The effect of training on two-way shape memory effect of binary NiTi and NiTi based ternary high temperature shape memory alloys

Materials Science & Engineering A 560 (2013) 653–666

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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

The effect of training on two-way shape memory effect of binary NiTi and NiTi based ternary high temperature shape memory alloys K.C. Atli a, I. Karaman a,b,n, R.D. Noebe c, D. Gaydosh c a

Department of Mechanical Engineering, Texas A&M University, College Station, TX 77843, USA Materials Science and Engineering Program, Texas A&M University, College Station, TX 77843, USA c NASA Glenn Research Center, Structures & Materials Division, Cleveland, OH 44135, USA b

a r t i c l e i n f o

abstract

Article history: Received 3 August 2012 Received in revised form 2 October 2012 Accepted 4 October 2012 Available online 12 October 2012

The propensity for various high-temperature shape memory alloys (HTSMA), i.e., Ni28.5Ti50.5Pt21, Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5, to exhibit two-way shape memory effect (TWSME) was compared to that of a conventional binary Ni49.9Ti50.1 shape memory alloy (SMA). Thermomechanical training in the form of thermal cycling under constant stress levels was employed to induce two-way shape memory behavior in the various materials. The resulting TWSME was characterized for its magnitude and stability under stress-free conditions, while parameters such as training stress and upper cycle temperature during training were investigated for their influence on this phenomenon. For Ni49.9Ti50.1, a negative correlation was found between an increasing training stress, from 80 MPa to 200 MPa, and the magnitude of the resulting TWSM strain, while a positive correlation was observed for Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5. The stability of the TWSME for the Ni49.9Ti50.1, measured by the strain evolution of the cold (martensitic) and hot (austenitic) shapes of the samples upon stress-free thermal cycling, was found to depend on the stress and temperature interval during training. Conversely, the stability of the NiTiPd based HTSMAs was much greater and less sensitive to these parameters over the stress and temperature intervals investigated. No TWSME was seen in Ni28.5Ti50.5Pt21 due to the higher upper cycle temperatures required during thermal cycling, which resulted in the recovery of any favorable dislocation structures generated during training. & 2012 Elsevier B.V. All rights reserved.

Keywords: Shape memory alloys Martensitic transformation Thermomechanical training Two-way shape memory effect Thermal stability

1. Introduction NiTi based shape memory alloys (SMAs) have been exploited in a wide range of product forms such as eyeglass frames, orthodontic wires, valves, pipe couplings, and switches due to their unique functional properties [1]. In addition, the fact that the shape memory effect can be used to do work against a load has led to the development of SMAs as compact, solid-state actuators. Compared to D.C. motors or their pneumatic counterparts, these actuators have several advantages such as light weight, reduction in total part count, ease of inspection and higher energy densities [2]. SMA actuators mostly operate based on the one-way shape memory effect (OWSME) combined with a biasing force to reset the SMA after each actuation cycle. However, it would be advantageous in many designs to eliminate the need to mechanically reset the actuator. In this respect, the twoway shape memory effect (TWSME) renders it possible for an

n Corresponding author at: Department of Mechanical Engineering, Texas A&M University, College Station, TX 77843, USA. Fax: þ1 979 862 2418. E-mail address: [email protected] (I. Karaman).

0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.10.009

actuator to remember both its low-temperature and hightemperature shapes without the need for a re-biasing force. Unlike OWSME, TWSME is not an inherent characteristic of SMAs, but rather it is obtained after thermomechanical treatments (training), such as stress or temperature cycling. Recently, a need for compact, high-temperature actuation sources has emerged in several aerospace, oil exploration, and automotive applications. One drawback of NiTi based binary SMAs is the relatively low transformation temperatures (o100 1C), limiting their use when high temperature actuation is required. To overcome this restriction, NiTi can be alloyed with Pd, Pt, Au, Hf and Zr, to raise transformation temperatures [3]. Among these ternary NiTi based high-temperature SMAs (HTSMAs), the NiTiPd system has recently attracted considerable attention due to its adequate workability, low thermal hysteresis, and good strain recovery under constrained and stress-free conditions [4–11]. Early studies on the TWSME mostly concentrated on conventional Cu-based [12–20] and binary NiTi SMAs [21–38]. These studies focused on the explanation of the TWSME mechanisms, generation of TWSME using different training procedures, and the effects of training parameters on the magnitude and stability of the TWSME. Similar studies were also conducted on ternary NiTi based SMAs, such

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as NiTiCu [38–41], NiTiFe [35], and NiTiNb [42,43]. However, there is very limited data on the TWSME characterization of HTSMAs. The only studies in this area for NiTi based HTSMAs were performed on NiTiHf alloys [44,45]. In the first study [44], Ni49Ti36Hf15 (at%) plates were trained via bending in martensite followed by unconstrained recovery for up to 30 cycles at different temperatures. The highest two-way shape memory (TWSM) strain of only 0.88% was achieved after a bending strain of 7.1% at room temperature. But even at this limited strain level, the stability of the TWSME was poor. This was attributed to the low strength of the martensite, which permitted the introduction of dislocations during the TWSME cycles, relaxing the oriented stress fields. Furthermore, the already limited TWSM strain decreased by 50% in just 10 stress-free cycles. In an effort to improve the stability of the TWSME, the same training procedure was applied to a precipitation-hardened Ni50.6Ti29.4Hf20 (at%) HTSMA [45]. While the stability of the TWSME improved, the alloy still exhibited a 20% decrease in the TWSM strain after 30 stress-free cycles and the magnitude of the initial TWSM strain was only slightly increased compared to the solutionized material. To the authors’ knowledge, there has not been a systematic study on the TWSME of promising HTSMA systems such as NiTiPd and NiTiPt. In this study, a series of NiTi based HTSMAs are compared to a conventional binary Ni49.9Ti50.1 SMA in terms of the magnitude and stability of the TWSME as a result of a 100cycle thermomechanical training procedure. The effects of Sc microalloying to NiTiPd are also investigated in an effort to enhance the TWSME. The effects of training parameters on the resulting TWSME, such as training stress and the upper cycle temperature (UCT) are also considered. Since most of the emerging actuator applications require the SMA to do work against a load, it is also useful to characterize the work output capability of the TWSME and its stability during actuation, which initially has been addressed for a Ni24.5Ti50.5Pd25 HTSMA [46]. However, this broader study is focused on optimization of the TWSME and thus characterization of this phenomenon under stress-free conditions is sufficient and will provide direction for the selection, processing, and design of HTSMA actuators exploiting the TWSME.

been shown that excessive residual strain in the form of generalized plastic deformation and the subsequent formation of retained martensite have undesirable effects on the TWSME [14]. Thus, the approach for obtaining maximum TWSM strain is to induce just enough localized plasticity that will yield the highest magnitude of stable oriented stress fields for nucleation of single variant martensite [28]. 1.2. Thermomechanical training procedures to obtain TWSME Different training techniques have successfully been implemented to obtain TWSME in various SMA systems. The most common training techniques are: deformation in martensite followed by constrained or unconstrained recovery (OWSME cycling) [24,26,37,40,42,44]; stress cycling above the Af temperature (superelastic cycling) [12]; temperature cycling through the martensitic transformation under a constant stress level (isobaric thermal cycling) [13,14,20,23–25,27,28,32,36] or a combination of the latter two methods [29]. Common to all these techniques is the repeated growth and shrinkage of particular martensitic variants, which is responsible for the generation of stable dislocation arrays [47–49]. Another technique, which is different in principle than the aforementioned techniques, is aging under constraint. This method has been used to obtain TWSME in Ni-rich NiTi SMAs. With this method, coherent precipitates are formed in preferred orientations under applied stress, resulting in oriented internal stress fields that bias the formation of singlevariant martensite [30]. Among the training procedures mentioned above, thermal cycling through the martensitic transformation under a constant stress level has been shown to yield satisfactory results in terms of the magnitude and stability of the TWSME generated. This method has been proven to be an efficient procedure, and depending on how it is implemented, can involve relatively low stresses, which can result in minimal plastic deformation compared to other training procedures [28]. Consequently, this approach was adopted in the current study.

1.1. Origin of TWSME 2. Experimental procedures Two mechanisms for the TWSME have been proposed in the literature. The first mechanism attributes the TWSME to the oriented residual stress fields of the dislocation arrays generated during thermomechanical training [12]. These residual stress fields are able to induce the same variants of martensite in the absence of an external stress as the ones that are generated by the external training stress, thus resulting in the TWSME [20]. The magnitude and stability of the TWSME depends to a great extent on the magnitude of these stress fields and how they can be maintained through repeated thermal cycling. The same mechanism has also been explained from a thermodynamics point of view [15]. According to this view, the dislocation arrays generated during thermomechanical training create low energy configurations in the repeatedly induced martensite variants, while a relatively higher energy configuration is induced in the less frequently induced variants. As a result, the growth of martensite variants with the low energy configuration is favored even when the external stress is removed, resulting in TWSME. The second mechanism is based on the local stabilization of martensite, retained above the austenite finish (Af) temperature. Similar to the residual stress fields in the first mechanism, retained martensite plates are attributed with influencing the nucleation of certain preferred martensite variants during the transformation process, resulting in TWSME [12]. Common to these two mechanisms is the prerequisite for some amount of plasticity to occur in the material during training, which manifests itself as residual strain [12]. However, it has also

2.1. Materials Four different material compositions were chosen for this study: conventional binary Ni49.9Ti50.1, Ni24.5Ti50.5Pd25, Ni24.5Ti50Pd25Sc0.5 and Ni28.5Ti50.5Pt21 (all in at%). It should be noted that all compositions fall on the Ti-rich side of stoichiometry so that precipitation is not a complicating factor in any of the materials in this study. Binary Ni49.9Ti50.1 (at%) was acquired from Special Metals, New Hartford, NY. in the form of 10 mm diameter rods in the hot-rolled and hotstraightened condition. This lot of material has been characterized in some detail and properties as diverse as thermal expansion [50], elastic moduli [51], basic shape memory behavior [52,53] and textural evolution during thermomechanical testing are available in the literature [52,54]. Ingots of the high-temperature ternary and quaternary alloys were prepared by vacuum induction melting of high-purity elemental constituents (99.95% Ti, 99.98% Ni, 99.995% Pd, 99.995% Pt and 99.95% Sc). Each ingot was homogenized at 1050 1C for 72 h under vacuum and allowed to furnace cool. Subsequent to the homogenization, ingots were placed in mild-steel extrusion cans and extruded at 900 1C with an area reduction ratio of 7:1. Rectangular and cylindrical dog-bone shaped tension specimens with gage dimensions of 8 mm  3 mm  1.5 mm and 3.81 mm Ø  16.4 mm, respectively, were cut from the hot-rolled and hot-extruded rods for thermomechanical training and TWSME characterization.

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Binary Ni49.9Ti50.1 was chosen to represent the baseline behavior of a conventional SMA without complications from off-stoichiometry strengthening and precipitation, as occurs in Ni-rich compositions. Ni24.5Ti50.5Pd25 was selected as an HTSMA appealing to applications requiring intermediate transformation temperatures of around 200 1C [3,5,7–10]. Ni24.5Ti50Pd25Sc0.5 has transformation temperatures about 10 1C lower compared to the ternary alloy. However, as shown in a previous study [4], this HTSMA displays slightly improved functional behavior in the form of better dimensional stability and smaller thermal hysteresis during isobaric thermal cycling compared to the ternary alloy. One question to be answered is whether these improvements will also extend to the TWSME response. Ni28.5Ti50.5Pt21 was chosen as a representative alloy for high-temperature applications, capable of actuation around 350 1C. Constant stress thermal cycling behavior of a similar composition, Ni30Ti50Pt20 has previously been studied by Noebe et al. [55,56] and this material exhibited good work output and dimensional stability with transformation temperatures around 300 1C. 2.2. Thermomechanical training and TWSME characterization The choice of training procedure to induce the TWSME in the currently studied materials was thermal cycling under constant stress for 100 cycles. Training was carried out under tensile stresses of 80 MPa, 150 MPa and 200 MPa for all alloys. Onehundred thermal cycles was selected, since it was expected to yield a reasonably stabilized material response with minimal changes in shape and transformation temperatures upon further thermal cycling [8,9]. For the Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5 HTSMAs, training was carried out on a custom-built constant-load testing frame using the small rectangular tensile samples. Heating and cooling was done at a rate of 571 1C/min. Samples were heated by conduction from the grips, which were in turn heated by radiation through the use of an environmental furnace equipped with four 1 kW halogen lamps. For cooling, water was circulated around the grips flowing through copper tubing. Temperature was controlled with a K-Type thermocouple attached to the middle of the sample gage section. To minimize the radiation heat transfer on the sample surface (which will lead to erroneous temperature readings that are not representative of the bulk specimen), the sample was shielded with a 1 mm thick reflective aluminum foil. A capacitive displacement probe (Capacitecs HPC-75) with a linear range of 0–1.5 mm was attached to the grips to measure the displacements during the training process. Axial strain was calculated by dividing the change in length to the initial gage section length. Following thermomechanical training, samples were unloaded and TWSME was characterized using a separate MTS testing frame for 10 stress-free cycles to assess its magnitude and stability. Strain was measured using an MTS high-temperature extensometer with a gage length of 12.7 mm and a  20/þ20% strain range. Samples were heated through conduction from the grips with heating bands. Cooling was achieved through conduction by flowing liquid nitrogen in copper tubes wrapped around the grips. Heating and cooling rate of the samples was 10 72 1C/min. Similar to the training procedure, temperature was measured using a K-type thermocouple attached to the middle of the sample gage section. For the Ni49.9Ti50.1 and Ni28.5Ti50.5Pt21 samples, both training and TWSME characterization were performed on an MTS servohydraulic load frame using the cylindrical dog-bone samples. Strain was measured using an MTS high-temperature extensometer with a gage-length of 12.7 mm and a strain range of 10/ þ20%. Heating of the samples was achieved through the use

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of a 7.5 kW induction heater. Temperature was measured with a K-type thermocouple, which was spot-welded to the middle of the gage section. Samples were heated and cooled at rates of 30 72 1C/min and 2072 1C/min, respectively. Additional experimental details related to this test setup can be found in [8].

3. Results and discussion Fig. 1a illustrates a typical strain vs. temperature evolution during the 100-cycle thermomechanical training procedure employed in this study. The results are for the binary Ni49.9Ti50.1 trained under 150 MPa. Training is initiated by loading the sample at room temperature in martensite and subsequently heating above the Af temperature to the upper cycle temperature (UCT), which in this example was 165 1C, to obtain austenite under the applied load. The first complete and all subsequent thermal cycles are comprised of cooling from the UCT to the lower cycle temperature (LCT) and heating of the sample back to the UCT, e.g., through the forward and reverse transformations, respectively. A common characteristic for all these materials is that the 1st training cycle was completed with a relatively high value of residual strain, eres , as compared to subsequent cycles. We use the term eres for residual strain per cycle to denote that it is strain that was not recovered during a given thermal cycle and not to imply any given mechanism for the strain generation. With increasing number of cycles, the material obtains an almost stable shape memory response, demonstrating a subtly different strain

Fig. 1. (a) An example of the thermomechanical training procedure used and the resulting strain evolution due to repeated isobaric thermal cycling. Data is for the Ni49.9Ti50.1 alloy at 150 MPa. The 1st and the 100th cycles are highlighted. Total residual strain,etotal res , was measured as the cumulative residual strain per cycle, eres, after 100 cycles. (b) 10 stress-free TWSME cycles following the training regimen shown in (a). TWSM strain, eTWSM, was calculated as the strain difference between the cold and hot shapes for a given thermal cycle. Note that the stress-free strain– temperature cycles were rezeroed for easier comparison between conditions and various materials.

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vs. temperature curve at the end of the 100th cycle that has a nearly negligible eres per cycle. During the course of training, the material accumulates a total residual strain, etotal res , which is the total strain increase in the sample after 100 isobaric thermal cycles, measured in the austenite state at the UCT. Following training, the sample is unloaded at room temperature and heated above the stress-free Af temperature before the 1st stress-free TWSME cycle, recovering a fraction of the strain associated with the detwinned, post-trained microstructure (Fig. 1b). During TWSME cycling, degradation in stability of the TWSME is encountered, evidenced by a decrease in cold and hotshape strains (Fig. 1b). This is the opposite behavior of what is observed in Fig. 1a during training, where both the cold and hotshape strains increase with training cycles. In this study, the stability of the TWSME is also quantified by the change in the two-way shape memory strain, eTWSM, which is calculated as the difference between the cold and hot-shape strains for a given stress-free cycle, as demonstrated in Fig. 1b.

extracted from the 80 MPa data along with 150 MPa and 200 MPa training results and summarized in Table 1. The 80 MPa training results, shown in Fig. 2, reveal that all materials exhibited a relatively small transient response during the very first heating cycle. After this initial transient response, a more consistent erec was observed that evolved much more subtly during the course of training. The highest erec values were seen in Table 1 A summary of recovered transformation strain, erec, during the 1st and 100th after 100 cycles, determined training cycles and total accumulated strain, etotal res from thermomechanical training tests conducted under different stress levels and upper cycle temperatures (UCTs). Material (at%)

Ni49.9Ti50.1

3.1. Thermomechanical training and the evolution of shape memory behavior Fig. 2 is a compilation of the strain vs. temperature response for binary Ni49.9Ti50.1, Ni24.5Ti50.5Pd25, Ni24.5Ti50Pd25Sc0.5 and Ni28.5Ti50.5Pt21 during the 100-cycle training process at 80 MPa. To better understand the generation of TWSM behavior in each trained material, pertinent information such as etotal and the res recovered transformation strain, erec for the 1st and 100th training cycles, measured during the reverse transformation, were

Ni24.5Ti50.5Pd25

Ni28.5Ti50.5Pt21

Ni24.5Ti50Pd25Sc0.5

Training stress (MPa)

80 150 150 200 80 150 150 200 80 150 200 80 150 200

Training UCT (1C)

165 165 200 165 280 280 320 280 500 500 500 280 280 280

Rec strain Rec strain (erec) 1st (erec) cycle (%) 100th cycle (%)

Total residual strain,

2.04 3.62 3.82 3.76 1.68 2.28 2.43 2.71 0.32 1.34 2.09 1.80 2.49 2.73

3.68 3.39 3.76 3.44 2.40 2.55 2.73 2.69 0.55 2.24

4.29 7.15 12.38 10.11 1.10 2.26 2.36 3.00 1.17 4.53

2.41 2.74 2.75

1.12 1.74 2.80

(etotal res ) (%)

Fig. 2. 80 MPa, 100-cycle thermomechanical training results for (a) Ni49.9Ti50.1, (b) Ni24.5Ti50.5Pd25, (c) Ni24.5Ti50Pd25Sc0.5 and (e) Ni28.5Ti50.5Pt21.

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Ni49.9Ti50.1 at all stress levels investigated, while Ni28.5Ti50.5Pt21 had the lowest values, which was attributed to a higher stress needed to reorient martensitic variants in this material [55]. All samples accumulated strain during the thermomechanical training procedure. At the end of the 80 MPa, 100-cycle training, the lowest etotal levels occurred in the Ni24.5Ti50.5Pd25 and res Ni24.5Ti50Pd25Sc0.5 HTSMAs, while Ni49.9Ti50.1 exhibited the highest values (Table 1). As expected, an increasing training stress led to an increased value of etotal res for all materials, though for a given stress level, the values were always larger for binary Ni49.9Ti50.1 (Table 1 and Fig. 3). It is also interesting to see that Ni28.5Ti50.5Pt21 accumulated a relatively higher value of etotal res under 150 MPa compared to Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5, while it had almost the same etotal at 80 MPa. This is most likely the result of creep res deformation due to the increased stress level and the need to cycle to a much higher UCT than the other materials, which was necessary to complete the reverse transformation. For this same reason, Ni28.5Ti50.5Pt21 could not endure 100 thermal cycles under 200 MPa and failed at the 60th cycle generating a etotal res of 8.3%. Fig. 4 illustrates the evolution of the TWSME during the subsequent 10 stress-free cycles after training at 80 MPa (Fig. 4a, c, e and g), as well as the changes in eTWSM as a function of cycle count after training at all three stress levels (Fig. 4b, d, f and h). Table 2 lists 1st and 10th cycle eTWSM values; efficiency factor for each trained material, defined as the ratio of the 1st cycle eTWSM and the erec of the 100th training cycle; the amount of degradation in eTWSM as well as the cold and hot-shape strain values during 10 TWSME cycles. With the exception of Ni28.5Ti50.5Pt21, efficiency factors for materials trained under 80 MPa were found to be quite similar, around 0.8, which indicated that the 1st cycle eTWSM values were only  20% less than the erec values of the 100th training cycle (an efficiency factor of 1 indicates that 1st cycle eTWSM is equal to the erec of the 100th training cycle). In terms of TWSME stability, Ni49.9Ti50.1 demonstrated poor performance with large degradations in eTWSM (10%) and cold and hot-shape strains (a decrease in strain of 0.6% and 0.3%, respectively) during TWSME cycling. Ni24.5Ti50.5Pd25 and Ni24.5Ti50 Pd25Sc0.5 HTSMAs exhibited superior stability with less degradation in eTWSM and cold and hot-shape strains. 3.2. Effect of training stress on the TWSME As mentioned previously, TWSME arises from the presence of oriented internal stress fields, which bias the formation of specific

Fig. 3. Total residual strain (etotal res ), measured in the austenite, after 100 thermomechanical cycles as a function of stress level for the different SMA systems.

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martensite variants during the martensitic transformation, while actually inhibiting the formation of non-favorable variants [15]. The oriented stress fields are usually imposed in the material through the presence of defects such as dislocations that are generated during thermomechanical training. Thus, some amount of residual strain is expected during training; yet overstressing the material might result in an increase in the plastic deformation and facilitate the formation of retained martensite. This in turn, would result in a decrease in the TWSME due to a decrease in the amount of transforming volume. Thus, it would be anticipated that there is an ideal stress or range of stresses for optimizing TWSME through thermomechanical training.

3.2.1. Ni49.9Ti50.1 For Ni49.9Ti50.1, training stress had a complex effect on the magnitude and stability of the TWSME. The magnitude of the 10th cycle eTWSM varied inversely proportional to the training stress, with increasing training stress resulting in lower 10th cycle eTWSM (Fig. 5). However, while the magnitude of the eTWSM decreased with increasing training stress, the stability of the eTWSM increased as evident from the smaller changes between successive cycles (Fig. 4b). For instance, upon unloading the 80 MPa-trained material and heating above the Af temperature, more than 4% strain was recovered (Fig. 4a). Only 73% of this value, or about 3% strain, was carried over to the 1st TWSME cycle and the eTWSM degraded by a further 10% in 10 thermal cycles (Table 2). The eTWSM for the 150 MPa-trained material decreased from 2.6% to 2.5% in 10 cycles, while the eTWSM for the 200 MPatrained material was found to be very stable during cycling, and may have actually increased slightly by the end of the 10th cycle (Fig. 4b). At this point, it should be noted that a stable eTWSM is not necessarily a sign of a stable TWSME. eTWSM might be quite stable while there is a significant change in overall dimensions of the sample due to an almost equal amount of degradation in cold and hot-shape strains. On the other hand, minimal changes in cold and hot-shape strains will by definition result in a stable eTWSM. NiTi was found to display poor TWSME stability by exhibiting large changes in cold and hot-shape strains upon TWSME cycling even though the changes in eTWSM were sometimes very minor (Table 2). For example, after 10 cycles, the hot shape of the 200 MPa-trained material changed by 0.7% strain even though the magnitude of eTWSM changed by only 0.1% strain. For the 150 MPa-trained Ni49.9Ti50.1, etotal was around 7% res (Fig. 3). Stress-free thermal cycling altered the post-trained dislocation structure in a way that each cycle caused a progressive relaxation of the oriented internal stresses. This inevitably caused a decrease in the cold-shape strain due to the formation of more self-accommodated martensite instead of heterogeneously nucleated single-variant martensite (Fig. 4a). At the end of each cycle, it was noticed that the hot-shape strain did not match the strain value of the previous cycle, i.e., there was also a decrease in the hot-shape strain. The primary reason for the decrease in hotshape strain upon stress-free thermal cycling is that after a total deformation of 7%, there is probably retained martensite in the material at the UCT [54]. But upon stress-free thermal cycling and gradual relaxation of the internal stresses some of this oriented retained martensite is able to transform back to austenite resulting in a decrease in the hot-shape strain as well. It is also possible that back stresses due to the dislocation structures developed during training were relaxed, which would also result in a slight strain recovery. TWSM behaviors of 80 MPa and 200 MPa-trained NiTi can be described based on the same scenarios. The notable difference between these conditions is the amount of etotal res generation during training. This variation could result in different amounts of retained martensite and dislocation densities in the

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Fig. 4. The evolution of strain vs. temperature behavior during no-load thermomechanical cycling used to determine TWSME after training at 80 MPa for (a) Ni49.9Ti50.1, (c) Ni24.5Ti50.5Pd25, (e) Ni24.5Ti50Pd25Sc0.5 and (g) Ni28.5Ti50.5Pt21. Note that the strain–temperature cycles were rezeroed after training for easy comparison. The variation of the TWSM strain during stress-free thermal cycling for (b) Ni49.9Ti50.1, (d) Ni24.5Ti50.5Pd25, (f) Ni24.5Ti50Pd25Sc0.5 and (h) Ni28.5Ti50.5Pt21 after training at various stress levels.

trained materials, which alter the volume fraction of oriented martensite variants that cause the TWSME. For instance, the peculiar TWSM behavior of the 200 MPa-trained material, i.e.,

stable eTWSM but unstable axial dimension upon cycling, can be attributed to a recovery of a large volume fraction of retained martensite. For this material, the decrease in the hot-shape strain

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Table 2 Magnitude and stability of the TWSME for the various SMA systems after thermomechanical training under different stress levels and upper cycle temperatures (UCTs).

Ni49.9Ti50.1

Ni24.5Ti50.5Pd25

Ni28.5Ti50.5Pt21

Ni24.5Ti50Pd25Sc0.5

Training stress (MPa)

Training UCT (1C)

TWSME UCT (1C)

1st cycle TWSM strain (eTWSM) (%)

10th cycle TWSM strain (eTWSM) (%)

Efficiency factora

TWSM strain degradationb

Hot-shape strain Cold-shape degradationc strain degradationc

80 150

165 165 200

200 80 150

165 280 280 320 280 500 500 500 280 280 280

165 165 165 200 165 280 280 280 280 500 500

3.06 2.60 2.95 2.43 2.34 2.12 2.41 2.34 2.58 -0.04 0.13

2.75 2.53 2.82 2.09 2.36 1.99 2.28 2.20 2.46 -0.08 0.00

0.83 0.77 0.78 0.65 0.68 0.88 0.95 0.86 0.96 -0.08 0.06

0.10 0.03 0.04 0.14 -0.01 0.06 0.05 0.06 0.05 -1.00 1.00

0.33 0.68 0.46 0.64 0.69 0.10 0.17 0.13 0.08 0.04 0.03

0.60 0.64 0.49 0.91 0.58 0.24 0.31 0.27 0.23 0.09 0.09

280 280 280

2.11 2.46 2.56

2.04 2.38 2.46

0.88 0.90 0.93

0.03 0.03 0.04

0.20 0.13 0.10

0.26 0.22 0.22

200 80 150 200 80 150 200

a Efficiency factor was calculated as the ratio of the two-way shape memory strain (eTWSM) of the 1st stress-free cycle to the recoverable transformation strain (erec) of the 100th training cycle, i.e., (1st cycle eTWSM/100th cycle erec). b Degradation in eTWSM upon stress-free thermal cycling was calculated as the ratio of the difference between the 1st and 10th stress-free cycle eTWSM values to the eTWSM of the 1st stress-free cycle, i.e., ((1st cycle eTWSM – 10th cycle eTWSM)/(1st cycle eTWSM)). c Degradations in cold and hot-shape strains were calculated as the difference between the percent shape strain values of the 1st and 10th stress-free cycles.

Fig. 5. (a) TWSM strain values after 10 stress-free cycles for the tested materials as a function of training stress. The 10th no-load thermal cycles for (b) Ni49.9Ti50.1 and (c) Ni24.5Ti50.5Pd25 after training under 80 MPa, 150 MPa and 200 MPa.

was very similar, if not larger, than the decrease in the cold-shape strain, resulting in a very stable or even slightly increasing eTWSM, while resulting in a significant change in the axial dimension of the sample.

3.2.2. Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5 Compared to Ni49.9Ti50.1, much lower etotal values were res recorded at all training stress levels for both Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5 HTSMAs (Fig. 3). While etotal was 1.1% for res

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training at 80 MPa for Ni24.5Ti50.5Pd25, it doubled to 2.3% when the stress was increased to 150 MPa and at 200 MPa, etotal res was 3.0%. Marginally smaller values for etotal were observed at each stress res level for Ni24.5Ti50Pd25Sc0.5 (Table 1). These values were much less than that observed for Ni49.9Ti50.1. Thus the NiTiPd(Sc) HTMSAs were much more resistant to the development of plastic deformation at the stress levels investigated than Ni49.9Ti50.1, even though the NiTiPd(Sc) HTSMAs were thermally cycled to much higher temperatures. In addition, the transformation strain for the NiTiPd(Sc) HTSMAs increased with training stress. Thus an increasing transformation strain with increasing training stress naturally led to an increasing eTWSM with training stress. Furthermore, eTWSM of the trained NiTiPd(Sc) HTSMAs were found to be very close to the maximum recoverable strain levels under stress, resulting in high training efficiencies (Table 2), which actually increased with increasing training stress level (Fig. 5), while the opposite behavior was observed in binary Ni49.9Ti50.1. For example, the strain recovered during the initial heating of the 150 MPa-trained material above the Af temperature was very close to the 1st cycle eTWSM. Out of the 2.55% strain recovered during the last training cycle, 95%, or 2.4% strain, was carried over to the subsequent TWSME cycle and only degraded by about 5% at the end of 10 stress-free thermal cycles for all training stress levels. In terms of TWSME stability, cold and hot-shape strains of Ni24.5Ti50.5Pd25 were very stable and changed much less upon stress-free thermal cycling than Ni49.9Ti50.1 (Fig. 4c and Table 2). The hot-shape strain change was less than 30% of that observed in Ni49.9Ti50.1 for a given stress level while the change in cold-shape dimensions was less than half of that observed in the binary alloy. This would indicate that the dislocation structures developed in Ni24.5Ti50.5Pd25 during training were more resistant to relaxation and recovery of the internal stresses than in Ni49.9Ti50.1, in spite of cycling to higher temperatures. But there was also a consistent anisotropy in the cold versus hot-shape strain change, which can be seen in Fig. 4c and quantified in Table 2. In general, the hot-shape strain change was about half of that observed for the change in cold shape strain during the ten stress-free thermal cycles. As discussed above, the change in shape of the sample measured at the LCT in the martensite phase is due to a slight reversion of oriented martensite back to self-accommodated martensite. The change in hot-shape dimensions measured at the UCT is presumably due to reversion of retained martensite to austenite. This would indicate that in Ni24.5Ti50.5Pd25 either very little retained martensite was generated during training or the retained martensite that was formed was very stable. Ni24.5Ti50Pd25Sc0.5 displayed very similar results to Ni24.5Ti50.5Pd25, with similar or slightly improved eTWSM and cold and hot-shape strain stability (Table 2).

3.2.3. Ni28.5Ti50.5Pt21 The high transformation temperatures of Ni28.5Ti50.5Pt21 necessitated thermal cycling to much higher temperatures during the training process, i.e., 500 1C compared to just 280 1C for the NiTiPd(Sc) HTSMAs. Heating of the 80 MPa-trained sample above the Af temperature resulted in a strain recovery of 0.5% (Fig. 4g), which is almost the same as the erec of the last training cycle (Table 2). Further stress-free cycling, however, did not yield any useful eTWSM. At the end of 10 stress-free thermal cycles, the eTWSM of Ni28.5Ti50.5Pt21 was almost non-existent regardless of the training stress. Although Ni28.5Ti50.5Pt21 exhibited reasonable recoverable and residual strains during training for an alloy with such high transformation temperatures [57], it did not show significant evidence of TWSME. This can be attributed to the high UCT required for thermal cycling. Therefore, regardless of the training

stress used, the first stress-free cycle at 500 1C relaxed any beneficial dislocation structure that might have developed during training [56], eliminating the driving force for the TWSME (Fig. 4g and h).

3.3. Effects of composition on the TWSME The most obvious change in the TWSME characteristics between the various SMA systems is the basic effect of composition on transformation temperatures and the role that cycling above the various transformation temperatures plays on subsequent microstructural development. Fig. 6 shows the 10th TWSME cycles for all materials previously trained under 150 MPa. The substitution of 25 at% Pd for Ni in NiTi results in an increase in the transformation temperatures of Ni24.5Ti50.5Pd25 to around 200 1C. The small quaternary alloying addition of Sc to Ni24.5Ti50.5Pd25 subsequently decreased the resulting transformation temperatures slightly, by about 10 1C, consistent with the isobaric thermal cycling results and differential scanning calorimetry (DSC) analysis reported previously [4]. Pt has a greater effect per unit of alloying addition on transformation temperatures than Pd, such that the addition of 21% Pt to NiTi raised the transformation temperatures of Ni28.5Ti50.5Pt21 to between 350 1C and 400 1C. Thus cycling above this temperature to ensure complete transformation under stress, i.e., 500 1C, was high enough to allow complete recovery of any dislocation structure that may have developed during training, removing any driving force for subsequent TWSME. Noebe et al. [56] have previously shown that recovery processes begin at approximately 450 1C in a Ni29.5Ti50.5Pt20 HTSMA. It thus remains a technological challenge to obtain stable TWSME response in HTSMAs when transformation temperatures are close to the recovery/recrystallization temperatures of the particular alloy. Substitution of Pd with Ni not only increases the transformation temperatures, but also acts as a solid-solution strengthener and results in a change in the martensite structure from monoclinic B19’, to orthorhombic B19, resulting in improved compatibility [4] between transforming phases (cubic, B2 austenite transforming to orthorhombic martensite instead of monoclinic martensite). This improved crystallographic compatibility coupled with increased strength levels, results in a much more stable thermomechanical cycling response compared to that of Ni49.9Ti50.1. This is evidenced by smaller etotal res values during training under all stress levels (Fig. 3) and smaller changes in cold and hot-shape strains during TWSME

Fig. 6. Effect of composition on the TWSME for materials previously trained at 150 MPa for 100 thermal cycles. The strain–temperature curves represent the TWSM response for the 10th load-free thermal cycle. The right axis represents the efficiency of the eTWSM compared to the erec during training.

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cycles (Fig. 4c), even though the sample was cycled to a higher UCT. For these same reasons, there was also a decrease in energy dissipation during transformation, contributing to a significant reduction of the thermal hysteresis during both training and the TWSME cycles. Microalloying with Sc further improved the compatibility of the transforming phases in Ni24.5Ti50.5Pd25 and increased the material strength due to solid-solution strengthening [4]. As a result, Ni24.5Ti50Pd25Sc0.5 had slightly smaller thermal hysteresis compared to Ni24.5Ti50.5Pd25 during both training and TWSME cycles. Recently, the second eigenvalue, l2, of the transformation stretch tensor that maps the austenite lattice to the martensite lattice, has been related to the thermal hysteresis associated with the martensitic transformation [58]. It has been shown that as the l2 value gets closer to 1, compatibility between transforming phases increases, leading to a smaller thermal hysteresis. Current results for the l2 of binary NiTi (0.9663) [59], Ni24.5Ti50.5Pd25 (1.0171) [4] and Ni24.5Ti50Pd25Sc0.5 (1.0158) [4] are consistent with this finding for the thermal hysteresis associated with the TWSME, as well. In terms of eTWSM, Ni49.9Ti50.1 exhibited a higher absolute value at the end of 10 TWSME cycles compared to Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5 (Fig. 6). However, when the TWSM values were normalized with respect to the highest amount of recoverable strain that could be obtained from each material at a stress level of 150 MPa, it was observed that Ni24.5Ti50.5Pd25 and Ni24.5Ti50 Pd25Sc0.5 HTSMAs outperformed Ni49.9Ti50.1 with values very close to 1 (Fig. 6). This indicates that for a specific thermomechanical treatment, the NiTiPd(Sc) HTMSAs responded to training more efficiently and have nearly perfect TWSME relative to their load-biased behavior during training. A high value of training efficiency indicates that dislocation structures and local stress fields generated by cycling are very effective in biasing the same martensite variants that were formed through load-biased thermal cycling, during the subsequent thermal cycling under zero stress [28]. On the other hand, Ni28.5Ti50.5Pt21 had a value close to 0, indicating that this material is not suitable for TWSME applications. 3.4. Effect of training upper cycle temperature (UCT) on the TWSME It is commonly accepted that for a chosen SMA system, transformation strains generated during thermomechanical training will be representative of the resultant eTWSM. This is the case for the Ni49.9Ti50.1 SMA and NiTiPd(Sc) HTSMAs studied here. Thus, any training parameter that can potentially increase the erec of the last training cycle will most likely increase the eTWSM of the following stress-free cycles. While erec is a function of applied stress, it also can be affected by the UCT the material is exposed to during cycling [53]. While a high UCT is more likely to trigger global plasticity due to the decrease in critical shear stress (CSS) for slip, it is also possible that higher UCT values may reduce the amount of retained martensite present in the alloy by heating further above the Af temperature. The initial choice of UCT for thermomechanical training was made to ensure complete transformation of martensite to austenite under the applied stress. Since all materials have different transformation temperatures, different UCTs were used during training. Ni49.9Ti50.1 was thermally cycled to 165 1C, whereas NiTiPd(Sc) HTSMAs were cycled to 280 1C and Ni28.5Ti50.5Pt21 was heated to 500 1C. Padula et al. [53] have already demonstrated that UCT affects the thermomechanical response of the Ni49.9Ti50.1 SMA used in this study, at least for a relatively low number of thermal cycles. They have shown that the material cycled using a higher UCT, up to some maximum value depending on the stress level, exhibited higher transformation strains

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reaching a maximum erec at a UCT of around 250 1C under 200 MPa [60]. It is of interest to determine whether similar trends will be seen after a relatively high number of thermal cycles and how this may affect the resulting TWSME. Consequently, Ni49.9Ti50.1 and Ni24.5Ti50.5Pd25 were trained using two different UCT values under 150 MPa. Ni49.9Ti50.1 was cycled to UCTs of 165 1C and 200 1C, while the Ni24.5Ti50.5Pd25 was cycled using UCTs of 280 1C and 320 1C.

3.4.1. Ni49.9Ti50.1 The strain vs. temperature response of the Ni49.9Ti50.1 during thermomechanical training under 150 MPa is shown in Fig. 7a and b with UCTs of 165 1C and 200 1C, respectively. At a first glance, it is clear that etotal generated during training increased res from 7.2% to 12.4% due to an increase of UCT from 165 1C to 200 1C. In spite of the large difference in etotal there are also res similarities in the strain vs. temperature evolution for these two cases. For both materials, the 1st training cycle is characterized by a relatively large eres and thermal hysteresis. However, during the course of training, the materials started to show a more stable behavior characterized by smaller shifts in transformation temperatures during cycling and smaller values of eres per cycle. Also, a smaller thermal hysteresis was observed, indicative of slightly less dissipation of elastic stored energy during later transformation cycles, which could be due to a reduced interaction between martensite variant pairs or to the development of low energy dislocation structures with cycling. There are also obvious differences between these two conditions in terms of erec . Compared to the 165 1C-UCT trained material, the 100th cycle of the 200 1C-UCT trained material had a noticeably higher erec (3.8% vs. 3.4%) and higher transformation temperatures. For the 165 1C-UCT trained material erec decreased from 3.6% to 3.4% after 100 thermal cycles at 150 MPa. For the 200 1C-UCT trained material, erec remained quite stable at around 3.8% during the entire 100 cycle training regime. As plasticity was more prominent at the higher UCT value (Fig. 7a vs. b), one would anticipate a reduced transforming volume. Regardless, the higher transformation strain of the 200 1C-UCT trained material can be explained by a retained martensite argument. While a higher UCT results in higher values of etotal during training, the amount of res retained martensite is probably diminished compared to the lower UCT case, as evidenced by the increased erec . In addition, while plasticity was greater in the 200 1C-UCT trained material as measured by etotal res , it is possible that the transforming volume was actually relatively unchanged or even greater than that for the 165 1C-UCT trained material, because excess dislocations could have been annihilated during cycling or rearranged into low energy networks, given the additional thermal energy introduced by cycling to 200 1C. The resultant TWSME obtained from these two cases is also different in terms of stability and magnitude. While a UCT of 200 1C during thermomechanical training resulted in a higher value of etotal res , it also yielded a higher initial eTWSM of around 3.0% (compared to 2.6% when a UCT of 165 1C was used). The degradation levels associated with cold and hot-shape strains are also different. After 10 TWSME cycles, the material trained with a UCT of 165 1C exhibited a eTWSM of around 2.5% (Fig. 7e) with about 0.7% decrease in both cold and hot-shape strains (Fig. 7c). Fig. 7d illustrates the TWSM response of the same material trained with a UCT of 200 1C. Compared to the material trained with a UCT of 165 1C, the decrease in both cold and hotshape strains are less by about 0.2% and the resulting eTWSM was 2.8% at the end of 10 TWSME cycles. The larger decrease in the cold-shape strain, in the material trained with a UCT of 165 1C, is probably due to a larger relaxation of internal stresses, and thus

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Fig. 7. The evolution of strain vs. temperature response of Ni49.9Ti50.1 trained under 150 MPa using a UCT of (a) 165 1C and (b) 200 1C. The evolution of the TWSME during the subsequent 10 stress-free cycles for the material trained with UCT of (c) 165 1C and (d) 200 1C. (e) A comparison of 10th TWSME cycles for the Ni49.9Ti50.1 alloy trained with different UCTs.

the formation of more self-accommodated martensite at the expense of the preferred oriented martensite. The larger etotal res of 12.4% in the 200 1C-UCT trained material is likely to have produced a more stable, lower energy dislocation network, which would be more resistant to relaxation of internal stresses in this material, leading to an improved TWSME stability. Finally, the smaller degradation in the hot-shape strain of the same material is possibly associated with smaller amounts of retained martensite present in the microstructure that reverts back to austenite during stress-free cycling.

3.4.2. Ni24.5Ti50.5Pd25 In contrast to Ni49.9Ti50.1, a 40 1C increase in the UCT used during training of Ni24.5Ti50.5Pd25 resulted in no significant difference in etotal levels (Table 1). The sample, which was res thermally cycled under 150 MPa using a UCT of 280 1C had a etotal res of 2.3% (Fig. 8a), while the same thermomechanical training

with a UCT of 320 1C resulted in a etotal value of 2.4% (Fig. 8b). res When the 1st and 100th cycles are inspected for both cases, it is clear that the evolution in strain vs. temperature responses are unaffected by the choice of UCT. Unlike Ni49.9Ti50.1, the 1st and 100th cycles for Ni24.5Ti50.5Pd25 are very comparable in terms of erec and the shape of the hysteresis loops. These are all indications of a shape memory response that is not affected by the choice of UCT, at least over the temperature range investigated. In addition, cycling to either UCT resulted in an approximately 16 1C shift in transformation temperatures during the course of training compared to a 7 1C shift in Ni49.9Ti50.1. This is an indication that larger residual stresses are generated in the Ni24.5Ti50.5Pd25 alloy, assisting the transformation. The initial stress-free heating curves after the training procedure follow almost the same strain–temperature path for the two cases (Fig. 8c and d) indicating similar values of internal stresses in both materials regardless of the UCT used. Consequently, the changes in cold and hot-shape strains are indistinguishable and are minimal

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Fig. 8. The evolution of strain vs. temperature response of Ni24.5Ti50.5Pd25 trained under 150 MPa using a UCT of (a) 280 1C and (b) 320 1C. The evolution of the TWSME during the subsequent 10 stress-free cycles for the material trained with UCT of (c) 280 1C and (d) 320 1C. (e) A comparison of 10th TWSME cycles for the Ni24.5Ti50.5Pd25 alloy trained with different UCTs.

compared to Ni49.9Ti50.1. The 10th TWSME cycle for each case is shown in Fig. 8e. The two strain vs. temperature curves are nearly identical with similar transformation strain, hysteresis, and transformation temperatures. The 0.1% difference in eTWSM for the 10th stress-free cycle (Fig. 8e) is essentially within the uncertainty of the strain measurements and is hardly significant. 3.5. Effect of TWSME cycling upper cycle temperature (UCT2W) on the TWSME In the previous section, the significance of the UCT used for thermomechanical training cycles was demonstrated for the binary Ni49.9Ti50.1 alloy (while the Ni24.5Ti50.5Pd25 alloy was insensitive to this effect over the narrow range of temperatures investigated). Consequently, the selection of UCT for the TWSME cycles also deserves attention. To avoid any confusion with the UCT used during the training process, the UCT used during the TWSME cycles will be denoted as UCT2W. To show the effect of

UCT2W on the stability and magnitude of TWSME, the binary Ni49.9Ti50.1 alloy, trained under 150 MPa using a UCT of 200 1C, was thermally cycled stress-free using a UCT2W of 165 1C and 200 1C. As expected for this alloy, a higher UCT2W used during stress-free cycling severely affected both the magnitude and stability of TWSME. Out of the 3.8% erec for the last training cycle, only 2.4% (vs. 3.0% with a UCT2W of 165 1C) could be carried over to the 1st TWSME cycle (Table 2). The eTWSM decreased by 0.3% (vs. 0.1% with a UCT2W of 165 1C) in 10 cycles and the largest decrease in cold-shape strain was observed for this condition. This increased degradation in stability is undoubtedly due to the rapid relaxation of oriented internal stresses at this higher cycling temperature. 3.6. Factors that affect the magnitude and stability of TWSME For a review of the factors that can affect the magnitude and stability of TWSME, the schematic in Fig. 9 is introduced.

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Fig. 9. Schematic showing the parameters involved in the generation of TWSME, its magnitude and stability. See text for details. (?: dislocations, s: applied stress during training, UCT: upper cycle temperature during either training or TWSME cycles, TRIP: transformation-induced plasticity, RM: retained martensite, etotal res : total residual strain measured in the austenite state at the UCT, OM: oriented martensite (responsible for the TWSME), SAM: self-accommodated martensite.).

The figure is divided into five main parts, each summarizing relevant microstructural parameters, mechanisms, or testing conditions that can directly or indirectly influence the magnitude and stability of the TWSME. Part I illustrates the key parameters that are likely to affect the microstructural evolution during thermomechanical training: the applied stress (s), the choice of UCT, and the number of cycles. In this study, only s and UCT were chosen as variables. The most evident effect of an increase in either s (Fig. 3) or UCT (Figs. 7 and 8) was an increase in etotal res generated during training. The effect of the number of cycles on the TWSME was not directly investigated in this study, with the number of cycles being kept constant at 100. But it is clear from the data that etotal res increases with cycle count, in a manner that appears to saturate at some higher number of cycles. Though the effect of cycle count on the TWSME was not directly investigated, it can be postulated that as the cycle count increases and the strain–temperature response saturates with cycling, a more stable TWSME should be developed. In addition to the martensitic phase transformation in SMAs it is possible for a number of additional inelastic and predominantly non-recoverable deformation processes to occur concurrently, as illustrated in Part II of Fig. 9. These are what give rise to the major component of the etotal generated during thermomechanical res cycling or training. Depending on the temperature and applied stress under which the transformation occurs, several deformation mechanisms can come into play. The most critical, because it will occur regardless of the stress–temperature conditions, is accommodation of the volume mismatch during the martensitic transformation by dislocation generation at austenite/martensite interfaces, a process also known as transformation-induced plasticity (TRIP). Evidence for this behavior in NiTi, through detailed TEM studies, has been presented by a number of investigators [47,49,61]. In addition, general plasticity is also likely as internal stresses or stress concentrations in the material reach the critical resolved shear stress for slip at first locally [52] and then more or

less generally throughout the material [62–65]. Finally, the effects dislocation relaxation, recovery and even creep deformation cannot be ignored during thermal cycling under stress at even moderate temperatures [66] as seen in the Ni28.5Ti50.5Pt21 case. Consequently, thermomechanical cycling during the training process results in an accumulation of defects in the microstructure. These defects are usually either dislocations or deformation twins [47–49,61–65,67] generated due to the mechanisms summarized in Part II. In addition to their contribution to plasticity, and potentially acting as a source for preferred nucleation of oriented martensite (OM), these defects also form a barrier against reverse transformation, stabilizing some of the martensitic variants such that they cannot revert back to austenite during transformation. Part III of Fig. 9 represents the formation of dislocations (?) and retained martensite (RM) as a result of training. The horizontal arrow is used between ? and RM to represent the interplay between dislocation structures and martensite stabilization. A consequence of both dislocation generation and to a much lesser extent the formation of RM is an increase in the etotal res and change in the hot-shape strain of the sample, which can be quantified from the thermomechanical testing results (Fig. 3). Some amount of etotal res is probably necessary for the generation of TWSME, yet excessive deformation leads to a decrease in the TWSME, as seen in binary Ni49.9Ti50.1 (Fig. 5b). Although there is no unique relationship between etotal res and the TWSME, dislocation structures and RM determine the amount of transforming volume available after training. Since it is not exactly known how dislocation structures contribute to the amount of transforming volume, the connection between ? and transforming volume has been indicated with a dashed arrow. Due to the presence of oriented internal stresses induced during training, a fraction of the transforming volume will be comprised of OM, which is responsible for the TWSME and directly controls the magnitude of the eTWSM. The remainder will transform into self-accommodated martensite

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(SAM). The difference between the observed eTWSM and the theoretical maximum transformation strain for the polycrystalline aggregate is then due to this SAM and any RM, which does not participate in the transformation process. Obviously, microstructural features such as crystallographic compatibility between transforming phases, texture, grain size and the initial dislocation density also affect microstructural evolution during training (Part IV of Fig. 9). Higher strength levels through finer grain size, higher initial dislocation density, precipitate strengthening, or preferred textures/orientations where slip is severely restricted, as well as improved compatibility between the transforming phases through compositional changes should decrease defect generation during training, resulting in a relatively lower dislocation density and possibly less RM. Other than improved compatibility between transforming phases as exemplified by comparing the response of the binary NiTi and NiTiPd(Sc) systems [4], the effects of these microstructural parameters on the TWSME were beyond the scope of this study, but are included in Part IV of Fig. 9 for completeness. The actual stress-free cycling of the trained material, used to investigate the stability of the TWSME, can cause additional microstructural evolution. First, a higher UCT2W will tend to cause rearrangement and annihilation of the dislocation structures generated during training, relaxing the internal stress fields and affecting the proportion of OM that is formed. Furthermore, it may assist in destabilization of any RM in the structure. The magnitude of the eTWSM and the stability of the TWSME depend directly on the relative amounts of all types of martensite structures (SAM, OM and RM) during stress-free cycling. Thus relaxation of the dislocation structures that affect the volume fractions of OM and the fraction of SAM retained during the transformation process will result in a change in eTWSM and instability of the TWSME due to changes in cold and hot-shape strains of the sample (Fig. 7c and d). Hence the selection of UCT during the TWSME cycling, as well as the other microstructural parameters mentioned above (Part IV of Fig. 9), will have a strong influence on the magnitude and stability of the TWSME.

4. Summary and conclusions In this study, a binary Ni49.9Ti50.1 SMA and Ni28.5Ti50.5Pt21, Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5 HTSMAs were characterized in terms of the stability and magnitude of the two-way shape memory effect (TWSME). The TWSME was induced in these materials through a thermo-mechanical training procedure consisting of 100 thermal cycles under various stress levels with different upper cycle temperatures (UCTs). In the process of training, the evolutionary behavior of the strain–temperature response of the materials as a function of stress was also determined. Subsequently, the stability and magnitude of the TWSME was assessed by running 10 stress-free thermal cycles following the training cycles. A summary of the results and conclusions that could be derived from this study are: 1. In general Ni49.9Ti50.1 exhibited poor TWSME stability evidenced by large degradations in cold (martensite) and hot (austenite) shape strains upon stress-free thermal cycling. Ni24.5Ti50.5Pd25 and Ni24.5Ti50Pd25Sc0.5 had superior stability characterized by minimal shape changes. 2. After training at 150 MPa, Ni49.9Ti50.1 exhibited a eTWSM of 2.5% at the end of 10 stress-free cycles. This value was about 60% of the recoverable strain level for the material under biased conditions at the end of training. On the other hand, NiTiPd(Sc) HTSMAs had much more efficient responses to training. These

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alloys had eTWSM values around 2.4%, which corresponded to almost 90% of the recoverable strain that could be obtained from these materials during 150 MPa thermal cycling. Due to the inherently high transformation temperatures of the Ni28.5Ti50.5Pt21, stress-free thermal cycling was performed at temperatures above the recovery range for the alloy, relaxing or recovering any dislocation structures in the material. Consequently, no significant TWSME was observed. 3. Both the stress level and UCT used during thermo-mechanical training have significant but varied effects on the magnitude and the stability of the resulting TWSME. UCT2W, the upper cycle temperature selected for stress-free thermal cycling during measurement of the TWSM response, was equally important for the stability of the TWSME. A higher UCT2W was found to result in a faster degradation of the TWSME in Ni49.9Ti50.1. 4. NiTiPd(Sc) HTSMAs have been shown to be attractive candidates for TWSME applications and are the only currently viable option for use at higher temperatures (to about 200 1C). These alloys generated relatively small amounts of residual strain during thermomechanical training, over the range of UCTs and stress levels investigated. The resultant TWSME was shown to have excellent stability, including relatively stable cold and hot-shape strains, and adequate eTWSM. Acknowledgments The authors gratefully acknowledge insightful discussions with Ji Ma (Graduate research assistant at Texas A&M University) and the Shape Memory Alloy and Active Structures Group at NASA Glenn Research Center. Particular thanks to Brian E. Franco (Graduate research assistant at Texas A&M University) for his help with the construction of the thermomechanical test setup and training of samples. This study has been supported by the NASA Fundamental Aeronautics Program, Subsonic Fixed Wing Project through Cooperative Agreement No. NNX07AB56A, with additional support from the Aeronautical Sciences Project. IK also acknowledges the support from the US Air Force Office of Scientific Research, Grant No. FA9550-12-1-0218. References [1] K.N. Melton, in: T.W. Duerig, K.N. Melton, D. Stockel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, Butterworth-Heinemann, New York, 1990, pp. 21–35. [2] C. Mavroidis, Res. Nondestr. Eval. 14 (2002) 1–32. [3] J. Ma, I. Karaman, R.D. Noebe, Int. Mater. Rev. 55 (2010) 257–315. [4] K.C. Atli, I. Karaman, R.D. Noebe, A. Garg, Y. Chumlyakov, I. Kireeva, Metall. Mater. Trans. A 41 (2010) 2485–2497. [5] K.C. Atli, I. Karaman, R.D. Noebe, H.J. Maier, Scr. Mater. 64 (2011) 315–318. [6] K.C. Atli, I. Karaman, R.D. Noebe, A. Garg, Y. Chumlyakov, I. Kireeva, Acta Mater. 59 (2011) 4747–4760. [7] G. Bigelow, R.D. Noebe, S.A. Padula, A. Garg, in: B. Berg, M.R. Mitchell, J. Proft (Eds.), Proceedings of the International Conference on Shape Memory and Superelastic Technologies, 2006, pp. 113–132. [8] G. Bigelow, S.A. Padula, A. Garg, D. Gaydosh, R.D. Noebe, Metall. Mater. Trans. A 41 (2010) 3065–3079. [9] G. Bigelow, D. Gaydosh, A. Garg, S.A. Padula, R.D. Noebe, in: S. Miyazaki (Ed.), Proceedings of the International Conference on Shape Memory and Superelastic Technologies, 2007, pp. 83–92. [10] G. Bigelow, S.A. Padula, A. Garg, R.D. Noebe, in: M.J. Dapino (Ed.), Proceedings of SPIE: Behavior and Mechanics of Multifunctional and Composite Materials, 2007, pp. 2B1-2B12. [11] B. Kockar, K.C. Atli, J. Ma, M. Haouaoui, I. Karaman, M. Nagasako, R. Kainuma, Acta Mater. 58 (2010) 6411–6420. [12] J. Perkins, R.O. Sponholz, Metall. Trans. A 15 (1984) 313–321. [13] R. Stalmans, J.V. Humbeeck, L. Delaey, Acta Metall. Mater. 40 (1992) 501–511. [14] R. Stalmans, J.V. Humbeeck, L. Delaey, J. Phys. IV 1 (1991) 403–408. [15] R. Stalmans, J.V. Humbeeck, L. Delaey, Acta Metall. Mater. 40 (1992) 2921–2931. [16] R. Stalmans, J.V. Humbeeck, L. Delaey, in: R. Yamamoto, E. Furubayashi, Y. Doi, R. Fang, B. Liu (Eds.), Advanced Materials ‘93, Vol. V: Pt. A: Ecomaterials; Pt. B: Shape Memory Materials and Hydrides, 1994, pp. 927–930.

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