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Thermo-mechanical processing and the shape memory effect in an Fe–Mn–Si-based shape memory alloy
Materials Science and Engineering A 422 (2006) 352–359
Thermo-mechanical processing and the shape memory effect in an Fe–Mn–Si-based shape memory alloy N. Stanford ∗ , D.P. Dunne Faculty of Engineering, University of Wollongong, Northfields Ave., Wollongong 2522, Australia Received 28 October 2005; accepted 14 February 2006
Abstract The effect of cold rolling and annealing on the shape memory effect (SME) in an Fe–Mn–Si-based alloy has been studied. It has been found that the SME in these alloys can be significantly increased by the appropriate thermo-mechanical processing (TMP). The optimum conditions were found to be 15% cold rolling followed by annealing at 800 ◦ C for 15 min. This produced a total strain recovery of 4.5%. TEM showed that this processing schedule produces a microstructure of evenly spaced, and well defined stacking faults throughout the parent phase. It is shown for the first time that samples processed in this way produce a larger fraction of martensite compared to samples in the as-austenitized condition. It is concluded that the stacking faults induced by TMP act as nucleation sites for martensite formation during deformation. The SME is improved primarily as a result of the increased amount of martensite that is formed in this condition. © 2006 Elsevier B.V. All rights reserved. Keywords: Shape memory; Training; Ferrous alloys; Cold rolling; Annealing
1. Introduction The Fe–Mn–Si-based alloy system produces a hexagonal martensite, known as ε, in response to deformation. This martensite is produced by the formation and overlap of stacking faults in the parent phase, and can revert back to austenite by heating. This imparts the alloy with shape memory behaviour [1]. It is known that optimum shape memory is achieved in Fe–Mn–Sibased alloys by using thermo-mechanical processing (TMP) to “train” the alloy [2]. The most common method is to train the specimen by tensile strain followed by a recovery anneal, and then to repeat the cycle for up to about five times. The shape memory is also tested in tension. The effects of repeated training have been studied in some detail. As well as markedly increasing the strain recovery, training is known to significantly decrease the “yield stress” of the alloy [3,4], which is identified as the stress required to form martensite. Training also increases the uniformity of deformation along the sample length, and the relative uniformity continues to increase with repeated training cycles [5]. The density of stacking faults retained within the parent phase is also known
∗
Corresponding author. Tel.: +61 2 4221 3457; fax: +61 2 4221 5474. E-mail address: [email protected] (N. Stanford).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.009
to increase with training cycles [6]. Observations of the sample surface have clearly shown that, at least qualitatively, more ε is formed during deformation after the training regime [5,7]. Atomic force microscopy of the free surface has shown that the morphology of the martensite is affected by training also. The martensite plates become thinner, and are more likely to form with a single variant within each parent grain [8,9]. In recent work [10], we have shown that TMP of the alloy by rolling and annealing can produce a marked increase in the shape memory effect (SME) measured in bending. Rolling and annealing is a much more convenient method of processing than tension, and lends itself to large-scale industrial production. The tedious tensile deformation and annealing of shape memory alloys are not a practical method of producing large products such as, for example, pipe joining collars. In this paper we report results of a systematic investigation of the effectiveness of cold rolling and annealing as a SME enhancing process, with the aim of optimizing the processing conditions.
2. Method The base stainless alloy used in this study was made in a 2.6 kg ingot by induction melting under an argon atmosphere. The composition of the alloy in wt.% was Fe–13Mn–5Si–9Cr–
N. Stanford, D.P. Dunne / Materials Science and Engineering A 422 (2006) 352–359
7Ni–0.03C. The as-cast alloy was hot rolled at 900 ◦ C to a thickness of ∼1 mm. The rolled strip was sectioned and machined into tensile test coupons or bend test strips before being solution treated at 1000 ◦ C for 1 h under flowing argon atmosphere, followed by immediate cold water quench. This is considered the starting condition, and will hereafter be referred to as the as-austenitized condition. As-austenitized samples were deformed at room temperature by cold rolling to strains between 5% and 39%. After the deformation step, samples were recovery annealed at temperatures between 600 and 800 ◦ C. The shape memory of each sample was measured in bending over a range of pre-strains between 1% and 6%. The pre-strain was taken as the maximum tensile/compressive strain, and is determined by the equation: ε=
1 (2R/ h) + 1
(1)
where ε = conventional strain, R = bend radius, h = sample thickness. The shape was recovered by annealing in a muffle furnace at 400 ◦ C for 15 min. After the samples had been recovery annealed the residual strain (εr ) remaining due to incomplete recovery was calculated using Eq. (1). The percentage recovery was determined by the equation: % recovery =
(ε − εr ) ε
(2)
Tensile specimens were deformed using an Instron tensile tester at a speed of 1 mm/min. The starting gage length of tensile specimens was 30 mm, and the starting gage width was 5 mm. Microstructural examination of samples was also carried out by transmission electron microscopy (TEM). The TEM foils were made using a Tenupol jet-polisher with a solution of 5% perchloric acid in acetic acid at 30 V. TEM observations were made using a JEOL JEM 2010 TEM at 200 kV. The transformation behaviour of the samples was investigated using a TA Q100 differential scanning calorimeter (DSC) at a heating rate of 10 ◦ C/min and a cooling rate of 5 ◦ C/min. X-ray diffraction (XRD) was carried out on a Philips PW 1730 XRD at a scan rate of 0.5 ◦ C/min. Prior to XRD analysis, samples were electropolished in a solution of 5% perchloric in acetic acid for 1 min at 30 V to remove the deformed layer induced by mechanical polishing. The crystallographic preferred orientation of the samples was measured using electron backscattered diffraction (EBSD) equipment supplied by HKL Technology. This system is fitted to a Leica 440 scanning electron microscope (SEM). Samples were prepared for EBSD by standard metallographic grinding and polishing techniques followed by a final polish with colloidal silica (Struers OPS) for 10 min. 3. Results 3.1. Orientation of bend testing The effect of the orientation of the test specimen on SME was investigated. A length of the as-austenitized strip was solution
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Fig. 1. Illustration of the four geometries tested in bending with respect to the rolling direction (RD) of the thermo-mechanical treatments. The shaded boxes represent the longitudinal (L), transverse (T), perpendicular longitudinal (PL) and perpendicular transverse (PT) samples. The orientations in which these samples were tested in bending are represented by corresponding grey shading. The direction of tensile stress on the outer sample edge during the bend test is also shown.
treated, and then processed by cold rolling to 15% reduction followed by a recovery anneal at 800 ◦ C for 15 min. After rolling, the strip was sectioned in four different bending geometries; longitudinal and transverse within the rolling plane, longitudinal and transverse perpendicular to the rolling plane (Fig. 1). Samples for bend tests in the rolling plane had a thickness of ∼1 mm, and those cut perpendicular to the rolling plane had a thickness of ∼0.8 mm. Shape memory tests of all four geometries were also carried out on a sample that was in the as-austenitized condition. The SME for samples tested in bending in the four different geometries is shown in Fig. 2. As can be seen, in the as-austenitized condition the perpendicular samples perform slightly better than those tested in the rolling plane. After being cold rolled and annealed, the samples exhibit a marked increase in the SME, compared to the samples in the austenitized condition. An increase of approximately 30% recovery occurred after TMP compared to the as-austenitized condition. The absolute strain recovered for the perpendicular– transverse and perpendicular–longitudinal samples is shown in Fig. 3. The as-austenitized samples plateau at a strain of around 2.5%, whilst the TMP samples have a maximum recoverable strain of 4.5% followed by a steep drop. 3.2. Effect of strain and annealing temperature on SME A series of experiments were carried out to optimize the cold rolling reduction and recovery annealing conditions for good SME. A group of austenitized strips were cold rolled to between 6% and 39% reduction. These were sectioned and annealed at 800 ◦ C for 15 min, 700 ◦ C for 30 min and 600 ◦ C for 1 h. The shape memory for each of these rolling reductions and annealing temperatures was tested in bending. All samples were tested in the same orientation with respect to the rolling direction, and this orientation is shown in Fig. 1 as section PT, the transverse section, perpendicular to the rolling plane. The shape memory of each sample at a pre-strain of 2% and 4% was estimated from the shape memory curves, and these estimations were used to examine the effect of cold rolling strain and recovery temperature on shape memory (Fig. 4). The best shape memory behaviour was
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Fig. 3. Total recovered strain for samples in the as-austenitized condition, and after cold rolling to 15% strain and annealing at 800 ◦ C (dashed line represents 100% strain recovery). Samples were tested in the perpendicular–transverse and perpendicular–longitudinal orientations (PT and PL in Fig. 1).
Fig. 2. Shape memory of samples measured in bending in different orientations with respect to the rolling direction. (a) Samples tested in as-austenitized condition. (b) Samples tested after TMP by cold rolling to 15% reduction and annealing for 15 min at 800 ◦ C.
obtained after cold rolling to 15% strain followed by recovery annealing at 800 ◦ C. Above this strain the shape memory decreases. At the highest strain of 39%, the shape recovery drops slightly below that of the austenitized condition. In all instances the minimum shape recovery occurred for the samples annealed at 600 ◦ C, and in this case the best rolling strain prior to annealing was 6%. For an annealing temperature of 700 ◦ C the best rolling strain was between 6% and 10%. TEM micrographs of selected samples are shown in Fig. 5. After rolling and annealing at 600 ◦ C the microstructure was typically dislocated, Fig. 5(a). After annealing at higher temperatures the deformed microstructure recovers into well-defined stacking faults that are dispersed throughout the parent phase. In the case of the optimum conditions for SME (15% cold rolling followed by 15 min at 800 ◦ C) there is a spacing between stacking faults of approximately 1 m (Fig. 5(b)). Between stacking faults, the microstructure is clear of dislocation tangles.
Fig. 4. Effect of cold rolling and annealing on shape recovery, measured in bending to a pre-strain of (a) 2% and (b) 4%.
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Fig. 5. TEM micrographs of samples deformed to (a) 17% strain and recovered at 600 ◦ C, (b) 15% strain and recovered for 15 min at 800 ◦ C, and (c) 39% strain and recovered at 800 ◦ C. All micrographs shown at same magnification.
However, even after annealing at 800 ◦ C the samples rolled to the highest strain of 39% did not form well-defined stacking faults and retained a significant amount of dislocation tangles (Fig. 5(c)). There were also some pockets of well-recovered microstructure and these did not contain any stacking faults. 3.3. Effect of TMP on martensite transformation Differential scanning calorimetry (DSC) was used to examine the transformation behaviour of the samples shown in Fig. 4 that were cold rolled to between 0% and 39% strain, and annealed at 800 ◦ C for 15 min. Before DSC testing, samples were first immersed in liquid nitrogen, then heated from room temperature to 300 ◦ C at 10 ◦ C/min to measure the reversion of martensite to austenite. The cooling cycle was run at a rate of 5 ◦ C/min from 100 to −50 ◦ C, which is the limit of the instrument used. Fig. 6 shows that the austenite start/finish temperatures (As and Af ) were not significantly altered by the cold rolling strain (Fig. 6(a)). However, the heat of transformation (Q) of the martensite to austenite reversion was markedly affected by strain (Fig. 6(b)). The samples tested after austenitizing and cooling in liquid nitrogen barely showed a distinguishable martensite to austenite transformation peak. With increasing strain up to 15% the peaks became more defined and exhibited a larger value of Qε→γ . Above 15% strain the size of the peaks decreased again, and returned to a slightly lower Qε→γ than in the as-austenitized condition. The martensite start temperature (Ms ) could not be measured in all samples, but for those that showed a distinct peak the Ms was around −20 ◦ C (Fig. 6(a)). It is not inferred that the other samples did not exhibit an Ms peak, but rather that the diffuse peaks with small Q shown by this alloy are easily masked by the baseline deviation that occurs because these low temperatures are close to the limit of the cooling range of the instrument (−50 ◦ C). Samples were analyzed using XRD to estimate the proportion of martensite in the microstructure in two different conditions: after cooling in a liquid nitrogen bath to produce thermal martensite, and after 5% cold rolling to produce stress-induced martensite, Fig. 7. An as-austenitized sample and a thermo-
Fig. 6. (a) Transformation temperatures measured using DSC for samples cold rolled to the strain shown followed by annealing for 15 min at 800 ◦ C. Ms : martensite start temperature, As : austenite start temperature, Af : austenite finish temperature, Ap : temperature at maximum heat flow. (b) Heat evolved during transformation (Q), measured using DSC, for samples cold rolled to the strain shown, followed by annealing for 15 min at 800 ◦ C.
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Fig. 8. Stress–strain curves for samples in the as-austenitized condition, and after TMP by 15% cold rolling followed by annealing at 800 ◦ C for 15 min.
3.4. Effect of TMP on mechanical behaviour Two tensile test coupons were machined from 1.3 mm thick hot rolled plate. These specimens were both austenitized at 1000 ◦ C for 1 h followed by water quenching. One specimen was then cold rolled in the tensile direction to 15% strain followed by annealing at 800 ◦ C for 15 min. Each specimen was then tensile tested to failure, Fig. 8. The as-austenitized sample was quite ductile, exhibiting a strain to failure of 80% elongation. TMP by cold rolling and annealing raised the yield point and tensile strength, and decreased the elongation to failure. Both samples exhibited continuous yielding. Fig. 7. (a) Example of XRD scan of as-austenitized sample (i) after immersion in liquid nitrogen and (ii) after cold rolling to 5% strain. (b) The relative amount of martensite to austenite compared by the ratio of counts for the martensite peak divided by the total number of counts for both peaks. Measurements are for the two pre-processing conditions: as-austenitized and TMP (15% cold roll, 15 min at 800 ◦ C); with each case being subjected to (i) cooling in liquid N2 and (ii) 5% cold rolling.
mechanically processed sample (15% cold rolling, 800 ◦ C for 15 min) were investigated in each of these conditions. To examine the relative changes in the ratio of martensite to austenite, ¯ ε and the (2 0 0) γ peaks were measured and an only the (1011) example is shown in Fig. 7(a). This analysis showed there was only a small amount of martensite present after the N2 bath treatment, and that the sample that had been thermo-mechanically processed had a larger fraction of martensite compared to the as-austenitized condition. For both pre-processing conditions (austenitized and TMP) more martensite was present after 5% cold rolling compared to the N2 bath treatment. Furthermore, more martensite was present in the TMP sample than in the asaustenitized sample for both N2 bath treatment and after 5% cold rolling.
3.5. Crystallographic preferred orientation The texture of the as-austenitized material was measured using EBSD and is shown in Fig. 9 as a series of (1 1 1) and (0 1 1) pole figures. The texture is consistent with the recrystallization texture of other low stacking fault energy alloys, such as brass [11]. The same measured texture is represented four times, once for each of the bend test geometries shown schematically in Fig. 1. It can be seen from the four texture representations that the testing orientations that perform best (the perpendicular transverse and perpendicular longitudinal) have the (1 1 0) normal aligned perpendicular to the bending normal direction (BND) and bending tensile direction (TenD) (i.e. in the centre of the pole figure). 4. Discussion 4.1. Orientation of bend test sample The SME measured in bending is affected by the orientation of the test strip in relation to the rolling direction (Fig. 2). This effect is evident in samples in the as-austenitized condition and
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Fig. 9. Texture of as-austenitized strip measured using EBSD, shown as (1 1 1) and (1 1 0) pole figures. The measured texture is represented four times, once for each of the bend test geometries shown schematically in Fig. 1. Each pole figure has the bending normal direction (BND) vertical, and the bending tensile direction (TenD) horizontal. Contour lines represent 1.25, 1.54, 1.83, . . ., 3.58 times random, f(g)max = 4.2 times random intensity.
also in those that have been processed by cold rolling and annealing. It is probably a result of preferred orientation of the parent grains. Stacking faults are crystallographic, and grains that are correctly orientated for these to easily nucleate and grow into martensite in response to an applied strain will be beneficial for SME. ε martensite forms by the movement of Shockley partial dislocations on alternate close packed (1 1 1) planes of the parent phase. This produces stacking faults, that can overlap to form hexagonal ε martensite plates [1]. On each of the four {1 1 1} planes there are three possible 1 1 2 directions in which the partials can move, resulting in 12 possible variants of ε from each individual parent orientation. In single crystals, the number of variants formed is know to be strongly dependant on the orientation of the crystal, and that formation of one variant only is the optimum condition for good shape memory [12]. It has been reported that for a 4 1 4 single crystal deformed by bending, the total true strain recovered was over 8% [1]. This represents close to twice the value achieved by polycrystalline shape memory alloys such as those investigated here (Fig. 3). It is likely that sharpening the texture could significantly increase the SME. Matsumura et al. [13] have investigated the effect of texture on shape memory in an Fe–Mn–Si-Cr alloy. They conclude that the SME will be maximized if the tensile axis is orientated at approximately 1 1 0, which is only 10◦ from the 4 1 4 orientation which is known to be optimal in the case of single crystals. The work of Matsumura et al. [13] supports the view that a strengthening of the appropriate texture will increase the SME, and it is planned to investigate this aspect of shape memory in more detail in a forth-coming paper. The difference in shape recovery as a function of orientation of bending (Fig. 2) highlights the inherent difficulties in comparing experimental data. Given this variability, it is difficult to
compare one set of results to another. It would also be helpful for the geometry in which the samples are tested to be explicitly stated in the text of published works. 4.2. Optimal TMP for shape memory A range of cold rolling strains and recovery temperatures have been tested (Fig. 4). The best SME was found after cold rolling to 15% strain and annealing for 15 min at 800 ◦ C. As the annealing temperature was decreased the optimum strain also decreased. At 600 ◦ C the optimum strain was 6%, and at 700 ◦ C the optimum strain was 6–10%. TEM observations showed that the samples with good SME corresponded to a microstructure of well defined stacking faults evenly spaced throughout the parent phase. The SME was lower in samples that exhibited a microstructure free of stacking faults, and was also lower in samples that showed dislocation tangles. Tensile testing (Fig. 8) showed that these alloys do not have a well-defined plateau in their stress–strain curve. In NiTi alloys this plateau indicates the strain regime in which martensite is forming in preference to slip [14]. It was hoped that a similar inflection in the stress–strain curved for these alloys would indicate how the prior thermo-mechanical processing affects the onset of the martensite transformation. Although there were no indications of this inflection for the alloys tested here, the mechanical testing did demonstrate that the samples that were cold rolled and annealed had a higher yield point than those in the austenitized condition. This is opposite to the mechanical behaviour described by Reyhani and McCormick [3] in tension, which showed that the yield point decreases with a training cycle. This difference is probably the result of different deformation geometries—Reyhani and McCormick [3] tensile deformed, annealed and then tensile tested their sample, whereas we have rolled, annealed and then tensile tested. It can be argued
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that increasing the yield point inhibits slip, and promotes ε formation; on the other hand, it can be argued that a lower yield point indicates that the stress to induce martensite has been decreased, and that this also aids SME. Otsuka et al. [7] indicate that training actually does both, and inspection of their Fig. 8 in Ref. [7] suggests that the effect of training on the yield point will depend on the Ms and temperature at which the test is carried out. Either way, whether training raises or lowers yield stress at room temperature, the effect on SME is still the same, it is enhanced. Therefore, the important factor is not the apparent yield stress but how yielding is affected: by production of ε martensite or irreversible plastic flow. 4.3. Martensite transformation DSC showed that the samples that produce the best SME also have large values of Q. This suggests that there is a larger volume fraction of martensite reverting to austenite in these samples. XRD confirmed that after a liquid nitrogen bath, the as-austenitized samples remained predominantly austenite, with only a minor increase in martensite. In comparison, the samples that were thermo-mechanically processed (TMP: 15% cold roll, 800 ◦ C for 15 min) showed a slightly larger increase in martensite formation after the nitrogen bath. After deformation the same trend is observed: the TMP samples have more deformation-induced martensite than samples in the as-austenitized condition. This shows that in the austenitized condition, it is difficult for martensite to form, even when subjected to temperatures far below the Ms . However, after being processed the alloy produced more martensite compared to the as-austenitized condition. This will clearly have a dramatic impact on SME. Alloy “training” by repeated tensile deformation and annealing cycles is well studied and the amount of information available in the literature is exhaustive e.g. Refs. [2–9]. It has been shown that training causes a decrease in the spacing of martensite plates and an increase in the likelihood of single-variant colonies [8,9]. Although these findings that pertain to the morphology of the martensite are not disputed, we have shown here that the amount of martensite formed, both thermal and stress induced, is markedly increased by TMP. It has been shown that after TMP, evenly spaced stacking faults are retained within the parent phase. It follows that these stacking faults in the parent phase act as nucleation sites for the formation of martensite during deformation. These nucleation sites allow the formation of more martensite compared to the as-austenitized condition, and it is this increase in the volume fraction of martensite that is the cause of the marked increase in shape memory exhibited by samples that have been thermo-mechanically processed. The quantitative result that a larger volume of martensite can form in samples that have been appropriately processed, compared to the as-austenitized condition, is the first reported evidence of this kind. It has also been shown that a high recoverable strain can be achieved by cold rolling to 15% strain followed by annealing at 800 ◦ C; and that this optimized TMP eliminates the need for the repetitive tensile training that has been commonly adopted for these alloys.
5. Conclusions The following conclusions are drawn for the Fe–Mn–Sibased alloy investigated: • The shape memory measured by bend testing is dependant on the orientation of the bend test strip with respect to the rolling direction. • The shape recovery can be increased by over 30% by cold rolling and annealing. This is commensurate with a total strain recovery increase from 2.5% to 4.5%. • The optimum processing conditions for shape memory are cold rolling to 15% strain followed by annealing at 800 ◦ C for 15 min. This produces an austenitic microstructure with stacking faults that are evenly spaced throughout the parent phase. • Below the optimum strain of 15% there are fewer stacking faults and larger areas of clear parent phase that are free of faults. Above the optimum strain (15%) the annealing treatment produces a complicated microstructure of dislocation tangles and stacking faults that do not recover sufficiently to provide good SME. • The TMP does not significantly affect the As or Ms transformation temperatures. • The optimum TMP increases the amount of martensite that forms both during cooling below the Ms , and also during deformation. • The optimum shape memory is achieved when there is an evenly spaced array of stacking faults throughout the parent microstructure. These stacking faults provide nucleation sites for the formation of martensite, and it is primarily an increase in the amount of martensite formed that is responsible for the increased SME exhibited by samples in this condition. Acknowledgements The research described in this paper was funded by the Australian Research Council Discovery Grant scheme. The authors would also like to thank Dr. David Wexler for his assistance with TEM, Mr. Greg Tillman for his metallographic expertise, and Dr. Chris Lukey for his assistance with calorie counting. References [1] A. Sato, E. Chishima, K. Soma, T. Mori, Acta Metall. 30 (1982) 117–1183. [2] H. Otsuka, M. Murakami, S. Matsuda, in: M. Doyama, S. S¯omiya, R. Chang (Eds.), Proceedings of the MRS International Meeting on Advanced Materials, vol. 9, Materials Research Society, 1988, pp. 451–456. [3] M.M. Reyhani, P.G. McCormick, Mater. Sci. Eng. A 160 (1993) 57–61. [4] C.Y. Chung, C. Shuchuan, T.Y. Hsu, Mater. Charact. 37 (1996) 227– 236. [5] Y. Watanabe, Y. Mori, A. Sato, J. Mater. Sci. 28 (1993) 1509–1514. [6] M.M. Reyhani, P.G. McCormick, Scripta Metall. Mater. 31 (1994) 875–878. [7] H. Otsuka, H. Yamada, T. Maruyama, H. Tanhashi, S. Matsuda, M. Murakami, ISIJ Int. 30 (1990) 674–679.
N. Stanford, D.P. Dunne / Materials Science and Engineering A 422 (2006) 352–359 [8] N. Bergeon, S. Kajiwara, T. Kikuchi, Acta Mater. 48 (2000) 4053–4064. [9] D.Z. Liu, S. Kajiwara, T. Kikuchi, N. Shinya, Phil. Mag. 83 (2003) 2875–2897. [10] N. Stanford, D.P. Dunne, J. Mater. Sci., in press. [11] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, 1995.
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[12] A. Sato, E. Chishima, Y. Yamaji, T. Mori, Acta Metall. 32 (1984) 539–547. [13] O. Matsumura, S. Furusako, T. Furukawa, H. Otsuka, ISIJ Int. 36 (1996) 1103–1108. [14] T. Saburi, in: K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge University Press, 1998, p. 58.
Thermo-mechanical processing and the shape memory effect in an Fe–Mn–Si-based shape memory alloy N. Stanford ∗ , D.P. Dunne Faculty of Engineering, University of Wollongong, Northfields Ave., Wollongong 2522, Australia Received 28 October 2005; accepted 14 February 2006
Abstract The effect of cold rolling and annealing on the shape memory effect (SME) in an Fe–Mn–Si-based alloy has been studied. It has been found that the SME in these alloys can be significantly increased by the appropriate thermo-mechanical processing (TMP). The optimum conditions were found to be 15% cold rolling followed by annealing at 800 ◦ C for 15 min. This produced a total strain recovery of 4.5%. TEM showed that this processing schedule produces a microstructure of evenly spaced, and well defined stacking faults throughout the parent phase. It is shown for the first time that samples processed in this way produce a larger fraction of martensite compared to samples in the as-austenitized condition. It is concluded that the stacking faults induced by TMP act as nucleation sites for martensite formation during deformation. The SME is improved primarily as a result of the increased amount of martensite that is formed in this condition. © 2006 Elsevier B.V. All rights reserved. Keywords: Shape memory; Training; Ferrous alloys; Cold rolling; Annealing
1. Introduction The Fe–Mn–Si-based alloy system produces a hexagonal martensite, known as ε, in response to deformation. This martensite is produced by the formation and overlap of stacking faults in the parent phase, and can revert back to austenite by heating. This imparts the alloy with shape memory behaviour [1]. It is known that optimum shape memory is achieved in Fe–Mn–Sibased alloys by using thermo-mechanical processing (TMP) to “train” the alloy [2]. The most common method is to train the specimen by tensile strain followed by a recovery anneal, and then to repeat the cycle for up to about five times. The shape memory is also tested in tension. The effects of repeated training have been studied in some detail. As well as markedly increasing the strain recovery, training is known to significantly decrease the “yield stress” of the alloy [3,4], which is identified as the stress required to form martensite. Training also increases the uniformity of deformation along the sample length, and the relative uniformity continues to increase with repeated training cycles [5]. The density of stacking faults retained within the parent phase is also known
∗
Corresponding author. Tel.: +61 2 4221 3457; fax: +61 2 4221 5474. E-mail address: [email protected] (N. Stanford).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.009
to increase with training cycles [6]. Observations of the sample surface have clearly shown that, at least qualitatively, more ε is formed during deformation after the training regime [5,7]. Atomic force microscopy of the free surface has shown that the morphology of the martensite is affected by training also. The martensite plates become thinner, and are more likely to form with a single variant within each parent grain [8,9]. In recent work [10], we have shown that TMP of the alloy by rolling and annealing can produce a marked increase in the shape memory effect (SME) measured in bending. Rolling and annealing is a much more convenient method of processing than tension, and lends itself to large-scale industrial production. The tedious tensile deformation and annealing of shape memory alloys are not a practical method of producing large products such as, for example, pipe joining collars. In this paper we report results of a systematic investigation of the effectiveness of cold rolling and annealing as a SME enhancing process, with the aim of optimizing the processing conditions.
2. Method The base stainless alloy used in this study was made in a 2.6 kg ingot by induction melting under an argon atmosphere. The composition of the alloy in wt.% was Fe–13Mn–5Si–9Cr–
N. Stanford, D.P. Dunne / Materials Science and Engineering A 422 (2006) 352–359
7Ni–0.03C. The as-cast alloy was hot rolled at 900 ◦ C to a thickness of ∼1 mm. The rolled strip was sectioned and machined into tensile test coupons or bend test strips before being solution treated at 1000 ◦ C for 1 h under flowing argon atmosphere, followed by immediate cold water quench. This is considered the starting condition, and will hereafter be referred to as the as-austenitized condition. As-austenitized samples were deformed at room temperature by cold rolling to strains between 5% and 39%. After the deformation step, samples were recovery annealed at temperatures between 600 and 800 ◦ C. The shape memory of each sample was measured in bending over a range of pre-strains between 1% and 6%. The pre-strain was taken as the maximum tensile/compressive strain, and is determined by the equation: ε=
1 (2R/ h) + 1
(1)
where ε = conventional strain, R = bend radius, h = sample thickness. The shape was recovered by annealing in a muffle furnace at 400 ◦ C for 15 min. After the samples had been recovery annealed the residual strain (εr ) remaining due to incomplete recovery was calculated using Eq. (1). The percentage recovery was determined by the equation: % recovery =
(ε − εr ) ε
(2)
Tensile specimens were deformed using an Instron tensile tester at a speed of 1 mm/min. The starting gage length of tensile specimens was 30 mm, and the starting gage width was 5 mm. Microstructural examination of samples was also carried out by transmission electron microscopy (TEM). The TEM foils were made using a Tenupol jet-polisher with a solution of 5% perchloric acid in acetic acid at 30 V. TEM observations were made using a JEOL JEM 2010 TEM at 200 kV. The transformation behaviour of the samples was investigated using a TA Q100 differential scanning calorimeter (DSC) at a heating rate of 10 ◦ C/min and a cooling rate of 5 ◦ C/min. X-ray diffraction (XRD) was carried out on a Philips PW 1730 XRD at a scan rate of 0.5 ◦ C/min. Prior to XRD analysis, samples were electropolished in a solution of 5% perchloric in acetic acid for 1 min at 30 V to remove the deformed layer induced by mechanical polishing. The crystallographic preferred orientation of the samples was measured using electron backscattered diffraction (EBSD) equipment supplied by HKL Technology. This system is fitted to a Leica 440 scanning electron microscope (SEM). Samples were prepared for EBSD by standard metallographic grinding and polishing techniques followed by a final polish with colloidal silica (Struers OPS) for 10 min. 3. Results 3.1. Orientation of bend testing The effect of the orientation of the test specimen on SME was investigated. A length of the as-austenitized strip was solution
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Fig. 1. Illustration of the four geometries tested in bending with respect to the rolling direction (RD) of the thermo-mechanical treatments. The shaded boxes represent the longitudinal (L), transverse (T), perpendicular longitudinal (PL) and perpendicular transverse (PT) samples. The orientations in which these samples were tested in bending are represented by corresponding grey shading. The direction of tensile stress on the outer sample edge during the bend test is also shown.
treated, and then processed by cold rolling to 15% reduction followed by a recovery anneal at 800 ◦ C for 15 min. After rolling, the strip was sectioned in four different bending geometries; longitudinal and transverse within the rolling plane, longitudinal and transverse perpendicular to the rolling plane (Fig. 1). Samples for bend tests in the rolling plane had a thickness of ∼1 mm, and those cut perpendicular to the rolling plane had a thickness of ∼0.8 mm. Shape memory tests of all four geometries were also carried out on a sample that was in the as-austenitized condition. The SME for samples tested in bending in the four different geometries is shown in Fig. 2. As can be seen, in the as-austenitized condition the perpendicular samples perform slightly better than those tested in the rolling plane. After being cold rolled and annealed, the samples exhibit a marked increase in the SME, compared to the samples in the austenitized condition. An increase of approximately 30% recovery occurred after TMP compared to the as-austenitized condition. The absolute strain recovered for the perpendicular– transverse and perpendicular–longitudinal samples is shown in Fig. 3. The as-austenitized samples plateau at a strain of around 2.5%, whilst the TMP samples have a maximum recoverable strain of 4.5% followed by a steep drop. 3.2. Effect of strain and annealing temperature on SME A series of experiments were carried out to optimize the cold rolling reduction and recovery annealing conditions for good SME. A group of austenitized strips were cold rolled to between 6% and 39% reduction. These were sectioned and annealed at 800 ◦ C for 15 min, 700 ◦ C for 30 min and 600 ◦ C for 1 h. The shape memory for each of these rolling reductions and annealing temperatures was tested in bending. All samples were tested in the same orientation with respect to the rolling direction, and this orientation is shown in Fig. 1 as section PT, the transverse section, perpendicular to the rolling plane. The shape memory of each sample at a pre-strain of 2% and 4% was estimated from the shape memory curves, and these estimations were used to examine the effect of cold rolling strain and recovery temperature on shape memory (Fig. 4). The best shape memory behaviour was
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Fig. 3. Total recovered strain for samples in the as-austenitized condition, and after cold rolling to 15% strain and annealing at 800 ◦ C (dashed line represents 100% strain recovery). Samples were tested in the perpendicular–transverse and perpendicular–longitudinal orientations (PT and PL in Fig. 1).
Fig. 2. Shape memory of samples measured in bending in different orientations with respect to the rolling direction. (a) Samples tested in as-austenitized condition. (b) Samples tested after TMP by cold rolling to 15% reduction and annealing for 15 min at 800 ◦ C.
obtained after cold rolling to 15% strain followed by recovery annealing at 800 ◦ C. Above this strain the shape memory decreases. At the highest strain of 39%, the shape recovery drops slightly below that of the austenitized condition. In all instances the minimum shape recovery occurred for the samples annealed at 600 ◦ C, and in this case the best rolling strain prior to annealing was 6%. For an annealing temperature of 700 ◦ C the best rolling strain was between 6% and 10%. TEM micrographs of selected samples are shown in Fig. 5. After rolling and annealing at 600 ◦ C the microstructure was typically dislocated, Fig. 5(a). After annealing at higher temperatures the deformed microstructure recovers into well-defined stacking faults that are dispersed throughout the parent phase. In the case of the optimum conditions for SME (15% cold rolling followed by 15 min at 800 ◦ C) there is a spacing between stacking faults of approximately 1 m (Fig. 5(b)). Between stacking faults, the microstructure is clear of dislocation tangles.
Fig. 4. Effect of cold rolling and annealing on shape recovery, measured in bending to a pre-strain of (a) 2% and (b) 4%.
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Fig. 5. TEM micrographs of samples deformed to (a) 17% strain and recovered at 600 ◦ C, (b) 15% strain and recovered for 15 min at 800 ◦ C, and (c) 39% strain and recovered at 800 ◦ C. All micrographs shown at same magnification.
However, even after annealing at 800 ◦ C the samples rolled to the highest strain of 39% did not form well-defined stacking faults and retained a significant amount of dislocation tangles (Fig. 5(c)). There were also some pockets of well-recovered microstructure and these did not contain any stacking faults. 3.3. Effect of TMP on martensite transformation Differential scanning calorimetry (DSC) was used to examine the transformation behaviour of the samples shown in Fig. 4 that were cold rolled to between 0% and 39% strain, and annealed at 800 ◦ C for 15 min. Before DSC testing, samples were first immersed in liquid nitrogen, then heated from room temperature to 300 ◦ C at 10 ◦ C/min to measure the reversion of martensite to austenite. The cooling cycle was run at a rate of 5 ◦ C/min from 100 to −50 ◦ C, which is the limit of the instrument used. Fig. 6 shows that the austenite start/finish temperatures (As and Af ) were not significantly altered by the cold rolling strain (Fig. 6(a)). However, the heat of transformation (Q) of the martensite to austenite reversion was markedly affected by strain (Fig. 6(b)). The samples tested after austenitizing and cooling in liquid nitrogen barely showed a distinguishable martensite to austenite transformation peak. With increasing strain up to 15% the peaks became more defined and exhibited a larger value of Qε→γ . Above 15% strain the size of the peaks decreased again, and returned to a slightly lower Qε→γ than in the as-austenitized condition. The martensite start temperature (Ms ) could not be measured in all samples, but for those that showed a distinct peak the Ms was around −20 ◦ C (Fig. 6(a)). It is not inferred that the other samples did not exhibit an Ms peak, but rather that the diffuse peaks with small Q shown by this alloy are easily masked by the baseline deviation that occurs because these low temperatures are close to the limit of the cooling range of the instrument (−50 ◦ C). Samples were analyzed using XRD to estimate the proportion of martensite in the microstructure in two different conditions: after cooling in a liquid nitrogen bath to produce thermal martensite, and after 5% cold rolling to produce stress-induced martensite, Fig. 7. An as-austenitized sample and a thermo-
Fig. 6. (a) Transformation temperatures measured using DSC for samples cold rolled to the strain shown followed by annealing for 15 min at 800 ◦ C. Ms : martensite start temperature, As : austenite start temperature, Af : austenite finish temperature, Ap : temperature at maximum heat flow. (b) Heat evolved during transformation (Q), measured using DSC, for samples cold rolled to the strain shown, followed by annealing for 15 min at 800 ◦ C.
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Fig. 8. Stress–strain curves for samples in the as-austenitized condition, and after TMP by 15% cold rolling followed by annealing at 800 ◦ C for 15 min.
3.4. Effect of TMP on mechanical behaviour Two tensile test coupons were machined from 1.3 mm thick hot rolled plate. These specimens were both austenitized at 1000 ◦ C for 1 h followed by water quenching. One specimen was then cold rolled in the tensile direction to 15% strain followed by annealing at 800 ◦ C for 15 min. Each specimen was then tensile tested to failure, Fig. 8. The as-austenitized sample was quite ductile, exhibiting a strain to failure of 80% elongation. TMP by cold rolling and annealing raised the yield point and tensile strength, and decreased the elongation to failure. Both samples exhibited continuous yielding. Fig. 7. (a) Example of XRD scan of as-austenitized sample (i) after immersion in liquid nitrogen and (ii) after cold rolling to 5% strain. (b) The relative amount of martensite to austenite compared by the ratio of counts for the martensite peak divided by the total number of counts for both peaks. Measurements are for the two pre-processing conditions: as-austenitized and TMP (15% cold roll, 15 min at 800 ◦ C); with each case being subjected to (i) cooling in liquid N2 and (ii) 5% cold rolling.
mechanically processed sample (15% cold rolling, 800 ◦ C for 15 min) were investigated in each of these conditions. To examine the relative changes in the ratio of martensite to austenite, ¯ ε and the (2 0 0) γ peaks were measured and an only the (1011) example is shown in Fig. 7(a). This analysis showed there was only a small amount of martensite present after the N2 bath treatment, and that the sample that had been thermo-mechanically processed had a larger fraction of martensite compared to the as-austenitized condition. For both pre-processing conditions (austenitized and TMP) more martensite was present after 5% cold rolling compared to the N2 bath treatment. Furthermore, more martensite was present in the TMP sample than in the asaustenitized sample for both N2 bath treatment and after 5% cold rolling.
3.5. Crystallographic preferred orientation The texture of the as-austenitized material was measured using EBSD and is shown in Fig. 9 as a series of (1 1 1) and (0 1 1) pole figures. The texture is consistent with the recrystallization texture of other low stacking fault energy alloys, such as brass [11]. The same measured texture is represented four times, once for each of the bend test geometries shown schematically in Fig. 1. It can be seen from the four texture representations that the testing orientations that perform best (the perpendicular transverse and perpendicular longitudinal) have the (1 1 0) normal aligned perpendicular to the bending normal direction (BND) and bending tensile direction (TenD) (i.e. in the centre of the pole figure). 4. Discussion 4.1. Orientation of bend test sample The SME measured in bending is affected by the orientation of the test strip in relation to the rolling direction (Fig. 2). This effect is evident in samples in the as-austenitized condition and
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Fig. 9. Texture of as-austenitized strip measured using EBSD, shown as (1 1 1) and (1 1 0) pole figures. The measured texture is represented four times, once for each of the bend test geometries shown schematically in Fig. 1. Each pole figure has the bending normal direction (BND) vertical, and the bending tensile direction (TenD) horizontal. Contour lines represent 1.25, 1.54, 1.83, . . ., 3.58 times random, f(g)max = 4.2 times random intensity.
also in those that have been processed by cold rolling and annealing. It is probably a result of preferred orientation of the parent grains. Stacking faults are crystallographic, and grains that are correctly orientated for these to easily nucleate and grow into martensite in response to an applied strain will be beneficial for SME. ε martensite forms by the movement of Shockley partial dislocations on alternate close packed (1 1 1) planes of the parent phase. This produces stacking faults, that can overlap to form hexagonal ε martensite plates [1]. On each of the four {1 1 1} planes there are three possible 1 1 2 directions in which the partials can move, resulting in 12 possible variants of ε from each individual parent orientation. In single crystals, the number of variants formed is know to be strongly dependant on the orientation of the crystal, and that formation of one variant only is the optimum condition for good shape memory [12]. It has been reported that for a 4 1 4 single crystal deformed by bending, the total true strain recovered was over 8% [1]. This represents close to twice the value achieved by polycrystalline shape memory alloys such as those investigated here (Fig. 3). It is likely that sharpening the texture could significantly increase the SME. Matsumura et al. [13] have investigated the effect of texture on shape memory in an Fe–Mn–Si-Cr alloy. They conclude that the SME will be maximized if the tensile axis is orientated at approximately 1 1 0, which is only 10◦ from the 4 1 4 orientation which is known to be optimal in the case of single crystals. The work of Matsumura et al. [13] supports the view that a strengthening of the appropriate texture will increase the SME, and it is planned to investigate this aspect of shape memory in more detail in a forth-coming paper. The difference in shape recovery as a function of orientation of bending (Fig. 2) highlights the inherent difficulties in comparing experimental data. Given this variability, it is difficult to
compare one set of results to another. It would also be helpful for the geometry in which the samples are tested to be explicitly stated in the text of published works. 4.2. Optimal TMP for shape memory A range of cold rolling strains and recovery temperatures have been tested (Fig. 4). The best SME was found after cold rolling to 15% strain and annealing for 15 min at 800 ◦ C. As the annealing temperature was decreased the optimum strain also decreased. At 600 ◦ C the optimum strain was 6%, and at 700 ◦ C the optimum strain was 6–10%. TEM observations showed that the samples with good SME corresponded to a microstructure of well defined stacking faults evenly spaced throughout the parent phase. The SME was lower in samples that exhibited a microstructure free of stacking faults, and was also lower in samples that showed dislocation tangles. Tensile testing (Fig. 8) showed that these alloys do not have a well-defined plateau in their stress–strain curve. In NiTi alloys this plateau indicates the strain regime in which martensite is forming in preference to slip [14]. It was hoped that a similar inflection in the stress–strain curved for these alloys would indicate how the prior thermo-mechanical processing affects the onset of the martensite transformation. Although there were no indications of this inflection for the alloys tested here, the mechanical testing did demonstrate that the samples that were cold rolled and annealed had a higher yield point than those in the austenitized condition. This is opposite to the mechanical behaviour described by Reyhani and McCormick [3] in tension, which showed that the yield point decreases with a training cycle. This difference is probably the result of different deformation geometries—Reyhani and McCormick [3] tensile deformed, annealed and then tensile tested their sample, whereas we have rolled, annealed and then tensile tested. It can be argued
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that increasing the yield point inhibits slip, and promotes ε formation; on the other hand, it can be argued that a lower yield point indicates that the stress to induce martensite has been decreased, and that this also aids SME. Otsuka et al. [7] indicate that training actually does both, and inspection of their Fig. 8 in Ref. [7] suggests that the effect of training on the yield point will depend on the Ms and temperature at which the test is carried out. Either way, whether training raises or lowers yield stress at room temperature, the effect on SME is still the same, it is enhanced. Therefore, the important factor is not the apparent yield stress but how yielding is affected: by production of ε martensite or irreversible plastic flow. 4.3. Martensite transformation DSC showed that the samples that produce the best SME also have large values of Q. This suggests that there is a larger volume fraction of martensite reverting to austenite in these samples. XRD confirmed that after a liquid nitrogen bath, the as-austenitized samples remained predominantly austenite, with only a minor increase in martensite. In comparison, the samples that were thermo-mechanically processed (TMP: 15% cold roll, 800 ◦ C for 15 min) showed a slightly larger increase in martensite formation after the nitrogen bath. After deformation the same trend is observed: the TMP samples have more deformation-induced martensite than samples in the as-austenitized condition. This shows that in the austenitized condition, it is difficult for martensite to form, even when subjected to temperatures far below the Ms . However, after being processed the alloy produced more martensite compared to the as-austenitized condition. This will clearly have a dramatic impact on SME. Alloy “training” by repeated tensile deformation and annealing cycles is well studied and the amount of information available in the literature is exhaustive e.g. Refs. [2–9]. It has been shown that training causes a decrease in the spacing of martensite plates and an increase in the likelihood of single-variant colonies [8,9]. Although these findings that pertain to the morphology of the martensite are not disputed, we have shown here that the amount of martensite formed, both thermal and stress induced, is markedly increased by TMP. It has been shown that after TMP, evenly spaced stacking faults are retained within the parent phase. It follows that these stacking faults in the parent phase act as nucleation sites for the formation of martensite during deformation. These nucleation sites allow the formation of more martensite compared to the as-austenitized condition, and it is this increase in the volume fraction of martensite that is the cause of the marked increase in shape memory exhibited by samples that have been thermo-mechanically processed. The quantitative result that a larger volume of martensite can form in samples that have been appropriately processed, compared to the as-austenitized condition, is the first reported evidence of this kind. It has also been shown that a high recoverable strain can be achieved by cold rolling to 15% strain followed by annealing at 800 ◦ C; and that this optimized TMP eliminates the need for the repetitive tensile training that has been commonly adopted for these alloys.
5. Conclusions The following conclusions are drawn for the Fe–Mn–Sibased alloy investigated: • The shape memory measured by bend testing is dependant on the orientation of the bend test strip with respect to the rolling direction. • The shape recovery can be increased by over 30% by cold rolling and annealing. This is commensurate with a total strain recovery increase from 2.5% to 4.5%. • The optimum processing conditions for shape memory are cold rolling to 15% strain followed by annealing at 800 ◦ C for 15 min. This produces an austenitic microstructure with stacking faults that are evenly spaced throughout the parent phase. • Below the optimum strain of 15% there are fewer stacking faults and larger areas of clear parent phase that are free of faults. Above the optimum strain (15%) the annealing treatment produces a complicated microstructure of dislocation tangles and stacking faults that do not recover sufficiently to provide good SME. • The TMP does not significantly affect the As or Ms transformation temperatures. • The optimum TMP increases the amount of martensite that forms both during cooling below the Ms , and also during deformation. • The optimum shape memory is achieved when there is an evenly spaced array of stacking faults throughout the parent microstructure. These stacking faults provide nucleation sites for the formation of martensite, and it is primarily an increase in the amount of martensite formed that is responsible for the increased SME exhibited by samples in this condition. Acknowledgements The research described in this paper was funded by the Australian Research Council Discovery Grant scheme. The authors would also like to thank Dr. David Wexler for his assistance with TEM, Mr. Greg Tillman for his metallographic expertise, and Dr. Chris Lukey for his assistance with calorie counting. References [1] A. Sato, E. Chishima, K. Soma, T. Mori, Acta Metall. 30 (1982) 117–1183. [2] H. Otsuka, M. Murakami, S. Matsuda, in: M. Doyama, S. S¯omiya, R. Chang (Eds.), Proceedings of the MRS International Meeting on Advanced Materials, vol. 9, Materials Research Society, 1988, pp. 451–456. [3] M.M. Reyhani, P.G. McCormick, Mater. Sci. Eng. A 160 (1993) 57–61. [4] C.Y. Chung, C. Shuchuan, T.Y. Hsu, Mater. Charact. 37 (1996) 227– 236. [5] Y. Watanabe, Y. Mori, A. Sato, J. Mater. Sci. 28 (1993) 1509–1514. [6] M.M. Reyhani, P.G. McCormick, Scripta Metall. Mater. 31 (1994) 875–878. [7] H. Otsuka, H. Yamada, T. Maruyama, H. Tanhashi, S. Matsuda, M. Murakami, ISIJ Int. 30 (1990) 674–679.
N. Stanford, D.P. Dunne / Materials Science and Engineering A 422 (2006) 352–359 [8] N. Bergeon, S. Kajiwara, T. Kikuchi, Acta Mater. 48 (2000) 4053–4064. [9] D.Z. Liu, S. Kajiwara, T. Kikuchi, N. Shinya, Phil. Mag. 83 (2003) 2875–2897. [10] N. Stanford, D.P. Dunne, J. Mater. Sci., in press. [11] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, 1995.
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[12] A. Sato, E. Chishima, Y. Yamaji, T. Mori, Acta Metall. 32 (1984) 539–547. [13] O. Matsumura, S. Furusako, T. Furukawa, H. Otsuka, ISIJ Int. 36 (1996) 1103–1108. [14] T. Saburi, in: K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge University Press, 1998, p. 58.