CLASSES OF MATERIALS
1.05 Cermets
1.05.2 Microstructure of Ti-Based Cermets .1 Instability Issues of TiC-Based Cermets
In transitional metal carbides, the interface stability between carbide and the binder metal is an important issue in determining the microstructure of cermets. Especially, it is so with TiC and Ti(CN). The tendency to reduce surface energy is a major driving force in the sintering of a single-phase system. Therefore, grain morphologies are typically spherical or at leastflat. However, in some special cases, negative curvatures are also found. This phenomenon is referred to as chemically induced grain migration (CIGM) and especially in the liquid-phase sintering, it is called instability of the solid–liquid interface (ISLI).
Many reports have appeared on the driving force for CIGM (Ahn, Kim, & Kang, 2001; Hillert, 1983; Lee, Baik, &
Yoon, 1987; Park, Park, Kang, & Kim, 1998; Song & Yoon, 1984; Yoon, 1995). These studies indicate that a large part of the total driving force for CIGM can be attributed to the presence of strain energy at the coherent interface layer.Chae, Chun, Kim, Baik, and Eun (1990)reported ISLI in a TiC–Fe system. LaterChun, Kim, and Eun (1993) also studied ISLI using a TiC–Mo–Ni system. They concluded that the strain was induced by the formation of a coherent layer, which is a solid-solution rim formed from the carbide and binder phases. However, CIGM is not always observed in all systems.
It was found that CIGM, narrowly speaking, ISLI is not commonly observed in carbide–metal systems during sintering. It is generally thought that the formation of a solid-solution rim is a necessary condition for ISLI.
Kwon and Kang (2006) investigated the microstructures of various carbide–Ni composites in terms of the parameters in Hume–Rothery rules such as atomic size, electronegativity, electron valence, and crystal structure.
Average size and manufacturers of initial powders are listed in Table 9. Of fourth period transition-metal carbides, TiC is the only carbide that exhibits a strong ISLI with negative curvatures in molten Ni (Kwon &
Kang, 2006). No ISLI was observed for other carbides of thefifth and sixth period elements.
Table 8 Chemical analysis of synthesized powders (in wt% balanced by Ti and W)
Compositions Processing temperature (C) Carbon (wt%) Nitrogen (wt%) Oxygen (wt%)
(Ti,15Mo)(CN)–Ni 1200 10.37 2.70 0.12
1300 5.64 1.59 0
(Ti,10Mo,5Nb)(CN)–Ni 1200 12.00 5.68 0.27
1300 11.31 3.90 0
(Ti,15W)C–Ni 1200 14.99 0.13 0.05
1300 12.14 0.52 1.14
(Ti,15W)(CN)–Ni 1300 11.45 5.58 0.36
(Ti,15W,5Mo)(CN) 1300 7.61 1.65 0.05
(Ti,15W,5Nb)(CN) 1300 15.22 4.28 0.05
(Ti,30W)C–Ni 1300 10.65 0.38 1.08
(Ti,30W)(CN)–Ni 1300 10.39 3.66 0.08
Ti(C0.7N0.3)–15WC–20Ni Commercial 10.10 4.44 0.81
Ti(C0.7N0.3)–30WC–20Ni Commercial 9.00 3.39 0.83
Table 9 Average particle size and manufacturers of initial powders (Kwon & Kang, 2006)
Particle size (mm) Manufacturer
TiC 1.00 Treibacher
ZrC 2.40 H. C. Starck
HfC 1.61 Kennametal
VC 1.10 HCST
NbC 1.75 Treibacher
TaC 1.12 Treibacher
Cr3C2 1.75 Kennametal
Mo2C 1.42 Kennametal
WC 1.88 Xiamen
Ni 4.00 INCO
Cermets 157
The origin of ISLI is strain developed at the interface between the carbides and the newly formed solid so- lutions. The difference in the size of the atoms involved can be used to predict the formation of a carbide–Ni solid solution when the Hume–Rothery rules are applied. Aside from the size factor, other factors in the rules are not found effective in predicting this phenomenon. Figure 14 shows the microstructures of the specimens without ISLI, which were obtained after infiltration and heat treatment at 1510C for 1 h. All those were from fifth period metal-carbide systems. Similar microstructures were obtained from sixth period carbides.
Figure 15shows the microstructures of the specimens obtained from the fourth period transitional-metal carbides. In these systems, carbide particles (TiC and Cr3C2) with negative curvatures can be seen. The IVB, VB, and VIB transitional-metal carbides in the fourth period were found to grow much faster (30–100mm) than the other carbides in thefifth and sixth periods (5–10mm). An XRD analysis was done to better understand the tendency toward solid-solution formation by measuring the shifts in the main peaks. As summarized inTable 10, all carbide systems demonstrate clear peak shifts. The peak shifts to high diffraction angles (2q) were observed for TiC, Cr3C2, and WC. The shifts to high angles, which indicate the decrease of lattice parameters, could be interpreted as a tendency of the carbides to form solid solutions or compounds readily with solutes such as Ni in these cases. The TiC–Ni system, which demonstrates a strong ISLI, shows a decrease in lattice constants as the result of forming a solid solution, (Ti,Ni)C. Other carbides such as VC, ZrC, NbC, HfC, TaC, and Mo2C showed peak shifts to low angles. This implies a low affinity between these transition metals and Ni in their crystal structure.
InTable 11, the differences in atom size between carbides and nickel are shown under a size factor. Other factors are also listed for comparison. The size factor was determined by the following formula: (atomic size of the metal in the carbide–atomic size of nickel)/(atomic size of the metal in the carbide). The size difference between impurity and host atoms must be underw15% for a solid solution to form. Thefifth and sixth period metal atoms (Zr, Hf, Mo, Nb, Ta, and W) have size factors of over 15% with nickel while fourth period metal atoms (Ti, V and Cr) have size factors underw15%. VC has the best conditions among the carbides of this study to form a solid solution with Ni. However, the tendency of VC to form a solid solution with Ni was small from XRD data and this system failed to show ISLI. Main peak of VC in XRD data is from (400) plane and (400) plane’s lattice parameter is smaller than that of other planes. This can be why the tendency to form a solid solution with Ni is not big as much as the case of TiC, Cr3C2and ISLI is not observed in VC–Ni system. However, all these factors including electronegativity, electron valence and crystal structure do not provide clear predictive information on the formation of solid solutions.
(a) (b)
(c)
10 µm 10 µm
100 µm
Figure 14 SEM micrographs of (a) ZrC–50Ni, (b) NbC–50Ni, and (c) Mo2C–50Ni (vol%), heat treated at 1510C for 1 h.
(a)
(c)
(b)
50 µm 50 µm
50 µm
Figure 15 SEM micrographs of (a) TiC–50Ni, (b) VC–50Ni, and (c) Cr3C2–50Ni (vol%) heat treated at 1510C for 1 h under vacuum.
Table 10 Peak shifts of major peaks of various carbide–metal systems (2q) (Kwon & Kang, 2006)
Carbides A B C D
JCPDS Raw powder 1510C heat treatment w/o Ni (C-B) 1510C sintering with Ni (D-C)
TiC 41.710 41.673 41.400 (0.273) 41.657(D0.257)
VC 43.490 43.415 43.360 (0.055) 43.340 (0.020)
Cr3C2 50.077 50.051 50.027 (0.024) 50.054(D0.027)
ZrC 33.040 33.054 33.034 (0.020) 32.755 (0.279)
NbC 34.730 34.822 34.633 (0.189) 34.458 (0.175)
Mo2C 39.392 39.530 39.466 (0.064) 39.229 (0.237)
HfC 33.437 33.456 33.213 (0.243) 33.212 (0.001)
TaC 34.855 34.927 35.065 (0.138) 34.607 (0.458)
WC 48.266 48.355 48.118 (0.237) 48.550(D0.432)
Table 11 Atomic radii and size factors of 4, 5, and 6 group transition-metal atoms and nickel atom (Kwon and Kang, 2006) Atomic radius
(nm) (A) Most common
valence (V) Crystal structure of carbide
Electronegativity
(E) DA(nm) DV DE Atomic size
factor (%)
Ni 0.149 2 FCC 1.91 – – – –
Ti 0.176 4 B1 1.54 0.027 2 0.37 15.34
Zr 0.206 4 B1 1.33 0.057 2 0.58 27.67
Hf 0.208 4 B1 1.3 0.059 2 0.61 28.37
V 0.171 5 B1 1.63 0.022 3 0.28 12.87
Nb 0.198 5 B1 1.6 0.049 3 0.31 24.75
Ta 0.200 5 B1 1.5 0.051 3 0.41 25.5
Cr 0.166 3 Orthorhombic 1.66 0.017 1 0.25 10.24
Mo 0.190 4 Hexagonal 2.16 0.041 2 0.25 21.58
W 0.193 4 Hexagonal 2.36 0.044 2 0.45 22.8
Cermets 159
1.05.2.2 Dissolution/Precipitation: Ti(CN)–MC–Ni
TiC- and Ti(C,N)-based cermets possess afine and stable microstructure. Such microstructure, when formed during liquid-phase sintering shows a typical core/rim structure where the cores are partially dissolved raw material particles on which the rim structure has grown through a dissolution–precipitation process (Ahn &
Kang, 1998; Ettmayer, Kolaska, Lengauer & Dreyer, 1995; Ettmayer & Lengauer, 1989; Qi & Kang, 1998).
Figure 16is a schematic of a scanning electron microscopy (SEM) microstructure for commercial cermets. In general, most of the particles in the microstructure consist of black TiC or Ti(C,N) cores and gray (Ti,W,.)C or (Ti,W,.)(C,N) solid-solution rims.
WC, secondary carbide, is a necessary carbide component in a cermet system which serves to improve me- chanical and functional properties. The general role of added WC in a cermet system is to enhance density via improving wetting and sinterability (Matsubara, Shin, & Sakuma, 1992). It has also been reported that the addition of WC to a TiC–Ni or Ti(C,N)–Ni system improves toughness and decreases particle growth rate (Suzuki et al., 1983; Suzuki & Matsubara, 1986). Finally, it has been noted that WC, which forms a solid so- lution with TiC or Ti(C,N), remains as an independent phase with increasing nitrogen content of a Ti(C,N)–Ni system. However, to date, few studies of the dissolution and precipitation behavior of Ti(C,N) and WC in a Ti(C,N)–Ni system have been reported (Rynemark, 1991). In the following, the discussion will be limited to Ti(CN) as an exemplary major carbide component.
1.05.2.2.1 Core/Rim Structure
There are numerous studies on core/rim microstructure of Ti(CN)-based cermets. As an example,Ahn and Kang (2000)investigated the core/rim structures of Ti(C,N)–xWC–20Ni in order to determine the effect of WC and nitrogen content on the microstructure of the system. In addition, the relative dissolution rate of WC to Ti(C,N) in the system was studied by analyzing the rim compositions which, to a large extent, reflects the equilibrium reaction during liquid-phase sintering.
Variations in WC content had a much less influence, as an additive, on the rim microstructure than that of other carbides (Ahn & Kang, 2000; Yang & Lee, 1996). This is attributed to the low dissolution rate of Ti(C,N) as compared to that of WC in the Ni melt. The known effect of particle refinement as the result of added WC to Ti(C,N)–Ni cermets was not evident sometimes (Matsubara et al., 1992).
1.05.2.2.2 Nitrogen Effect on Dissolution of Ti(CN)
The nitrogen content in Ti(C1xNx), however, had a significant effect on the microstructure by changing the stability of Ti(C1xNx), resulting in different rates of dissolution. It has been reported that the dissolution of WC in the Ni binder phase can be limited if the total nitrogen concentration is high in a cermet (Doi et al., 1985). This can be attributed to the low chemical affinity between W and N.Figure 17 shows the microstructures of the systems composed of the Ti(C0.7N0.3), Ti(C0.5N0.5) and Ti(C0.3N0.7) phases, respectively, and it demonstrates the
Figure 16 An SEM schematic of Ti(C,N)-based cermets.
influence of N in Ti(C1xNx) on the microstructure. For other systems (Ahn & Kang, 1998; Fukuhara & Mitani, 1982; Jung et al., 1999; Mun & Kang, 1999), it has been reported that the microstructure is strongly affected by the nitrogen content of Ti(C,N). Somewhat different results were reported byFukuhara and Mitani (1982)for a Ti(C,N)–Mo–Ni system via the P/M technique and byQi and Kang (1998)from a Ti(C,N)–NbC–Ni system via infiltration techniques. Both investigators showed that when Mo or NbC were added to Ti(C,N)–Ni, the Ti(C0.5N0.5) phase exhibited the slowest dissolution rate among the various Ti(C1xNx) phases. Based on ther- modynamic calculations for Ti(C,N) stability (Jung et al., 1999), it has been reported that Ti(C0.3N0.7) is the most stable compound in the 1400–1600C range and it is consistent with the result fromFigure 17.
Figure 17 SEM/BSE micrographs of Ti(C1xNx)–10WC–20Ni systems sintered at 1510C in vacuum for 1 h: (a)x¼0.3 (b)x¼0.5 (c)x¼0.7.
Cermets 161
1.05.2.2.3 Dissolution Rates of Secondary Carbides
The cores of undissolved Ti(C,N) particles act as nucleation sites for the rim structure. The rims are forming when oversaturated solutes precipitate out from the melt on the Ti(C,N) particles (Gee, Reece, & Roebuck, 1992). Thus, the equilibrium reactions in the dissolution and precipitation of the hard phase during sintering can be inferred from the composition analysis of the core/rim structure as long as the solute content is within solubility limit in a carbide system (Doi et al., 1985; Mun & Kang, 1999). It has been regarded that the inner rim forms during the heating stage from the onset temperature for liquid forming (w1300C), whereas the outer rim forms at the sintering temperature (1510C) (Doi et al., 1985; Gee et al., 1992).
The dissolution rate of WC was measured to be approximately two andfive times faster than that of Ti(C,N) in the system at 1300 and 1510C, respectively (Ahn & Kang, 2000). The results were interpreted in terms of phase stability and precipitation phenomena. Ahn and Kang (2001) reported the dissolution behaviors of Ti(C0.7N0.3) with various secondary carbides such as HfC, TaC and WC. The average dissolution rate of Ti(C0.7N0.3) in the HfC-containing system is foundw1.6 and 1.9 times higher than those for Ti(C0.7N0.3) in the TaC- and WC-containing systems at the same sintering conditions.Figure 18shows the SEM/backscattered electron micrographs of Ti(C0.7N0.3)–10 MC–30Ni systems sintered at 1510C in vacuum for 1 h. MC is (a) HfC, (b) TaC, and (c) WC, which are the basis for the measurement of dissolution rate.
1.05.2.2.4 Factors Influencing Dissolution Behavior
The particle dissolution is an interesting issue, which is complicated by various thermodynamics and kinetics factors. In practice, the dissolution phenomena of a compound in a liquid melt can be considered as two separate steps: if the dissolution of the particle is limited to its surface or in the neighborhood of the surface, (1) the elemental dissociation of a compound and (2) mixing between the dissociated species and the surrounding materials, such as a liquid melt. Thefirst step, the dissociation of a compound particle, can be described as below:
J ¼ A n vexpðQ=RTÞexp
DGf=RT
; (25)
whereJis theflux of dissociating species from the particle to the melts,Ais the probability of available jump sites in the melt for the elements,nis the number of possible unit compounds to be dissociated,nis the Debye
Figure 18 SEM/BSE micrographs of Ti(C0.7N0.3)–10 MC–30Ni systems sintered at 1510C in vacuum for 1 h. MC is (a) HfC, (b) TaC, and (c) WC.
frequency,Qis the activation energy for particle dissociation, andDGfis the energy of formation of the com- pound from the elements. In the case of carbide dissolution in Ni melts, the reverse reaction can be assumed to be negligible. InEqn (25), the probability of available jump sites,A, is closely related to the solubility of the dissolving species in the liquid melt. This is, in turn, determined by interactions (coefficients) among species in the melts as the number of elements involved in the system increases. This issue has been well documented for the effect of secondary solutes on the solubility of a primary solute in thefield of steel production (Pehlke &
Elliot, 1960). A similar approach as shown in Eqn (25) can be found to describe the migration of grain boundaries (Crank, 1956; Shewman, 1989). The equation indicates that little dissociation would occur due to the high energy of formation term unless it is in a liquid melt.
The driving force for the dissolution of dissociated elements in the melt is the reduction in total free energy of the system, as the result of the entropy of mixing and solid–liquid instability. Thus, for the case of (1) the free energy of mixing,DGM, and (2) strain energy,DGε, due to the instability of the solid–liquid interface and (3) a surface-related term,DGg(Park et al., 1998; Chun et al., 1993; Rhee and Yoon, 1987) are additional thermo- dynamic factors that are needed to describe the dissolution behavior. Thus,Eqn (25)can be expressed as below:
J ¼ aA n vexpð Q=RTÞexp
DGf=RT
; (26)
where DGrxn¼DGfþDGMþDGgþDGεþ. and a is a constant. The surface energy not only determines particle morphology but also influences the level of strain energy at the solid–liquid interface in a complicated manner. Different shapes of a compound particle in a solution with slightly different compositions are not uncommon.
Further, the transport of the species is often described by the mobility and a gradient of activity co- efficient of the diffusing species with respect to concentrations. The product of these factors determines the diffusivity of the species at a given temperature. Therefore, at a given temperature, the rate-limiting step in the transport for dissolution–precipitation or coarsening is generally determined by a domi- nating factor between (1) the thermal stability of the compounds to be dissolved (interface-controlled) and (2) difficulties in transport due to interactions among the diffusing species and the melts (diffusion- controlled).
The above factors show that the dissolution phenomena involve many intricate thermodynamic and kinetic factors and that the accurate quantification of these factors is not an easy task.
1.05.2.2.5 Onset of Core/Rim Structure Formation
A number of studies related to the microstructure of Ti(C,N)-based cermets have appeared (Suzuki et al., 1971;
Zhang, 1993). However, issues regarding thefinal microstructures, especially regarding the onset of formation, remain ambiguous and controversial. Some studies suggest that thefinal rim structure is formed during thefinal sintering step via dissolution of constituent carbides into the liquid binder and its reprecipitation in the form of a solid solution onto the Ti(C,N) particles (Ahn et al., 2001; Ahn & Kang, 2000; Ettmayer et al., 1995). Early studies have proposed that the inner rims are formed by solid-state reactions or during an early stage of sintering (w1300C) while the outer rims form thereafter via a dissolution/reprecipitation process (Doi et al., 1985;
Lihndahl, Gustafson, Rolander, Stals, & Andrén, 1999; Zackrisson & Andrén, 1999). An extensive investigation was made on this issue to clarify the onset time for formation of the final rim structure with Ti(C0.7N0.3)–xWC–20Ni (x¼5–25 wt%) systems (Kim, Min, & Kang, 2003). The study concluded that the inner rim forms at sintering temperature and the outer rim forms thereafter during the sintering or cooling stage of sintering.
It would be anticipated that rim structures start forming on the Ti(C,N) particles when liquid Ni becomes saturated with solutes (Pehlke & Elliot, 1960). Assuming that the rim phase forms at compositional equilib- rium, the solid solution of the rim should have the same Ti/W ratio in composition as that in the liquid Ni.
Thus, reactions occurring during the dissolution/reprecipitation process could also be inferred from the composition analysis of the core/rim structure (Figure 19(b)) (Ahn et al., 2001; Ahn & Kang, 2000; Doi et al., 1985; Lihndahl et al., 1999; Qi & Kang, 1998; Zackrisson & Andrén, 1999). In thefigure, 1510C/0 h stands for immediate cooling after furnace temperature reaches to 1510C.
The increase in the solubility of W in liquid Ni with temperature is consistent with thermodynamic prin- ciples and phase diagrams for W–Ni and WC–Ni(Co) (Exner, 1979; Ettmayer et al., 1995). The relation be- tween sintering temperature and W content in the rim phase, as shown inFigure 19, reveals that the solid solution in the outer rim could be formed at lower temperatures than that in the inner rim or at the same sintering temperature by the dissolution of Ti(CN) cores and inner rim phase at the sintering temperature. In Cermets 163
addition, from a location point of view, the outer rim forms at a later time than the inner rim. Based on these two facts, it can be concluded that the onset of inner rim formation in“thefinal microstructure”takes place at the sintering temperature, 1510C. A portion of outer rim or the entire outer rim forms thereafter near or at the cooling stage, depending on the initial WC content. The unusually high amount of W in the Figure 19 (a) TEM micrograph of a Ti(C0.7N0.3)–15WC–20Ni system; (b,c) TEM/EDS concentrations of tungsten, for
Ti(C0.7N0.3)–xWC–20Ni systems for various sintering temperatures: (b) inner rim (c) outer rim.
Ti(C,N)–5WC–20Ni sintered at 1510C for 0 h inFigure 19might be due to the combined effect of the short time for carbide dissolution and the high dissolution rate of WC compared to that of Ti(C,N) (Ahn & Kang, 2000; Shewman, 1989).
It is possible for the solid solution to form inner rim at any temperature. However, the rim phases formed at low temperatures in early stages become thermodynamically unstable and tend to redissolve at high temper- atures. Thus, the rim phases formed by solid-state diffusion or during the early stage of sintering (w<1400C) with the initial liquid Ni cannot constitute the final rim structure. Figure 20 is the microstructure of a Ti(C,N)–25WC–20Ni system infiltrated at 1510C for 0, 10, 30, 60 min (Kim et al., 2003). The change in the inner rim compositions of the infiltrated samples is shown inFigure 21. Based on the information inFigure 22, the coalescence of Ti(C,N) is likely to occur at the sintering temperature to a great extent prior to the formation of thefinal rim phases.
Figure 20 SEM/BSE micrographs of Ti(C0.7N0.3)–25WC–20Ni systems infiltrated at 1510C for (a) 0 min, (b) 10 min, (c) 30 min, and (d) 60 min in vacuum.
Figure 21 TEM/EDS concentrations of tungsten, for Ti(C0.7N0.3)–xWC–20Ni systems for various sintering times, (a) inner rim and (b) outer rim.
Cermets 165