Regular Article
Solidification Microstructure and Mechanical Properties of B1-type TiC in Fe–Ti–C Ternary Alloys
2024 Volume 64 Issue 11 Pages 1723-1731
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2024 Volume 64 Issue 11 Pages 1723-1731
The microstructure of the B1-type TiC formed during solidification and its mechanical properties were investigated using arc-melted Fe–Ti–C ternary alloys. The TiC formed at relatively high temperatures in the liquid as the primary phase exhibited a dendritic shape. With decreasing temperature and/or decreasing Ti and C content in the liquid, the morphology of the TiC changed to a cubic shape with a {001}TiC habit plane, a plate shape with a {011}TiC habit plane, and a needle shape with a preferential growth direction of <001>TiC. The morphology of the TiC was characterized by the anisotropy of its surface energy and its growth rate. The cubic shape with a {001}TiC habit plane was formed as a result of the {001}TiC surface exhibiting the lowest surface energy among the TiC surfaces. However, the plate shape with a {011}TiC habit plane and the needle shape with a <001>TiC preferential growth direction likely formed because the slowest and fastest growth rates corresponded to the <011>TiC and <001>TiC directions, respectively. At room temperature, the alloy with dendritic TiC was fractured in the elastic deformation region because TiC exhibited no plastic deformation. However, the results obtained at 800°C suggested that the TiC exhibited plastic deformability and that the alloy with the dendritic TiC was also plastically deformed.
B1-type MX ceramics have the NaCl structure, where the main components M and X are transition-metal and nonmetal elements, respectively. One of these ceramics, TiC, has a high melting point (3067°C), low density (4.93 g/cm3), and high hardness, making it a promising strengthening phase for high-temperature materials.1,2) However, one of the drawbacks of TiC is its poor toughness. The room-temperature fracture toughness of TiC with a stoichiometric composition is ~3 MPa
When B1-type MX ceramics are in equilibrium with a metal phase, various metallic elements of the equilibrium metal phase are substituted at the M site and several nonmetallic elements such as carbon, nitrogen, and oxygen are substituted at the X site, which introduces structural vacancies because of off-stoichiometry. Such deviations from stoichiometric compositions can alter the physical properties of MX ceramics, which can, in turn, affect their mechanical properties.10,11,12) Solidified MX ceramics are likely to have a large off-stoichiometry because of undercooling and other effects. In addition, during solidification, the anisotropy of the solid/liquid interfacial energy can promote or inhibit the growth of certain planes, resulting in a solidified MX ceramic with a unique morphology. Because the morphology of MX ceramics is expected to be one of the factors that govern the toughness of alloys, controlling the solidification microstructure of MX ceramics is critical for designing alloys with high strength without losing toughness and deformability, even when large amounts of MX ceramics are introduced. However, in the sintering process, which is the main manufacturing process for cermets, the ceramic and metal phases are usually nonequilibrium, making the introduction of off-stoichiometric effects and control of the morphology difficult. For example, the microstructure of Mo–TiC eutectic alloys fabricated by casting completely differ from that of Mo–TiC eutectic alloys fabricated by sintering.4,13,14) Alloys with a fine eutectic lamellar microstructure consisting of the Mo phase and TiC formed by the casting process exhibit high-temperature strength and ductility superior to those of the corresponding alloys prepared by the sintering process.13,14) Similar behavior has been reported for Mo–ZrC eutectic alloys.15)
The literature includes several reports on the solidification morphology of TiC. TiC in Fe-based alloys formed by the eutectic reaction of L → Fe (body-centered cubic (BCC) structure) + TiC has a habit plane.16,17,18) TiC with a similar habit plane has also been observed in FeAl-based composites,19,20) TiAl alloys,21,22) and Al composites.23) However, the formation of such a habit plane has not been reported for TiC in Mo-based alloys formed by the eutectic reaction of L → Mo (BCC structure) + TiC.24,25) Here, the M–Ti–C ternary system can be divided into two patterns in terms of the off-stoichiometry of TiC: cases in which TiC in equilibrium with the metal or liquid phase has substantial off-stoichiometry, and cases in which it does not. Examples of the former are Mo–Ti–C and Nb–Ti–C ternary systems, in which M is mainly an element of group 4–6.24,26) Examples of the latter include Fe–Ti–C and Ni–Ti–C ternary systems.27,28)
In the present study, we investigate the Fe–Ti–C ternary system to discuss the solidification morphology of TiC, which exhibits almost no off-stoichiometry. The TiC in equilibrium with the metal or liquid phase in this system is an almost stoichiometric composition irrespective of temperature, and the solidification microstructure of the TiC exhibits little compositional change during solidification. Therefore, in the present study, we investigated the solidification microstructure and mechanical properties of TiC using Fe–Ti–C ternary model alloys.
The atomic ratios of the investigated alloys were Fe-15Ti-15C (15TiC alloy), Fe-5Ti-5C (5TiC alloy), and Fe-7.5Ti-2.5C (5Ti-2.5TiC alloy). According to the Fe–Ti–C ternary liquidus projection, all of the alloys form primary TiC and the liquidus temperature of the 15TiC alloy is higher than those of the 5TiC and 5Ti-2.5TiC alloys.27) These alloys were prepared as 9–10 cm3 ingots by arc melting under an Ar atmosphere using pure Fe (99.99 wt%), Ti (99.9 wt%), and TiC (99 wt%) as raw materials. Hereafter, all compositions are expressed in atomic ratio. The alloy compositions were analyzed by inductively coupled plasma atomic emission spectroscopy; the analyzed chemical compositions of the as-cast alloys are shown in Table 1. The chemical and nominal compositions did not substantially differ. Field-emission scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were used for microstructural observations. TEM disk specimens with a 3 mm diameter were cut by electrical-discharge machining, mechanically polished to a thickness of approximately 50–70 μm, then twin-jet electropolished in an ethanolic solution of 10 vol% perchloric acid. Part of the SEM specimens was deep etched by immersion in 4% nital solution for 30 min for three-dimensional observation of the TiC. Compositional analyses of the phases were performed by field-emission electron probe microanalyzer (EPMA) equipped with a wavelength-dispersive X-ray spectroscopy (WDX) operated 10 kV and 5.0 × 10−8 A. The orientation was determined by electron backscatter diffraction (EBSD) analysis. The specimen for tensile testing was prepared by electrical-discharge machining into a dog-bone shape with a gage length of 5 mm, a width of 2 mm, and a thickness of 1 mm. The strain rate for the tensile test was 1.0 × 10−3 s−1.
| Sample | Nominal composition/at% | Chemical composition/at% | ||||
|---|---|---|---|---|---|---|
| Fe | Ti | C | Fe | Ti | C | |
| 15TiC | 70.0 | 15.0 | 15.0 | 70.4 | 14.5 | 15.1 |
| 5TiC | 90.0 | 5.0 | 5.0 | 91.0 | 4.6 | 4.4 |
| 5Ti-2.5TiC | 90.0 | 7.5 | 2.5 | 90.5 | 7.2 | 2.3 |
Figure 1 shows a backscattered electron image (BEI), the kernel average misorientation (KAM) map of TiC, and the inverse pole figure (IPF) map with respect to the normal direction of the sample surface of the as-cast 15TiC alloy. The other IPF maps in this paper also show the orientation with respect to the normal direction of the sample surface. In the BEI, a relatively large primary TiC with dark contrast, an Fe phase (BCC structure at room temperature) with bright contrast surrounding the primary TiC, and a eutectic microstructure containing the Fe phase and TiC are observed (Fig. 1(a)). The volume fraction of TiC in the 15TiC alloy was 30%. Almost no orientation change was observed in the primary TiC, whereas a small orientation change was observed in the eutectic TiC (Fig. 1(b)). The primary TiC exhibited mainly a dendritic morphology; however, the tips of the dendrite arms were faceted.
To study the solidification process of this primary TiC, we calculated the Fe fraction at the Ti sites (
| (1) |
where
Table 2 shows the compositions and
| Sample | Composition/at% | |||
|---|---|---|---|---|
| Fe | Ti | C | ||
| Dendric TiC | 0.6 | 52.9 | 46.5 | 1.1 |
| Faceted TiC | 0.7 | 52.6 | 46.7 | 1.3 |
Figure 2 shows the as-cast microstructure of the 5TiC alloy. The primary TiC was not dendritic but had faceted habit planes (Fig. 2(a)). The volume fraction of TiC in the 5TiC alloy was 8%. As in the 15TiC alloy, almost no orientation change was observed in the primary TiC, whereas a small orientation change was observed in the eutectic TiC (Fig. 2(b)). The habit plane of the primary TiC in the 5TiC alloy was {001}TiC, which is the same as that of the facets of the primary TiC in the 15TiC alloy. The habit plane of TiC around the primary TiC is considered to be {011}TiC because it is always parallel to the trace of the {011}TiC. A secondary electron image (SEI) of the deep-etched specimen shows that the primary TiC was cubic shaped and that plate-shaped TiC grew from the primary TiC toward the liquid phase (Fig. 2(d)). The plate-shaped TiC had the same orientation as the cubic primary TiC, suggesting that it was formed by growth from the primary TiC (Figs. 2(c) and 2(d)). In the eutectic microstructure, plate-shaped TiC and needle-shaped TiC grown from the tip of the plate-shaped TiC were observed (Fig. 2(e)). The habit plane of the plate-shaped TiC in the eutectic structure was {011}TiC, like that of the plate-shaped TiC grown from primary TiC (Fig. 2(f)). The angle between needle-shaped TiC particles was almost 90° (Fig. 2(e)). The dotted TiC was considered to be needle-shaped TiC grown in the direction normal to the sample surface (Figs. 2(c) and 2(f)), and the orientation of the needle-shaped TiC was close to {001}TiC (Fig. 2(c)). Therefore, the preferential growth direction of the needle-shaped TiC is considered to be <001>TiC.
Figure 3 shows the as-cast microstructure of the 5Ti-2.5TiC alloy. The primary TiC and eutectic TiC had faceted habit planes like those of 15TiC and 5TiC, and the volume fraction of TiC was 5% (Fig. 3(a)). However, the Fe matrix was very coarse. According to the liquidus surface projection,27) the Fe phase in the 5Ti-2.5TiC alloy has a BCC structure at the eutectic reaction temperature and no transformation occurs during cooling. Therefore, the Fe phase in the 5Ti-2.5TiC alloy became much coarser than that in the 5TiC and 15TiC alloys. However, the habit plane of the primary TiC was {001}TiC (Fig. 3(b)), which suggests that the morphology of the primary TiC in the 5Ti-2.5TiC alloy is cubic, like the primary TiC in the 5TiC alloy. The habit plane of the eutectic TiC was {011}TiC, which is the same as that of the plate-shaped TiC in the 5TiC alloy (Figs. 2(f), 3(c), and 3(d)).
The solidification of the TiC phase in Fe–Ti–C ternary alloys with primary TiC occurs, in order from the early stage of solidification, dendritic TiC, dendritic TiC with a habit plane of {001}TiC at the edge, cubic-shaped TiC with a habit plane of {001}TiC, plate-shaped TiC with a habit plane of {011}TiC, and needle-shaped TiC with a preferential growth direction of <001>TiC. The shape of the TiC was found to vary depending on the liquidus temperature and the liquidus composition at the time of the growth.
In general, the solidification microstructure morphology is controlled primarily by the crystallographic anisotropy of the surface energy and the growth rate. In particular, the relationship between surface energy and crystallographic morphology is discussed on the basis of the corresponding γ-plot.29,30) At relatively low temperatures, the effect of surface-energy anisotropy is substantial and the faceted plane is formed. In this case, the morphology is governed by the surface-energy anisotropy. Above a certain temperature (i.e., the roughening temperature), the surface-energy anisotropy weakens because of entropy effects; although anisotropic growth occurs, the crystals no longer have a habit plane. Similarly, when the growth rate dominates the morphology, the morphology is determined by the dependence of the growth rate of each surface.30) For example, the end of a dendrite of a metal with the FCC structure is surrounded by {111}FCC because {111}FCC is the densest surface and has the lowest surface energy.31,32,33) In this case, the predominant growth direction would be <001>FCC surrounded by four equivalent {111}FCC.34)
A similar phenomenon is considered to have occurred during the solidification process of B1-type TiC in Fe–Ti–C ternary alloys. In the early stage of solidification, if the temperature of the liquid phase is sufficiently high, the growth rate is relatively isotropic because of weak surface-energy anisotropy caused by entropy effects and the TiC is considered to be dendritic, with an irregular surface profile. According to the reported liquid surface projection, primary TiC is formed from ~2100°C in 15TiC and ~1600°C in 5TiC, with eutectic reactions occurring at ~1450°C.27) Therefore, TiC formed above ~1600°C is considered to be dendritic. In fact, TiC in Mo-based alloys that form TiC at temperatures greater than 1600°C is dendritic.24,25,35) The solidification rate and composition are factors that affect the solidification microstructure. However, dendritic-shaped TiC was formed both in the case of a relatively high cooling rate24,25) such as that in arc melting and in the case of a low cooling rate,35) suggesting that the solidification rate does not strongly affect the formation of dendritic-shaped TiC. The composition of TiC is also inferred to have no substantial effect on the formation of dendritic-shaped TiC because the TiC in the Mo-based alloys has a large off-stoichiometry, whereas the TiC in the present study exhibits almost no off-stoichiometry; however, both formed dendritic-shaped TiC at high temperatures.
As solidification progressed, dendritic-shaped primary TiC with a {001}TiC habit-plane edge or cubic-shaped primary TiC with a {001}TiC habit plane was formed. The crystal-growth behavior of B1-type TiC and TiN varied depending on the ratio of the {001} and {111} surface energies, and it has been reported to change from cubic with a {001} habit plane to 14-plane with both {001} and {111} habit planes and to octahedral with a {111} habit plane.36) Figure 4 shows the crystal structure of B1-type TiC and the atomic arrangement in the {011}TiC, {001}TiC, and {111}TiC planes. The {001}TiC is the densest plane of B1-type TiC and has been reported to have a lower surface energy than the {011}TiC and {111}TiC planes.37,38) As solidification progresses, the temperature of the liquid phase should gradually decrease and the composition of the liquid phase should change accordingly. These changes should gradually enhance the effect of the TiC surface energy and lead to the growth of TiC with a {001}TiC habit plane. Cubic-shaped TiC has also been reported to form in FeAl-based composites at approximately the same temperature at which cubic-shaped primary TiC in 5TiC was formed, although the habit plane has not been identified.20,21) However, tetradecahedral TiC has been reported to form in Al composites.23) The crystal growth behavior of these TiC phases is also attributable to the surface energy.
With a further decrease in the liquidus temperature and concentration of the solute elements in the liquidus phase, plate-shaped TiC with a {011}TiC habit plane was formed. The surface energy of {011}TiC has been reported to be higher than that of {001}TiC and {111}TiC.37,38) From this viewpoint, the formation of plate-shaped TiC might be strongly influenced by the anisotropy of the growth rate rather than by the surface energy. The growth rate is affected by elemental interactions: the Ti–Ti interaction force in Fe solvent is positive, whereas the Ti–C interaction force is strongly negative.39) This interaction suggests that, in the liquid phase, an attraction exists between Ti and C atoms and a repulsion exists between Ti atoms. Therefore, the growth rate in the direction formed solely from Ti should be slower than that in the direction where Ti and C are mixed. The <110>TiC direction is the only direction that consists of only Ti or only C. Therefore, the growth rate in the <110>TiC direction was likely the lowest, which might have resulted in the formation of plate-shaped TiC with a {011}TiC habit plane.
As solidification proceeds, the liquidus temperature decreases further. As a result, the degree of undercooling increases and the anisotropy of the growth rate is expected to become more substantial. Because the <001>TiC direction has Ti and C in nearest neighbor positions, the growth rate in this direction should be the highest. Therefore, we speculate that needle-shaped TiC with <001>TiC as the preferential growth direction might have formed in the final stage of solidification.
3.2. Mechanical Properties at Room Temperature and 800°CFigure 5 shows the nominal stress–plastic strain curves of 15TiC alloy, 5TiC alloy, and 5Ti-2.5TiC alloy at room temperature. The 15TiC alloy ruptured within the elastic deformation region, whereas the 5TiC alloy exhibited a yield stress of ~700 MPa, a plastic strain of ~1%, and a rupture stress of ~830 MPa. For the 5Ti-2.5TiC alloy, the yield stress decreased to ~250 MPa; however, the plastic strain was ~13% and the rupture stress was ~500 MPa. SEM images showing the fractography and cross section near the fracture surface after the room temperature tensile test are shown in Fig. 6. Most of the fracture surface of the 15TiC alloy was covered by denticulate fractures with smooth dendritic-shaped TiC (Fig. 6(a)). Therefore, it is considered that the cracks propagated mainly through the dendritic-shaped TiC. By contrast, the fracture surface of the 5TiC alloy showed cubic TiC without fracture (Fig. 6(b)), suggesting that interfacial delamination between the cubic TiC and the Fe phase occurred. However, most of the fracture surface of the 5Ti-2.5TiC alloy was occupied by the Fe phase with rever patterns (Fig. 6(c)). The cracks were wavy and secondary crack formation was also observed (Fig. 6(d)). Some of the cracks were observed to be blocked by plate-shaped TiC (Fig. 6(e)).
At room temperature, the 5TiC alloy was plastically deformed but the plastic elongation was small; by contrast, the 15TiC alloy did not exhibit plastic deformation. To investigate the deformation behavior of these alloys, high-temperature tensile tests were conducted. Figure 7 shows the nominal stress–plastic strain curves of 15TiC alloy and 5TiC alloy at 800°C. The 15TiC alloy showed a yield stress of ~200 MPa and a tensile strength of ~220 MPa at a plastic strain of 1.5%. The 5TiC alloy exhibited a yield stress of approximately 90 MPa and a tensile strength of approximately 105 MPa at a plastic strain of approximately 1.5%. The fracture strains of both alloys at 800°C were much higher than those at room temperature. Figure 8 shows the microstructure and KAM map of cross sections of the fracture surfaces of the 15TiC and 5TiC alloys after the tensile test at 800°C. The voids and cracks in the 15TiC alloy were mainly observed in the dendritic-shaped TiC (Fig. 8(a)). In the dendritic-shaped TiC, orientation changes, which were not observed in the as-cast specimen, occurred irrespective of the formation of cracks (Figs. 2(b) and 8(b)). Therefore, the results suggest that, at 800°C, the dendritic-shaped TiC and plate-shaped TiC in the 15TiC alloy deformed plastically, accompanied by plastic deformation of the Fe phase, and cracks formed in the TiC at the maximum tensile strength, resulting in a drastic decrease in stress. The cracks in the dendritic-shaped TiC were parallel to the {001}TiC trace, suggesting that the cracks in the TiC propagated mainly through {001}TiC (Fig. 8(c)). This interpretation is consistent with the reported cleavage planes of B1-type carbides.40,41) By contrast, in the 5TiC alloy, voids and cracks were observed in the Fe phase but not in the cubic-shaped TiC (Fig. 8(d)). In addition, as shown in the KAM map, a large orientation change was observed in the plate-shaped TiC but not in the cubic-shaped TiC (Figs. 2(b) and 8(e)). Therefore, plastic deformation of the 5TiC alloy at 800°C is considered to have occurred mainly in the Fe phase and in the plate-shaped TiC.
The 15TiC alloy with coarse dendritic-shaped TiC failed within the elastic deformation region at room temperature. This result indicates that the dendrite-like TiC lacks plastic deformability at room temperature. However, the 5TiC alloys and 5Ti-2.5TiC alloys, which do not contain dendritic-shaped TiC, exhibited plastic strain at room temperature. These results suggest that the plastic deformation in the 5TiC and 5Ti-2.5TiC alloys at room temperature is caused by plastic deformation of the Fe phase. At 800°C, the 15TiC alloy yielded at a higher stress than the 5TiC alloy and exhibited a plastic strain of ~10%. An orientation change was also observed within the dendritic- and plate-shape TiC, suggesting that TiC exhibits slight plastic deformability at 800°C. However, no orientation change was observed within the cubic-shaped TiC of the 5TiC alloy, indicating that the cubic-shaped TiC contributed little to the plastic deformation behavior.
The solidification microstructure and mechanical properties of TiC were investigated using Fe–Ti–C ternary model alloys prepared by the arc-melting technique. The conclusions are summarized as follows:
(1) The TiC formed at high temperatures exhibited a dendritic morphology. With decreasing liquidus temperature and concentration of solute elements in the liquidus, the morphology of the TiC changed to cubic-shaped with a {001}TiC habit plane, plate-shaped with a {011}TiC habit plane, and needle-shaped with a <001>TiC preferential growth direction.
(2) The morphology of TiC was determined by the anisotropy of the surface energy and the growth rate of TiC. Cubic-shaped TiC with a {001}TiC habit plane was formed because of the minimum surface energy of the {001}TiC surface, whereas plate-shaped TiC with a {011}TiC habit plane and needle-shaped TiC with a <001>TiC preferential growth direction were formed because they exhibit the slowest and fastest growth rates, respectively.
(3) At room temperature, TiC exhibited no plastic deformability and the 15TiC alloy with dendritic-shaped TiC ruptured in the elastic deformation region. However, the results suggested that TiC exhibited slight plastic deformability at 800°C.
This research was financially supported by The Iron and Steel Institute of Japan, Research Group for High-strengthening Theory in High-temperature Materials.
