lay, E. Pauty, Discussion of nonconventional effects in solid-state sintering of cemented carbides, J. Am. Ceram. Soc. 92 (2009) 1396–1402. [17] A. Delanoë, J.M. Missiaen, C.H. Allibert, S. Lay, E. Pauty, Effect of the C potential and of Cr doping on the densification of alloys WC–Co, in: G. Kneringer (Ed.), 16th International Plansee Seminar Reutte, Austria, Plansee holding AG, Reutte 2005, pp. 642–652. [18] J.M. Doh, J.R. Groza, J. Oakes, Solid state sintering of WC–Co based cemented carbides, in: J.J. Oakes, J.H. Reinshage (Eds.),Advances in Powder Metallurgy & Particulate Materials, MPIF, Princetown 1998, pp. 105–114. [19] K. Okada, M. Hisanaga, T. Tanase, Effect of duplex addition of VC and Cr3C2 on the sintering of cemented carbides, in: K. Kosuge, H. Nagai (Eds.),PM2000 2000, pp. 1312–1315 (Kyoto, Japan). [20] K. Okada, T. Taniuchi, T. Tanase, Effect of Cr3C2 addition on the sintering of cemented carbide, in: J.J. Oakes, J.H. Reinshage (Eds.),Advances in Powder Metallurgy & Particulate Materials, MPIF, Princetown 1998, pp. 75–84. [21] G. Gille, G. Leitner, B. Roebuck, Sintering behaviour and properties of WC–Co hardetals in relation to the WC powder properties, Euro'PM96 European Conference on Advances in Hard Materials Production, EPMA, Stockholm 1996, pp. 195–210.

[22] K. Frisk, A. Markström, Effect of Cr and V on phase equilibria in Co–WC based hardmetals, Int. J. Mater. Res. 99 (2008) 287–293. [23] D.B. Williams, C.B. Carter, Transmission electron microscopy, Spectroscopy, IV, Plenum Press, New York, 1996. [24] M. Göthelid, S. Haglund, J. Ågren, Influence of O and Co on the early stages of sintering of WC–Co: a surface study by AES and STM, Acta Mater. 48 (2000) 4357–4362. [25] G. Leitner, K. Jaenicke-Rössler, C. Wagner, Process during dewaxing and sintering of hardmetals, International Conference on Advances in Hard Materials Production, Paper, 13, EPMA, Bonn 1992, pp. 1–10. [26] J. Angseryd, M. Zwinkels, Study of the chemical reactions during debinding of cemented carbide, in: D. Bouvard, F. Valdivieso (Eds.),Sintering'05: 4th International Conference on Science, Technology and Applications of Sintering, Grenoble, France, INPG 2005, pp. 327–330. [27] G. Leitner, T. Gestrich, Shrinkage, liquid phase formation and gaseous reactions during sintering of WC–Co hardmetals and correlation to the WC grain size, in: G. Kneringer, P. Rodhammer, H. Wilner (Eds.),14th Plansee Seminar, Plansee AG, Reutte 1997, pp. 86–99. [28] G.E. Spriggs, Importance of atmosphere control in hard-metal production, Powder Metall. 13 (1970) 369–393. [29] A. Markström, B. Sundman, K. Frisk, A revised thermodynamic description of the Co–W–C system, J. Phase Equilib. Diffus. 26 (2005) 152–160. [30] O. Lavergne, F. Robaut, F. Hodaj, C.H. Allibert, Mechanism of solid-state dissolution of WC in Co-based solutions, Acta Mater. 50 (2002) 1683–1692. [31] M. Elfwing, S. Norgren, Study of solid-state sintered fine-grained cemented carbides, Int. J. Refract. Met. Hard Mater. 23 (2005) 242–248. [32] M. Brieseck, M. Bohn, W. Lengauer, Diffusion and solubility of Cr in WC, J. Alloys Compd. 489 (2010) 408–414. [33] J. Weidow, S. Johansson, H.-O. Andrén, G. Wahnström, Transition metal solubilities in WC in cemented carbide materials, J. Am. Ceram. Soc. 94 (2011) 605–610. [34] S.A.E. Johansson, G. Wahnström, A computational study of thin cubic carbide films in WC/Co interfaces, Acta Mater. 59 (2011) 171–181. [35] S. Sundin, S. Haglund, A comparison between magnetic properties and grain size for WC/Co hard materials containing additives of Cr and V, Int. J. Refract. Met. Hard Mater. 18 (2000) 297–300. [36] G.K. Schwenke, J.V. Sturdevant, Magnetic saturation & coercivity measurements on chromium-doped cemented carbides, Int. J. Powder Metall. (Princeton, N. J.) 43 (2007) 21–31. [37] A. Henjered, M. Hellsing, H.O. Andren, H. Norden, Quantitative microanalysis of carbide/carbide interfaces in tungsten carbide-cobalt-base cemented carbides, Mater. Sci. Technol. 2 (1986) 847–855. [38] J. Weidow, H.-O. Andrén, Grain and phase boundary segregation in WC–Co with small V, Cr or Mn additions, Acta Mater. 58 (2010) 3888–3894. [39] M. Christensen, G. Wahnström, Strength and reinforcement of interfaces in cemented carbides, Int. J. Refract. Met. Hard Mater. 24 (2006) 80–88. [40] J. Weidow, S. Norgren, H.-O. Andrén, Effect of V, Cr and Mn additions on the microstructure of WC–Co, Int. J. Refract. Met. Hard Mater. 27 (2009) 817–822.