Effects of Carbon Content and Thermo-Mechanical Treatment on Fe 59 Mn 30 Si 6 Cr 5 C X (X ¼ 0:015{0:1 mass%) Shape Memory Alloys

Transcription

1 Materials Transactions, Vol. 49, No. 8 (2008) pp to 1857 #2008 The Japan Institute of Metals Effects of Carbon Content and Thermo-Mechanical Treatment on Fe 59 Mn 30 Si 6 Cr 5 C X (X ¼ 0:015{0:1 mass%) Shape Memory Alloys C. H. Yang 1, H. C. Lin 2; *, K. M. Lin 1 and W. H. Ho 3 1 Department of Materials Science and Engineering, Feng Chia University, Taichung, Taiwan, R.O. China 2 Department of Materials Science and Engineering, National Taiwan University, Taipei, Taiwan, R.O. China 3 Department of Product Development, Taiwan Textile Research Institute, Taipei, Taiwan, R.O. China Slight amounts of carbon (0:015{0:1 mass%) were added into the Fe-30Mn-6Si-5Cr shape memory alloy. The microstructure, precipitation behaviour and shape memory performance of these Fe 59 Mn 30 Si 6 Cr 5 C X alloys with various thermo-mechanical treatments were investigated. Experimental results show that two particles, the -phase and M 23 C 6 carbide, are observed within the grain-boundary phases. M 23 C 6 carbide only appears in an alloy with high carbon content, but the -phase can form in all alloys with low and high carbon contents. The deformationinduced defects in the 10% pre-deformed specimens are preferential sites for nucleation of precipitates (the -phase and M 23 C 6 carbide) in the following aging treatment. After 550{800 C aging, the rearranged uniform defects and homogeneous precipitates within the grain interior will strengthen the matrix and increase the shape recovery ability. After 850{900 C aging, both the uniform defects and homogeneous precipitates disappear due to recrystallization, and hence the specimen s shape recovery ratios are significantly reduced. [doi: /matertrans.mra ] (Received April 7, 2008; Accepted May 20, 2008; Published July 2, 2008) Keywords: iron-manganese-silicon-chromium shape memory alloy, thermo-mechanical treatment, precipitation hardening, shape memory performance 1. Introduction In light of their low cost and excellent workability, the Fe-Mn-Si-based shape memory alloys have attracted much attention recently. Their SME arises from the reverse transformation of stress-induced " martensite (HCP structure) into parent austenite (FCC structure) upon heating. 1) In the past decades, extensive studies of the Fe-Mn-Sibased alloys were made focusing on the transformation behaviours, 2 4) physical properties, 3 5) effects of thermomechanical training 6 10) and composition dependence of SME and corrosion resistance ) Also, effort is made to increase the use of these alloys, especially the heat-toshrink pipe coupling. 20) It is well known that the Cr addition in Fe-Mn-Si alloys will increase their shape recovery ability 21) and corrosion resistance. 12) In a previous paper, 22) the grain-boundary phase in the Fe-30Mn-6Si-5Cr alloy would occur during annealing at 600{750 C, but was not observed at annealing temperatures exceeding 800 C or lower than 550 C. This grain-boundary phase has been analyzed to have an ordered BCC structure with a lattice parameter of about nm. But, it is still unknown that if carbon atoms have contribution to the formation of this grain-boundary phase or not. Hence, by adding various amounts of carbon into the Fe-30Mn-6Si-5Cr alloy in the present study, we aim at investigating systematically the effects of carbon content on the formation of grain-boundary phase. Meanwhile, the thermo-mechanical treatments, involving the pre-deformation and subsequent aging at various temperatures, are also carried out for these Fe 59 Mn 30 Si 6 Cr 5 C X alloys. Their effects on the formation of grain-boundary phase (or other aging precipitates) and shape memory performance of these alloys will also be discussed. *Corresponding author, hclinntu@ntu.edu.tw 2. Experimental A vacuum melting technique was employed to prepare the Fe 59 Mn 30 Si 6 Cr 5 C X alloys, including the Fe 59 Mn 30 - Si 6 Cr 5 C 0:015,Fe 59 Mn 30 Si 6 Cr 5 C 0:06 and Fe 59 Mn 30 Si 6 Cr 5 C 0:1 (mass%). The as-cast ingots were hot-rolled into plates with 2 mm-thickness, fully annealed at 1150 C for 1 hour and water quenched. Some of these as-quenched plates were then directly aged at 550{900 C for 1 hour and the others were subjected to cold-rolling of 10% prior to the subsequent aging treatment at 550{900 C for 1 hour. The Fe 59 Mn 30 - Si 6 Cr 5 C X alloys were analyzed by using a JEOL JXA-8500F EPMA. Their measured chemical compositions, being presented in Table 1, are similar to those nominal ones. The alloys transformation temperatures, being measured by using a DuPont 2000 DSC thermal analyzer, are presented in Table 2. As shown in Table 2, the alloys with higher carbon contents will have much lower transformation temperatures, especially for the Ms and Mf. The specimen s microstructure was observed by using a HITACHI S3000 scanning electron microscope (SEM) and the element colour mapping was carried out by using the EPMA. The chemical compositions of aged particles were also determined by EPMA analysis. For each kind of particle, the chemical compositions were averaged from at least 10 test readings. Specimens for transmission electron micro- Table 1 The EPMA measured chemical compositions of Fe 59 Mn 30 - Si 6 Cr 5 C X alloys. No. Nominal compositions Fe Mn Si Cr C A Fe 59 Mn 30 Si 6 Cr 5 C 0: B Fe 59 Mn 30 Si 6 Cr 5 C 0:06 balance C Fe 59 Mn 30 Si 6 Cr 5 C 0: C content was measured by using an Optical Emission Spectrometer.

2 1854 C. H. Yang, H. C. Lin, K. M. Lin and W. H. Ho Table 2 The DSC measured transformation temperatures of Fe 59 Mn 30 - Si 6 Cr 5 C X alloys. No. Alloys Ms ( C) Mf ( C) As ( C) Af ( C) A Fe 59 Mn 30 Si 6 Cr 5 C 0: : B Fe 59 Mn 30 Si 6 Cr 5 C 0:06 44:8 69: C Fe 59 Mn 30 Si 6 Cr 5 C 0:1 < 100 These transformation temperatures can not be exactly measured in this study. scope (TEM) were prepared by jet electro-polishing at 25 C with an electrolyte consisting of 3% HClO 4 and 97% C 2 H 5 OH by volume. TEM observation was carried out by using a JEOL JEM-2010 microscope at an operating voltage of 200 kv. The shape memory performance was examined by a bending test with a surface strain of 4% and recovery heating temperature of 500 C. The testing system and calculation of shape recovery have been presented in a previous paper. 23) The tensile tests were carried out at room temperature on a multi-functional MTS (Model: CY- 6040A4, made in Taiwan) tester. 3. Results and Discussion 3.1 Analysis of grain-boundary phase for Fe 59 Mn 30 - Si 6 Cr 5 C X alloys Figures 1 1(e) show the SEM micrographs of Fe 59 - Mn 30 Si 6 Cr 5 C 0:015 specimens quenched from 1150 C and then aged at 550{900 C for 1 hour, respectively. It can be seen that there is no appearance of grain-boundary phase for specimens aged at both 550 C and 900 C, as shown in Figs. 1 and 1(e). Namely, the grain-boundary phase will mainly occur in the temperature range of 650{850 C, as (e) (d) Fig. 1 SEM micrographs of Fe 59 Mn 30 Si 6 Cr 5 C 0:015 specimens quenched from 1150 C and then aged at 550{900 C for 1 hour, 550 C, 650 C, 750 C, (d) 850 C, (e) 900 C. particle 2 particle 1 Fig. 2 SEM micrographs of Fe 59 Mn 30 Si 6 Cr 5 C 0:015, Fe 59 Mn 30 - Si 6 Cr 5 C 0:06, Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimens quenched from 1150 C and then aged at 750 C for 1 hour. shown in Figs. 1, 1, and 1(d). Figures 2 2 show the SEM micrographs of Fe 59 Mn 30 Si 6 Cr 5 C 0:015,Fe 59 Mn 30 - Si 6 Cr 5 C 0:06 and Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimens, respectively, quenched from 1150 C and then aged at 750 C for 1 hour. These SEM micrographs show that the quantity of grain-boundary phase increases with increasing carbon content within the alloys. Meanwhile, two kinds of particles (particles 1 and 2) within the grain-boundary phase in Fig. 2 can be clearly observed for the Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimen, but only one kind of particle (particle 1) is obvious within the grain-boundary phases in Figs. 2 and 2 for the Fe 59 Mn 30 Si 6 Cr 5 C 0:015 and Fe 59 Mn 30 Si 6 Cr 5 C 0:06 specimens, respectively. This phenomenon demonstrates that the carbon content within the alloys will also influence the formation of different grain-boundary phases. To enlarge the formation of these grain-boundary phases for easier analysis, these alloys were 10% pre-cold-rolled and then aged at 750 C for 24 hours. Figures 3 and 4 present the results of EPMA element colour mapping for the as-treated Fe 59 Mn 30 Si 6 Cr 5 C 0:015 and Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimens, respectively. The results of EPMA element colour mapping for Fe 59 Mn 30 Si 6 Cr 5 C 0:06, being similar to those of Fe 59 - Mn 30 Si 6 Cr 5 C 0:015, hence are omitted. In Figs. 3 and 4, it can be seen that the particles not only occur at the grain boundaries, but also within the grain interior. This feature is ascribed to the high density of deformation defects, such as dislocations, jogs, stacking faults and twins, introduced by the 10% pre-deformation prior to the subsequent aging. These particles can nucleate on both preferential sites of deformation defects and grain boundaries, and then grow up during the long-time aging. Carefully examining the element colour mapping in these two figures, only one kind of particle (particle 1) occurs for the Fe 59 Mn 30 Si 6 Cr 5 C 0:015 specimen (Fig. 3), but two kinds of particles, as-distinguished by big and small particles (corresponding to particles 1 and 2), are observed for the Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimen (Fig. 4). The chemical compositions of these particles are measured and the results are presented in Table 3. Based on these chemical compositions, the particles 1 and 2 are determined to be -phase and M 23 C 6 carbide, respectively. Their crystal

3 Effects of Carbon Content and Thermo-Mechanical Treatment on Fe 59 Mn 30 Si 6 Cr 5 C X (X ¼ 0:015{0:1 mass%) Shape Memory Alloys 1855 Table 3 The EPMA measured chemical compositions of particle 1 (-phase) and particle 2 (M 23 C 6 carbide). phases particle 1 particle 2 elements (-phase) (M 23 C 6 ) Fe Mn Si Cr C Zone axis = [011] BCC Fig. 3 EPMA element colour mapping for Fe 59 Mn 30 Si 6 Cr 5 C 0:015 specimen with 10% pre-cold-rolling and subsequent aging at 750 C for 24 hours Zone axis = [011] FCC Fig. 5 TEM analysis of Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimen with 10% pre-coldrolling and subsequent aging at 750 C for 24 hours. Bright field image, SADP of particle 1, SADP of particle 2. Fig. 4 EPMA element colour mapping for Fe 59 Mn 30 Si 6 Cr 5 C 0:1 specimen with 10% pre-cold-rolling and subsequent aging at 750 C for 24 hours. structures will be more discussed using the TEM analysis in the next paragraph. Figures 5 5 shows the TEM analysis of Fe 59 Mn 30 - Si 6 Cr 5 C 0:1 specimen with 10% pre-cold-rolling and subsequent aging at 750 C for 24 hours. As shown in Fig. 5, particles 1 and 2 can be observed simultaneously in the TEM bright field image. The chemical compositions of particles 1 and 2, being measured by using the TEM-EDS, are similar to those presented in Table 2. Based on the analysis of SADP shown in Fig. 5, particle 1 has a BCC structure with lattice constant of 0.88 nm. This structure exhibits that particle 1 is the as-called -phase. 24) The particle 2, being analyzed from the SADP shown in Fig. 5, has a FCC structure with lattice constant of 1.05 nm. This indicates that particle 2 is M 23 C 6 carbide. The single -phase has also been observed for the Fe-Mn-Si-Cr 22) and Fe-14Mn-6Si- 10Cr-7Ni-0.026C 24) alloys. The single M 23 C 6 carbide has

4 1856 C. H. Yang, H. C. Lin, K. M. Lin and W. H. Ho 800 Stress, σ / MPa (d) Strain, ε / % Fig. 6 The stress-strain curves of Fe 59 Mn 30 Si 6 Cr 5 C 0:015 specimens. Curve 1: as-quenched from 1150 C, curve 2: quenched from 1150 C and then aged at 750 C for 1 hour, curve 3: 10% pre-cold-rolled and then aged at 750 C for 1 hour. Fig. 7 SEM micrographs of Fe 59 Mn 30 Si 6 Cr 5 C 0:015 specimens 10% precold-rolled and then aged at 700 C, 750 C, 800 C and (d) 850 C for 1 hour. also been observed for the Fe-15Mn-5Si-8Cr-4Ni-0.18C and Fe-13.53Mn-4.86Si-8.16Cr-3.82Ni-0.16C alloys ) However, the results shown in Figs. 2, 4, and 5 are the first time to simultaneously observe both -phase and M 23 C 6 carbide in one alloy. Meanwhile, it is also understood in the present study that the M 23 C 6 carbide only appears in an alloy with high carbon content, but the -phase can form in all alloys with low and high carbon contents. 3.2 Effects of pre-cold-rolling on Fe 59 Mn 30 Si 6 Cr 5 C X alloys Figure 6 shows the stress-strain curves of Fe 59 Mn 30 Si 6 - Cr 5 C 0:015 specimens with different thermo-mechanical treatments (TMT). Curves 1, 2 and 3 exhibit the results for specimens with the TMT conditions of as-quenched from 1150 C, quenched from 1150 C and then aged at 750 C for 1 hour, and 10% pre-cold-rolled and then aged at 750 C for 1 hour, respectively. The yield strengths presented in these stress-strain curves have the sequence as curve 3 > curve 2 curve 1. This results demonstrate that the strengthening effect occurs for the specimen with 10% precold-rolling and subsequent aging at 750 C, but not for the as-quenched or directly aged specimens. This feature, being ascribed to both the deformation-induced defects and the homogeneously distributed precipitates, will be more discussed in next paragraph. Figures 7 7(d) shows the SEM micrographs of Fe 59 - Mn 30 Si 6 Cr 5 C 0:015 specimens with 10% pre-cold-rolling and subsequent aging at C, respectively. In Fig. 2 for the directly aged specimens, all precipitates occur along the grain boundaries. However, as shown in Figs. 7 and 7 for specimens with 10% pre-cold-rolling and subsequent aging at 700 and 750 C, the precipitates occur homogeneously within the grain interior as well as the grain boundary. This indicates that the deformation-induced defects in the 10% pre-cold-rolled specimens, such as dislocations, jogs, stacking faults and twins, are preferential sites for nucleation of the precipitates if aged at 700 and 750 C. In Fig. 7 for specimen aged at 800 C for 1 hour, it can be seen that the recrystallization begins to occur in addition to the formation of precipitates. This feature is consistent with that reported by Matsumura et al. 28) They found that the recrystallization temperature is in between C for the Fe-28Mn-6Si-5Cr alloy. If the aging temperature is increased to 850 C, the recrystallized grains grow up and a less quantity of precipitates occur mainly along the grain boundaries. Figure 8 shows the shape recovery ratios as a function of aging temperatures for the Fe 59 Mn 30 Si 6 Cr 5 C X alloys with 10% pre-cold-rolling prior to aging. In Fig. 8, one can obviously find that the as-cold-rolled specimens have a worse shape recovery for all alloys. This feature is reasonable because the deformation-induced defects with high internal stress will reside non-uniformly within the grains and inhibit the subsequent " $ transformation during the SME test. After C aging, these Fe 59 Mn 30 Si 6 Cr 5 C X alloys can exhibit quite high shape recovery ratios. The maximum shape recovery ratios can be higher than 90% for the C aged Fe 59 Mn 30 Si 6 Cr 5 C 0:015 specimens. Whilst, the shape recovery ratios of the C aged specimens are significantly reduced. These results can be explained as below. After the C aging, these nonuniform defects are partially annealed out or rearranged into a uniform state within the grains. These uniform defects with a small internal stress are expected to be preferential sites for the formation of stress-induced " martensite during the SME test. Besides, the large quantity of homogeneous precipitates within the grain interior will strengthen the matrix and increase the shape recovery ability. But after the C aging, the recrystallization occurs and the uniform defects are annealed out. Meanwhile, a less quantity of precipitates form only at the recrystallized grain boundaries and can not have obvious strengthening effect on the matrix. The disappearance of both uniform defects and homogeneous precipitates within the grain interior will significantly reduce the shape recovery of C aged specimens. In Fig. 8, it can also be seen that the shape recovery ratios are different for Fe 59 Mn 30 Si 6 Cr 5 C X alloys with different carbon contents. Namely, the specimen with higher carbon content will exhibit lower shape recovery ratios. This feature may be

5 Effects of Carbon Content and Thermo-Mechanical Treatment on Fe 59 Mn 30 Si 6 Cr 5 C X (X ¼ 0:015{0:1 mass%) Shape Memory Alloys 1857 Shape recovery ratio (%) ascribed to that the specimens with higher carbon contents will have lower Ms temperatures, as presented in Table 2. If the Ms temperature is far below room temperature, less amount of strain-induced " martensite can be formed during the SME test. Hence, their shape recovery ratios are decreased. 4. Conclusions alloy A alloy C as-quenched as-rolled alloy B Temperature, T / o C Fig. 8 The shape recovery ratios as a function of aging temperatures for the Fe 59 Mn 30 Si 6 Cr 5 C X alloys with 10% pre-cold-rolling prior to aging. The Fe 59 Mn 30 Si 6 Cr 5 C X alloys will form the grain-boundary phase after aging at C. The quantity of grainboundary phase increases with increasing carbon content within the alloys. Two particles, the -phase and M 23 C 6 carbide, are observed within the grain-boundary phases. M 23 C 6 carbide only appears in an alloy with high carbon content, but the -phase can form in all alloys with low and high carbon contents. The -phase is identified to exhibit a BCC structure with lattice constant of 0.88 nm, and M 23 C 6 carbide has an FCC structure with lattice constant of 1.05 nm. The as-rolled specimens have a worse shape recovery because the deformation-induced defects will inhibit the " $ transformation during SME test. After C aging for specimens with 10% pre-deformation, the rearranged uniform defects and homogeneous precipitates within the grain interior will strengthen the matrix and increase the shape recovery ability. After C aging, the recrystallization occurs and a less quantity of precipitates form mainly at the grain boundaries. This will make both the uniform defects and homogeneous precipitates disappear, and hence the shape recovery of C aged specimens is significantly reduced. The shape recovery ratios are different for Fe 59 Mn 30 Si 6 Cr 5 C X alloys with different carbon contents. Specimens with high carbon content will exhibit low shape recovery ratios in this study. Acknowledgement The authors are pleased to acknowledge the financial support of this research by the National Science Council (NSC), Republic of China, under Grant No. NSC E and NSC E REFERENCES 1) A. Sato, E. Chishima, K. Soma and T. Mori: Acta Metall. 30 (1982) ) M. Murakami, H. Otsuka, H. Suzuki and S. Matsuda: Trans. ISIJ 27 (1987) B-88. 3) H. C. Lin and K. M. Lin: Scripta Metall. Mater. 34 (1996) ) H. C. Lin, K. M. Lin and T. S. Chou: Scripta Metall. Mater. 35 (1996) ) A. Sato, Y. Yamaji and T. Mori: Acta Metall. 34 (1986) ) K. Tsuzaki, M. Ikegami, Y. Tomota and T. Maki: ISIJ 30 (1990) ) B. H. Jiang, T. Tadaki, H. Mori and T. Y. Hsu: Mat. Trans., JIM 38 (1997) ) B. H. Jiang, X. A. Qi, W. M. Zhou, Z. L. Xi and T. Y. Hsu: Scripta Metall. Mater. 34 (1996) ) R. D. Xia, G. W. Liu and T. Liu: Mater. Lett. 32 (1997) ) H. C. Lin, K. M. Lin S. K. Wu, T. P. Wang and Y. C. Hsia: Mater. Sci. Eng. A (2006) ) X. X. Wang and L. C. Zhao: Scripta Metall. Mater. 26 (1992) ) H. C. Lin, K. M. Lin, C. S. Lin and T. M. Ouyang: Corrosion Science 44 (2002) ) T. Maki and I. Tamura: Proc. ICOMAT-86 (1986) ) M. Murakami, H. Otsuka, G. Suzuki and M. Masuda: Proc. ICOMAT- 86 (1986) ) A. Sato, K. Takagaki, S. Horie, M. Kato and T. Mori: Proc. ICOMAT- 86 (1986) ) O. Soderberg, X. W. Liu, P. G. Yakovenko, K. Ullakko and V. K. Lindroos: Mater. Sci. Eng. A (1999) ) V. V. Bliznuk, V. G. Gavriljuk, B. D. Shanina, A. A. Konchits and S. P. Kolesnik: Acta Materialia 51 (2003) ) J. C. Li, W. Zheng and Q. Jiang: Mater. Lett. 38 (1999) ) H. Kubo, K. Nakamura, S. Farjami and T. Maruyama: Mater. Sci. Eng. A 378 (2004) ) H. Tanahashi, T. Maruyama and H. Kubo: Trans. Mat. Soc. Japan 18B (1994) ) L. J. Rong, D. H. Ping, Y. Y. Li and C. X. Shi: Scripta Metallurgica et Materialia 132 (1995) ) H. C. Lin, C. S. Lin, K. M. Lin and Y. C. Chuang: Journal of Alloys and Compounds 319 (2001) ) H. C. Lin and S. K. Wu: Scripta Metall. Mater. 26 (1992) ) N. Stanford, D. P. Dunne and B. J. Monaghan: Journal of Alloys and Compounds 430 (2007) ) Y. H. Wen, M. Yan and N. Li: Materials Letters 58 (2004) ) Y. H. Wen, M. Yan and N. Li: Scripta Materialia 50 (2004) ) Y. H. Wen, W. Zhang, N. Li, H. B. Peng and L. R. Xiong: Acta Materialia 55 (2007) ) O. Matsumura, S. Furusako, T. Sumi, T. Furukawa and H. Otsuka: Materials Science and Engineering A 272 (1999) 459.