Morphology evolution of Ir–Nb–X (X = Hf, Ta, or Ti) ternary alloys

Journal of Alloys and Compounds 428 (2007) 220–229

Morphology evolution of Ir–Nb–X (X = Hf, Ta, or Ti) ternary alloys C. Huang a,∗ , Y. Yamabe-Mitarai b , H. Harada b a

b

China Institute of Atomic Energy (CIAE), P.O. Box 275-108, Beijing 102413, China National Institute for Materials Science (NIMS), Sengen 1-2-1, Tsukuba, Ibaraki 305-0047, Japan Received 13 December 2005; accepted 27 January 2006 Available online 7 July 2006

Abstract The microstructure evolution of nine samples from three Ir-base ternary systems, Ir–Nb–Hf, Ir–Nb–Ta, and Ir–Nb–Ti, was investigated by microstructure observation using scanning electron microscopy (SEM), composition map-analysis using electron probe microscopy analysis (EPMA), and phase determination using X-ray diffraction (XRD) patterns. The fcc/L12 two-phase structure was detected in all the samples. Lattice misfits between fcc and L12 phases were calculated. Ir–Nb–Ta and Ir–Nb–Ti alloys exhibited a microstructure quite similar to that of Ni-base superalloys, and the cuboidal L12 precipitates in Ir–Nb–Ta and Ir–Nb–Ti alloys could maintain up to 1900 ◦ C. © 2006 Published by Elsevier B.V. Keywords: Intermetallics; High-temperature alloys; Microstructure

1. Introduction There is a great deal of interest in the development of a new generation of high-temperature alloys that will exceed Ni-base superalloys. A coherent fcc/L12 two-phase structure in which cuboidal L12 precipitates embed in the fcc matrix is believed to be the starting point for Ni-base superalloys to become a fascinating high-temperature material. Accordingly, one alternative method is to develop a novel superalloy with a microstructure that is analogous to that of Ni-base superalloys by applying high-melting-point elements to obtain higher high-temperature capabilities. Ir, which belongs to platinum group metals (PGMs), has potential as such a matrix element because of its fcc structure (the same structure as Ni), high melting point (2447 ◦ C), and excellent environmental resistance. Ir-base binary superalloys, Ir–Zr, Ir–Hf, Ir–Nb, Ir–Ta, and Ir–Ti, have been proposed and investigated in our group [1,2]. A coherent fcc/L12 two-phase structure was found in all of these systems. Among them, the Ir–Nb alloys presented cuboidal precipitates, whose microstructure was the most analogous to that of Ni-base superalloys. The lattice misfits between fcc and L12 phases in these alloys were calculated to be about 2.2, 1.9, 0.4, 0.3, and 0.1% for Ir–Zr, Ir–Hf, Ir–Nb, Ir–Ta, and Ir–Ti alloys, respectively. Lately, the study on ∗

Corresponding author. Tel.: +86 10 69358135; fax: +86 10 69358126. E-mail address: [email protected] (C. Huang).

0925-8388/$ – see front matter © 2006 Published by Elsevier B.V. doi:10.1016/j.jallcom.2006.01.111

Ir–Nb–Zr ternary alloys by combing Ir–Nb and Ir–Zr binary alloys by Yamabe-Mitarai et al. [3,4] prevailed the formation of fcc/L12 two-phase structure as well as stronger creep-resistance than either of Ir–Nb or Ir–Zr binary alloys. This prompts the present investigation of three ternary alloys by combining Ir–Nb with other binary systems (Ir–Hf, Ir–Ta, or Ir–Ti). In this paper, a primary investigation on these three ternary alloy systems, Ir–Nb–Hf, Ir–Nb–Ta, and Ir–Nb–Ti, is carried out to evaluate the microstructure evolution. Hf, Ta, and Ti, as the third elements added to the Ir–Nb system, are selected according to the following considerations. First, the fcc/L12 twophase structure exists in Ir–Hf, Ir–Ta, and Ir–Ti binary alloys, and the differences in the lattice parameters among Ir3 Hf, Ir3 Ta, Ir3 Ti, and Ir3 Nb are small [5]. This implies that a continuous solid solution between the two-pair L12 phases (Ir3 Nb–Ir3 Hf, Ir3 Nb–Ir3 Ta, and Ir3 Nb–Ir3 Ti) can form and, furthermore, that the fcc/L12 structure in the ternary systems can be obtained. Second, Ir–Hf, Ir–Ta, and Ir–Ti binary alloys have been previously investigated, and the resulting data can be used for comparison. 2. Experimental procedure Ir (99.9 wt.% purity) powder, Nb (99.9 wt.% purity) powder, Hf (99.4 wt.% purity) powder, Ta (99.99 wt.% purity) plate, and Ti (99.9 wt.% purity) powder were prepared. Nine button ingot samples, 15 g each, were obtained by arcmelting the above starting materials in an Ar atmosphere. The composition of the alloys is shown in Table 1. Two binary compositions of Ir–17%Nb and Ir–15%Hf (or Ir–18%Ta or Ir–15%Ti) in three proportions, 25:75, 50:50, and 75:25, were

C. Huang et al. / Journal of Alloys and Compounds 428 (2007) 220–229 Table 1 Compositions of the samples investigated

Table 2 Average area composition obtained by EPMA

Sample no.

Sample no.

Composition (at.%) Mixing proportion

Nominal composition

Ir–Nb–Hf Hf1 Hf2 Hf3

Ir–17Nb:Ir–15Hf = 25:75 Ir–17Nb:Ir–15Hf = 50:50 Ir–17Nb:Ir–15Hf = 75:25

84.5Ir–4.25Nb–11.25Hf 84Ir–8.5Nb–7.5Hf 83.5Ir–12.75Nb–3.75Hf

Ir–Nb–Ta Ta1 Ta2 Ta3

Ir–17Nb:Ir–18Ta = 25:75 Ir–17Nb:Ir–18Ta = 50:50 Ir–17Nb:Ir–18Ta = 75:25

Ir–Nb–Ti Ti1 Ti2 Ti3

Ir–17Nb:Ir–15Ti = 25:75 Ir–17Nb:Ir–15Ti = 50:50 Ir–17Nb:Ir–15Ti = 75:25

Region A (at.%)

221

Region B (at.%)

Ir

Nb

Third element

Ir

Ir–Nb–Hf Hf1 Hf2 Hf3

83.78 84.00 84.96

4.79 4.72 12.78

11.05 6.73 2.26

91.02 89.55 88.01

1.93 4.41 9.17

7.05 6.04 2.81

82.25Ir–4.25Nb–13.5Ta 82.5Ir–8.5Nb–9Ta 82.75Ir–12.75Nb–4.5Ta

Ir–Nb–Ta Ta1 Ta2 Ta3

82.28 78.80 83.52

3.62 8.64 11.61

14.10 12.56 4.87

85.93 85.24 85.52

4.15 7.71 11.37

9.92 7.06 3.11

84.5Ir–4.25Nb–11.25Ti 84Ir–8.5Nb–7.5Ti 83.5Ir–12.75Nb–3.75Ti

Ir–Nb–Ti Ti1 Ti2 Ti3

87.06 85.75 84.36

3.86 8.21 12.37

9.08 6.03 3.28

84.41 83.48 83.38

4.34 8.93 12.88

11.24 7.59 3.74

mixed to form the ternary alloys. These binary compositions of Ir–17%Nb, Ir–15%Hf, Ir–18%Ta, and Ir–15%Ti were selected because an fcc/L12 twophase structure with a 50% volume fraction of L12 phase could form in these composition ranges according to the binary phase diagrams. Cylindrical samples, 3 mm in diameter and 5–6 mm in height, were cut by an electro-discharge machine (EDM) and heat-treated at four temperatures of 1500, 1700, 1900, and 2000 ◦ C. All the heat treatments were conducted in a vacuum furnace for 168 h (168 h heat-treatment was applied in order to get the equilibrium phases of Ir-base alloys according to our experience), and the samples were then cooled within the furnace. The cooling rate was not a certain value, which was larger at higher temperature, but decreased at lower temperature. The average cooling rate was about 30 ◦ C/min. The cooling rate is an important factor for the microstructure evolution of an alloy that is treated by solid solution heattreatment. However, for Ir-base alloys, the solid solution temperature is higher than 2000 ◦ C that is very difficult to reach. Although different cooling rates have some effect on the precipitate microstructure and size for Ir-base alloys, but not as important as that for the alloys with lower melting temperature (such as Ni-base superalloys). Accordingly at the first stage of this investigation, the emphasis was not put on the cooling rate effect, but on the phase constituents. The samples under as-cast and heat-treated stages were polished and electrolytically etched with a solution of 5% HCl in ethyl alcohol for microstructure observation by SEM. One face of the heat-treated samples was polished and characterized by XRD (filtered Cu K␣ radiation) using a RICOH RINT 2500 with a voltage of 40 kV and a current of 300 mA. In order to confirm the phase type, EPMA analysis using a JEOL JAX-8900R was performed. Because the same phase constituents were found in the samples heat-treated at different temperatures from the results of SEM observation and XRD patterns analysis, only the samples heat-treated at 1700 ◦ C were diamondpolished for phase composition analysis by EPMA which was conducted under an accelerating voltage of 20 kV and a current of 5 × 10−8 A. Ir–L␣ and Nb–L␣ were selected for analyzing the content of Ir and Nb.

3. Results 3.1. Microstructure and phase constituents The heat-treated samples appeared similar microstructure evolution at the different heat-treatment temperatures, although the consisted phases became coarse along with temperature increment. The typical microstructures of alloys under as-cast and heat-treated stages (1900 ◦ C for 168 h) are shown in Fig. 1. As the phases in each sample were too fine, the phase composition could not be obtained by point-analysis by EPMA. In

Nb

Third element

this case, map-analysis was conducted instead. The area average composition value obtained by averaging several point values randomly picked in regions A and B is listed in Table 2. 3.1.1. Ir–Nb–Hf alloys A spheroid structure with an interdendritic line was observed in all three alloys under as-cast condition. Within the spheroid structure, different contrast (brighter in the center and blacker at the edge) was observed. This might be due to the heterogeneous element distribution. For alloy Hf1 with the largest Hf content (84.5Ir–4.25Nb–11.25Hf), three regions with clear boundaries were found under the as-cast stage (Fig. 1-1(a)). From the EPMA results, in which the Ir content in region B was higher than that in region A while Nb and Hf showed more concentration in region A, it was believed that region A was an L12 matrix and region B, an fcc matrix. The region C with a eutectic-like fcc/L12 twophase structure filled the area between regions A and B. After it was heat-treated, the fcc phase with a rod-shape precipitated from the matrix in region A, and the L12 phase with a cuboidal shape formed within the fcc matrix in region B. The boundary between each region was still obvious (Fig. 1-1(a )). The morphology was slightly different in alloy Hf2 with a moderate Hf content (84Ir–8.5Nb–7.5Hf) (Fig. 1-1(b and b )). Although three regions existed in the as-cast sample, the primary L12 phase (region A) occupied a larger area; on the other hand, region B, which was the primary fcc phase, was limited to a smaller size than Hf1 alloy. After heat-treatment, the fcc/L12 two-phase structure formed in both regions A and B. The microstructure underwent further changes in alloy Hf3 with the lowest Hf content (83.5Ir–12.75Nb–3.75Hf). Region A increased in size, while region B became further smaller (Fig. 1-1(c and c )). The boundary between each area was less evident. Similarly, after it was heat-treated, the fcc/L12 two-phase structure formed in regions A and B. One obvious phenomenon observed from EPMA map-analysis was that Hf showed an extremely high content around the edge of the dendritic arm, especially in Hf1, where the Hf content was the highest among the three alloys. This implied that the diffusion of Hf within Ir–Nb was difficult.

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Fig. 1. (1) Microstructures of Ir–Nb–Hf alloys under as-cast stage (a–c), and heat-treated at 1900 ◦ C for 168 h (a–c ). (a and a ) Hf1 (84.5Ir–4.25Nb–11.25Hf); (b and b ) Hf2 (84Ir–8.5Nb–7.5Hf); (c and c ) Hf3 (83.5Ir–12.75Nb–3.75Hf). (2) Microstructures of Ir–Nb–Ta alloys under as-cast stage (a–c), and heat-treated at 1900 ◦ C for 168 h (a , a , b , c , c ). (a, a , and a ) Tal (82.25Ir–4.25Nb–13.5Ta); (b and b ) Ta2 (82.5Ir–8.5Nb–9Ta); (c, c , and c ) Ta3 (82.75Ir–12.75Nb–4.5Ta). (a and c ) are enlarged pictures of precipitates in (a and c ). (3) Microstructures of Ir–Nb–Ti alloys under as-cast stage (a–c), and heat-treated at 1900 ◦ C for 168 h (a –c ). (a and a ) Ti1 (84.5Ir–4.25Nb–11.25Ti); (b and b ) Ti2 (84Ir–8.5Nb–7.5Ti); (c and c ) Ti3 (83.5Ir–12.75Nb–3.75Ti).

3.1.2. Ir–Nb–Ta alloys Under as-cast stage, the two-region structure was clearly observed in all the three alloys. In Table 2, the Ir content in region B was higher than that in region A. Nb was slightly concentrated in region A, and Ta was rich in region A. This implied that the matrix phase in region A was L12 and the matrix phase in region B was fcc. The difference among the three Ta alloys was in the size of the two regions in each one. The B area became smaller as the Ta content within the alloys decreased from Ta1, Ta2, to Ta3. Region A, on the other hand, increased as the Ta content decreased (Fig. 1-2(a–c)). The evolution in which the separation of each region became less obvious along with the decrement of the third element content in the Ta series, was similar to that of the Hf series. After heat treatment, the fcc phase formed within the Ll2 matrix in region A, and the Ll2 phase precipitated from the fcc phase in region B (Fig. 1–2(a –c )).

However, the microstructures of the fcc phase formed from the Ll2 matrix were different among these alloys. Ta1 and Ta2 presented strip-like fcc precipitates; however, Ta3 showed cuboidal precipitates, as shown in Fig. 1-2(a and c ). 3.1.3. Ir–Nb–Ti alloys Although two regions existed, region B was very limited (Fig. 1-3(a–c)). Different from the Ir–Nb–Hf and Ir–Nb–Ta alloys, the Ir–Nb–Ti alloy did not show any obvious change in the size of the region as the Ti content changed under ascast stage. All three alloys showed a homogeneous distribution of Ir and Nb (Fig. 2 shows the map-analysis results of Ti2). However, the content of Ti was slightly higher on the grain boundary region (region B). Region A was thought to be a primary fcc phase, and region B was believed to be an Ll2 phase. After they were heat-treated, a fine cuboidal Ll2 phase

C. Huang et al. / Journal of Alloys and Compounds 428 (2007) 220–229

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Fig. 1. (Continued ) .

precipitated homogeneously from the fcc matrix in all three alloys. 3.2. XRD patterns The cross section (about 3 mm in diameter) of the samples was polished and analyzed by XRD. All the samples heat-treated

at different temperatures presented the similar peaks on XRD profiles and these peaks were identified as from the fcc and L12 phases. The representative XRD patterns of the alloys heattreated at 1700 ◦ C are shown in Fig. 3. The peaks of fcc and L12 phases from the same Miller indices could be clearly separated in all three Ir–Nb–Hf alloys. However, the separation was not as evident in the Ir–Nb–Ta alloys as it

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Fig. 1. (Continued ) .

was in the Ir–Nb–Hf alloys. For Ir–Nb–Ti alloys, the peaks of fcc and L12 completely overlapped at lower diffraction angles. This indicated that the lattice misfits between the fcc and L12 phases were the largest in the Ir–Nb–Hf alloys and the smallest in the Ir–Nb–Ti alloys. According to the XRD patterns, the lattice parameters of the fcc and L12 phases were measured. Then the lattice misfit (δ) between fcc and L12 phase was calculated according to the following equation: δ = (αf − αL )/αf (where αf and αL are the lattice parameters of the fcc and the L12 phases, respectively). For the Ir–Nb–Ti alloys, larger samples (about 5 mm in diameter on cross section) and higher degree diffraction angles (from 80◦ to 120◦ ) were applied in order to get clearly separated peaks. Table 3 lists the results, and the lattice misfit change as a function of the third element concentration is shown in Fig. 4. The lattice misfit in Ir–Nb–Hf and Ir–Nb–Ta alloys increased as the Hf or Ta concentration increased; however, the lattice misfit decreased with the increment in the Ti concentration. Because the continuous solid solution between each pair-L12

Table 3 Lattice parameters of fcc and L12 phases in each alloy, and lattice misfit between the two phases Sample no.

Lattice parameter (nm)

Lattice misfit, δ = (aL12 − afcc )/afcc (%)

fcc phase

L12 phase

Ir–Nb–Hf Hf1 Hf2 Hf3

0.3862 0.3860 0.3866

0.3923 0.3913 0.3911

1.6 1.4 1.2

Ir–Nb–Ta Ta1 Ta2 Ta3

0.3860 0.3864 0.3869

0.3891 0.3890 0.3892

0.8 0.7 0.6

Ir–Nb–Ti Ti1 Ti2 Ti3

0.3849 0.3851 0.3842

0.3856 0.3861 0.3854

0.2 0.3 0.3

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Fig. 2. Map-analysis result of alloy Ti2 after heat-treated at 1700 ◦ C for 168 h. (a–c) Distribution map of Ir, Nb, and Ti, respectively; (d) back scattered image (BSI).

phase (Ir3 Nb–Ir3 Hf, Ir3 Nb–Ir3 Ta, and Ir3 Nb–Ir3 Ti) could form, the lattice parameter value of the L12 phase in each ternary alloy was between the lattice parameter values of two pure L12 phases. All the lattice parameter values of the L12 phases in Ir–Nb–Hf, Ir–Nb–Ta, and Ir–Nb–Ti alloys were between the lattice parameter of Ir3 Nb (0.3892 nm) and the lattice parameters of Ir3 Hf (0.3933 nm), Ir3 Ta (0.3884 nm), and Ir3 Ti (0.3840 nm) [5], respectively. 3.3. Stability of precipitates at high temperature Although the fcc/L12 two-phase structure existed in all the alloys, along with the increment of heat-treatment temperature, the precipitates presented different stability. The microstructure evolution of precipitates along with the heat-treatment temperature (1700, 1900, and 2000 ◦ C) in the three kinds of samples is representatively exhibited in Fig. 5. The cuboidal precipitates in Ir–Nb–Ta and Ir–Nb–Ti alloys showed stronger microstructure stability than the Ir–Nb–Hf alloys by presenting sharply aligned cubes up to 1900 ◦ C (Fig. 5(b, b , a, and a )). The sizes of the precipitates were about 300–400 nm at 1900 ◦ C. The Ir–Nb–Ti alloy was more stable, as it could maintain a aligned cube up to 2000 ◦ C (Fig. 5(c )). By this temperature, the Ir–Nb–Ta alloys became unstable and a lamella structure formed (Fig. 5(b )).

However, the lamella structure formed in Ir–Nb–Hf alloys at as low as 1700 ◦ C (Fig. 5(a)). 4. Discussions 4.1. Continuous solid solution of the L12 compound in ternary system An fcc/L12 two-phase structure was found in all the alloys in this study. This was obtained from designing the ternary alloys by combining two binary alloys, each of which had an fcc/L12 two-phase structure. In the study by Yamabe-Mitarai and Harada, the fcc/L12 two-phase structure also formed in Ir–Nb–Zr ternary alloys that had been designed in the same way [3]. This provides a sensible method for designing a ternary alloy in which the fcc/L12 two-phase structure is expected. The key to obtain the fcc/L12 two-phase structure in the ternary system is whether the two L12 phases in each binary system could form a continuous solid solution. It is well known that the capability for forming solid solutions in metallic systems is limited by differences in the atom sizes of the constituents. The first attempt was made by HumeRothery et al. [6,7], who introduced the concept of favorable size-factor as bounded by limits of ±15%. Kornilov gave a

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Fig. 4. Lattice misfit change as a function of third element concentration.

Fig. 3. XRD patterns of the samples heat-treated at 1700 ◦ C.

qualitative expression on the size-factor on continuous solid solutions between intermetallic compounds, i.e. the compounds should have close values of the lattice parameters. In an earlier study in our group, the continuous solid solution between the two L12 phases could be obtained when their lattice parameters were close enough, at least less than about 8% in a quaternary system [8]. For example, the two L12 compounds of Ir3 Nb and Pt3 Al could form the continuous solid solution L12 –Ir3 Nb(Pt3 Al) because their lattice parameter difference was 1.18% ((aIr3Nb − aPt3Al )/aPt3Al ). However, the continuous solid solutions could not form between L12 –Ir3 Nb and L12 –Ni3 Al and between L12 –Ir3 Ta and L12 –Ni3 Al, as the size differences were large (8.94 and 8.73% for Ir3 Nb–Ni3 Al and Ir3 Ta–Ni3 Al, respectively). Among the three ternary sys-

tems investigated in this research, all the pair-L12 phases, Ir3 Nb–Ir3 Hf, Ir3 Nb–Ir3 Ta, and Ir3 Nb–Ir3 Ti, have close lattice parameters. The largest difference value existing between Ir3 Nb–Ir3 Ti is 1.26% ((aIr3Nb − aIr3Ti )/aIr3Ti ). Therefore, continuous solid solutions of L12 phases formed in these systems. In conclusion, in order to get an fcc/L12 two-phase structure in an Ir-base ternary system, the obvious method is to combine two binary systems, each binary system should have an fcc/L12 two-phase structure, and the lattice parameters of the two L12 crystals should be close. Moreover, it is known that the lattice parameter changes linearly with the content of the solvent in the case of continuous solid solution in dilute binary and ternary systems. The dependence of the change of the lattice parameter of the L12 phase in Ir–Nb–Hf, Ir–Nb–Ta, and Ir–Nb–Ti ternary alloys on the addition of a third element could not be demonstrated in this study because the phase composition could not be obtained. However, Table 3 shows that lattice parameter values of ternary alloys existed between the lattice parameters of the two pure L12 phases. For example, the lattice parameters of Ir–Nb–Hf alloys, 0.3911–0.3923 nm, are between the lattice parameters of pure Ir3 Nb (0.3892 nm) and pure Ir3 Hf (0.3933 nm). Accordingly, the lattice parameter of the L12 phase of the Ir-base ternary alloys can be roughly estimated when the continuous solid solution of L12 phase could form. 4.2. Microstructure and lattice misfit The lattice misfit has an effect on the morphology of the precipitates in Ni-base superalloys and Ir-base superalloys [9]. Studies on Ir-base binary alloys by Yamabe-Mitarai et al. [10] indicated that the shape of the L12 precipitates strongly depended on the lattice misfit. The precipitate shape was irregular in the Ir–Ti alloy with the smallest lattice misfit (0.1%), cuboidal in the Ir–Nb and Ir–Ta alloys with a moderate lattice misfit (0.4 and 0.3% for Ir–Nb and Ir–Ta, respectively), and plate-like in the Ir–Zr and Ir–Hf alloys with the largest lattice misfit (2.2 and 1.9% for Ir–Zr and Ir–Hf, respectively). The present study of Ir–Nb–X (X = Hf, Ta, or Ti) ternary alloys showed that cuboidal precipitates existed in all the alloys. This

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Fig. 5. Evolution of precipitate morphology with the heat-treated temperature: (a) Hf1 (84.5Ir–4.25Nb–11.25Hf); (b) Ta3 (82.75Ir–12.75Nb–4.5Ta); (c) Ti3 (83.5Ir–12.75Nb–3.75Ti); (a–c) heat-treated at 1700 ◦ C for 168 h; (a –c ) heat-treated at 1900 ◦ C for 168 h; (a –c ): heat-treated at 2000 ◦ C for 168 h.

suggested that the shape of precipitate in Ir–Nb–X (X = Hf, Ta, or Ti) ternary alloys did not only depend on the lattice misfit. This phenomena was also found in Ir–Nb–Zr ternary alloys which presented cuboidal precipitates, however possessed large lattice misfit value of 1.3–1.8% [4]. The lattice misfit also affects the coarsening processing of L12 precipitates in Ni-base superalloys. It was reported that the coarsening rate of the cuboidal L12 –Ni3 Al phase in Ni–Al–Mo alloys decreased as the lattice misfit was reduced [9]. The results obtained in this study also showed the similar effect of the lattice misfit on the coarsening of precipitates: the larger of the lattice misfit value the more profound of precipitate coarsening. Among the three kinds of alloys, Ir–Nb–Hf has the largest lattice misfit value of 1.2–1.6%, Ir–Nb–Ta presents the moderate value 0.6–0.8%, and Ir–Ir–Nb–Ti has the smallest one of 0.2–0.3%. Accordingly, the cuboidal precipitates showed the most stability in Ir–Nb–Ti alloys; however, they were easier to coarsen in Ir–Nb–Ta alloys, and presented the serious coarsening in Ir–Nb–Hf alloys. This difference of the structure stability of the three kinds of ternary alloys was believed to be mainly due to the difference of their lattice misfit values. This was because

larger lattice misfit would induce higher interfacial energy which caused higher driving force of coarsening. 4.3. Microstructure control Under as-cast stage, a heterogeneous microstructure involving two or three regions each with a different morphology was detected in Ir–Nb–Hf and Ir–Nb–Ta alloys. However, the regional separation phenomenon was not observed within the Ir–Nb–Ti alloys, and a homogeneous microstructure was found. The reason for the difference was attributed to the solidifying processes of the alloys, which can be briefly interpreted according to their binary phase diagrams [11]. As shown in Fig. 6, the common point among Ir–17%Nb, Ir–15%Hf, and Ir–18%Ta is that all these compositions are in fcc and L12 two-phase regions and that the primary phase formed from the liquid is L12 . This caused the formation of the first region (region A in all Ir–Nb–Hf and Ir–Nb–Ta alloys). During the further cooling, the second region with an fcc phase formed in Ir–Nb–Hf and Ir–Nb–Ta alloys (region B in Ir–Nb–Hf and Ir–Nb–Ta alloys), and the fcc/L12 eutectic structure formed by eutectic reaction

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Fig. 6. Partial binary phase diagrams of: (a) Ir–Nb; (b) Ir–Hf; (c) Ir–Ta; (d) Ir–Ti.

in Ir–Nb–Hf alloys (region C in Ir–Nb–Hf alloys). In contrast, the composition of Ir–15%Ti lies in the range where fcc is the primary phase formed from liquid. From the microstructure observation, the primary phase formed in all the Ir–Nb–Ti alloys investigated in this study was the fcc phase (region A in Ir–Nb–Ti alloys), and the second region was very limited. This indicated that the combining Ir–Nb with Ir–Ti caused the ternary alloy composition to be almost in a single fcc phase region, which induced a homogeneous microstructure. 4.4. The potential of Ir-base ternary alloys for the development of novel high-temperature materials The coherent fcc/L12 two-phase structure with cuboidal L12 precipitates formed after solid solution treatment is believed to be the main reason for the excellent creep-resistance of the Nibase superalloy. In this study, Ir–Nb–Ta and Ir–Nb–Ti alloys presented a very analogous microstructure to Ni-base superalloys at a much higher temperature than that Ni-base superalloys could be used. Ir–Nb–Ta and Ir–Nb–Ti ternary systems show great potential for the development of novel materials that could be served at a higher temperature than Ni-base superalloys. In a comparison of these two systems of Ir–Nb–Ta and Ir–Nb–Ti, Ir–Nb–Ti alloys exhibited some superior characteristics. First, Ir–Nb–Ti alloys showed a higher temperature capability than the Ir–Nb–Ta alloys. The L12 precipitates within Ir–Nb–Ti alloys maintained the cuboidal shape to higher temperature. The structure stability is beneficial to high-temperature creep property. Second, the microstructure was more homogeneous in the Ir–Nb–Ti alloys than that in the Ir–Nb–Ta alloys.

Ir–Nb–Ti appears to be a more promising system for ultra hightemperature alloy development. 5. Conclusions The microstructure evolution of the alloys in Ir–Nb–Hf, Ir–Nb–Ta, and Ir–Nb–Ti ternary systems was investigated. The fcc/L12 two-phase structure was detected from the Ir–Nb side to the Ir–Hf (or Ir–Ta or Ir–Ti) side. The lattice misfits between the fcc and Ll2 phases were calculated to be 1.2–1.6, 0.6–0.8, and 0.2–0.3% for Ir–Nb–Hf, Ir–Nb–Ta, and Ir–Nb–Ti alloys, respectively. The stability of precipitates in these alloys decreased with increasing of lattice misfits. The phase constituents, precipitate morphology, and microstructure stability were discussed in terms of the lattice parameter difference and lattice misfit. Ir–Nb–Ti with the smallest lattice misfit value is found to be a promising system where fcc/L12 two-phase structure with cuboidal precipitates could remained up to 2000 ◦ C. Acknowledgements The authors would like to thank Mr. E. Bannai for preparing the ingots and Mr. K. Nishida for EPMA analysis. References [1] Y. Yamabe, Y. Koizumi, H. Murakami, Y. Ro, T. Maruko, H. Harada, Scr. Mater. 35 (1996) 211–215. [2] Y. Yamabe-Mitarai, Y. Ro, T. Maruko, H. Harada, Metall. Trans. A 29A (1998) 537–549. [3] Y. Yamabe-Mitarai, H. Harada, J. Alloys Compd. 361 (2003) 169–179.

C. Huang et al. / Journal of Alloys and Compounds 428 (2007) 220–229 [4] Y. Yamabe-Mitarai, Y. Gu, H. Harada, Metall. Trans. A 34A (2003) 2207–2215. [5] Powder Diffraction File, Inorganic Phases, International Center for Diffraction Data, 1991. [6] W. Hume-Rothery, Acta Metall. 14 (1966) 17–20. [7] W. Hume-Rothery, R.E. Smallman, C.W. Haworth, The Structure of Metals and Alloys, The Institute of Metals, London, 1969, p. 125.

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[8] C. Huang, Y. Yamabe-Mitarai, X.H. Yu, H. Harada, Intermetallics 12 (2004) 619–623. [9] J.G. Conley, M.E. Fine, J.R. Weertman, Acta Metall. 37 (1989) 1251–1263. [10] Y. Yamabe-Mitarai, Y. Ro, T. Maruko, H. Harada, Intermetallics 7 (1999) 49–58. [11] T.B. Massalski (Ed.), Binary Alloy Phase Diagrams, second ed., ASM, 1992.