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A comparative study of the high temperature deformation behavior of Fe-25Al and Fe-25Al-10Ti alloys
Scripta mater. 42 (2000) 905–910 www.elsevier.com/locate/scriptamat
A COMPARATIVE STUDY OF THE HIGH TEMPERATURE DEFORMATION BEHAVIOR OF Fe-25Al AND Fe-25Al-10Ti ALLOYS Su-Ming Zhu, Kazushi Sakamoto, Makoto Tamura and Kunihiko Iwasaki Japan Ultra-high Temperature Materials Research Institute, 573-3 Okiube, Ube, Yamaguchi 755-0001, Japan (Received December 8, 1999) (Accepted December 22, 1999) Keywords: Iron aluminide; High temperature compression; Strain rate sensitivity; Deformation microstructure
Introduction Iron aluminides based on Fe3Al are of interest as structural intermetallics in view of their low cost, low density, remarkable corrosion resistance and environmental friendliness. However, commercialization of these intermetallics has been very limited due to their poor ductility and impact resistance at ambient temperature as well as inadequate creep strength at elevated temperatures. The best application of the Fe3Al alloy to date is the use in porous hot-gas filters [1]. Recent studies by Hawk et al. [2,3] have demonstrated that the addition of Ti to Fe3Al has a positive influence on their tribological properties. For example, the alloy with 10 at. % Ti substituting for Fe in Fe3Al shows a 40% decrease in volume wear compared with Fe3Al. This inspiring result suggests that iron aluminides may find their application in some tribological circumstances, especially where the oxidation or sulfidation is also a major concern. A systematic study has been conducted by the present authors on the processing and characterization of Ti alloyed Fe3Al [4,5]. In this communication, the compressive deformation behavior of Fe-25Al10Ti (all compositions in at. %) alloy in the temperature range of 873–1273 K is reported and compared with that of Fe-25Al alloy. The fundamental data provided are expected to be useful for the engineering processing of Ti alloyed Fe3Al-based alloys.
Experimental The Fe-25Al and Fe-25Al-10Ti alloys (simply termed as 0Ti and 10Ti alloys hereafter) used in this study were prepared by arc-melting 99.99 wt. % elemental wires into buttons in Ar. The buttons were remelted several times to ensure homogeneity. After melting, the buttons were further homogenized in vacuum for 14.4 ks at 1323 K. Rectangular compression specimens in dimensions of 3 ⫻ 3 ⫻ 6 mm were sectioned from the buttons by electro-discharge machining, with the long axis parallel to the solidification direction. These specimens were then tested over a range of strain rate (4.2 ⫻ 10⫺5–1.4 ⫻ 10⫺3 s⫺1) and temperature (873–1273 K) in a Shimadzu EU-1-3 high temperature materials testing system. All tests were carried out in air with 15 min hold at the testing temperature prior to deformation. 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(00)00312-2
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Figure 1. Optical micrographs of (a) Fe-25Al and (b) Fe-25Al-10Ti alloys prior to deformation. The solidification direction is vertical.
To minimize friction effects, the specimen-die interface was lubricated with silicon carbide. Microstructures of the alloys before and after deformation were examined by optical microscopy. Metallography samples were etched in a solution containing 5% HNO3, 5% HF, 75% glycerol and 15% H2O. Results Figure 1 shows the optical micrographs of the 0Ti and 10Ti alloys prior to deformation. It can be seen that the 0Ti alloy consists mainly of coarse columnar grains in the solidification direction (Fig. 1(a)), whilst the 10Ti alloy is composed of relatively fine and equiaxed grains (Fig. 1(b)). This indicates that the Ti addition decreases the grain size of the Fe-25Al alloy. Similar results have also been reported by Sikka et al. in FeAl-based alloys [6]. Furthermore, the Ti addition leads to the precipitation of second phase particles. The structure of these precipitates is not defined at the present stage. Examination of the deformed specimens reveals that the 0Ti alloys can be deformed without surface cracking over the whole temperature and strain rate range used. However, the 10Ti alloy tends to exhibit surface cracking especially at temperatures below 1073 K. The cross sectional microstructures of the
Figure 2. Optical micrographs of Fe-25Al alloy after deformation at (a) 873 K and (b) 1273 K.
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Figure 3. Optical micrographs of Fe-25Al-10Ti alloy after deformation at (a) 873 K and (b) 1273 K.
0Ti and 10Ti alloys after high temperature deformation are shown in Figs. 2 and 3, respectively. For the 0Ti alloy, formation of subgrains is typical at 873 K (Fig. 2(a)), and refined grains are observed in specimens deformed at temperatures above 1073 K (Fig. 2(b)). For the 10Ti alloy, cavitation is visible in specimens deformed below 1073 K (Fig. 3(a)). Such cavitation occurs preferentially at grain boundaries, especially at the triple junctions of the grain boundaries, indicating that grain boundary sliding plays an important role in the deformation process. Formation of new grains is evidenced in specimens deformed at temperatures above 1173 K (Fig. 3(b)). These results reveal that the 10Ti alloy has a relatively narrower processing window than the 0Ti alloy. The hot working temperature of the 10Ti alloy should be above 1173 K. Figure 4 shows representative true stress-strain curves of the 0Ti and 10Ti alloys in the temperature range of 873–1273 K. It is evident that temperature and Ti addition have significant effects on the flow behavior. In general, the flow stress decreases greatly with increasing temperature for both of the alloys. The 10Ti alloy shows much higher flow stress than the 0Ti one, and this tendency is more pronounced at higher temperatures, indicating that the Ti addition improves the high temperature strength of the Fe-25Al alloy. Moreover, different features are observed in the stress-strain curves of the 0Ti and 10Ti alloys. The 0Ti alloy is characterized by a sharp or no peak stress at low strain, followed by a near steady-state flow behavior. The 10Ti alloy exhibits a marked peak stress at relatively high strain, followed by gradual flow softening until attaining a steady state. With increasing temperature, the peak stress shifts from high strain to low strain.
Figure 4. Representative stress-strain curves of Fe-25Al and Fe-25Al-10Ti alloys (strain rate ⫽ 1.4 ⫻ 10⫺3 s⫺1).
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Figure 5. Strain rate dependence of compressive flow stress of Fe-25Al and Fe-25Al-10Ti alloys.
To analyze the temperature and strain rate dependence of flow stress, the flow stress at true strain of 0.2 is used in this work. Figure 5 shows the strain rate dependence of the compressive flow stress for the 0Ti and 10Ti alloys. From the plots, the strain rate sensitivity (m) values are determined. The results show that the m value generally increases with increasing temperature. For the 0Ti alloy, especially, the m value reaches to 0.33 at 1073 K and 0.32 at 1273 K. Similar m values have been reported by Lin et al. [7] and Chu et al. [8] in Fe3Al based alloys subject to high temperature tensile deformation. In both studies, superplastic behavior have been observed. According to Nieh et al. [9], m ⫽ 0.3 is a critical value for intermetallics to show superplastic behavior. Thus it is expected the 0Ti alloy shows superplastic behavior at temperatures above 1073 K. The 10Ti alloy, however, exhibits much lower m values than the 0Ti alloy, implying that superplastic deformation may not be possible in the present temperature and strain rate ranges. The activation energy for flow deformation, Q, can be determined by the following equation [7,8] Q⫽n䡠R䡠
⭸ln , ⭸共1/T兲
(1)
where n ⫽ 1/m, R is the universal gas constant, is the flow stress and T is the absolute temperature. Based on this equation, ln is plotted against 1/T for the 0Ti and 10Ti alloys around the temperature of 1073 K, as shown in Fig. 6. The calculated flow activation energy is 245 kJ/mol for the 0Ti alloy and 210 kJ/mol for the 10Ti alloy. The values are very close to those reported by Lin et al. [7]. In their work, values of 263 kJ/mol and 191 kJ/mol are obtained for the Fe-28Al and Fe-28Al-2Ti alloys, respectively. It appears that Ti additions tend to decrease the flow activation energy of Fe3Al based alloys. This decrease is speculated to be associated with the grain refinement by Ti additions, as grain boundary plays an important role in the flow deformation of polycrystalline alloys. On the other hand,
Figure 6. Plot of ln against 1/T for Fe-25Al and Fe-25Al-10Ti alloys. The flow activation energy is calculated.
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these activation energies are much lower than the creep activation energy of 300 –350 kJ/mol for Fe3Al-based alloys [10]. This suggests that the flow deformation is not controlled by the lattice diffusion, but by grain boundary related mechanism. Discussion The present results reveal that the 10Ti alloy can only be hot-deformed without surface cracking in a relatively narrower processing range than the 0Ti alloy. Two factors are considered to account for this difference. Firstly, Ti additions have been found to raise the D03 3 B2 transition temperature appreciably, thus enhancing the stability of D03 ordered structure at high temperatures [11,12]. It is well known that D03 structure has poorer deformability than the B2 structure. Secondly, Ti additions tend to decrease the tensile ductility of Fe3Al, though the yield strength is greatly increased, especially at high temperatures [13]. Limited tensile ductility is believed to be the direct cause for the surface cracking of Fe3Al [14]. In the present work, fine-grained microstructure can be acquired after deformation in certain temperature range for both the 0Ti and 10Ti alloy. In the former case, superplastic behavior is expected to occur. Several mechanisms for the grain refinement by superplastic deformation in Fe3Al have been proposed [7–9]. In view of the steady state flow behavior and the absence of grain boundary sliding, the continuous grain boundary migration mechanism proposed by Chu et al. [8] is believed to be dominant for the grain refinement of the present 0Ti alloy. For the 10Ti alloys, the refined grains are considered to result from dynamic recrystallization, as dynamic recrystallization causes not only grain refinement but also flow softening [15]. The flow softening is evidenced for the present 10Ti alloys deformed at temperatures above 1173 K (Fig. 3(b)). Conclusions The compressive deformation behavior of Fe-25Al and Fe-25Al-10Ti was studied in the range of strain rate (4.2 ⫻ 10⫺5–1.4 ⫻ 10⫺3 s⫺1) and temperature (873–1273 K). The results show that Ti addition tend to reduce the strain rate sensitivity and activation energy for flow deformation of Fe-25Al alloy. The Fe-25Al alloy can be deformed without cracking at temperatures above 873 K. However, the Ti-added alloy has a relatively narrow processing window. The hot working temperature of Fe-25Al10Ti alloy should be above 1173 K. Acknowledgements Financial support by the New Energy and Industrial Technology Development Organization (NEDO), Japan, is greatly acknowledged. References 1. 2. 3. 4. 5. 6.
V. K. Sikka, in Proceedings of the International Symposium on Nickel and Iron Aluminides: Processing, Properties, and Applications, ed. S. C. Deevi et al., p. 361, The Materials Information Society (1997). J. A. Hawk, D. E. Alman, and R. D. Wilson, Mater. Res. Soc. Symp. Proc. 364, 243 (1995). J. A. Hawk and D. E. Alman, Mater. Sci. Eng. A239 –240, 899 (1987). S.-M. Zhu and K. Iwasaki, Mater. Sci. Eng. A270, 170 (1999). S.-M. Zhu and K. Iwasaki, Mater. Trans. JIM. 40, 1311 (1999). V. K. Sikka, D. Wilkening, J. Liebetrau, and B. Mackey, Mater. Sci. Eng. A258, 229 (1998).
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7. 8. 9. 10. 11. 12. 13.
14. 15.
HIGH TEMPERATURE DEFORMATION
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D. Lin, A. Shan, and M. Chen, Intermetallics. 4, 489 (1996). J. P. Chu, I. M. Liu, J. H. Wu, W. Kai, J. Y. Wang, and K. Inoue, Mater. Sci. Eng. A258, 236 (1998). T. G. Nieh, J. Wadsworth, and O. D. Sherby, Superplasticity in Metals and Ceramics, p. 125, Cambridge University Press, Cambridge (1997). D. H. Sastry and R. S. Sundar, in Proceedings of the International Symposium on Nickel and Iron Aluminides: Processing, Properties, and Applications, ed. S. C. Deevi et al., p. 123, The Materials Information Society (1997). R. T. Fortnum and D. E. Mikkola, Mater. Sci. Eng. 91, 223 (1987). Y. Nishino, S. Asano, and T. Ogawa, Mater. Sci. Eng. A234 –236, 271 (1997). C. Testani, A. Di Gianfrancesco, O. Tassa, and D. Pocci, in Proceedings of the International Symposium on Nickel and Iron Aluminides: Processing, Properties, and Applications, ed. S. C. Deevi et al., p. 213, The Materials Information Society (1997). B. Voyzelle and J. D. Boyd, Mater. Sci. Eng. A258, 243 (1998). H. Fukutomi, C. Hartig, and H. Mecking, Z. Metallkd. 81, 272 (1990).
A COMPARATIVE STUDY OF THE HIGH TEMPERATURE DEFORMATION BEHAVIOR OF Fe-25Al AND Fe-25Al-10Ti ALLOYS Su-Ming Zhu, Kazushi Sakamoto, Makoto Tamura and Kunihiko Iwasaki Japan Ultra-high Temperature Materials Research Institute, 573-3 Okiube, Ube, Yamaguchi 755-0001, Japan (Received December 8, 1999) (Accepted December 22, 1999) Keywords: Iron aluminide; High temperature compression; Strain rate sensitivity; Deformation microstructure
Introduction Iron aluminides based on Fe3Al are of interest as structural intermetallics in view of their low cost, low density, remarkable corrosion resistance and environmental friendliness. However, commercialization of these intermetallics has been very limited due to their poor ductility and impact resistance at ambient temperature as well as inadequate creep strength at elevated temperatures. The best application of the Fe3Al alloy to date is the use in porous hot-gas filters [1]. Recent studies by Hawk et al. [2,3] have demonstrated that the addition of Ti to Fe3Al has a positive influence on their tribological properties. For example, the alloy with 10 at. % Ti substituting for Fe in Fe3Al shows a 40% decrease in volume wear compared with Fe3Al. This inspiring result suggests that iron aluminides may find their application in some tribological circumstances, especially where the oxidation or sulfidation is also a major concern. A systematic study has been conducted by the present authors on the processing and characterization of Ti alloyed Fe3Al [4,5]. In this communication, the compressive deformation behavior of Fe-25Al10Ti (all compositions in at. %) alloy in the temperature range of 873–1273 K is reported and compared with that of Fe-25Al alloy. The fundamental data provided are expected to be useful for the engineering processing of Ti alloyed Fe3Al-based alloys.
Experimental The Fe-25Al and Fe-25Al-10Ti alloys (simply termed as 0Ti and 10Ti alloys hereafter) used in this study were prepared by arc-melting 99.99 wt. % elemental wires into buttons in Ar. The buttons were remelted several times to ensure homogeneity. After melting, the buttons were further homogenized in vacuum for 14.4 ks at 1323 K. Rectangular compression specimens in dimensions of 3 ⫻ 3 ⫻ 6 mm were sectioned from the buttons by electro-discharge machining, with the long axis parallel to the solidification direction. These specimens were then tested over a range of strain rate (4.2 ⫻ 10⫺5–1.4 ⫻ 10⫺3 s⫺1) and temperature (873–1273 K) in a Shimadzu EU-1-3 high temperature materials testing system. All tests were carried out in air with 15 min hold at the testing temperature prior to deformation. 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(00)00312-2
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Figure 1. Optical micrographs of (a) Fe-25Al and (b) Fe-25Al-10Ti alloys prior to deformation. The solidification direction is vertical.
To minimize friction effects, the specimen-die interface was lubricated with silicon carbide. Microstructures of the alloys before and after deformation were examined by optical microscopy. Metallography samples were etched in a solution containing 5% HNO3, 5% HF, 75% glycerol and 15% H2O. Results Figure 1 shows the optical micrographs of the 0Ti and 10Ti alloys prior to deformation. It can be seen that the 0Ti alloy consists mainly of coarse columnar grains in the solidification direction (Fig. 1(a)), whilst the 10Ti alloy is composed of relatively fine and equiaxed grains (Fig. 1(b)). This indicates that the Ti addition decreases the grain size of the Fe-25Al alloy. Similar results have also been reported by Sikka et al. in FeAl-based alloys [6]. Furthermore, the Ti addition leads to the precipitation of second phase particles. The structure of these precipitates is not defined at the present stage. Examination of the deformed specimens reveals that the 0Ti alloys can be deformed without surface cracking over the whole temperature and strain rate range used. However, the 10Ti alloy tends to exhibit surface cracking especially at temperatures below 1073 K. The cross sectional microstructures of the
Figure 2. Optical micrographs of Fe-25Al alloy after deformation at (a) 873 K and (b) 1273 K.
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Figure 3. Optical micrographs of Fe-25Al-10Ti alloy after deformation at (a) 873 K and (b) 1273 K.
0Ti and 10Ti alloys after high temperature deformation are shown in Figs. 2 and 3, respectively. For the 0Ti alloy, formation of subgrains is typical at 873 K (Fig. 2(a)), and refined grains are observed in specimens deformed at temperatures above 1073 K (Fig. 2(b)). For the 10Ti alloy, cavitation is visible in specimens deformed below 1073 K (Fig. 3(a)). Such cavitation occurs preferentially at grain boundaries, especially at the triple junctions of the grain boundaries, indicating that grain boundary sliding plays an important role in the deformation process. Formation of new grains is evidenced in specimens deformed at temperatures above 1173 K (Fig. 3(b)). These results reveal that the 10Ti alloy has a relatively narrower processing window than the 0Ti alloy. The hot working temperature of the 10Ti alloy should be above 1173 K. Figure 4 shows representative true stress-strain curves of the 0Ti and 10Ti alloys in the temperature range of 873–1273 K. It is evident that temperature and Ti addition have significant effects on the flow behavior. In general, the flow stress decreases greatly with increasing temperature for both of the alloys. The 10Ti alloy shows much higher flow stress than the 0Ti one, and this tendency is more pronounced at higher temperatures, indicating that the Ti addition improves the high temperature strength of the Fe-25Al alloy. Moreover, different features are observed in the stress-strain curves of the 0Ti and 10Ti alloys. The 0Ti alloy is characterized by a sharp or no peak stress at low strain, followed by a near steady-state flow behavior. The 10Ti alloy exhibits a marked peak stress at relatively high strain, followed by gradual flow softening until attaining a steady state. With increasing temperature, the peak stress shifts from high strain to low strain.
Figure 4. Representative stress-strain curves of Fe-25Al and Fe-25Al-10Ti alloys (strain rate ⫽ 1.4 ⫻ 10⫺3 s⫺1).
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Figure 5. Strain rate dependence of compressive flow stress of Fe-25Al and Fe-25Al-10Ti alloys.
To analyze the temperature and strain rate dependence of flow stress, the flow stress at true strain of 0.2 is used in this work. Figure 5 shows the strain rate dependence of the compressive flow stress for the 0Ti and 10Ti alloys. From the plots, the strain rate sensitivity (m) values are determined. The results show that the m value generally increases with increasing temperature. For the 0Ti alloy, especially, the m value reaches to 0.33 at 1073 K and 0.32 at 1273 K. Similar m values have been reported by Lin et al. [7] and Chu et al. [8] in Fe3Al based alloys subject to high temperature tensile deformation. In both studies, superplastic behavior have been observed. According to Nieh et al. [9], m ⫽ 0.3 is a critical value for intermetallics to show superplastic behavior. Thus it is expected the 0Ti alloy shows superplastic behavior at temperatures above 1073 K. The 10Ti alloy, however, exhibits much lower m values than the 0Ti alloy, implying that superplastic deformation may not be possible in the present temperature and strain rate ranges. The activation energy for flow deformation, Q, can be determined by the following equation [7,8] Q⫽n䡠R䡠
⭸ln , ⭸共1/T兲
(1)
where n ⫽ 1/m, R is the universal gas constant, is the flow stress and T is the absolute temperature. Based on this equation, ln is plotted against 1/T for the 0Ti and 10Ti alloys around the temperature of 1073 K, as shown in Fig. 6. The calculated flow activation energy is 245 kJ/mol for the 0Ti alloy and 210 kJ/mol for the 10Ti alloy. The values are very close to those reported by Lin et al. [7]. In their work, values of 263 kJ/mol and 191 kJ/mol are obtained for the Fe-28Al and Fe-28Al-2Ti alloys, respectively. It appears that Ti additions tend to decrease the flow activation energy of Fe3Al based alloys. This decrease is speculated to be associated with the grain refinement by Ti additions, as grain boundary plays an important role in the flow deformation of polycrystalline alloys. On the other hand,
Figure 6. Plot of ln against 1/T for Fe-25Al and Fe-25Al-10Ti alloys. The flow activation energy is calculated.
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these activation energies are much lower than the creep activation energy of 300 –350 kJ/mol for Fe3Al-based alloys [10]. This suggests that the flow deformation is not controlled by the lattice diffusion, but by grain boundary related mechanism. Discussion The present results reveal that the 10Ti alloy can only be hot-deformed without surface cracking in a relatively narrower processing range than the 0Ti alloy. Two factors are considered to account for this difference. Firstly, Ti additions have been found to raise the D03 3 B2 transition temperature appreciably, thus enhancing the stability of D03 ordered structure at high temperatures [11,12]. It is well known that D03 structure has poorer deformability than the B2 structure. Secondly, Ti additions tend to decrease the tensile ductility of Fe3Al, though the yield strength is greatly increased, especially at high temperatures [13]. Limited tensile ductility is believed to be the direct cause for the surface cracking of Fe3Al [14]. In the present work, fine-grained microstructure can be acquired after deformation in certain temperature range for both the 0Ti and 10Ti alloy. In the former case, superplastic behavior is expected to occur. Several mechanisms for the grain refinement by superplastic deformation in Fe3Al have been proposed [7–9]. In view of the steady state flow behavior and the absence of grain boundary sliding, the continuous grain boundary migration mechanism proposed by Chu et al. [8] is believed to be dominant for the grain refinement of the present 0Ti alloy. For the 10Ti alloys, the refined grains are considered to result from dynamic recrystallization, as dynamic recrystallization causes not only grain refinement but also flow softening [15]. The flow softening is evidenced for the present 10Ti alloys deformed at temperatures above 1173 K (Fig. 3(b)). Conclusions The compressive deformation behavior of Fe-25Al and Fe-25Al-10Ti was studied in the range of strain rate (4.2 ⫻ 10⫺5–1.4 ⫻ 10⫺3 s⫺1) and temperature (873–1273 K). The results show that Ti addition tend to reduce the strain rate sensitivity and activation energy for flow deformation of Fe-25Al alloy. The Fe-25Al alloy can be deformed without cracking at temperatures above 873 K. However, the Ti-added alloy has a relatively narrow processing window. The hot working temperature of Fe-25Al10Ti alloy should be above 1173 K. Acknowledgements Financial support by the New Energy and Industrial Technology Development Organization (NEDO), Japan, is greatly acknowledged. References 1. 2. 3. 4. 5. 6.
V. K. Sikka, in Proceedings of the International Symposium on Nickel and Iron Aluminides: Processing, Properties, and Applications, ed. S. C. Deevi et al., p. 361, The Materials Information Society (1997). J. A. Hawk, D. E. Alman, and R. D. Wilson, Mater. Res. Soc. Symp. Proc. 364, 243 (1995). J. A. Hawk and D. E. Alman, Mater. Sci. Eng. A239 –240, 899 (1987). S.-M. Zhu and K. Iwasaki, Mater. Sci. Eng. A270, 170 (1999). S.-M. Zhu and K. Iwasaki, Mater. Trans. JIM. 40, 1311 (1999). V. K. Sikka, D. Wilkening, J. Liebetrau, and B. Mackey, Mater. Sci. Eng. A258, 229 (1998).
910
7. 8. 9. 10. 11. 12. 13.
14. 15.
HIGH TEMPERATURE DEFORMATION
Vol. 42, No. 9
D. Lin, A. Shan, and M. Chen, Intermetallics. 4, 489 (1996). J. P. Chu, I. M. Liu, J. H. Wu, W. Kai, J. Y. Wang, and K. Inoue, Mater. Sci. Eng. A258, 236 (1998). T. G. Nieh, J. Wadsworth, and O. D. Sherby, Superplasticity in Metals and Ceramics, p. 125, Cambridge University Press, Cambridge (1997). D. H. Sastry and R. S. Sundar, in Proceedings of the International Symposium on Nickel and Iron Aluminides: Processing, Properties, and Applications, ed. S. C. Deevi et al., p. 123, The Materials Information Society (1997). R. T. Fortnum and D. E. Mikkola, Mater. Sci. Eng. 91, 223 (1987). Y. Nishino, S. Asano, and T. Ogawa, Mater. Sci. Eng. A234 –236, 271 (1997). C. Testani, A. Di Gianfrancesco, O. Tassa, and D. Pocci, in Proceedings of the International Symposium on Nickel and Iron Aluminides: Processing, Properties, and Applications, ed. S. C. Deevi et al., p. 213, The Materials Information Society (1997). B. Voyzelle and J. D. Boyd, Mater. Sci. Eng. A258, 243 (1998). H. Fukutomi, C. Hartig, and H. Mecking, Z. Metallkd. 81, 272 (1990).