Other Cemented Carbides and Cermets

INTRODUCTION TO HARDMETALS

1.03 Microstructure and Morphology of Hardmetals

1.03.3 Other Cemented Carbides and Cermets

MC cubic carbide additions are used to increase the hardness and the wear resistance of WC–Co alloys for appli- cations like high-speed machining of steels. The metals M belonging to the group IV (Ti, Zr, Hf) or to the group V (V, Nb, Ta) of the periodic table were considered because the corresponding carbides are harder than WC (Exner, 1979).

Among them, TiC is the hardest and mainly mixtures of WC–TiC carbides with additions of (Ta,Nb)C were selected.

Besides, much harder TiC-based cemented carbides like TiC–Mo₂C–Ni alloys were developed for their higher tool life in specific operations likefine turning (Pastor, 1987). Then further improvements mainly on toughness and wear resistance of hardmetals were obtained by introducing TiN in TiC–Ni/Mo alloys and thus developing titanium carbonitride-based cermets (Pastor, 1999). The available data of the literature on the wetting behavior and on the interface energies in these systems will befirst briefly presented before describing the microstructure.

1.03.3.1 Wetting of TiC and TiN

Compared to WC, the wetting angle of the transition-metal carbides by Co is found higher especially for TiC what reflects the more difficult densification of TiC–Co alloys (Warren & Waldron, 1972c) (Table 1). In early studies of cemented carbides, the poorer strength of the TiC/Co interface was pointed out by a lower value of the Figure 26 Example of the local formation of abnormal grains by the carburization of previously formedh-phase. (a) Interrupted at 1350C; (b) 1450C, 1 h (fromAdorjan et al., 2006).

Microstructure and Morphology of Hardmetals 113

work of adhesion (3.64 J m²) relative to WC/Co. The analysis of the electronic structure of TiC/Co and WC/Co interfaces indicates that strong covalent Co–C bonds form at both interfaces (Christensen et al., 2002; Dudiy &

Lundqvist, 2001). However, the enhanced strength of WC/Co interface originates from a larger contribution of metal–metal Co–W bonds. In agreement with the experiments, a lower value of the separation work is obtained for the TiC/Co interface (W_sep¼3.25–3.45 J m²) (Dudiy & Lundqvist, 2001). The poorer wettability of TiC by Co results in the formation of pores and worsens the mechanical properties of WC–TiC–Co alloys (Bhaumik, Upadhyaya, & Vaidya, 1991). In TiC-based grades, nickel proved to be more suitable than cobalt as a binder (Brookes, 1992) and wetting was improved by the addition of Mo or Mo₂C (Exner, 1979).

Compared to WC and TiC carbides, scarce experimental data on the wetting behavior of TiN by Co or Ni are available although it is admitted that wetting is not so good but does not limit the densification of these alloys.

The adhesion properties of the Co/TiN interface were investigated by means of DFT calculations (Dudiy &

Lundqvist, 2001). It was found that the adhesion of the interface is provided by strong covalent Co–N bonds comparable to Co–C bonds occurring at TiC/Co interface. However, the extra electron of the N atom induces a reduction of the strength of these bonds resulting in a weaker adhesion in the Co/TiN case. This analysis is consistent with the practice that a higher sintering temperature is needed as the TiN content increases to get full densification of the cermets (Ettmayer, Kolaska, Lengauer, & Dreyer, 1995).

1.03.3.2 WC–TiC–Co Alloys

As TiC carbide is much more brittle than WC, the TiC content of WC–TiC–Co alloys usually does not exceed 18 wt% (Brookes, 1992). After sintering, the microstructure consists in a mixture of angular WC grains and TiC rounded grains embedded in the Co-rich binder. The TiC grains, called g phase, usually show a core–rim structure (Figure 27), with the core having the composition of the initial TiC or (Ti,W)C powder and the rim being a (Ti,W)C phase with a higher amount of W. Solubility of WC in TiC is significant, whereas it is negligible for TiC in WC. At 1400C, the equilibrium composition of the cubic carbide varies from W_0.35Ti_0.65C inh containing alloys to W_0.40Ti_0.60C when graphite is present (Kruse, Jansson, & Frisk, 2001). The formation of the rim around TiC grains arises from the partial dissolution of TiC and WC in the binder/liquid followed by the epitaxial precipitation of the (Ti,W)C solid solution onto undissolved TiC grains (Andrén, 2001). For low TiC additions, all TiC raw grains are dissolved in the binder and only (Ti,W)C grains are present (Weidow, Zack- risson, Jansson, & Andrén, 2009). The distribution of the elements in the core–rim structure of thegphase was accurately determined using atom probefield ion microscopy (Rolander & Andrén, 1988). It was found that the diffusion distance of W atoms from the rim to the core is only a few nanometers because of the low mobility of W atoms in TiC (Figure 28). The Ti content of the binder is low and depends on the carbon content during sintering. For example, at 1450C, the solubility of Ti is expected to be 2.4 wt% in carbon-rich alloys and 1.0 wt

% in tungsten-rich alloys (Frisk et al., 2001). As for WC–Co alloys, a low carbon content was found in the binder after sintering.

The inspection of the microstructure (Figure 27) shows the formation of (Ti,W)C/WC phase boundaries and (Ti,W)C/(Ti,W)C grain boundaries where segregated cobalt was detected (Henjered et al., 1986; Weidow &

Andrén, 2011). These observations confirm the formation of a WC–MC skeleton (Andrén, 2001). The overall microstructure is however highly dependent on the amount of TiC and on the WC/TiC grain size ratio and

Figure 27 Microstructure of 70(25TiC–75WC)–30Co (wt%) samples sintered at 1450C for (a) 2 h and a starting size of powder of 1.3mm (b) sintered for 5 h with a starting size of 4.1mm, (A) TiC core, (B) (Ti,W)C grain, (C) WC grain, (D) Co-based matrix (Yoon, Lee,

& Kang, 2005).

finally greatly influences the mechanical properties (Lee, Cha, Kim, & Hong, 2006). In some grades containing high percentages of TiC, nickel is used as an entire or partial replacement of cobalt.

1.03.3.3 TiC–Mo₂C–Ni Alloys

The addition of Mo₂C to TiC–Ni alloys induces the formation of the characteristic core–rim structure of the carbide grains, the rim having the same crystallographic orientation than the core. The composition of the core corresponds to the initial TiC powder while the rim is a (Ti,Mo)C solid solution and a rather abrupt change in composition occurs at the core–rim interface (Figure 29) (Yamamoto, Jaroenworaluck, Ikuhara, & Sakuma, 1999). The mechanism of the rim formation is not completely assessed (Pastor, 1999). The rim formation would start at the solid state by the diffusion of Mo into TiC grains. Moreover, a (Ti,Mo)C solid solution likely precipitates on TiC grains from the eutectic liquid with a composition in equilibrium with the melt. The Figure 28 Atom probefield ion microscopy (APFIM) composition profile through a core–rim boundary in an 85WC–7.7(Ti,W)C–7.3Co (wt%) material. Some impurity nitrogen was found in the core (Andrén, 2001).

Figure 29 (a) Bright-field TEM image in TiC–20 wt% Mo2C–20 wt% Ni sintered at 1370C for 1 h, showing the typical core–rim structure. The white arrows indicate Ni phases included into carbide grains during sintering. (b) Plots of Ti and Mo weight ratio estimated from EDS analysis across the interface. The weight ratio shown is calculated from the ratio of Ti or Mo to (TiþMo) (Yamamoto et al., 1999).

Microstructure and Morphology of Hardmetals 115

composition of the liquid binder depends on the carbon content of the alloy. For specimens sintered between 1300C and 1450C, the Ti and Mo binder content was found to range from about 3 wt% Ti and 0 wt% Mo for high-carbon alloys to about 10 wt% Ti and 6 wt% Mo for low-carbon alloys (Suzuki, Hayashi, & Terada, 1971).

The Mo enrichment of the rim at the surface of the TiC grains is likely to be responsible for the better wetting of the carbide grains when Mo₂C is present in the alloy. Moreover, the addition of Mo can limit grain growth (Barranco & Warenchak, 1989). Development of this grade of alloys by changing the composition is still a topicalfield of research (Guo, Xiong, Yang, & Jiang, 2008).

1.03.3.4 Ti(C,N)-Based Cermets

Cermets are prepared from a mixture of TiC and TiN hard phases or from the Ti(C,N) compound, with nickel or/

and cobalt as a binder. This class of materials also contains WC or Mo₂C to improve the wetting of the hard phase and increase the mechanical performance of the alloys interfaces (Brookes, 1992; Ettmayer et al., 1995).

The microstructures are usually complex, depending on the composition, starting powders and sintering treatment. The hard-phase particles are embedded in the binder and show the characteristic core–rim structure while the rim may be divided into the inner and outer rims that have different compositions (Figure 30). The microstructure has been carefully characterized in various cermets and at different stages of the sintering cycle and processes leading to the formation of the complex structure could be identified. Among the variety of starting materials, the Ti(C,N)–TiN–Co–WC or Ti(C,N)–Ni–WC model alloys are chosen to illustrate the microstructure evolution on sintering and to get a better insight on the core/rim structure formation (Ahn &

Kang, 2000; Zackrisson, Rolander, & Andrén, 2001; Kim, Min, & Kang, 2003).

On heating, the hard-phase particles start to dissolve in the binder and the (Ti,W)(C,N) inner rim with a high amount of tungsten precipitates epitaxially onto undissolved Ti(C,N) particles. The thickness of the inner rim is not uniform what could be related to the difficulty for this phase to nucleate. The outer rim forms after melting of the binder by precipitation of a (Ti,W)(C,N) solid solution from the liquid. The outer rim has a smaller tungsten content and higher nitrogen content than the inner rim. The composition of the core of the car- bonitride grains evolves during sintering due to the diffusion of carbon and nitrogen atoms. They do not contain tungsten atoms except at the level of crystal defects where the diffusion is enhanced. The onset of formation of the inner rim is not fully established. It could form on heating at the solid state by diffusion through the solid binder with a composition close to the equilibrium (Zackrisson et al., 2001). The difference between inner and outer rims would be due to a change of the chemical equilibrium as the temperature in- creases and as the nitrogen activity is enhanced due to the closure of the porosity. On the other hand, it is also assumed that the inner rim forms at the onset of binder melting and the outer rim when sintering proceeds (Ahn

& Kang, 2000). The composition change would be related to the difference in dissolution rate of the WC and Ti(C,N) phases at higher temperature. The dissolution rate is higher for WC and the relative rates vary as a function of the temperature. As the solute content of the binder determines the composition of the rim, the inner rim is enriched in W atoms compared to the outer rim.

Figure 30 (a) SEM micrograph of a Ti(C,N)–TiN–Co–WC cermet sintered at 1430C for 90 min showing the typical core/rim structure, where the inner rim close to the core has a bright contrast and the outer rim a darkest contrast (Zackrisson et al., 2001). (b) Schema of the microstructure of a Ti(C,N)-based cermet (Ahn & Kang, 2000).

Very similar microstructures are observed in Mo₂C-containing cermets with the formation of the core/rims structure, the core and the rims having the Ti(C,N) and the (Ti,Mo)(C,N) composition, respectively, the inner rim being richer in Mo atoms than the outer rim (Andrén, 2001; Lindahl, Gustafson, Rolander, Stals, &

Andrén, 1999).

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