WC–Co (Ni) Alloys

INTRODUCTION TO HARDMETALS

1.03 Microstructure and Morphology of Hardmetals

1.03.2 WC–Co (Ni) Alloys

1.03.2.1 Microstructure of the Co-Rich (Ni) Binder 92

1.03.2.2 Interface and Grain Morphology 93

1.03.2.2.1 Wetting of WC ₉₃

1.03.2.2.2 Interface Energy ₉₄

1.03.2.2.3 WC Grain Shape Quantification 98

1.03.2.2.4 Effect of Transition Metals Addition 99

1.03.2.3 Microstructure of the WC Hard Phase 100

1.03.2.3.1 Orientation and Structure of WC/WC Grain Boundaries ₁₀₀

1.03.2.3.2 Grain Boundary Energy ₁₀₃

1.03.2.3.3 Segregation of 3D-Transition Metals to WC/WC Grain Boundaries ₁₀₄

1.03.2.3.4 Stability of WC/WC Grain Boundaries and Infiltration by Co ₁₀₅

1.03.2.3.5 Mean Free Path in the Binder Phase ₁₀₇

1.03.2.3.6 Contiguity ₁₀₇

1.03.2.4 Grain Growth 110

1.03.2.4.1 Mechanisms and Kinetics of Grain Growth in WC–Co Alloys ₁₁₀

1.03.2.4.2 Effect of C Content ₁₁₃

1.03.2.4.3 Effect of Inhibitors ₁₁₃

1.03.3 Other Cemented Carbides and Cermets 113

1.03.3.1 Wetting of TiC and TiN 113

1.03.3.2 WC–TiC–Co Alloys 114

1.03.3.3 TiC–Mo₂C–Ni Alloys 115

1.03.3.4 Ti(C,N)-Based Cermets 116

References 117

1.03.1 Introduction

Mechanical properties of hard materials are closely related to their microstructure, especially the phase composition, phase distribution and grain size. The phase distribution is to a large extent monitored by the relative energy of the grain boundaries and interfaces. In what follows, interfacial structures and energetics will be the connecting thread to get a comprehensive overview of the microstructure in cemented carbides. These quantities will be described using available experimental or calculated data of the literature. While few experimental data are available, atomistic calculations are helpful to interpret the microstructure and predict the effect of additional elements.

This chapter will focus mainly on WC–Co alloys which were studied extensively for nearly one century at different scales. The microstructure of the binder phase will first be discussed. Then properties of WC/Co interfaces and WC/WC grain boundaries will be described together with their consequences on the grain morphology and on the structure of the hard-phase skeleton. Finally, the state of the art about mechanisms and microstructural evolution related to grain coarsening will be presented. In an additional section, the microstructure of TiC-containing cemented carbides and of Ti(C,N)-based cermets will be briefly addressed.

1.03.2 WC–Co (Ni) Alloys

WC-based cemented carbides were mainly developed with cobalt as binder phase because it gives rise to a dense material with excellent mechanical properties. Nickel and iron were also considered as alternative binders but cobalt is the most common binder. The microstructure of WC–Co (Ni) alloys consists of facetted WC grains

Comprehensive Hard Materials, Volume 1 http://dx.doi.org/10.1016/B978-0-08-096527-7.00003-9 91

embedded in a Co(Ni)-rich binder as illustrated inFigure 1. WC grains are not completely surrounded by the binder but always form contacts with other WC grains. The microstructure depends to a large extent on the interfacial energy relationships. In what follows, the binder phase will befirst described. Then available data on WC/binder interfaces and WC/WC grain boundaries will be given. Finally, the contiguity which quantifies the relative amount of grain boundaries and interfaces will be discussed.

1.03.2.1 Microstructure of the Co-Rich (Ni) Binder

The mechanical properties of the binder are controlled by its composition and its distribution in the alloy after sintering (Almond & Roebuck, 1988). There is no solubility of cobalt in WC while cobalt dissolves a high amount of WC at high temperature. The solubility of W and C atoms in cobalt depends both on the temperature and on the total content of carbon in the material (Akesson, 1982). At 1425C, the C and W amount in the binder in equilibrium with WC ranges from about 9Cþ13W to 16Cþ6W (at%) when the carbon content increases from theh-phase limit to the graphite limit. On cooling, the solubility in the cobalt is reduced and W and C atoms reprecipitate on WC grains. The diffusion rate of W atoms decreases, so a gradient of concentration forms over a distance of about 0.1mm close to the WC grains (Andrén, 2001). For a slow cooling rate, the W amount remaining in the cobalt at room temperature reflects the composition of the binder at approximately 1000C. On the other hand, nearly no carbon is found in the binder likely due to the fast diffusion of this element in the cobalt.

The crystallographic structure of the Co-rich binder is mostly face-centered cubic (fcc) while some hexag- onal close-packed (hcp) pools resulting from the fcc–hcp transformation are also present. The fcc high- temperature form of Co is probably stabilized by residual stresses and the presence of dissolved atoms (Almond & Roebuck, 1982). However, a systematic characterization of the binder using the electron back- scatter diffraction (EBSD) technique indicates that in some specimens, the hcp form is predominant while no relationship with the composition or with the WC grain size could be established (Mingard, Roebuck, Marshall, & Sweetman, 2011). Although Co grains are interpenetrated, they can be delimited according to their orientation and they attain sizes much larger than the WC grain size (Figure 2). This size is influenced by the cooling rate, the amount of W dissolved in the binder and the cobalt content in the alloy (Weidow & Andrén, 2010a). It is also dependent on the WC grain size as cobalt likely nucleates on the surface of some WC grains, adopting a particular orientation relationship (Bounhoure, Lay, Loubradou, & Missiaen, 2008). Small fcc cobalt inclusions are also found in the WC grains with a shape and an orientation that minimize the interface energy (Mohan & Strutt, 1996).

Less data are available for the Ni-based binder since WC–Ni alloys represent a minor part of the cemented carbides. At the sintering temperatures, the two-phasefield (WCþbinder) is less extended than for WC–Co alloys and the solubility of WC is reduced compared to Co (Gabriel, Pastor, Deo, Basu, & Allibert, 1986;

Wittmann, Schubert, & Lux, 2002). Concerning the size of the Ni-based pools, microstructural investigations reported in the literature indicate that both small and large binder grains relative to the WC structure are encountered in such alloys (Mingard et al., 2011).

Figure 1 (a) Scanning electron microscopy (SEM) micrograph showing the typical microstructure of a WC–12Co alloy (at%) prepared from a WC powder with a median size of 0.85mm and sintered at 1450C for 1 h (WC in bright grey and Co in dark). (b) SEM image of a high-carbon WC–10 wt% Ni alloy sintered for 1 h at 1480C from a starting WC grain size of 0.35mm (Wittmann et al., 2002).

1.03.2.2 Interface and Grain Morphology

The facetted shape of WC grains in WC–Co alloys is due to the development of low-energy habit planes. The stability of these facets is related to their relative energy. In what follows, present knowledge on WC/Co in- terfaces is reviewed with a special emphasis on the energies. The energy calculations are achieved using density functional theory (DFT) that is based on the determination of the ground-state electron density of the system.

This approach may be very useful to compare different atomic configurations.

1.03.2.2.1 Wetting of WC

During liquid-phase sintering, the full densification can be achieved only if a good wetting of the hard phase by the binder occurs. The cohesion of the resulting WC–Co alloys is also related to the wetting behavior through the work of adhesion (Eqn (3)). The degree of wetting results from the energy balance between the surfaces of the system. It can be represented by the contact angle q of the liquid on the solid surface under thermody- namical equilibrium:

g_SV ¼ g_SLþg_LV cosq (1)

whereg_SVandg_LVrefer to the surface energies of the solid and liquid, respectively, andg_SLto the solid–liquid interface energy (Figure 3).

The magnitude of the contact angle can be obtained from the geometry of the drop. A small contact angle characterizes a good wetting of the solid by the liquid phase. Although sintering of a powder mixture requires an optimal wetting, few experimental data are available on wetting angles, surface and interface energies for car- bides and metals used in the fabrication of cemented carbides. Only the surface energy of the liquid can be measured directly by different methods and the value determined for Co at 1420C is 1.91 J m²and for Ni at 1380C is 1.81 J m²(Ramqvist, 1965).

Figure 2 (a) EBSD orientation map of a WC–11Co (wt%) alloy and (b) corresponding fcc Co binder EBSD map where black lines delineate Co regions with the same orientation (scale marker¼100mm) (Mingard et al., 2011).

Figure 3 Contact angleqof a liquid drop deposited on a solid for a solid–liquid–vapor equilibrium.

Microstructure and Morphology of Hardmetals 93

The wetting angle of WC by molten Co was experimentally found to be close to zero, which indicates a complete wetting of WC grains by Co (Ramqvist, 1965) and participated to the success of this metal as a binder in the cemented carbide industry. The prevalence of Co as a binder is not only related to its wetting behavior.

Other binders like Fe or Ni or combined with Co also offer a good wetting of WC (Table 1) and similar or superior mechanical properties were found for (Co–Ni–Fe)–WC alloys (Almond & Roebuck, 1988; Gille, Bredthauer, Gries, Mende, & Heinrich, 2000; Prakash, Holleck, Thümmler, & Walter, 1981). The main obstacle in the use of these alternative binders is the higher accuracy needed to control the carbon content of such alloys and to avoid the formation of complex carbides (Pastor, 1999). However, this difficulty should be overcome by choosing suitable compositions (Uhrenius, Pastor, & Pauty, 1997).

The wettability of a solid can also be quantified by the work of adhesion,W_ad. This quantity represents the external work required to separate the two crystals and create free surfaces of solid and liquid phases as expressed by the Dupré equation:

W_ad ¼ g_LVþg_SVg_SL (2)

By combiningEqns (1) and (2), the work of adhesion can be conveniently determined from the contact angle and the surface energy of the liquid:

W_ad ¼ g_LVð1þcosqÞ (3)

Hence, in early studies of cemented carbides, the work of adhesion of the WC/Co interface was estimated to 3.82 J m² (Ramqvist, 1965). This high value emphasizes the good wettability of the carbide by the cobalt.

A slightly lower value is found for the WC/Ni interface (Table 1).

Bonding at an interface can be quantified using DFT calculations by the work of separationW_septhat is the work to separate the interface between two solid phases into two free surfaces, without relaxation processes (Finnis, 1996, p. 422). According to the definition, the calculated values should be higher than the work of adhesion. As a consequence, the comparison betweenW_adandW_sepshould be only qualitative. DFT calcula- tions were used to evaluate the work of separation of the WC{100}/Co{100} interface assuming the cubic NaCl crystallographic structure for WC, that is stable only at very high temperatures (Christensen, Dudiy, & Wahn- ström, 2002). A value close to 3.7 J m²was determined in the same range as the adhesion work found for this interface.

1.03.2.2.2 Interface Energy

WC has a hexagonal unit cell with the P-6m² space group and close lattice parameters a¼0.2906 nm, c¼0.2837 nm. Its stability is due to the formation of strong nearest-neighbor W–C bonds (Mattheiss &

Hamann, 1984). Tungsten and carbon atoms lie at (000) and at (1/3, 2/3, 1/2) positions, respectively, so the structure of WC consists of alternate layers of tungsten and carbon atoms along [0001] or [10-10] directions (Figure 4). Each carbon atom lies in the centre of a trigonal prism delineated by six tungsten atoms and reversely. As seen on the drawing, carbon atoms occupy thegsites and thebsites are empty. Consequently, the structure is not centrosymmetric such as (10-10) and (-1010) surfaces are not equivalent. Experiments including etching, sessile-drop wetting and hardness behavior have revealed this polar character of the prismatic surfaces of WC grains (French, 1969).

Table 1 Experimental values of wetting anglesq() and adhesion energiesWad(J m²) for a range of materials used in the cemented carbide industry. The wetting experiments inRamqvist (1965)were conducted at 1420C for Co, and at 1380C for Ni, inWarren and Waldron (1972a)between 1380C and 1450C

WC TiC TaC VC NbC TiC–22WC Mo2C

Co wetting qCo() 0^a 25^a

262^b 13^a 13^a

13^b 14^a

11.51^b 21^a

24.5^b 0^a

Wad^a 3.82 3.64 3.77 3.77 3.76 3.69 3.82

Ni wetting q_Ni()^a 0 23 16 17 18 0

Wad^a 3.62 3.47 3.55 3.54 3.53 3.62

aRamqvist (1965).

bWarren and Waldron (1972a).

In WC–Co alloys, WC grains adopt the shape of a triangular prism with truncated corners delimited by {0001} basal and {10-10} prismatic facets (Exner, 1979) (Figure 4(c)). The grain shape is influenced by the carbon potential in the alloy, showingflat facets in C-rich alloys and some rounded parts and steps at WC/Co interfaces in W-rich alloys (Wang, Heusch, Lay, & Allibert, 2002) (Figure 5) (Kim, Han, Park, & Kim, 2003).

The basal and prismatic planes play a major role in the microstructure formation of WC–Co alloys. According to the crystallographic structure of WC, two kinds of basal surfaces can arise, they can be carbon or tungsten terminated. On the other hand, there are two types of WC(10-10) surface, denoted I and II. For type I surface, the closest atoms lie at an interplane spacing of 0.084 nm while for type II, they lie at 0.168 nm (Figure 4(b)).

Each type of facet can be carbon or tungsten terminated which leads to four candidate WC(10-10)/Co interfaces.

This faceted shape of carbide grains is caused by the anisotropy in WC/Co interface energy. The equilibrium shape corresponds to the minimal interface energy of the grain and reflects the relative stability of the habit planes (Herring, 1951). A theoretical estimation of the energyg_SLof the WC/Co interface was carried out on the basis of thermodynamic considerations and a value of 0.5 J m² was deduced (Warren, 1980) (Table 2).

Quantities like interface energies are difficult to get experimentally. For example, deducing the interface energy from wetting experiments requires the knowledge of the solid surface energy (Eqn (1)). However, unlike liquid surface energies, solid surface energies cannot be measured directly and only estimates are available. Moreover, in the case of ideal wetting in the WC/Co system, a lower bound can only be obtained. Estimates of surface energy were determined for several carbides using wetting data but these later caused controversies (Livey &

Murray, 1956). Warren (Warren, 1976) has also determined lower values of the average energiesg_SV,g_SLand of the average grain boundary energy g_SS in systems with a continuous WC skeleton (Section 1.03.2.3.2). Ac- cording to this study, a lower bound of 0.575 J m²is obtained for the WC/Co interface energy, which is a little higher than the thermodynamic calculation (Table 2). All these values do not account for the anisotropy of the interface energies.

Figure 4 (a) Hexagonal unit cell of WC, (b) atomic projection along [0001] showing the four types of (10-10) prismatic facets denoted IW, IC, IIW and IIC on the drawing, and (c) morphology of WC grains in WC–Co alloys.

(a) (b)

Figure 5 Morphology of WC grains extracted from specimens (a) WC–35%Co–0.7%C (1500C, 2 h) and (b) WC–35%Co–0.7%C (1500C, 5 h). The compact in (b) corresponds to the compact of (a) embedded in a pack of carbon-black powder and resintered in order to increase the carbon content of the alloy (Kim, Han, et al., 2003).

Microstructure and Morphology of Hardmetals 95

The measurements of the interface energy of each facet would be extremely tricky. DFT calculations were used for assessing the energy of the different types of habit planes. At the sintering temperature, the binder is liquid so cobalt atoms are mobile and the atomic structure of the interface is constantly changing. In order to get reasonable quantities for the energy, an upper and a lower limit were defined and inside the interval of energy, the parameterawas used to characterize the degree of coherency at the interface (Christensen, Wahnström, Lay,

& Allibert, 2007):

g ¼ g_minþaðg_maxg_minÞwith 0<a<1 (4) Fora¼0, the energy is minimum and the interface is assumed to be coherent, while fora¼1, the interface is completely incoherent. In these calculations, the effect of carbon chemical potential was also taken into account.

The four types of (10-10) facets, denoted I and II, terminated by W or carbon atoms, were examined (Christensen et al., 2007; Christensen, Wahnström, Allibert, & Lay, 2005). For each type of facet, three atomic configurations at the interface were considered (Figure 6) and their energy was calculated. Values in the range of 0.7–2.9 J m² were found (Figure 7(a)) (Table 2). Most stable interfaces are determined to be W-terminated whatever the carbon content of the alloy. This is due to the formation of metal–- metal bonding across the interface. Moreover, type I facets have a lower energy. According to these re- sults, the set of most developed facets are expected to be of type I. Both types of facets are W-terminated independently on the carbon content. High-resolution transmission electron microscopy (HRTEM) characterization of the prismatic facets is in agreement with these calculations (Lay, Donnadieu, &

Loubradou, 2010).

The same approach was used to model the WC(0001)/Co interface (Christensen et al., 2007). C- and W-terminated basal facets were considered and in each case, four atomic configurations were assumed. An upper and a lower limit of the energy were deduced as a function of the coherency parameter in a W- or C-rich alloy.

Table 2 Solid surface energy (gSV), interface energy (gSL) and grain boundary energy (gSS) estimates (expressed in J m²) based on experimental measurements (exp), thermodynamic calculations (therm) or calculated by DFT techniques as well as grain boundary energy variation due to Co segregation (DgSS) (J m²) and dihedral angle (fexp) available for WC–Co and MC–Co alloys

WC TiC TaC VC

gSV exp >2.470.2^a 1.190.35^b

2.210.2^a 1.290.39^b

2.400.2^a 1.670.5^b

gSL exp >0.570.2^a 0.44^c 0.47^c 0.46^c

g_{SL therm} 0.5^d

gSL DFT WC(10-10)/Co WC(0001)/Co

0.7–2.9^e 0.7–2.3^e

gSS exp >0.990.3^a 0.880.3^a 0.960.3^a

gSS DFT

clean boundary 0.03–7.41 1.9 (mean value) (S¼2 twist)^f 2.1–2.6 (S¼2 tilt)^g 2.7 (mean value) (generic)^f DgSS DFT

Co segregation 0.4 (S¼2 tilt)^g

fexp() z60^h 25–27ⁱ 6–8ⁱ

aWarren (1976).

bLivey and Murray (1956).

cWarren and Waldron (1972b).

dWarren (1980).

eChristensen et al. (2007).

fChristensen and Wahnstro¨m (2003).

gChristensen and Wahnstro¨m (2004).

hDeshmukh and Gurland (1986).

iWarren and Waldron (1972a).

The energy of the basal facets is in the same range as the prismatic facets, between 0.7 and 2.3 J m² (Figure 7(b);Table 2). In the W-rich alloy, the most stable interface is W-terminated while in the C-rich limit, the basal facet is C- or W-terminated depending on the degree of coherency. These energy values were used as shown in the following to predict the WC grain morphology as a function of the carbon content.

Interface energy values deduced from DFT calculations are significantly higher than the values obtained by Warren (1976)(Table 2).

Figure 7 Interface energy of W- or C-terminated (a) prismatic and (b) basal facets as a function of the coherency parameterafor a WC–Co alloy lying in the C-rich or W-rich limit. For sake of clarity, an average value for type I and II facets is plotted for prismatic interfaces (Christensen et al., 2007).

Figure 6 Models of WC(10-10)/Co facets terminated by tungsten atoms used to calculate the interface energy. Three relative positions (a-c) of the cobalt atoms at the interface were considered.

Microstructure and Morphology of Hardmetals 97

1.03.2.2.3 WC Grain Shape Quantification

In the condition of an equilibrium shape, the energy ratio between the facets can be deduced from the shape of the WC grains. This requires a large volume of Co to decrease the effect of WC/WC contacts and long sintering times to get the equilibrium shape. Two parameters were defined to characterize the facetted shape of WC grains:

the truncation and the elongation factors. The truncation factorris the ratio between the lengths of the two types of facets: it expresses the anisotropy between the two sets of prismatic facets. The elongation factorkis the ratio of the thickness of the prism,t, over the height,h, of the truncated triangle. It measures the anisotropy between the prismatic facets and the basal facet (Figure 8).

r ¼ X

a_short=X

a_long (5)

k ¼ t=h (6)

If the shape is governed by the minimization of the energy, the shape parameters are related to the interface energies:

r ¼

2gPlong

gPshort 1 2_g^gPlong

Pshort

(7)

k ¼ 2g_Bð2rþ1Þ

3g_Plongðrþ1Þ (8)

whereg_Plong(g_Pshort) is the energy of the largest (smallest) prismatic facets andg_Bthe energy of the basal facet.

Using the calculated energies of the facets, the shape parameters were determined including the effect of carbon chemical potential for WC–Co alloys (Christensen et al., 2007) (Figure 9). For simplicity, the samea value was chosen for basal and prismatic interfaces. The calculation results in a truncation and elongation factor close to 0.2 and 0.7, respectively, in the incoherent limit whatever the carbon content. However, the atomic

Figure 8 Parameters used for the determination of the shape factors.

Figure 9 (a) Truncation parameter and (b) elongation parameter as a function of the parametera. Solid and dashed lines correspond to W-rich and C-rich WC–Co alloys, respectively (Christensen et al., 2007).