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Role of microstructure on phase transformation behavior in Ni–Ti–Fe shape memory alloys during thermal cycling
Journal of Alloys and Compounds 652 (2015) 459e469
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Role of microstructure on phase transformation behavior in NieTieFe shape memory alloys during thermal cycling Ritwik Basu a, *, M.A. Mohtadi-Bonab b, Xu Wang b, Mostafa Eskandari b, c, Jerzy A. Szpunar b a b c
Department of Mechanical Engineering, The NORTHCAP University, Gurgaon, 122017, India Department of Mechanical Engineering, University of Saskatchewan, Saskatoon, S7N5A9, Canada Department of Materials Science & Engineering, Faculty of Engineering, Shahid Chamran University, Ahvaz, Iran
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 July 2015 Received in revised form 27 August 2015 Accepted 28 August 2015 Available online 1 September 2015
Two different microstructures of NieTieFe shape memory alloys processed through different thermomechanical treatments with nearly similar grain size and in-grain misorientation but different crystallographic textures were subjected to series of thermal cycles without external loading. The microstructures and the phase transformation behavior of the samples were examined after every seventy cycles. The experiment involved treating the samples with liquid nitrogen (LN2) for complete martensitic (B190 ) transformation and then heating it back to parent austenite (B2) condition. Thermal cycling introduced significant differences in microstructural parameters especially the grain boundary nature, stored elastic energy and the misorientation or defect densities. These microstructural alterations during thermal cycling were related to changes in transformation temperatures and enthalpy. Thermal cycling also brought out changes in crystallographic orientations of austenite grains. The present study aims to address the role of microstructure on the thermal fatigue behavior in NieTi based shape memory alloys during cyclic transformations. © 2015 Elsevier B.V. All rights reserved.
Keywords: NieTieFe shape memory alloys Thermal cycling Electron back scattered diffraction (EBSD) Texture X-ray diffraction (XRD)
1. Introduction Shape memory alloys of NieTi, commonly called as NITINOL possess the two most important characteristics, the shape memory and the pseudoelasticity, that makes it an important candidate for a broad range of engineering applications [1]. In many of these applications the components fabricated from these alloys are subjected to repeated thermal or mechanical transformation cycles. Examples include, (1) actuators that are designed to experience repeated thermal cycles throughout a wide range of transformation temperatures that are typically determined by the start and finish temperature of austenite4martensite transformation [2,3] and (2) endovascular stents of self-expanding types that undergo several pseudoelastic cycles that results from cardiac pressure pulses [4]. The thermal cycles could be performed either under external loading or zero load conditions. Both thermal and mechanical cycles involve a reversible, first-order, cubic-to-monoclinic (B2/B190 ) thermoelastic martensitic transformation [5,6]. The B2
* Corresponding author. E-mail address: [email protected] (R. Basu). http://dx.doi.org/10.1016/j.jallcom.2015.08.239 0925-8388/© 2015 Elsevier B.V. All rights reserved.
and B190 phases have been referred as austenite and martensite respectively in the present study. The temperatures and the stresses to initiate the transformation from B2/B190 are significantly affected during successive thermal and mechanical cycling [7e13]. Thermoelastic martensites differ from ferrous martensites through presence of internal twin variants that accommodates elastic strains during cooling from high temperature B2 phase. Further deformation causes the detwinning of these micro-twin structures leading to several new martensite variants. Upon heating these different variants revert back to the original high temperature B2 structure. The shape accommodation of martensite in steels occurs by classical slip mechanism and this is irreversible in nature [14]. The effects of thermal cycling on the transformation characteristics and shape memory properties have been studied in the past by several researchers [15e30]. Some of the important findings are summarized as follows: (1) Thermal cycling in absence of external stresses tends to decrease the start temperature of both marteniste and austenite; the Ms and As respectively. This is also followed by widening of thermal hysteresis [22,23]. In addition, repeated cycling show characteristic R-phase (intermediate phase) transformation due to precipitation of the metastable intermetallic
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Ni4Ti3 prior to the martensitic transformation [22,26]. (2) Thermal cycling under external stress increases the start temperatures of transformation [26]. This behavior is attributed partly due to the alignment of dislocation structures in the parent B2 matrix. However a complete understanding on the transformation characteristics of these alloys during cycling under constant stress is still not comprehensible. Thus it can be stated that the observed changes in the functional characteristics as a result of cycling are largely associated to subtle changes in the microstructure [22,30,31]. The available literature on improved cyclic stability of NieTi SMAs reports precipitation strengthening and partial recovery annealing after cold working [22,32,33] as important factors. Although thermal cycling of the SMA components in absence of external stresses introduces functional fatigue [34,35], it has also been shown that these components subjected to thermal cycles at zero stress prior to application (pre-cycling) can improve the fatigue resistance largely in these alloys [36e38]. The dislocation accumulation as a result of precycling facilitates hardening of the thermally cycled alloy, thus improving the shape memory effect. The available literature on thermal cycling of shape memory alloys with or without external loading discusses the dislocations generated during thermal cycling [22,39]. A gap remains in deriving a correlation in the observed microstructural quantities to the cyclic transformation pattern. The present study brings out a relationship between microstructures to the calorimetric signatures of phase transformation (B2/B190 ) in NieTieFe shape memory alloy during stress-free thermal cycling. The study further attempts to addresses the microstructural causes of thermal fatigue. The microstructure and texture investigation have been analyzed through electron back scattered diffraction (EBSD) and x-ray diffraction (XRD) techniques and calorimetric measurements were obtained through differential scanning calorimetry (DSC). The study presents a systematic investigation on the role of grain boundary (GB) characteristics, misorientation, stored energy and crystallographic texture on the transformation parameters; the temperature of transformation and the enthalpy. It has been shown that temperatures of both forward and reverse transformation bear a strong correlation with the fraction of high angle grain boundaries (HAGBs). Increase or decrease in the HAGB fraction reflected through similar changes in the martensite start temperature of transformation. The enthalpy of transformation was a clear function of the stored energy. Thermal cycling caused increase in the ingrain misorientation which is a signature of increased defect density in the samples. It was also brought out in the study that crystallographic orientation of parent austenite grains in the microstructure decides the thermal fatigue response in shape memory materials. Identifying these uncertainties in the transformation sequence and its correlation to the microstructure of the alloy will form a baseline study for manufacturing of actuator devices. Ternary Ni47eTi50eFe3 alloy has been used in the present research as it is widely used in many engineering applications such as heat shrinkable couplings, self-healing gaskets, safety valves, cryogenic thermal switches etc. [40e42].
annealed sample here onwards has been referred as P. Another 10% reduction in thickness was attained in sample P through cold rolling followed by final annealing at 730 C for 0.5 h. This sample here has been referred as A. Details about the microstructure developments through aforementioned deformation routes is beyond the scope of present research; a separate field of study reported in Ref. [43]. Samples for EBSD and bulk crystallographic texture measurements have been prepared thorough conventional metallographic polishing techniques to bring down the thickness of the sample to half. The surface of interest in this study was chosen as the rolling plane. The sample surface was further treated with colloidal silica (0.04 mm) solution to generate a strain free surface for carrying out microstructural measurements [43]. For the calorimetric measurements, specimens were cut using slow speed diamond wafer blade to avoid deformation in the microstructure. The thermal cycling was conducted by holding the samples in liquid nitrogen for martensitic transformation and then heating back for reverse austenite transformation. The microstructures and calorimetric properties were measured after every seventy thermal cycles. 2. 2. Experimental procedures Characterization through EBSD techniques were performed in a SU 6600 Hitachi field emission scanning electron microscope (FESEM) coupled with an Oxford Instruments Norlys Nano EBSD detector. The AZTEC 2.0 data acquisition software was used for acquiring diffraction patterns from the samples which was synchronized with the beam scan. An acceleration voltage of 20 keV was maintained during the scan. The number of bands used for pattern identification were set to six and the data acquisition rate of 20 frames/s was used. A typical step size of 0.35 mm was used for EBSD measurements. The scan parameters and SEM conditions were kept identical for all scans performed in this research. Oxford Instrument's Channel 5 post processing software was used for processing the raw EBSD data. Grain boundaries (GBs) were identified as continuous misorientation region, where misorientation angle > 5 were considered. GBs greater that 15 were considered as high angle grain boundaries (HAGBs). Grain size diameter was determined from diameter of a circle that has the same area of the grain. The area of a grain equals to the number of data points in the grain multiplied by their pixel size. In-grain misorientations were measured as average point-to-point misorientation inside an identified grain: local average misorientation [43]. Local stored energy associated to grain boundary was estimated using the misorientation criterion. Here boundary misorientations between 5 and 15 are categorized as low-angle boundaries (LAGBs) while misorientations 15 are high-angle boundaries (HAGBs). The energy contribution from grain boundary misorientation is estimated from the ReadeShockley equation [44] such that:
gðqÞ ¼ gm
2. Experimental 2.1. Materials and preparations
q qm
q 1 ln ; q qm qm
and
gðqÞ ¼ gm ; q qm Alloy of Ni47eTi50eFe3 was supplied by Hlmet Co., Ltd, China in hot rolled conditions. Further cold rolling was carried out in a lab scale rolling mill to 50% reduction in thickness in three passes with intermediate annealing at 800 C for 25 min followed by final annealing treatments at the same temperature for 1.5 h. The post
(1)
(2)
where gm ¼ 0.706 Jm2 [45] is the energy per unit area of a HAGB, is the boundary misorientation and qm ¼ 15 is the misorientation angle above which the energy per unit area is independent of misorientation angle. The average boundary energy is calculated
R. Basu et al. / Journal of Alloys and Compounds 652 (2015) 459e469
from equations (1) and (2) considering all the boundary misorientations between 5 q 62.8 via summation:
g¼
62:8 X
½gðqÞf ðqÞ
5
(3)
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where f(q) represents the boundary fraction for a given misorientation. The stored energy per unit volume due to a grain boundary dislocation (Eb) is obtained by multiplying the average energy per unit boundary area (g) to the area per unit volume (SV d 3 ) ECD where, dECD is the average grain size measured from equivalent circle diameter considering all misorientations between
Fig. 1. Microstructures represented through inverse pole figure (IPF) map of (a) sample A and (b) sample P under pre-cycled conditions. The IPF color triangle presents important orientations corresponding to the colors. Sample A shows increased presence of ND//<111> grains. Sample P had ND//<110> grains along with few other orientations. Orientation distribution functions (ODFs) of 42 ¼ 0 and 45 sections estimated from EBSD measurements for (c) sample A and (d) sample P shows clear differences in crystallographic textures. The important texture components are illustrated in the ODF plots. (e) Grain size and (f) misorientation distributions were nearly identical for samples A and P. The inset shows the most probable values.
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5 q 62.8 . Thus the stored energy due to boundary dislocations (Eb) can thus be represented in the following form:
Eb ¼ SV g ¼
3g dECD
(4)
It is to be noted that Eb values do not account for in-grain misorientations or dislocations.
Measurements for bulk crystallographic textures were performed in a Bruker AXS General Area Detector system equipped with a Cr-ka radiation. Two pole figures (PF), (110) and (200) were measured for calculation of orientation distribution function (ODF) using Textool programs [46]. The long wavelength of Cr-Ka X-rays imposed difficulties in measuring a third pole figure for ODF calculations. The Cr-Ka X-rays offers advantage over Co or Cu X-rays due to much lower defocus effect because pole figures are
Fig. 2. Differential scanning calorimetric (DSC) plots showed characteristic transformation temperatures and the enthalpy of transformation for (a) sample A and (b) sample P. (c) Martensite (Ms) and (d) austenite start (As) temperatures as a function of thermal cycle show a drop; drop being more clear in sample P. Enthalpy of (e) forward and (f) reverse transformation as a function of thermal cycle is presented. Sample P exhibited a steady decrease in transformational enthalpy. Changes were less significant in sample A. A drop in transformation temperature and enthalpy with multiple thermal cycle is a valid signature of thermal fatigue. Error bars represent typical accuracy of estimated enthalpy and transformation temperatures within 5% of the mean values.
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measured above 2q ¼ 70 . The results on texture for austenite NieTi phase was verified by comparing with Cu (3 PFs). The texture results were nearly identical. The crystal orientations are represented through Euler angles (41, F, 42) [47]. The sample reference is chosen as the rolling direction (RD) and the normal to the rolling plane (ND). For measurement of initiation temperatures and enthalpies of transformations, a TA Instrument's Q2000 differential scanning calorimetry (DSC) unit equipped with a LN2 cooling attachment was used. An inert atmosphere was maintained by using Ar. The heatingecooling cycles for all thermally cycled specimens were performed in the temperature span of 50 C to 50 C at the rate of 3 C/min. 3. Results Fig. 1(aef) bring out comprehensible differences in the microstructures of the samples A and P under investigation prior to subjecting to thermal cycles. The EBSD measured inverse pole figure (IPF) colored orientation maps illustrated in Fig. 1a and b shows distinctly different microstructures. The colors in the IPF map represent different crystallographic orientation of individual grains parallel to the normal direction (ND) of the samples. Readers are advised to refer the standard colored stereo-triangle for interpretation of orientations. The textures represented through the orientation distribution function (ODF) sections of 42 ¼ 0 and 45 in Fig. 1c shows clear presence of g-fiber orientations with highest
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intensities around {111} <123> and {111} <112> in sample A. The ODF plot for Sample P in Fig. 1d shows a relatively high intensity around {211} <111> and a partial spread of orientation along the ND//<110> fiber. However the XRD estimated ODF for sample P also shows a week spread of Goss and cube texture which will be discussed shortly. The grain size and in-grain misorientation distributions with their most probable values presented in Fig. 1 d and e respectively were nearly identical for both the samples A and P. The results on the calorimetric signatures of phase transformation are illustrated in Fig. 2a-f. Fig. 2a and b are representative DSC plots for samples A and P respectively subjected to thermal cycles. The austenite4martensite transformation show a decrease in martensite (Ms) and austenite start (As) temperatures with repeated thermal cycling, see Fig. 2c and d. . The decreasing trend of the start temperatures during thermal cycling has been reported by several researchers in the past [48e50]. This has been attributed primarily to the introduction of transformation-induced dislocations [48] that tend to increase the elastic strain energy in each cycle leading to stabilization of austenite [51,52]. The rate of decrease in transformation temperatures (TT) measured in sample P is higher than sample A. Sample A exhibits a nearly consistent Ms except that a dip is observed at 70th thermal cycle. Significant changes in the enthalpy of transformation during thermal cycles were noticeable. The enthalpy data of forward and reverse transformation as a function of temperature for each cycle is presented in Fig. 2e and f. The enthalpies of both transformation show a monotonic decrease with increasing number of cycles. The effect is
Fig. 3. Microstructural changes in samples A and P as a function of thermal cycles. (a) High angle grain boundary (HAGB) fraction, (b) stored energy (SE) and (c) most probable misorientation are plotted against thermal cycles. Both HAGB and SE show a decrease with increase in thermal cycles. Sample P showed a steady drop in the values. Misorientation showed an increasing trend.
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more pronounced in sample P. The enthalpy of marteniste transformation in sample A shows no significant change. Lower transformation enthalpy is a valid signature of suppression of austenite4martensite transformation. The results on TTs and enthalpy suggest that sample P is more susceptible to functional fatigue and is not an ideal candidate for shape memory applications. Fig. 3a-c discusses the microstructural developments in terms of boundary fraction, stored energy of deformation and the most probable value misorientation at every stage of thermal cycling estimated from EBSD measurements. The measurements were obtained from three EBSD scans performed on the same sample after every thermal cycle. The standard deviations estimated from multiple EBSD measurements are represented through error bars. The fraction of HAGB and stored energy show a decrease with number of thermal cycles. Misorientation showed an opposite trend. More importantly the Ms Temperature is a clear function of the HAGB fraction, i.e an increase or decrease in the Ms temperature correlated well with the measured HAGB fraction. The enthalpy of transformation on the other hand bears a strong dependence on the stored energy estimated from EBSD. Fig. 4a and b brings out the changes in the austenite textures of samples A and P during thermal cycling represented through XRD measured ODFs. The texture ODFs show a wide scatter of orientations around the g-fiber (ND//<111>) line with the highest intensity located in the vicinity of {322} <258> and {322} <236> in the sample A in pre-cycled condition brought out in the 42 ¼ 45
section. The EBSD measured ODF of sample A in Fig. 1c presents similar results. With successive thermal cycling the g-fiber became pronounced along with strong presence of near {100} <011> orientations. The starting microstructure of Sample P exhibits a weak Goss {110} <001> followed by two secondary peaks around {100} <001>, cube and {211}<111> orientations. However the intensities around Goss orientation were too low to be detected in EBSD. Thermal cycling of sample P resulted in strengthening of ND// <110> fiber suggesting that these orientations have grown at the expense of Goss orientation. 4. Discussions The transformation of austenite to martensite proceeds with formation of glissile dislocation interfaces [51]. Any deviation in the glissile nature of the interface would retard the transformation. Several authors have established the presence of transformation induced dislocations that assist the formation of martensite nuclei [53e58]. However, the dislocation density progressively increases with thermal cycling and the glissile relationship between austenite and martensite is gradually lost. There exists no unified theory to explain the mechanism of dislocation accumulation during repeated thermal cycling. The dislocations that may be present in the material prior to thermal cycling as a result of deformation could act as a source of martensite nuclei formation, however with successive cycles additional dislocations contribute which could arise as a result of coalescence of different plates of
Fig. 4. Effects of thermal cycling on the texture developments. The X-ray estimated textures for the austenite phase are represented in standard 42 ¼ 0 and 45 ODF section. ODF plots of (a) sample A and (b) sample P showed differences in their initial and final textures. Sample A showed characteristic g-fiber texture which got strengthened with thermal cycling. Sample P shows development of ND//<110> fiber with increased thermal cycling. The maximum ODF intensities are clearly indicated in the figures.
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martensite variants, stress relaxation due to change in volume during transformation or interaction of dislocations with different martensite plate groups [41]. The transformation induced dislocations impede further transformation of austenite to martensite. The measured average misorientation in Fig. 3c show an increase with thermal cycles which scales nearly with defect densities. A decrease in the Ms Temperature implies difficulties in transformation. The increased presence of dislocations with thermal cycling is depicted in the band contrast map obtained from two scans from the same sample with and without cycling, see Fig. 5aed. The pre-cycled sample shows dislocations in the form of banded structures in few different grains. These different structures were apparent primarily in the near ND//<110> grains as seen in the IPF maps. The dark regions are regions of poor band contrast referring to low indexed points as a result of strain accumulation. There were also grains which were free of such micro features. Significant differences in the sub-structures were brought out in the 140th cycle with clear presence of increased dislocation arrays and bands. The GBs shown in blue are HAGBs. These structures show misorientation varying from 2 to 5 . It can be said that the structure formation occurs at the expense of the HAGBs. This led to drop in the fraction of HAGBs with successive thermal cycling which coincided with similar change in the Ms temperature. These dislocation structures were more significant in grains with ND//<110> orientations. Sample P showed higher presence of ND//<110>. Though the average misorientations of the starting samples were nearly the same, the distribution of misorientations were different. Sample P showed a peak widening which accounts for increased dislocation sub-structures primarily in the ND//<110> grains. The repeated thermal cycling incurred more defects in sample P at the expense of HAGBs. Thus a steady drop in the in the HAGB fraction was evident
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with thermal cycling reflecting through identical changes in the Ms temperature. Furthermore Sample A and P showed presence of fine recrystallized (Rex) grains with grain size varying between 3 and 5 mm. Such grains are characterized by high band contrast (white). Fig. 6a shows band contrast map obtained at 210th cycle for sample P with visibly small Rex Grains. These grains were more significant at the GBs, triple junctions and few in the grain interiors. Such grains do not show any possible signatures of undergoing phase transformation even after several thermal cycles. A phase transformation would have induced defects in these grains as discussed before. In contrast to rex grains, a grain with defect substructures as seen in Fig. 6b would show a wide scatter around the orientation in the IPF orientation triangle implying a significant orientation gradient. However small Rex Grains do not exhibit any spread of orientation implying that these grains did not transform. It is likely that dislocations will pile up around these small grains pinning the phase interfaces and thereby imparting greater resistance to phase transformation. It can therefore be concluded that Rex grains provided additional hindrance to phase transformation lowering the start transformation temperatures. The other microstructural change brought out in the present study was the stored energy during thermal cycling. According to the classical deformation theory, the stored energy is expected to increase with dislocation density, however in the present study the measured stored energy from EBSD show an inverse trend. The only plausible explanation available at this point is that the stored energy in a cyclic phase transformation process is strongly affected by the distribution and nature of defects rather than the density alone. The observed pattern of changes in the stored energy correlated to identical changes in transformational enthalpies, see Fig. 2c,d and
Fig. 5. (a) Band contrast (BC) and (b) IPF map of sample P without cycling. There were few grains primarily of near ND//<110> orientations which exhibited dislocation substructures seen in the BC map. Regions specific to these defect structures were difficult to index and showed up as dark regions or poor band contrast. White regions correspond to regions with high BC. These structures could arise as a result of previous deformation process. (c) BC and (d) IPF maps of sample P after 140th thermal cycle. Grains with ND//<110> showed increased dislocation structures.
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3b. The enthalpy of transformation during a thermoelastic martensitic transformation is the combined effect of three components; (i) the chemical molar enthalpy resulting from differences in entropy between austenite and martensite structures, i.e. the total available energy for transformation, (ii) a part of the chemical enthalpy that accounts for stored elastic energy associated to GBs and (iii) the other part, associated to defects in the microstructures. The enthalpy of both forward and reverse transformation show a decrease with thermal cycling, the drop being more significant in sample P. Fig. 7a and b shows enthalpy as a function of stored energy. No clear correlation exists between enthalpy of forward transformation and stored energy in sample A. However, sample P
shows a steady increase in the enthalpy with stored energy. The reverse transformational enthalpy showed an increase with stored energy for both samples. This of course raises a question about the possible interdependability of these quantities. Though crystallographic effects of deformation such as rolling, extrusion, and drawing processes are ignored many a times for the design of shape memory devices, the orientation of the parent austenite grains alter the initiation temperatures and stresses for martensitic phase transformation [59e62]. Few literature on martensitic transformation in steels suggest that orientation of reversed austenite during a thermal cycle process is not altered when compared to the original austenite orientation [63], the
Fig. 6. (a) BC map of sample P exhibiting fine recrystallized grains after 140th thermal cycle, encircled in white. These grains even after multiple thermal cycling did not accumulate dislocations/defects. This was verified by their high BC value. The orientations of few selected grains also did not show scatter, implying defect free structure. (b) BC map showing large grains n sample P after 140th cycle. The presence of defects increased the spread of orientation. The orientation of one such grain in the stereo triangle shows a large spread.
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Fig. 7. The enthalpy of (a) forward and (b) reverse transformation as a function of stored energy. Sample P exhibited no correlation with stored energy in the forward transformation path. An increase in enthalpy with stored energy was clear in sample A. The reverse transformation showed an increasing trend in both samples A and P.
present work clearly contradicts the findings. The thermal fatigue discussed here is a consequence of texture strengthening. Repeated thermal cycling caused g fiber strengthening in Sample A. Sample P showed texturing of ND//<110> fiber after the 210th cycle. Earlier works on texture strengthening in shape memory alloys subjected to stress-induced martensite transformation suggested rotation of parent austenite grains to favorable orientation that has the highest driving force for transformation; however no literature exists till date on the texture development as a result of thermal cycle. In absence of external loading, the possibility of crystal rotation must be ignored. Hence the texture development occurs during martensite/austenite reversion process. In general there are 24 self accommodating variants of martensite that can be formed from a single austenite during a thermally activated process under no load; however the reversion of martensite variants does not correspond to the exact parent austenite orientation. Thermal cycling introduces a complex dislocation arrangement between the self accommodating variants resulting in stress field between the plates which keeps increasing with thermal cycles. This residual stress disturbs the invariant plane stain condition to maintain the interface between parent and the product phases. In a previous study, it was reported that a phase boundary relationship (PBR) of 40 <001> between austenite and martensite is necessary to trigger the reverse transformation [64]. Any deviation from the exact PBR does not lead to the right parent austenite orientation. This deviation is largely due to residual stress that is developed as a result of incomplete strain accommodation. It could therefore be speculated that sample P that exhibited weaker Goss orientation in pre-cycled condition accumulated higher strain with successive thermal cycling primarily in the Goss ({110} <001>) oriented grains. The martensite variants originating from these orientations could not revert to the original austenite orientations due to increased stress field instead they collapsed to orientations of ND//<110> represented by the ND//<110> fiber. Such grains were also characterized by significant presence of remnant martensite that could not undergo reverse transformation to austenite as a result of residual stress which built up with cycling, see Fig. 8a,b. The fatigue effects due to g fiber texturing in sample A seems to be less severe. The underlying conclusion presented in this work is that microstructure similar to sample P exhibiting significant ND//<110> orientations are more liable to thermal fatigue. The clarification for thermal fatigue in NieTi SMAs based on few microstructural observations presented here is partly qualitative and speculative, but the results remain reproducible. A theoretical model would be required to explain the variant transformation during cycling and work is in progress.
5. Conclusions The present investigation brings out changes in the microstructures and transformation characteristics during thermal cycling in samples A and P exhibiting different starting microstructures. The conclusions are outlined as follows. The transformation temperatures (TTs) (Ms and As) decreased after multiple thermal cycles. Sample P showed a steady decrease in TT; however the change was less significant in sample A. Furthermore the transformation temperatures correlated to similar changes in the high angle grain boundary (HAGB) fractions. The decrease in HAGB fraction with thermal cycling was attributed to increased formation of dislocation or defect sub-structures in grain interiors. These defects between the interfaces of parent and product phases hindered the transformation process in successive cycles. This led to lowering of both Ms and As. Sample P showed higher rate of decrease in transformational enthalpy than sample A. The enthalpy had a strong bearing on
Fig. 8. Phase map of sample P after 210 thermal cycles. Red grains are parent austenite grains. Thermal cycling led to presence of remnant masrtensite (blue) which were significant in defect bands inside austenite grains of the ND//<110> orientations. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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the stored energy of the microstructures. Increase in stored energy related to increase in both forward and reverse transformational enthalpy. The change in stored energy has been attributed to distribution of dislocations/defects in the microstructures rather than dislocation densities alone. The observed result on calorimetric signatures of phase transformation suggests, sample P was more prone to thermal fatigue. Lastly repeated thermal cycles led to texturing in samples A and P. After prolong thermal cycles, texturing of g-fiber became pronounced in sample A. Sample P showed significant strengthening of ND//<110> fiber. ND//<110> grains exhibited remnant martensite that could not transform due to residual stresses that built up between different plate variants with thermal cycling. The observed patterns of crystallographic modifications suggests that grains with ND//<110> orientations (sample P) are more liable to thermal fatigue. Microstructures such as sample P should therefore be avoided for applications which require high fatigue resistance. Acknowledgments Authors would like to thank Prof. Supratim Ghosh, Department of Food and Bioproduct Sciences, University of Saskatchewan, Canada for extending help to carry out DSC measurements in his lab. References [1] T.W. Duerig, K.N. Melton, J.L. Proft, in: T.W. Duerig, K.N. Melton, D. Stockel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, Butterworth Heineman Ltd, Guildford (UK), 1990, p. 130. [2] N.B. Morgan, C.M. Friend, A review of shape memory stability in NiTi alloys, J. Phys. IV 11 (2001) 325e332. [3] H. Scherngell, A.C. Kneissl, Generation, development and degradation of the intrinsic two-way shape memory effect in different alloy systems, Acta Mater. 50 (2002) 327e341. [4] A.R. Pelton, V. Schroeder, M.R. Mitchell, X.-Y. Gong, M. Barney, S.W. Robertson, Fatigue and durability of Nitinol stents, J. Mech. Behav. Biomed. Mater. 1 (2008) 153e164. [5] K. Bhattacharya, Microstructure of Martensite, Oxford University Press, Oxford, 2003. [6] K. Otsuka, X. Ren, Physical metallurgy of TieNi-based shape memory alloys, Prog. Mater. Sci. 50 (2005) 511e678. [7] R. Heinen, K. Hackl, W. Windl, M.F. Wagner, Microstructural evolution during multi-axial deformation of pseudoelastic NiTi studied by first principles based micromechanical modeling, Acta Mater. 57 (2009) 3856e3867. €nemeyer, G. Eggeler, Design of a [8] M. Wagner, J. Richter, J. Frenzel, D. Gro medical non-linear drilling device: the influence of twist and wear on the fatigue behaviour of NiTi wires subjected to bending rotation, Mater. Werkst. 35 (2004) 320e325. € ckel, An overview of nitinol medical applications, [9] T. Duerig, A. Pelton, D. Sto Mater. Sci. Eng. A 273e275 (1999) 149e160. [10] A.R. Pelton, S.M. Russell, J. DiCello, The physical metallurgy of Nitinol for medical applications, J. Metal. 55 (2003) 33e37. € ver, M. Wagner, G. Eggeler, [11] L. Krone, J. Mentz, M. Bram, H.P. Buchkremer, D. Sto €ller, S. Esenwein, G. Muhr, O. Prymak, D. Christ, S. Reese, D. Bogdanski, M. Ko M. Epple, The potential of powder metallurgy for the fabrication of biomaterials on the basis of nickel-titanium: a case study with a staple showing shape memory behavior, Adv. Eng. Mater. 7 (2005) 613e619. [12] J. Mentz, J. Frenzel, M.F.-X. Wagner, K. Neuking, G. Eggeler, H.P. Buchkremer, €ver, Powder metallurgical processing of NiTi shape memory alloys with D. Sto elevated transformation temperatures, Mater. Sci. Eng. A 491 (2008) 270e278. [13] W. Huang, On the selection of shape memory alloys for actuators, Mater. Des. 23 (2002) 11e19. [14] C. Großmann, J. Frenzel, V. Sampath, T. Depka, A. Oppenkowski, Ch Somsen, K. Neuking, W. Theisen, G. Eggeler, Processing and property assessment of NiTi and NiTiCu shape memory actuator springs, Mater. Werkst. 39 (2008) 499e510. [15] C. Großmann, J. Frenzel, V. Sampath, T. Depka, G. Eggeler, Elementary transformation and deformation processes and the cyclic stability of NiTi and NiTiCu shape memory spring actuators, Metall. Mater. Trans. A 40 (2009) 2530e2544. [16] G.B. Olson, M. Cohen, Dislocation theory of martensitic transformations, in: F.R.N. Nabarro (Ed.), Dislocations in Solids, North-Holland, Amsterdam, 1986, pp. 384e387.
[17] B. Kockar, I. Karaman, J.I. Kim, Y. Chumlyakov, A method to enhance cyclic reversibility of NiTiHf high temperature shape memory alloys, Scr. Mater. 54 (2006) 2203e2208. [18] C.-H. Li, L.-J. Chiang, Y.-F. Hsu, W.-H. Wang, Cold Rolling-Induced Multistage Transformation in Ni-Rich NiTi Shape Memory Alloys, Mater. Trans. 49 (9) (2008) 2136e2140. [19] Y. Liu, J. Laeng, T.V. Chin, T.-H. Nam, Effect of incomplete thermal cycling on the transformation behaviour of NiTi, Mater. Sci. Eng. A 435e436 (2006) 251e257. [20] K. Gall, J. Tyber, G. Wilkesanders, S.W. Robertson, R.O. Ritchie, H.J. Maier, Effect of microstructure on the fatigue of hot-rolled and cold-drawn NiTi shape memory alloys, Mater. Sci. Eng. A 486 (2008) 389e403. [21] C. Urbina, S.D. la Flor, F. Ferrando, Effect of thermal cycling on the thermomechanical behaviour of NiTi shape memory alloys, Mater. Sci. Eng. A 501 (2009) 197e206. [22] S. Miyazaki, Y. Igo, K. Otsuka, Effect of thermal cycling on the transformation temperatures of TieNi alloys, Acta Metal. 34 (1986) 2045e2051. [23] P.G. McCormick, Y. Liu, Thermodynamic analysis of the martensitic transformation in NiTidII. Effect of transformation cycling, Acta Metall. Mater 42 (1994) 2407e2413. [24] A.R. Pelton, G.H. Huang, P. Moinec, R. Sinclair, Effects of thermal cycling on microstructure and properties in Nitinol, Mater. Sci. Eng. A 532 (2012) 130e138. [25] R. Basu, L. Jain, B.C. Maji, M. Krishnan, K.V. Mani Krishna, I. Samajdar, P. Pant, Origin of microstructural irreversibility in NidTi based shape memory alloys during thermal cycling, Metall. Mater. Trans. A 43 (2012) 1277e1287. [26] Y. Liu, P.G. McCormick, Factor influencing the development of two-way shape memory in NiTi, Acta Metall. Mater. 38 (1990) 1321e1326. [27] S.K. Giri, M. Krishnan, U. Ramamurty, Enhancement of fatigue life of NieTieFe shape memory alloys by thermal cycling, Mater. Sci. Eng. 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Wagner, Structural and functional fatigue of NiTi shape memory alloys, Mater. Sci. Eng. A 378 (2004) 24e33. [35] N.B. Morgan, C.M. Friend, A review of shape memory stability in NiTi alloys, J. Phys. IV C 11 (2001) 325. [36] E. Hornbogen, Thermo-mechanical fatigue of shape memory alloys, J. Mater. Sci. 39 (2004) 385e399. [37] E. Horbogen, A. Heckmann, Improved fatigue resistance of pseudo-elastic NiTi alloys by thermo-mechanical treatment, Mater. Wiss. Werk. 34 (2003) 464e468. [38] D.Y. Li, X.F. Wu, The effect of stress on soft modes for the phase transformation in a Ti-Ni alloy: II. Effects of ageing and thermal cycling on the phase transformation, Philos. Mag. A 63 (1991) 603e616. [39] J. Perkins, Shape memory behavior and thermoelastic martensitic transformations, Mater. Sci. Eng. 51 (1981) 181e192. [40] H. Funakubo, Shape Memory Alloys, Gordon and Breach Science Publishers, New York, 1987. [41] K.N. Melton, in: K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge University Press, Cambridge, 1998, p. 220. [42] V.B. Krishnan, Low Temperature NiTiFe Shape Memory Alloys: Actuator Engineering and Investigation of Deformation Mechanisms Using In-situ Neutron Diffraction at Los Alamos National Laboratory, University of Central Florida Orlando, 2007 (PhD thesis). [43] R. Basu, J. Szpunar, M. Eskandari, M.A. M-Bonab, Microstructural investigation on marforming and conventional cold deformation in NieTieFe-based shape memory alloys, Int. J. Mater. Res. 106 (8) (2015) 852e862, http://dx.doi.org/ 10.3139/146.111252. [44] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, second ed., Elsevier, Oxford, 2004. [45] K. Yamada, C.C. Koch, The influence of mill energy and temperature on the structure of the TiNi intermetallic after mechanical attrition, J. Mater. Res. 8 (1993) 1317. [46] H.R. Wenk, P. Van Houtte, Texture and anisotropy, Rep. Prog. Phys. 67 (8) (2004) 1367e1428. 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Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Role of microstructure on phase transformation behavior in NieTieFe shape memory alloys during thermal cycling Ritwik Basu a, *, M.A. Mohtadi-Bonab b, Xu Wang b, Mostafa Eskandari b, c, Jerzy A. Szpunar b a b c
Department of Mechanical Engineering, The NORTHCAP University, Gurgaon, 122017, India Department of Mechanical Engineering, University of Saskatchewan, Saskatoon, S7N5A9, Canada Department of Materials Science & Engineering, Faculty of Engineering, Shahid Chamran University, Ahvaz, Iran
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 July 2015 Received in revised form 27 August 2015 Accepted 28 August 2015 Available online 1 September 2015
Two different microstructures of NieTieFe shape memory alloys processed through different thermomechanical treatments with nearly similar grain size and in-grain misorientation but different crystallographic textures were subjected to series of thermal cycles without external loading. The microstructures and the phase transformation behavior of the samples were examined after every seventy cycles. The experiment involved treating the samples with liquid nitrogen (LN2) for complete martensitic (B190 ) transformation and then heating it back to parent austenite (B2) condition. Thermal cycling introduced significant differences in microstructural parameters especially the grain boundary nature, stored elastic energy and the misorientation or defect densities. These microstructural alterations during thermal cycling were related to changes in transformation temperatures and enthalpy. Thermal cycling also brought out changes in crystallographic orientations of austenite grains. The present study aims to address the role of microstructure on the thermal fatigue behavior in NieTi based shape memory alloys during cyclic transformations. © 2015 Elsevier B.V. All rights reserved.
Keywords: NieTieFe shape memory alloys Thermal cycling Electron back scattered diffraction (EBSD) Texture X-ray diffraction (XRD)
1. Introduction Shape memory alloys of NieTi, commonly called as NITINOL possess the two most important characteristics, the shape memory and the pseudoelasticity, that makes it an important candidate for a broad range of engineering applications [1]. In many of these applications the components fabricated from these alloys are subjected to repeated thermal or mechanical transformation cycles. Examples include, (1) actuators that are designed to experience repeated thermal cycles throughout a wide range of transformation temperatures that are typically determined by the start and finish temperature of austenite4martensite transformation [2,3] and (2) endovascular stents of self-expanding types that undergo several pseudoelastic cycles that results from cardiac pressure pulses [4]. The thermal cycles could be performed either under external loading or zero load conditions. Both thermal and mechanical cycles involve a reversible, first-order, cubic-to-monoclinic (B2/B190 ) thermoelastic martensitic transformation [5,6]. The B2
* Corresponding author. E-mail address: [email protected] (R. Basu). http://dx.doi.org/10.1016/j.jallcom.2015.08.239 0925-8388/© 2015 Elsevier B.V. All rights reserved.
and B190 phases have been referred as austenite and martensite respectively in the present study. The temperatures and the stresses to initiate the transformation from B2/B190 are significantly affected during successive thermal and mechanical cycling [7e13]. Thermoelastic martensites differ from ferrous martensites through presence of internal twin variants that accommodates elastic strains during cooling from high temperature B2 phase. Further deformation causes the detwinning of these micro-twin structures leading to several new martensite variants. Upon heating these different variants revert back to the original high temperature B2 structure. The shape accommodation of martensite in steels occurs by classical slip mechanism and this is irreversible in nature [14]. The effects of thermal cycling on the transformation characteristics and shape memory properties have been studied in the past by several researchers [15e30]. Some of the important findings are summarized as follows: (1) Thermal cycling in absence of external stresses tends to decrease the start temperature of both marteniste and austenite; the Ms and As respectively. This is also followed by widening of thermal hysteresis [22,23]. In addition, repeated cycling show characteristic R-phase (intermediate phase) transformation due to precipitation of the metastable intermetallic
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Ni4Ti3 prior to the martensitic transformation [22,26]. (2) Thermal cycling under external stress increases the start temperatures of transformation [26]. This behavior is attributed partly due to the alignment of dislocation structures in the parent B2 matrix. However a complete understanding on the transformation characteristics of these alloys during cycling under constant stress is still not comprehensible. Thus it can be stated that the observed changes in the functional characteristics as a result of cycling are largely associated to subtle changes in the microstructure [22,30,31]. The available literature on improved cyclic stability of NieTi SMAs reports precipitation strengthening and partial recovery annealing after cold working [22,32,33] as important factors. Although thermal cycling of the SMA components in absence of external stresses introduces functional fatigue [34,35], it has also been shown that these components subjected to thermal cycles at zero stress prior to application (pre-cycling) can improve the fatigue resistance largely in these alloys [36e38]. The dislocation accumulation as a result of precycling facilitates hardening of the thermally cycled alloy, thus improving the shape memory effect. The available literature on thermal cycling of shape memory alloys with or without external loading discusses the dislocations generated during thermal cycling [22,39]. A gap remains in deriving a correlation in the observed microstructural quantities to the cyclic transformation pattern. The present study brings out a relationship between microstructures to the calorimetric signatures of phase transformation (B2/B190 ) in NieTieFe shape memory alloy during stress-free thermal cycling. The study further attempts to addresses the microstructural causes of thermal fatigue. The microstructure and texture investigation have been analyzed through electron back scattered diffraction (EBSD) and x-ray diffraction (XRD) techniques and calorimetric measurements were obtained through differential scanning calorimetry (DSC). The study presents a systematic investigation on the role of grain boundary (GB) characteristics, misorientation, stored energy and crystallographic texture on the transformation parameters; the temperature of transformation and the enthalpy. It has been shown that temperatures of both forward and reverse transformation bear a strong correlation with the fraction of high angle grain boundaries (HAGBs). Increase or decrease in the HAGB fraction reflected through similar changes in the martensite start temperature of transformation. The enthalpy of transformation was a clear function of the stored energy. Thermal cycling caused increase in the ingrain misorientation which is a signature of increased defect density in the samples. It was also brought out in the study that crystallographic orientation of parent austenite grains in the microstructure decides the thermal fatigue response in shape memory materials. Identifying these uncertainties in the transformation sequence and its correlation to the microstructure of the alloy will form a baseline study for manufacturing of actuator devices. Ternary Ni47eTi50eFe3 alloy has been used in the present research as it is widely used in many engineering applications such as heat shrinkable couplings, self-healing gaskets, safety valves, cryogenic thermal switches etc. [40e42].
annealed sample here onwards has been referred as P. Another 10% reduction in thickness was attained in sample P through cold rolling followed by final annealing at 730 C for 0.5 h. This sample here has been referred as A. Details about the microstructure developments through aforementioned deformation routes is beyond the scope of present research; a separate field of study reported in Ref. [43]. Samples for EBSD and bulk crystallographic texture measurements have been prepared thorough conventional metallographic polishing techniques to bring down the thickness of the sample to half. The surface of interest in this study was chosen as the rolling plane. The sample surface was further treated with colloidal silica (0.04 mm) solution to generate a strain free surface for carrying out microstructural measurements [43]. For the calorimetric measurements, specimens were cut using slow speed diamond wafer blade to avoid deformation in the microstructure. The thermal cycling was conducted by holding the samples in liquid nitrogen for martensitic transformation and then heating back for reverse austenite transformation. The microstructures and calorimetric properties were measured after every seventy thermal cycles. 2. 2. Experimental procedures Characterization through EBSD techniques were performed in a SU 6600 Hitachi field emission scanning electron microscope (FESEM) coupled with an Oxford Instruments Norlys Nano EBSD detector. The AZTEC 2.0 data acquisition software was used for acquiring diffraction patterns from the samples which was synchronized with the beam scan. An acceleration voltage of 20 keV was maintained during the scan. The number of bands used for pattern identification were set to six and the data acquisition rate of 20 frames/s was used. A typical step size of 0.35 mm was used for EBSD measurements. The scan parameters and SEM conditions were kept identical for all scans performed in this research. Oxford Instrument's Channel 5 post processing software was used for processing the raw EBSD data. Grain boundaries (GBs) were identified as continuous misorientation region, where misorientation angle > 5 were considered. GBs greater that 15 were considered as high angle grain boundaries (HAGBs). Grain size diameter was determined from diameter of a circle that has the same area of the grain. The area of a grain equals to the number of data points in the grain multiplied by their pixel size. In-grain misorientations were measured as average point-to-point misorientation inside an identified grain: local average misorientation [43]. Local stored energy associated to grain boundary was estimated using the misorientation criterion. Here boundary misorientations between 5 and 15 are categorized as low-angle boundaries (LAGBs) while misorientations 15 are high-angle boundaries (HAGBs). The energy contribution from grain boundary misorientation is estimated from the ReadeShockley equation [44] such that:
gðqÞ ¼ gm
2. Experimental 2.1. Materials and preparations
q qm
q 1 ln ; q qm qm
and
gðqÞ ¼ gm ; q qm Alloy of Ni47eTi50eFe3 was supplied by Hlmet Co., Ltd, China in hot rolled conditions. Further cold rolling was carried out in a lab scale rolling mill to 50% reduction in thickness in three passes with intermediate annealing at 800 C for 25 min followed by final annealing treatments at the same temperature for 1.5 h. The post
(1)
(2)
where gm ¼ 0.706 Jm2 [45] is the energy per unit area of a HAGB, is the boundary misorientation and qm ¼ 15 is the misorientation angle above which the energy per unit area is independent of misorientation angle. The average boundary energy is calculated
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from equations (1) and (2) considering all the boundary misorientations between 5 q 62.8 via summation:
g¼
62:8 X
½gðqÞf ðqÞ
5
(3)
461
where f(q) represents the boundary fraction for a given misorientation. The stored energy per unit volume due to a grain boundary dislocation (Eb) is obtained by multiplying the average energy per unit boundary area (g) to the area per unit volume (SV d 3 ) ECD where, dECD is the average grain size measured from equivalent circle diameter considering all misorientations between
Fig. 1. Microstructures represented through inverse pole figure (IPF) map of (a) sample A and (b) sample P under pre-cycled conditions. The IPF color triangle presents important orientations corresponding to the colors. Sample A shows increased presence of ND//<111> grains. Sample P had ND//<110> grains along with few other orientations. Orientation distribution functions (ODFs) of 42 ¼ 0 and 45 sections estimated from EBSD measurements for (c) sample A and (d) sample P shows clear differences in crystallographic textures. The important texture components are illustrated in the ODF plots. (e) Grain size and (f) misorientation distributions were nearly identical for samples A and P. The inset shows the most probable values.
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5 q 62.8 . Thus the stored energy due to boundary dislocations (Eb) can thus be represented in the following form:
Eb ¼ SV g ¼
3g dECD
(4)
It is to be noted that Eb values do not account for in-grain misorientations or dislocations.
Measurements for bulk crystallographic textures were performed in a Bruker AXS General Area Detector system equipped with a Cr-ka radiation. Two pole figures (PF), (110) and (200) were measured for calculation of orientation distribution function (ODF) using Textool programs [46]. The long wavelength of Cr-Ka X-rays imposed difficulties in measuring a third pole figure for ODF calculations. The Cr-Ka X-rays offers advantage over Co or Cu X-rays due to much lower defocus effect because pole figures are
Fig. 2. Differential scanning calorimetric (DSC) plots showed characteristic transformation temperatures and the enthalpy of transformation for (a) sample A and (b) sample P. (c) Martensite (Ms) and (d) austenite start (As) temperatures as a function of thermal cycle show a drop; drop being more clear in sample P. Enthalpy of (e) forward and (f) reverse transformation as a function of thermal cycle is presented. Sample P exhibited a steady decrease in transformational enthalpy. Changes were less significant in sample A. A drop in transformation temperature and enthalpy with multiple thermal cycle is a valid signature of thermal fatigue. Error bars represent typical accuracy of estimated enthalpy and transformation temperatures within 5% of the mean values.
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measured above 2q ¼ 70 . The results on texture for austenite NieTi phase was verified by comparing with Cu (3 PFs). The texture results were nearly identical. The crystal orientations are represented through Euler angles (41, F, 42) [47]. The sample reference is chosen as the rolling direction (RD) and the normal to the rolling plane (ND). For measurement of initiation temperatures and enthalpies of transformations, a TA Instrument's Q2000 differential scanning calorimetry (DSC) unit equipped with a LN2 cooling attachment was used. An inert atmosphere was maintained by using Ar. The heatingecooling cycles for all thermally cycled specimens were performed in the temperature span of 50 C to 50 C at the rate of 3 C/min. 3. Results Fig. 1(aef) bring out comprehensible differences in the microstructures of the samples A and P under investigation prior to subjecting to thermal cycles. The EBSD measured inverse pole figure (IPF) colored orientation maps illustrated in Fig. 1a and b shows distinctly different microstructures. The colors in the IPF map represent different crystallographic orientation of individual grains parallel to the normal direction (ND) of the samples. Readers are advised to refer the standard colored stereo-triangle for interpretation of orientations. The textures represented through the orientation distribution function (ODF) sections of 42 ¼ 0 and 45 in Fig. 1c shows clear presence of g-fiber orientations with highest
463
intensities around {111} <123> and {111} <112> in sample A. The ODF plot for Sample P in Fig. 1d shows a relatively high intensity around {211} <111> and a partial spread of orientation along the ND//<110> fiber. However the XRD estimated ODF for sample P also shows a week spread of Goss and cube texture which will be discussed shortly. The grain size and in-grain misorientation distributions with their most probable values presented in Fig. 1 d and e respectively were nearly identical for both the samples A and P. The results on the calorimetric signatures of phase transformation are illustrated in Fig. 2a-f. Fig. 2a and b are representative DSC plots for samples A and P respectively subjected to thermal cycles. The austenite4martensite transformation show a decrease in martensite (Ms) and austenite start (As) temperatures with repeated thermal cycling, see Fig. 2c and d. . The decreasing trend of the start temperatures during thermal cycling has been reported by several researchers in the past [48e50]. This has been attributed primarily to the introduction of transformation-induced dislocations [48] that tend to increase the elastic strain energy in each cycle leading to stabilization of austenite [51,52]. The rate of decrease in transformation temperatures (TT) measured in sample P is higher than sample A. Sample A exhibits a nearly consistent Ms except that a dip is observed at 70th thermal cycle. Significant changes in the enthalpy of transformation during thermal cycles were noticeable. The enthalpy data of forward and reverse transformation as a function of temperature for each cycle is presented in Fig. 2e and f. The enthalpies of both transformation show a monotonic decrease with increasing number of cycles. The effect is
Fig. 3. Microstructural changes in samples A and P as a function of thermal cycles. (a) High angle grain boundary (HAGB) fraction, (b) stored energy (SE) and (c) most probable misorientation are plotted against thermal cycles. Both HAGB and SE show a decrease with increase in thermal cycles. Sample P showed a steady drop in the values. Misorientation showed an increasing trend.
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more pronounced in sample P. The enthalpy of marteniste transformation in sample A shows no significant change. Lower transformation enthalpy is a valid signature of suppression of austenite4martensite transformation. The results on TTs and enthalpy suggest that sample P is more susceptible to functional fatigue and is not an ideal candidate for shape memory applications. Fig. 3a-c discusses the microstructural developments in terms of boundary fraction, stored energy of deformation and the most probable value misorientation at every stage of thermal cycling estimated from EBSD measurements. The measurements were obtained from three EBSD scans performed on the same sample after every thermal cycle. The standard deviations estimated from multiple EBSD measurements are represented through error bars. The fraction of HAGB and stored energy show a decrease with number of thermal cycles. Misorientation showed an opposite trend. More importantly the Ms Temperature is a clear function of the HAGB fraction, i.e an increase or decrease in the Ms temperature correlated well with the measured HAGB fraction. The enthalpy of transformation on the other hand bears a strong dependence on the stored energy estimated from EBSD. Fig. 4a and b brings out the changes in the austenite textures of samples A and P during thermal cycling represented through XRD measured ODFs. The texture ODFs show a wide scatter of orientations around the g-fiber (ND//<111>) line with the highest intensity located in the vicinity of {322} <258> and {322} <236> in the sample A in pre-cycled condition brought out in the 42 ¼ 45
section. The EBSD measured ODF of sample A in Fig. 1c presents similar results. With successive thermal cycling the g-fiber became pronounced along with strong presence of near {100} <011> orientations. The starting microstructure of Sample P exhibits a weak Goss {110} <001> followed by two secondary peaks around {100} <001>, cube and {211}<111> orientations. However the intensities around Goss orientation were too low to be detected in EBSD. Thermal cycling of sample P resulted in strengthening of ND// <110> fiber suggesting that these orientations have grown at the expense of Goss orientation. 4. Discussions The transformation of austenite to martensite proceeds with formation of glissile dislocation interfaces [51]. Any deviation in the glissile nature of the interface would retard the transformation. Several authors have established the presence of transformation induced dislocations that assist the formation of martensite nuclei [53e58]. However, the dislocation density progressively increases with thermal cycling and the glissile relationship between austenite and martensite is gradually lost. There exists no unified theory to explain the mechanism of dislocation accumulation during repeated thermal cycling. The dislocations that may be present in the material prior to thermal cycling as a result of deformation could act as a source of martensite nuclei formation, however with successive cycles additional dislocations contribute which could arise as a result of coalescence of different plates of
Fig. 4. Effects of thermal cycling on the texture developments. The X-ray estimated textures for the austenite phase are represented in standard 42 ¼ 0 and 45 ODF section. ODF plots of (a) sample A and (b) sample P showed differences in their initial and final textures. Sample A showed characteristic g-fiber texture which got strengthened with thermal cycling. Sample P shows development of ND//<110> fiber with increased thermal cycling. The maximum ODF intensities are clearly indicated in the figures.
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martensite variants, stress relaxation due to change in volume during transformation or interaction of dislocations with different martensite plate groups [41]. The transformation induced dislocations impede further transformation of austenite to martensite. The measured average misorientation in Fig. 3c show an increase with thermal cycles which scales nearly with defect densities. A decrease in the Ms Temperature implies difficulties in transformation. The increased presence of dislocations with thermal cycling is depicted in the band contrast map obtained from two scans from the same sample with and without cycling, see Fig. 5aed. The pre-cycled sample shows dislocations in the form of banded structures in few different grains. These different structures were apparent primarily in the near ND//<110> grains as seen in the IPF maps. The dark regions are regions of poor band contrast referring to low indexed points as a result of strain accumulation. There were also grains which were free of such micro features. Significant differences in the sub-structures were brought out in the 140th cycle with clear presence of increased dislocation arrays and bands. The GBs shown in blue are HAGBs. These structures show misorientation varying from 2 to 5 . It can be said that the structure formation occurs at the expense of the HAGBs. This led to drop in the fraction of HAGBs with successive thermal cycling which coincided with similar change in the Ms temperature. These dislocation structures were more significant in grains with ND//<110> orientations. Sample P showed higher presence of ND//<110>. Though the average misorientations of the starting samples were nearly the same, the distribution of misorientations were different. Sample P showed a peak widening which accounts for increased dislocation sub-structures primarily in the ND//<110> grains. The repeated thermal cycling incurred more defects in sample P at the expense of HAGBs. Thus a steady drop in the in the HAGB fraction was evident
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with thermal cycling reflecting through identical changes in the Ms temperature. Furthermore Sample A and P showed presence of fine recrystallized (Rex) grains with grain size varying between 3 and 5 mm. Such grains are characterized by high band contrast (white). Fig. 6a shows band contrast map obtained at 210th cycle for sample P with visibly small Rex Grains. These grains were more significant at the GBs, triple junctions and few in the grain interiors. Such grains do not show any possible signatures of undergoing phase transformation even after several thermal cycles. A phase transformation would have induced defects in these grains as discussed before. In contrast to rex grains, a grain with defect substructures as seen in Fig. 6b would show a wide scatter around the orientation in the IPF orientation triangle implying a significant orientation gradient. However small Rex Grains do not exhibit any spread of orientation implying that these grains did not transform. It is likely that dislocations will pile up around these small grains pinning the phase interfaces and thereby imparting greater resistance to phase transformation. It can therefore be concluded that Rex grains provided additional hindrance to phase transformation lowering the start transformation temperatures. The other microstructural change brought out in the present study was the stored energy during thermal cycling. According to the classical deformation theory, the stored energy is expected to increase with dislocation density, however in the present study the measured stored energy from EBSD show an inverse trend. The only plausible explanation available at this point is that the stored energy in a cyclic phase transformation process is strongly affected by the distribution and nature of defects rather than the density alone. The observed pattern of changes in the stored energy correlated to identical changes in transformational enthalpies, see Fig. 2c,d and
Fig. 5. (a) Band contrast (BC) and (b) IPF map of sample P without cycling. There were few grains primarily of near ND//<110> orientations which exhibited dislocation substructures seen in the BC map. Regions specific to these defect structures were difficult to index and showed up as dark regions or poor band contrast. White regions correspond to regions with high BC. These structures could arise as a result of previous deformation process. (c) BC and (d) IPF maps of sample P after 140th thermal cycle. Grains with ND//<110> showed increased dislocation structures.
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3b. The enthalpy of transformation during a thermoelastic martensitic transformation is the combined effect of three components; (i) the chemical molar enthalpy resulting from differences in entropy between austenite and martensite structures, i.e. the total available energy for transformation, (ii) a part of the chemical enthalpy that accounts for stored elastic energy associated to GBs and (iii) the other part, associated to defects in the microstructures. The enthalpy of both forward and reverse transformation show a decrease with thermal cycling, the drop being more significant in sample P. Fig. 7a and b shows enthalpy as a function of stored energy. No clear correlation exists between enthalpy of forward transformation and stored energy in sample A. However, sample P
shows a steady increase in the enthalpy with stored energy. The reverse transformational enthalpy showed an increase with stored energy for both samples. This of course raises a question about the possible interdependability of these quantities. Though crystallographic effects of deformation such as rolling, extrusion, and drawing processes are ignored many a times for the design of shape memory devices, the orientation of the parent austenite grains alter the initiation temperatures and stresses for martensitic phase transformation [59e62]. Few literature on martensitic transformation in steels suggest that orientation of reversed austenite during a thermal cycle process is not altered when compared to the original austenite orientation [63], the
Fig. 6. (a) BC map of sample P exhibiting fine recrystallized grains after 140th thermal cycle, encircled in white. These grains even after multiple thermal cycling did not accumulate dislocations/defects. This was verified by their high BC value. The orientations of few selected grains also did not show scatter, implying defect free structure. (b) BC map showing large grains n sample P after 140th cycle. The presence of defects increased the spread of orientation. The orientation of one such grain in the stereo triangle shows a large spread.
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Fig. 7. The enthalpy of (a) forward and (b) reverse transformation as a function of stored energy. Sample P exhibited no correlation with stored energy in the forward transformation path. An increase in enthalpy with stored energy was clear in sample A. The reverse transformation showed an increasing trend in both samples A and P.
present work clearly contradicts the findings. The thermal fatigue discussed here is a consequence of texture strengthening. Repeated thermal cycling caused g fiber strengthening in Sample A. Sample P showed texturing of ND//<110> fiber after the 210th cycle. Earlier works on texture strengthening in shape memory alloys subjected to stress-induced martensite transformation suggested rotation of parent austenite grains to favorable orientation that has the highest driving force for transformation; however no literature exists till date on the texture development as a result of thermal cycle. In absence of external loading, the possibility of crystal rotation must be ignored. Hence the texture development occurs during martensite/austenite reversion process. In general there are 24 self accommodating variants of martensite that can be formed from a single austenite during a thermally activated process under no load; however the reversion of martensite variants does not correspond to the exact parent austenite orientation. Thermal cycling introduces a complex dislocation arrangement between the self accommodating variants resulting in stress field between the plates which keeps increasing with thermal cycles. This residual stress disturbs the invariant plane stain condition to maintain the interface between parent and the product phases. In a previous study, it was reported that a phase boundary relationship (PBR) of 40 <001> between austenite and martensite is necessary to trigger the reverse transformation [64]. Any deviation from the exact PBR does not lead to the right parent austenite orientation. This deviation is largely due to residual stress that is developed as a result of incomplete strain accommodation. It could therefore be speculated that sample P that exhibited weaker Goss orientation in pre-cycled condition accumulated higher strain with successive thermal cycling primarily in the Goss ({110} <001>) oriented grains. The martensite variants originating from these orientations could not revert to the original austenite orientations due to increased stress field instead they collapsed to orientations of ND//<110> represented by the ND//<110> fiber. Such grains were also characterized by significant presence of remnant martensite that could not undergo reverse transformation to austenite as a result of residual stress which built up with cycling, see Fig. 8a,b. The fatigue effects due to g fiber texturing in sample A seems to be less severe. The underlying conclusion presented in this work is that microstructure similar to sample P exhibiting significant ND//<110> orientations are more liable to thermal fatigue. The clarification for thermal fatigue in NieTi SMAs based on few microstructural observations presented here is partly qualitative and speculative, but the results remain reproducible. A theoretical model would be required to explain the variant transformation during cycling and work is in progress.
5. Conclusions The present investigation brings out changes in the microstructures and transformation characteristics during thermal cycling in samples A and P exhibiting different starting microstructures. The conclusions are outlined as follows. The transformation temperatures (TTs) (Ms and As) decreased after multiple thermal cycles. Sample P showed a steady decrease in TT; however the change was less significant in sample A. Furthermore the transformation temperatures correlated to similar changes in the high angle grain boundary (HAGB) fractions. The decrease in HAGB fraction with thermal cycling was attributed to increased formation of dislocation or defect sub-structures in grain interiors. These defects between the interfaces of parent and product phases hindered the transformation process in successive cycles. This led to lowering of both Ms and As. Sample P showed higher rate of decrease in transformational enthalpy than sample A. The enthalpy had a strong bearing on
Fig. 8. Phase map of sample P after 210 thermal cycles. Red grains are parent austenite grains. Thermal cycling led to presence of remnant masrtensite (blue) which were significant in defect bands inside austenite grains of the ND//<110> orientations. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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the stored energy of the microstructures. Increase in stored energy related to increase in both forward and reverse transformational enthalpy. The change in stored energy has been attributed to distribution of dislocations/defects in the microstructures rather than dislocation densities alone. The observed result on calorimetric signatures of phase transformation suggests, sample P was more prone to thermal fatigue. Lastly repeated thermal cycles led to texturing in samples A and P. After prolong thermal cycles, texturing of g-fiber became pronounced in sample A. Sample P showed significant strengthening of ND//<110> fiber. ND//<110> grains exhibited remnant martensite that could not transform due to residual stresses that built up between different plate variants with thermal cycling. The observed patterns of crystallographic modifications suggests that grains with ND//<110> orientations (sample P) are more liable to thermal fatigue. Microstructures such as sample P should therefore be avoided for applications which require high fatigue resistance. Acknowledgments Authors would like to thank Prof. Supratim Ghosh, Department of Food and Bioproduct Sciences, University of Saskatchewan, Canada for extending help to carry out DSC measurements in his lab. References [1] T.W. Duerig, K.N. Melton, J.L. Proft, in: T.W. Duerig, K.N. Melton, D. Stockel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, Butterworth Heineman Ltd, Guildford (UK), 1990, p. 130. [2] N.B. Morgan, C.M. Friend, A review of shape memory stability in NiTi alloys, J. Phys. IV 11 (2001) 325e332. [3] H. Scherngell, A.C. Kneissl, Generation, development and degradation of the intrinsic two-way shape memory effect in different alloy systems, Acta Mater. 50 (2002) 327e341. [4] A.R. Pelton, V. Schroeder, M.R. Mitchell, X.-Y. Gong, M. Barney, S.W. Robertson, Fatigue and durability of Nitinol stents, J. Mech. Behav. Biomed. Mater. 1 (2008) 153e164. [5] K. Bhattacharya, Microstructure of Martensite, Oxford University Press, Oxford, 2003. [6] K. Otsuka, X. Ren, Physical metallurgy of TieNi-based shape memory alloys, Prog. Mater. Sci. 50 (2005) 511e678. [7] R. Heinen, K. Hackl, W. Windl, M.F. Wagner, Microstructural evolution during multi-axial deformation of pseudoelastic NiTi studied by first principles based micromechanical modeling, Acta Mater. 57 (2009) 3856e3867. €nemeyer, G. Eggeler, Design of a [8] M. Wagner, J. Richter, J. Frenzel, D. Gro medical non-linear drilling device: the influence of twist and wear on the fatigue behaviour of NiTi wires subjected to bending rotation, Mater. Werkst. 35 (2004) 320e325. € ckel, An overview of nitinol medical applications, [9] T. Duerig, A. Pelton, D. Sto Mater. Sci. Eng. A 273e275 (1999) 149e160. [10] A.R. Pelton, S.M. Russell, J. DiCello, The physical metallurgy of Nitinol for medical applications, J. Metal. 55 (2003) 33e37. € ver, M. Wagner, G. Eggeler, [11] L. Krone, J. Mentz, M. Bram, H.P. Buchkremer, D. Sto €ller, S. Esenwein, G. Muhr, O. Prymak, D. Christ, S. Reese, D. Bogdanski, M. Ko M. Epple, The potential of powder metallurgy for the fabrication of biomaterials on the basis of nickel-titanium: a case study with a staple showing shape memory behavior, Adv. Eng. Mater. 7 (2005) 613e619. [12] J. Mentz, J. Frenzel, M.F.-X. Wagner, K. Neuking, G. Eggeler, H.P. Buchkremer, €ver, Powder metallurgical processing of NiTi shape memory alloys with D. Sto elevated transformation temperatures, Mater. Sci. Eng. A 491 (2008) 270e278. [13] W. Huang, On the selection of shape memory alloys for actuators, Mater. Des. 23 (2002) 11e19. [14] C. Großmann, J. Frenzel, V. Sampath, T. Depka, A. Oppenkowski, Ch Somsen, K. Neuking, W. Theisen, G. Eggeler, Processing and property assessment of NiTi and NiTiCu shape memory actuator springs, Mater. Werkst. 39 (2008) 499e510. [15] C. Großmann, J. Frenzel, V. Sampath, T. Depka, G. Eggeler, Elementary transformation and deformation processes and the cyclic stability of NiTi and NiTiCu shape memory spring actuators, Metall. Mater. Trans. A 40 (2009) 2530e2544. [16] G.B. Olson, M. Cohen, Dislocation theory of martensitic transformations, in: F.R.N. Nabarro (Ed.), Dislocations in Solids, North-Holland, Amsterdam, 1986, pp. 384e387.
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