- Email: [email protected]
Effect of low-temperature precipitation on the transformation characteristics of Ni-rich NiTi shape memory alloys during thermal cycling
Intermetallics 18 (2010) 1172e1179
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Effect of low-temperature precipitation on the transformation characteristics of Ni-rich NiTi shape memory alloys during thermal cycling M.F.-X. Wagner a, b, *, S.R. Dey a, b, H. Gugel a, J. Frenzel a, Ch. Somsen a, G. Eggeler a, b a b
Institute of Materials Science and Engineering, Chemnitz University of Technology, Erfenschlager Str. 73, 09125 Chemnitz, Germany Research Department IS3/HTM e Integrity of Small-Scale Systems/High Temperature Materials, Ruhr University Bochum, Universitaetsstr. 150, 44801 Bochum, Germany
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 December 2009 Accepted 26 February 2010 Available online 1 April 2010
Thermal cycling of NiTi shape memory alloys is associated with functional fatigue: the characteristic phase transformation temperatures decrease with increasing number of cycles, and the transformation behavior changes from a single- to a two-stage martensitic transformation involving the intermediate R-phase. These effects are usually attributed to a gradual increase of dislocation density associated with micro-plasticity during repeated cycling through the transformation range. Here, these changes are shown to increase at a higher maximum temperature (in the fully austenitic state) during differential scanning calorimetric cycling of a Ni-rich alloy. Additional thermal cycling experiments without repeated phase transformations, and post-mortem microstructural observations by transmission electron microscopy, demonstrate that a relevant portion of functional fatigue is due to the formation of nanoscale Ni-rich precipitates of type Ni4Ti3 even at temperatures relatively close to the austenite finish temperature. These results show that both dislocation generation during the diffusion-less phase transformation, and diffusion-controlled nucleation and growth of Ni4Ti3 precipitates, can interact and contribute to the evolution of functional properties during thermal cycling of Ni-rich NiTi. Ó 2010 Elsevier Ltd. All rights reserved.
Keywords: A: Multiphase intermetallics B: Martensitic transformations B: Precipitates B: Shape-memory effects
1. Introduction NiTi shape memory alloys are widely used in medical implants and devices, where the reversible martensitic transformation between a more symmetric high-temperature phase (austenite; cubic B2 type crystal structure) and a less symmetric low temperature phase (martensite; monoclinic B190 type crystal structure) is exploited to obtain pseudoelastic material behavior [1e4]. Future applications are also expected in the field of actuators and devices that use thermal shape memory during repeated heating and cooling [5,6]. Actuator forces are generated as the NiTi component is heated and contracts during the reverse transformation from martensite to austenite [7e9]. Differential scanning calorimetry (DSC) can be used to assess the transformation behavior of a NiTi alloy. By slowly cooling and heating a DSC specimen between a maximum temperature Tmax above the austenite finish temperature, Af, and a minimum temperature Tmin below the martensite finish temperature, Mf, the material is fully transformed to martensite and back to austenite. The characteristic phase * Corresponding author. Institute of Materials Science and Engineering, Chemnitz University of Technology, Erfenschlager Str. 73, 09125 Chemnitz, Germany. Tel.: þ49 371 531 38683; fax: þ49 371 531 838683. E-mail address: [email protected] (M.F.-X. Wagner). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.02.048
transformation temperatures (PTTs) can be determined from the heat flow data (see Fig. 1; the forward martensitic transformation is exothermic, whereas the reverse transformation is endothermic). It is well known that thermal cycling affects the transformation characteristics of NiTi measured by DSC [10e16]. As shown schematically in Fig. 1, one typically observes decreasing PTTs with increasing number of DSC cycles [10e12]. Moreover, the R-phase, an intermediate phase that is also formed in a reversible, martensitic transformation from B2 austenite, can occur, resulting in a two-step phase transformation (B2 / R / B190 ). In engineering applications, the martensitic forward and reverse transformations must be repeatable for many cycles [17,18]. While it is in theory fully reversible [19], the phase transformation involves nucleation and growth of martensite into the adjacent, austenitic material. As the growing martensite interacts with defects and interfaces in the surrounding matrix, some local formation of dislocations (and hence irreversible plastic deformation) is inevitable [20e23]. Micro-plasticity in NiTi can lead to classical structural fatigue, where it ultimately results in crack formation and growth during mechanical cycling [24e26]. Moreover, functional fatigue is associated with a gradual degradation of shape memory behavior when NiTi is subjected to several forward and reverse transformation cycles. Even during pure thermal cycling without mechanical loading, the functional characteristics
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
1173
nucleation of Ni4Ti3 occurs more easily than particle growth in the early stages of precipitation, low-temperature/short-time ageing results in a high number of very fine, nano-scale precipitates (with particle diameters <10 nm, [37]). Despite a substantial body of work on functional fatigue on the one hand, and on the effect of heat treatments on the other hand, there is no systematic study that takes into account how Tmax may affect the elementary microstructural processes that contribute to the evolution of functional properties during thermal cycling. In this work, we investigate how functional changes and precipitation are interrelated when Tmax approaches the temperature regime where precipitates can form in a Ni-rich NiTi alloy. By analyzing the DSC data of up to 30 thermal cycles through the transformation regime, we document that functional degradation is more severe at higher Tmax. From this unexpected temperature dependency, and from complementary TEM analysis, we conclude that Ni4Ti3 precipitation occurs even at the relatively low temperatures associated with thermal cycling. We show that both processes (formation of dislocations and precipitation) may simultaneously affect functional degradation, and that precipitation of nano-scale Ni4Ti3 particles may well be a dominant mechanism when thermal cycling is performed at higher temperatures. Fig. 1. Schematic representation of the evolution of DSC data during thermal cycling of NiTi. With increasing cycle number N, the transformation temperatures are shifted towards lower temperatures, and the transformation sequence changes from a singleto a two-step transformation, involving the intermediate R-phase, during cooling. An individual phase transformation can be characterized by the start, peak, and finish temperatures (indices s, p, and f, indicated here for the martensite peak).
often change considerably because of relatively subtle changes in the microstructure [10e12]. In the light of functional fatigue studies, the changes in functional behavior and transformation sequence during simple thermal cycling are generally rationalized by the accumulation of dislocations and their effect on subsequent transformation cycles [13e16]. However, it is not clear whether other elementary processes can also affect functional degradation during thermal cycling. Interestingly, similar effects on the functional properties are observed when Ni-rich alloys are heat treated at constant, elevated temperatures (typically: T > 350 C, [27]). Then, occurrence of R-phase and decreasing PTTs are attributed to precipitation of the meta-stable Ni4Ti3 phase. Such heat treatments are beneficial because particle hardening increases the yield strength of austenite, which in turn contributes to better functional stability as less dislocations are formed during mechanical or thermal cycling. Recently, Undisz et al. [28] reported that shorttime heat treatments (annealing times of 2e30 min) considerably affect the PTTs and the pseudoelastic behavior of Ni-rich NiTi wires. Moreover, Jiang et al. [29] provided a detailed study on how precipitation limits the stress- and temperature regimes for pseudoelastic and shape memory behavior of an alloy with 50.9 at.% Ni. Heterogeneous precipitation of Ni4Ti3 is also one of the key mechanisms underlying multi-stage transformations that have recently attracted considerable attention [30e32]. This interest has motivated detailed studies on the properties and precipitation behavior of Ni4Ti3 [33e36]. In the context of the present work, a careful investigation by Kim and Miyazaki [37] is especially important as it provides timeetemperatureetransformation data on, and a detailed microstructural analysis of, Ni4Ti3 precipitation in solution-annealed Ni-rich (50.9 at.% Ni) NiTi. They [37] found that precipitation occurs at temperatures as low as 200 C (but only after long ageing times >100 h), whereas precipitates could already be observed by transmission electron microscopy (TEM) after ageing at 300 C for 20 min. They documented that ageing at low temperatures (<330 C) is most effective for improving the shape memory stability of NiTi. Moreover they demonstrated that, as
2. Experiments An ingot of binary NiTi with a nominal composition of 50.7 at.% Ni was produced by vacuum induction melting. Details on the individual processing steps for high purity melting and thermomechanical processing have been published elsewhere [38]. For chemical homogenization, the ingot was annealed at 850 C for 1 h in argon and subsequently quenched in water. A rectangular piece was cut from the ingot and hot-rolled in several steps (with intermediate heating to 850 C) to a final thickness of 2 mm. The hotrolled sheet was then again annealed at 850 C for 1 h under argon atmosphere. The grain size of the sheet material was determined from optical micrographs after mechanical grinding and chemical etching with Beraha-I solution. The observations were made over the horizontal plane of the sample containing the rolling and transverse directions. Several optical micrographs were acquired and the grain size was statistically determined using the line intersection method. The as-processed microstructure contained slightly elongated grains along the rolling direction, with an average grain size of w42 mm. As the dimensions of the DSC specimens are w2.5 2.5 2 mm3 (this corresponds to a specimen mass of about 80 mg), the grain size is sufficiently small to ensure that the observed transformation behavior is fully representative of the polycrystalline sheet material. Thermal cycling was performed in a DSC instrument of type 2920 CE (TA Instruments) at a constant heating/cooling rate of 10 K/ min under an argon/helium atmosphere (to minimize oxidation in the long-term cycling experiments). DSC cycles were performed on three specimens with constant Tmin (100 C) but with varying Tmax (150, 200 and 250 C). 20 DSC cycles were carried out for Tmax ¼ 150 and 200 C, and 30 cycles for Tmax ¼ 250 C. The first cycle consisted of (1) heating the fully austenitic specimens from ambient temperature to Tmax, (2) holding the specimen temperature constant at Tmax for 3 min, (3) cooling the specimen to Tmin, (4) holding the specimen temperature constant at Tmin for 3 min, and (5) heating the specimen again to Tmax. Subsequent cycles were performed by repeating steps (2)e(4), see also Fig. 1. Note that using hold times (often up to 10 min) at Tmin and Tmax is a standard procedure during DSC testing, with the aim of ensuring thermal equilibrium of the specimen and the reference sample. The DSC data allows detailed analysis of PTTs, i.e., the martensite, austenite, and R-phase start and finish temperatures Ms, Mf, As, Af, Rs, Rf,
1174
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
respectively. Moreover, the general evolution of the transformation peaks as a function of the number of cycles was documented by plotting the temperatures of maximum/minimum heat flows, Mp, Ap, and Rp (which are referred to as peak temperatures hereafter). To separate the effects of conventional functional fatigue (due to cyclic transformation) and of precipitation, 19 “transformationfree” thermal cycles between Tmin ¼ 30 C (i.e., well above Af, ensuring that no transformation occurred during this experiment; see dashed arrows in Fig. 1) and Tmax ¼ 250 C were performed on an additional DSC specimen. This specimen was subsequently subjected to one complete thermal cycle to determine the transformation behavior after the transformation-free cyclic heat treatment. The as-processed material, and the thermally cycled DSC specimens, were carefully analyzed by TEM. The TEM foils were prepared by mechanical grinding, followed by electro-polishing at T ¼ 19 C and U ¼ 19 V, using 95% acetic acid and 5% perchloric acid. The TEM investigations were performed on a Philips CM 20 instrument operating at 200 kV. 3. Results and discussion 3.1. Effect of Tmax on functional stability during thermal cycling The DSC curves (exothermic heat flow dq/dT as a function of temperature) during thermal cycling with different Tmax-values are summarized in Fig. 2. Since the ranges of transition temperatures for NiTi in all cases lie in a temperature regime much narrower than that between Tmin and Tmax, the DSC data are shown from 75 C to 75 C only. Outside this range, heat flows were constant in all measurements. To better illustrate the general evolution of the transformation behavior for the different specimens, cycles 1, 5, 10, and 20 (and for Tmax ¼ 250 C, also cycle 30) are highlighted by full lines in Fig. 2. The as-processed material initially exhibits a one-stage transformation with single peaks on cooling and heating representing the B2 / B190 , and the B190 / B2 transformations, respectively. In all three graphs, one can observe a shift towards lower temperatures of the PTTs
associated with the initial, one-stage, transformation. Moreover, for all experiments, the transformation behavior changes to a two-step transformation during cooling. The occurrence of the intermediate R-phase (associated with the first, smaller peak during cooling) is first evident in the 9th cycle for Tmax ¼ 150 C; it is observed from the 6th cycle onwards for Tmax ¼ 200 C, and already from the 3rd cycle onward for Tmax ¼ 250 C. At the highest Tmax (Fig. 2c), a two-stage reverse transformation can also be observed during heating from the 9th cycle onwards. For further analysis, the evolution of PTTs during the three thermal cycling experiments is presented in Fig. 3. Fig. 3aec show the PTTs associated with the martensitic forward transformations during cooling, and Fig. 3def show the corresponding PTTs during heating, as a function of the number of cycles. With increasing cycle number, stronger cooling is generally required to trigger the formation of B190 martensite. As a consequence, the R-phase transformation becomes a thermodynamically favorable intermediate path. Fig. 3aec also demonstrate that the R-phase peak temperatures are shifted towards higher temperatures with increasing cycle number. This process results in an increasing gap between the B2 / R and R / B190 peaks (i.e., a clear separation of both phase transformations). Most importantly, the PTT data sets in Fig. 3 illustrate that there is a considerable effect of Tmax on the functional properties: with increasing maximum temperatures, the martensite peaks (B2 / B190 and R / B190 during cooling, and B190 / B2 and B190 / R during heating) are more strongly shifted towards lower temperatures. This effect is stronger during the forward transformation. At higher Tmax, functional degradation occurs faster and is therefore more pronounced after similar cycle numbers. This is, for instance, evident from the first occurrence of the R-phase. When an R-phase peak can be distinguished for the first time, the Rp temperature is quite similar in all three experiments (Rp ¼ 13.5 1.5 C). However, the first occurrence of the R-phase is observed earlier in experiments with higher Tmax. Moreover, as functional evolution is more pronounced at higher Tmax, Rp-values change more strongly with increasing cycle
Fig. 2. DSC data during thermal cycling between Tmin ¼ 100 C and different maximum temperatures Tmax: (a) 150 C, (b) 200 C, and (c) 250 C. The heat flow curves have been truncated at 75 C and remain constant outside this temperature range.
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
1175
Fig. 3. Transformation start, peak and finish temperatures as a function of the number of cycles, (a)e(c) during cooling, (d)e(f) during heating.
number; after 20 thermal cycles, Rp-values of approx. 11, 7, and 12 C are observed for Tmax ¼ 150, 200, and 250 C, respectively. Similarly, the other PTTs decrease more strongly with increasing Tmax. At Tmax ¼ 250 C, the shift of the martensite peaks is so pronounced that it even allows for a well-defined two-step reverse transformation during heating. The general trends observed here (decreasing PTTs associated with the martensite peaks and occurrence of the R-phase) are in full agreement with the literature [10e12]. The commonly accepted view behind this special type of shape memory degradation is that the formation and accumulation of defects (dislocations) occurs due to a local misfit between the phases during the cyclic phase transformation. In a recent study, it was shown that the crystallographic compatibility can be improved by alloying copper (substituting nickel), which results in a better functional stability because less defects are introduced during cycling [9]. Generally, the increase in dislocation density with increasing number of cycles hinders the martensitic transition, and therefore, stronger cooling is required to trigger the B2 / B190 transformation. Moreover, the formation of the R-phase (which is associated with smaller transformation strains than the formation of B190 martensite) is less affected by an increased dislocation density. The shift of the R-phase peaks towards higher temperatures with increasing cycle number may indicate that the internal stresses due to micro-plasticity further promote the nucleation of the R-phase at higher temperatures.
While this classical view fully rationalizes the general evolution of the DSC data presented in Figs. 2 and 3, we emphasize that it cannot account for the most prominent result of our study, namely, the effect of Tmax on the transformation behavior of the Ni-rich specimens studied here. If functional degradation during thermal cycling was simply related to an increasing dislocation density, there should be little or no effect of Tmax on the cyclic functional behavior at all, as long as all experiments ensured that the specimen was fully transformed in every thermal cycle (this is clearly the case for all data presented here, see Fig. 2). Indeed, one would at the most expect an opposite effect, where recovery (i.e., a decreasing dislocation density) during heating to higher temperatures might result in less pronounced functional fatigue; but given the relatively low temperatures used in the present study, such an effect is quite unlikely. Instead, our findings point out that an additional, temperature-dependent mechanism contributes to functional degradation. In the remainder of this paper, we demonstrate, both indirectly and by direct TEM observations, that precipitation of the Ni4Ti3 phase in the temperature interval near Tmax results in the functional changes observed in our DSC data. 3.2. Effect of transformation-free cycling and ageing heat treatment Some important indirect information on the microstructural mechanism behind the newly observed effect of Tmax e and further
1176
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
evidence that an increasing dislocation density alone cannot account for it e is provided by our additional, “transformation-free” DSC experiment. In this experiment, thermal cycling was performed in the temperature range near Tmax ¼ 250 C (using the same heating/cooling rates of 10 K/min, and hold times of 3 min, as before), but completely avoiding the martensitic phase transformation by increasing Tmin to 30 C, i.e., above Af (see also Figs. 1 and 2). While the effect of ‘cyclic ageing’ experiments with varying temperatures in the classical ageing temperature regime (w670e750 K) was investigated in [39], a cyclic experiment at the low temperatures used here has to our best knowledge not been performed before. Transformation-free thermal cycling in the fully austenitic state was performed for 19 cycles in the DSC apparatus. The DSC data (not presented here) simply showed constant heat flows during heating and cooling and thus indicated that no martensitic transformation occurred. During the 20th thermal cycle, the transformation behavior was recorded by first cooling to below Mf and then heating to above Af. The corresponding DSC curve is shown in Fig. 4, where it is compared to the initial material behavior prior to thermal cycling. Although the specimen is not subjected to cyclic martensitic transformations in the transformation-free cycling experiment, its transformation behavior is considerably altered. The DSC peaks after transformation-free cycling are not as well defined, but similar to the material behavior after several regular transformation cycles (see Figs. 2 and 3), a two-step transformation with relatively broad peaks is observed, and the PTTs of the martensite peaks are shifted towards lower temperatures. Again, since no martensitic transition takes place during the first 19 transition-free thermal cycles, under classical understanding of dislocationcontrolled functional degradation one should expect the 20th thermal cycle to be equivalent to the 1st regular thermal cycle for Tmax ¼ 250 C (dashed curve in Fig. 4). In contrast, an additional thermal process obviously contributes to microstructural evolution in the temperature regime near higher Tmax-values. Considering that the alloy studied here is Ni-rich, a likely candidate mechanism is precipitation of Ni4Ti3 particles, in particular in the light of previous results [37] that show that nano-scale precipitates can be formed even at the relatively low temperatures similar to Tmax in our cyclic DSC experiments.
Fig. 4. After 19 transformation-free cycles (Tmax ¼ 250 C), the transformation behavior of the Ni-rich NiTi specimen clearly exhibits functional fatigue. The martensite peak during cooling and the austenite peak during heating are shifted towards lower temperatures, and a second peak indicates the occurrence of the Rphase during cooling.
To further corroborate this hypothesis, one additional specimen was placed into an evacuated Quartz tube, using Ti foil as getter material to further reduce oxidation, and annealed in a furnace at 240 C for 4.5 h, followed by water-quenching. The reasoning behind this choice of heat treatment parameters is as follows: Precipitation is diffusion-controlled and will hence be most active in the temperature range close to Tmax. In our cyclic experiments with Tmax ¼ 250 C, the specimen temperature is in the temperature range between 220 and 250 C for 9 min in every thermal cycle (3 min each during heating and cooling at 10 K/min, and 3 min during holding time at 250 C). The average temperature in this time interval is 240 C, and a total ageing time of 4.5 h corresponds to 30 cycles (at 9 min per cycle). The DSC curve of the aged specimen is compared to the transformation behavior after 20 and 30 thermal cycles with Tmax ¼ 250 C in Fig. 5. Similar to the transformation-free cycling measurement, the transformation behavior of the aged specimen differs considerably from the as-processed material behavior. As no cyclic transformation, and hence no micro-plasticity, occurred during ageing, the DSC curve in Fig. 5 clearly indicates that Ni4Ti3 precipitation occurs in the studied alloy (with a Ni-concentration of 50.7 at.%) at T ¼ 240 C, which agrees well with the timeetemperaturee transformation diagram published in [37]. The peaks associated with the two-stage transformation during cooling and the single peak during heating are well-defined and the overall DSC curve very closely resembles the cyclic data, Fig. 2c. Interestingly, the aged specimen’s PTTs correspond best to the 20th thermal DSC cycle, as opposed to the 30th thermal cycle that was initially targeted with the heat treatment. This deviation may well result from our very simple, non-weighted averaging procedure used to estimate appropriate ageing conditions. But more importantly, the data presented in Fig. 5 clearly demonstrates that both processes, classical functional fatigue due to an increasing dislocation density during cyclic forward and reverse transformations, and precipitation of Ni4Ti3, contribute to the functional degradation at Tmax near 250 C. 3.3. Precipitation during thermal cycling of NiTi The relatively modest differences between cyclic measurements with Tmax ¼ 150 and 200 C (Figs. 2 and 3) show that the
Fig. 5. The effect of a representative heat treatment (T ¼ 240 C, t ¼ 4.5 h) is very similar to that of pure thermal cycling (Tmax ¼ 250 C). The functional changes indicate that precipitation occurs even at the relatively low temperatures used in the thermal cycling experiments.
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
evolution of functional characteristics due to cyclic transformation and an increasing dislocation density is more dominant than precipitation at lower temperatures, while nucleation of Ni4Ti3 considerably affects functional degradation near Tmax ¼ 250 C. Incidentally, Tmax ¼ 250 C is about one third of the melting temperature of NiTi (Tmelt ¼ 1310 C, [40]), and diffusioncontrolled processes generally become relevant near this temperature range and above [41]. Therefore, our DSC results indirectly suggest the diffusion-controlled thermal formation of Ni4Ti3 particles during thermal cycling with elevated Tmax. Motivated by these indirect findings, a careful TEM analysis was performed on the as-processed material and on the thermally cycled DSC specimens. Fig. 6aec show selected area electron diffraction patterns of the [210] zone of B2 austenite for the three most interesting material states (as-processed, after 20 cycles with Tmax ¼ 200 C, and after 30 cycles with Tmax ¼ 250 C). Fig. 6d shows a schematic representation of the expected diffraction image for B2 austenite also containing R-phase and Ni4Ti3 particles. All three specimens considered here (Fig. 6aec) are clearly austenitic. In addition to the austenitic [210] reflections, small, diffuse intensities can also be noticed in Fig. 6b and c: In Fig. 6b (specimen after 20 thermal cycles with Tmax ¼ 200 C), one can also identify 1/3 <120> reflections of the R-phase [42]. These spots are shown as black rings in Fig. 6d. The presence of some R-phase is expected at room temperature after thermal cycling (see also Fig. 2b). The diffraction pattern after thermal cycling with Tmax ¼ 250 C (Fig. 6c) also indicates the
1177
presence of the R-phase. The corresponding peaks are more clearly visible than in Fig. 6b. Moreover, the diffraction image also exhibits 1/7 <321> reflections that clearly indicate the presence of the Ni4Ti3 phase (the corresponding reflections are schematically shown as grey circles in Fig. 6d). Bright field images of the Ni-rich precipitates could not to be captured under conventional usage of TEM. This is in line with the observations of Kim and Miyazaki that, under the thermal conditions prevalent here, Ni4Ti3 precipitates with diameters below 15 nm are formed [37]. Our findings confirm that, despite their small size, these particles significantly influence the functional properties of NiTi during thermal cycling. Considering furthermore the cyclic data in Figs. 2 and 3, it is obvious that nucleation of the nano-scale precipitates already occurs during the very first thermal cycles. While longer ageing times are necessary for growth of the precipitates, the driving force for nucleation is present from the very beginning, and the fast nucleation of Ni4Ti3 particles results in the stronger decrease of PTTs and the earlier observation of the R-phase at higher Tmax. This fast nucleation after very short times indicates that the cyclic accumulation of defects during thermal cycling may well provide an increase of the number of preferential nucleation sites, and therefore speed up Ni4Ti3 precipitation compared to conventional ageing experiments. This interpretation also means that the effect of Tmax on the transformation characteristics during thermal cycling is expected to be even more dominant in alloys with higher Ni-concentrations, such as 50.9 at.%, which is commonly used to obtain pseudoelastic behavior.
Fig. 6. TEM diffraction patterns in the [210] zone: (a) as-processed material, (b) after 20 thermal cycles with Tmax ¼ 200 C, (c) after 30 thermal cycles with Tmax ¼ 250 C, (d) Schematic representation of the diffraction pattern in the [210] zone. Black circles represent reflexes of the B2 structure, open circles are associated with the 1/3 <120> reflexes of the R-phase, and grey circles represent reflexes of type 1/7 <321> associated with Ni4Ti3.
1178
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
4. Summary and conclusions Thermal cycling with different maximum temperatures (Tmax ¼ 150, 200 and 250 C) was carried out on a Ni-rich NiTi alloy (50.7 at.% Ni) in a differential scanning calorimeter. The evolution of phase transformation temperatures was analyzed for up to 30 cycles. In all cases, typical changes in phase transition temperatures (shifting of peaks towards lower temperatures) and a transition from single- to two-step transformations associated with the occurrence of the R-phase were observed. Most importantly, it was observed for the first time that the transformation characteristics also are a function of Tmax. Contrary to what one would expect from the classical studies on functional fatigue, functional degradation increases with increasing Tmax. Dislocation accumulation due to repeated cycling through the transformation range cannot account for this phenomenon. Indirect evidence from DSC measurements as well as complementary TEM analyses of thermally cycled specimens clearly demonstrate that precipitation of Ni4Ti3 particles occurs at increased Tmax and plays an important role for the transformation behavior in subsequent cycles. Traditionally, functional fatigue has been attributed to microplasticity during the repeated martensitic transformation, which, by definition, is diffusion-less. Here, we find that a rather substantial portion of the observed functional changes are in fact due to the diffusion-controlled mechanism of Ni4Ti3 precipitation even at (for heat treatments) relatively low maximum temperatures. The generation of dislocations, the precipitation of Ni4Ti3, and the interaction between these microstructural processes, have a combined effect on the functional properties of NiTi during thermal cycling. It is interesting to note that precipitation prior to thermal cycling results in an increase of functional stability, whereas precipitation during thermal cycling actually contributes to the effects typically associated with functional fatigue. While further work (for instance, comparing alloys with different Ni-concentrations, and high-resolution TEM to directly characterize the nano-scale precipitates) is necessary to obtain a more systematic understanding of this phenomenon, our first results presented here demonstrate that functional degradation during thermal cycling is more complex than previously thought. These observations must be taken into account for the interpretation (in terms of dislocation-based mechanisms, hysteresis, dissipation and fatigue) of fatigue data published previously in the literature. As thermal cycling experiments are also used to verify the functional stability of novel high-temperature shape memory alloys (for instance, ternary NiTi-based alloys with transformation temperatures above 120 C, [43,44]), similar effects may well become an important factor affecting functional fatigue of these materials. Acknowledgements This work was supported by the German Research Foundation (Deutsche Forschungsgemeinschaft, DFG) in the framework of the Collaborative Research Center SFB459 “Shape Memory Technology” and by funding through the Ruhr University’s Research Department “Integrity of small-scale systems/High Temperature Materials e IS3/ HTM”. MW gratefully acknowledges funding in the DFG’s Emmy Noether programme and by the North Rhine Westfalia Academy of Science. References [1] Heinen R, Hackl K, Windl W, Wagner MF-X. Microstructural evolution during multi-axial deformation of pseudoelastic NiTi studied by first principles based micromechanical modeling. Acta Mater 2009;57:3856e67.
[2] Wagner M, Richter J, Frenzel J, Grönemeyer D, Eggeler G. Design of a medical non-linear drilling device: the influence of twist and wear on the fatigue behaviour of NiTi wires subjected to bending rotation. Materialwiss Werkstofftech 2004;35:320e5. [3] Duerig T, Pelton A, Stöckel D. An overview of nitinol medical applications. Mater Sci Eng 1999;A273e275:149e60. [4] Pelton AR, Russell SM, DiCello J. The physical metallurgy of Nitinol for medical applications. JOM 2003;55:33e7. [5] Krone L, Mentz J, Bram M, Buchkremer HP, Stöver D, Wagner M, et al. 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 behaviour. Adv Eng Mater 2005;7:613e9. [6] Mentz J, Frenzel J, Wagner MF-X, Neuking K, Eggeler G, Buchkremer HP, et al. Powder metallurgical processing of NiTi shape memory alloys with elevated transformation temperatures. Mater Sci Eng 2008;A491:270e8. [7] Huang W. On the selection of shape memory alloys for actuators. Mater Des 2002;23:11e9. [8] Großmann C, Frenzel J, Sampath V, Depka T, Oppenkowski A, Somsen Ch, et al. Processing and property assessment of NiTi and NiTiCu shape memory actuator springs. Materialwiss Werkstofftech 2008;39:499e510. [9] Großmann C, Frenzel J, Sampath V, Depka T, Eggeler G. Elementary transformation and deformation processes and the cyclic stability of NiTi and NiTiCu shape memory spring actuators. Metall Mater Trans 2009; A40:2530e44. [10] Miyazaki Y, Igo Y, Otsuka K. Effect of thermal cycling on the transformation temperatures of TieNi alloys. Acta Metall 1986;34:2045e51. [11] Tadaki T, Nakata Y, Shimizu K. Thermal cycling effects in an aged Ni-rich TieNi shape memory alloy. Trans Japan Instit Metals 1987;28:883e90. [12] Matsumoto H. Transformation behaviour with thermal cycling in NiTi alloys. J Alloys Compd 2003;350:213e7. [13] Miyazaki S, Imai T, Igo Y, Otsuka K. Effectt of cyclic deformation on the pseudoelasticity characteristics of TieNi alloys. Met Trans 1986;A17: 115e20. [14] Gall K, Maier H. Cyclic deformation mechanisms in precipitated NiTi shape memory alloys. Acta Mater 2002;50:4643e57. [15] Hurley J, Ortega AM, Lechniak J, Gall K, Maier HJ. Structural evolution during the cycling of NiTi shape memory alloys. Z Metallkunde 2003;94:547e52. [16] Wagner MF-X, Nayan N, Ramamurty U. Healing of fatigue damage in NiTi shape memory alloys. J Phys D Appl Phys 2008;41:185408. [17] Morgan NB, Friend CM. A review of shape memory stability in NiTi alloys. J Physique IV 2001;11:325e32. [18] Scherngell H, Kneissl AC. Generation, development and degradation of the intrinsic two-way shape memory effect in different alloy systems. Acta Mater 2002;50:327e41. [19] Bhattacharya K, Conti S, Zanzotto G, Zimmer J. Crystal symmetry and the reversibility of martensitic transformations. Nature 2004;428:55e9. [20] Norfleet DM, Sarosi PM, Manchiraju S, Wagner MF-X, Uchic MD, Anderson PM, et al. Study of transformation-induced plasticity during pseudo-elastic deformation of NieTi microcrystals. Acta Mater 2009;57:3549e61. [21] Yawny A, Sade M, Eggeler G. Pseudoelastic cycling of ultra-fine-grained NiTi shape-memory wires. Z Metallkunde 2005;96:608e18. [22] Eggeler G, Hornbogen E, Yawny A, Heckmann A, Wagner M. Structural and functional fatigue of NiTi shape memory alloys. Mater Sci Eng 2004; A378:24e33. [23] Pfetzing J, Schaefer A, Somsen Ch, Wagner MF-X. Nanoindentation of pseudoelastic NiTi shape memory alloys: thermomechanical and microstructural aspects. Int J Mater Res 2009;100:936e42. [24] Wagner M, Sawaguchi TS, Kausträter G, Höffken D, Eggeler G. Structural fatigue of pseudoelastic NiTi shape memory wires. Mater Sci Eng 2004; A378:105e9. [25] Wagner MF-X, Eggeler G. New aspects of bending rotation fatigue in ultrafine-grained pseudo-elastic NiTi wires. Int J Mater Res 2006;97:1687e96. [26] Vaidyanathan R, Dunand DC, Ramamurty U. Fatigue crack-growth in shapememory NiTi and NiTieTiC composites. Mater Sci Eng 2000;A289:208e16. [27] Miyazaki S, Otsuka K. Development of shape memory alloys. ISIJ Int 1989;29:353e77. [28] Undisz A, Fink M, Rettenmayr M. Response of austenite finish temperature and phase transformation characteristics of thin medical-grade NieTi wire to short-time annealing. Scripta Mater 2008;59:979e82. [29] Jiang F, Liu Y, Yang H, Li L, Zheng Y. Effect of ageing treatment on the deformation behaviour of Ti-50.9 at.% Ni. Acta Mater 2009;57:4773e81. [30] Khalil-Allafi J, Dlouhy A, Eggeler G. Ni4Ti3-precipitation during aging of NiTi shape memory alloys and its influence on martensitic phase transformations. Acta Mater 2002;50:4255e74. [31] Dlouhy A, Khalil-Allafi J, Eggeler G. Multiple-step martensitic transformations in Ni-rich NiTi alloys e an in-situ transmission electron microscopy investigation. Phil Mag 2003;83:339e63. [32] Zheng Y, Jiang F, Li L, Yang H, Liu Y. Effect of ageing treatment on the transformation behaviour of Ti-50.9 at.% Ni alloy. Acta Mater 2008;56:736e45. [33] Wagner MF-X, Windl W. Elastic anisotropy of Ni4Ti3 from first principles. Scripta Mater 2009;60:207e10. [34] Tirry W, Schryvers D. High resolution TEM study of Ni4Ti3 precipitates in austenitic Ni51Ti49. Mater Sci 2004;A378:157e60. [35] Tirry W, Schryvers D. Quantitative determination of strain fields around Ni4Ti3 precipitates in NiTi. Acta Mater 2005;53:1041e9.
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179 [36] Tirry W, Schryvers D, Jorissen K, Lamoen D. Electron-diffraction structure refinement of Ni4Ti3 precipitates in Ni52Ti48. Acta Crystallogr 2006; B62:966e71. [37] Kim J, Miyazaki S. Effect of nano-scaled precipitates on shape memory behavior of Ti-50.9 at.%Ni alloy. Acta Mater 2005;53:4545e54. [38] Frenzel J, Zhang Z, Neuking K, Eggeler G. High quality vacuum induction melting of small quantities of NiTi shape memory alloys in graphite crucibles. J Alloys Compd 2004;385:214e23. [39] Jiang F, Li L, Zheng Y, Yang H, Liu Y. Cyclic ageing of Ti-50.8 at.% Ni alloy. Intermet 2007;16:394e8.
1179
[40] Murray J. NieTi (nickeletitanium). Binary alloy phase diagrams. Materials Park, OH, USA: ASM International; 1990. p. 2874e2876. [41] Nabarro F, de Villiers H. The physics of creep. London: Taylor & Francis; 1995. [42] Hara T, Ohba T, Okunishi E, Otsuka K. Structural study of R-phase in Ti50.23 at.%Ni and Ti-47.75 at.%Ni-1.50 at.%Fe alloys. Mater Trans JIM 1997;38:11e7. [43] Firstov G, van Humbeeck J, Koval Y. High-temperature shape memory alloys: some recent developments. Mater Sci Eng A 2004;378:2e10. [44] Firstov G, van Humbeeck J, Koval Y. High temperature shape memory alloys: problems and prospects. J Intell Mater Syst Struct 2006;17:1041e7.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Effect of low-temperature precipitation on the transformation characteristics of Ni-rich NiTi shape memory alloys during thermal cycling M.F.-X. Wagner a, b, *, S.R. Dey a, b, H. Gugel a, J. Frenzel a, Ch. Somsen a, G. Eggeler a, b a b
Institute of Materials Science and Engineering, Chemnitz University of Technology, Erfenschlager Str. 73, 09125 Chemnitz, Germany Research Department IS3/HTM e Integrity of Small-Scale Systems/High Temperature Materials, Ruhr University Bochum, Universitaetsstr. 150, 44801 Bochum, Germany
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 December 2009 Accepted 26 February 2010 Available online 1 April 2010
Thermal cycling of NiTi shape memory alloys is associated with functional fatigue: the characteristic phase transformation temperatures decrease with increasing number of cycles, and the transformation behavior changes from a single- to a two-stage martensitic transformation involving the intermediate R-phase. These effects are usually attributed to a gradual increase of dislocation density associated with micro-plasticity during repeated cycling through the transformation range. Here, these changes are shown to increase at a higher maximum temperature (in the fully austenitic state) during differential scanning calorimetric cycling of a Ni-rich alloy. Additional thermal cycling experiments without repeated phase transformations, and post-mortem microstructural observations by transmission electron microscopy, demonstrate that a relevant portion of functional fatigue is due to the formation of nanoscale Ni-rich precipitates of type Ni4Ti3 even at temperatures relatively close to the austenite finish temperature. These results show that both dislocation generation during the diffusion-less phase transformation, and diffusion-controlled nucleation and growth of Ni4Ti3 precipitates, can interact and contribute to the evolution of functional properties during thermal cycling of Ni-rich NiTi. Ó 2010 Elsevier Ltd. All rights reserved.
Keywords: A: Multiphase intermetallics B: Martensitic transformations B: Precipitates B: Shape-memory effects
1. Introduction NiTi shape memory alloys are widely used in medical implants and devices, where the reversible martensitic transformation between a more symmetric high-temperature phase (austenite; cubic B2 type crystal structure) and a less symmetric low temperature phase (martensite; monoclinic B190 type crystal structure) is exploited to obtain pseudoelastic material behavior [1e4]. Future applications are also expected in the field of actuators and devices that use thermal shape memory during repeated heating and cooling [5,6]. Actuator forces are generated as the NiTi component is heated and contracts during the reverse transformation from martensite to austenite [7e9]. Differential scanning calorimetry (DSC) can be used to assess the transformation behavior of a NiTi alloy. By slowly cooling and heating a DSC specimen between a maximum temperature Tmax above the austenite finish temperature, Af, and a minimum temperature Tmin below the martensite finish temperature, Mf, the material is fully transformed to martensite and back to austenite. The characteristic phase * Corresponding author. Institute of Materials Science and Engineering, Chemnitz University of Technology, Erfenschlager Str. 73, 09125 Chemnitz, Germany. Tel.: þ49 371 531 38683; fax: þ49 371 531 838683. E-mail address: [email protected] (M.F.-X. Wagner). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.02.048
transformation temperatures (PTTs) can be determined from the heat flow data (see Fig. 1; the forward martensitic transformation is exothermic, whereas the reverse transformation is endothermic). It is well known that thermal cycling affects the transformation characteristics of NiTi measured by DSC [10e16]. As shown schematically in Fig. 1, one typically observes decreasing PTTs with increasing number of DSC cycles [10e12]. Moreover, the R-phase, an intermediate phase that is also formed in a reversible, martensitic transformation from B2 austenite, can occur, resulting in a two-step phase transformation (B2 / R / B190 ). In engineering applications, the martensitic forward and reverse transformations must be repeatable for many cycles [17,18]. While it is in theory fully reversible [19], the phase transformation involves nucleation and growth of martensite into the adjacent, austenitic material. As the growing martensite interacts with defects and interfaces in the surrounding matrix, some local formation of dislocations (and hence irreversible plastic deformation) is inevitable [20e23]. Micro-plasticity in NiTi can lead to classical structural fatigue, where it ultimately results in crack formation and growth during mechanical cycling [24e26]. Moreover, functional fatigue is associated with a gradual degradation of shape memory behavior when NiTi is subjected to several forward and reverse transformation cycles. Even during pure thermal cycling without mechanical loading, the functional characteristics
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
1173
nucleation of Ni4Ti3 occurs more easily than particle growth in the early stages of precipitation, low-temperature/short-time ageing results in a high number of very fine, nano-scale precipitates (with particle diameters <10 nm, [37]). Despite a substantial body of work on functional fatigue on the one hand, and on the effect of heat treatments on the other hand, there is no systematic study that takes into account how Tmax may affect the elementary microstructural processes that contribute to the evolution of functional properties during thermal cycling. In this work, we investigate how functional changes and precipitation are interrelated when Tmax approaches the temperature regime where precipitates can form in a Ni-rich NiTi alloy. By analyzing the DSC data of up to 30 thermal cycles through the transformation regime, we document that functional degradation is more severe at higher Tmax. From this unexpected temperature dependency, and from complementary TEM analysis, we conclude that Ni4Ti3 precipitation occurs even at the relatively low temperatures associated with thermal cycling. We show that both processes (formation of dislocations and precipitation) may simultaneously affect functional degradation, and that precipitation of nano-scale Ni4Ti3 particles may well be a dominant mechanism when thermal cycling is performed at higher temperatures. Fig. 1. Schematic representation of the evolution of DSC data during thermal cycling of NiTi. With increasing cycle number N, the transformation temperatures are shifted towards lower temperatures, and the transformation sequence changes from a singleto a two-step transformation, involving the intermediate R-phase, during cooling. An individual phase transformation can be characterized by the start, peak, and finish temperatures (indices s, p, and f, indicated here for the martensite peak).
often change considerably because of relatively subtle changes in the microstructure [10e12]. In the light of functional fatigue studies, the changes in functional behavior and transformation sequence during simple thermal cycling are generally rationalized by the accumulation of dislocations and their effect on subsequent transformation cycles [13e16]. However, it is not clear whether other elementary processes can also affect functional degradation during thermal cycling. Interestingly, similar effects on the functional properties are observed when Ni-rich alloys are heat treated at constant, elevated temperatures (typically: T > 350 C, [27]). Then, occurrence of R-phase and decreasing PTTs are attributed to precipitation of the meta-stable Ni4Ti3 phase. Such heat treatments are beneficial because particle hardening increases the yield strength of austenite, which in turn contributes to better functional stability as less dislocations are formed during mechanical or thermal cycling. Recently, Undisz et al. [28] reported that shorttime heat treatments (annealing times of 2e30 min) considerably affect the PTTs and the pseudoelastic behavior of Ni-rich NiTi wires. Moreover, Jiang et al. [29] provided a detailed study on how precipitation limits the stress- and temperature regimes for pseudoelastic and shape memory behavior of an alloy with 50.9 at.% Ni. Heterogeneous precipitation of Ni4Ti3 is also one of the key mechanisms underlying multi-stage transformations that have recently attracted considerable attention [30e32]. This interest has motivated detailed studies on the properties and precipitation behavior of Ni4Ti3 [33e36]. In the context of the present work, a careful investigation by Kim and Miyazaki [37] is especially important as it provides timeetemperatureetransformation data on, and a detailed microstructural analysis of, Ni4Ti3 precipitation in solution-annealed Ni-rich (50.9 at.% Ni) NiTi. They [37] found that precipitation occurs at temperatures as low as 200 C (but only after long ageing times >100 h), whereas precipitates could already be observed by transmission electron microscopy (TEM) after ageing at 300 C for 20 min. They documented that ageing at low temperatures (<330 C) is most effective for improving the shape memory stability of NiTi. Moreover they demonstrated that, as
2. Experiments An ingot of binary NiTi with a nominal composition of 50.7 at.% Ni was produced by vacuum induction melting. Details on the individual processing steps for high purity melting and thermomechanical processing have been published elsewhere [38]. For chemical homogenization, the ingot was annealed at 850 C for 1 h in argon and subsequently quenched in water. A rectangular piece was cut from the ingot and hot-rolled in several steps (with intermediate heating to 850 C) to a final thickness of 2 mm. The hotrolled sheet was then again annealed at 850 C for 1 h under argon atmosphere. The grain size of the sheet material was determined from optical micrographs after mechanical grinding and chemical etching with Beraha-I solution. The observations were made over the horizontal plane of the sample containing the rolling and transverse directions. Several optical micrographs were acquired and the grain size was statistically determined using the line intersection method. The as-processed microstructure contained slightly elongated grains along the rolling direction, with an average grain size of w42 mm. As the dimensions of the DSC specimens are w2.5 2.5 2 mm3 (this corresponds to a specimen mass of about 80 mg), the grain size is sufficiently small to ensure that the observed transformation behavior is fully representative of the polycrystalline sheet material. Thermal cycling was performed in a DSC instrument of type 2920 CE (TA Instruments) at a constant heating/cooling rate of 10 K/ min under an argon/helium atmosphere (to minimize oxidation in the long-term cycling experiments). DSC cycles were performed on three specimens with constant Tmin (100 C) but with varying Tmax (150, 200 and 250 C). 20 DSC cycles were carried out for Tmax ¼ 150 and 200 C, and 30 cycles for Tmax ¼ 250 C. The first cycle consisted of (1) heating the fully austenitic specimens from ambient temperature to Tmax, (2) holding the specimen temperature constant at Tmax for 3 min, (3) cooling the specimen to Tmin, (4) holding the specimen temperature constant at Tmin for 3 min, and (5) heating the specimen again to Tmax. Subsequent cycles were performed by repeating steps (2)e(4), see also Fig. 1. Note that using hold times (often up to 10 min) at Tmin and Tmax is a standard procedure during DSC testing, with the aim of ensuring thermal equilibrium of the specimen and the reference sample. The DSC data allows detailed analysis of PTTs, i.e., the martensite, austenite, and R-phase start and finish temperatures Ms, Mf, As, Af, Rs, Rf,
1174
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
respectively. Moreover, the general evolution of the transformation peaks as a function of the number of cycles was documented by plotting the temperatures of maximum/minimum heat flows, Mp, Ap, and Rp (which are referred to as peak temperatures hereafter). To separate the effects of conventional functional fatigue (due to cyclic transformation) and of precipitation, 19 “transformationfree” thermal cycles between Tmin ¼ 30 C (i.e., well above Af, ensuring that no transformation occurred during this experiment; see dashed arrows in Fig. 1) and Tmax ¼ 250 C were performed on an additional DSC specimen. This specimen was subsequently subjected to one complete thermal cycle to determine the transformation behavior after the transformation-free cyclic heat treatment. The as-processed material, and the thermally cycled DSC specimens, were carefully analyzed by TEM. The TEM foils were prepared by mechanical grinding, followed by electro-polishing at T ¼ 19 C and U ¼ 19 V, using 95% acetic acid and 5% perchloric acid. The TEM investigations were performed on a Philips CM 20 instrument operating at 200 kV. 3. Results and discussion 3.1. Effect of Tmax on functional stability during thermal cycling The DSC curves (exothermic heat flow dq/dT as a function of temperature) during thermal cycling with different Tmax-values are summarized in Fig. 2. Since the ranges of transition temperatures for NiTi in all cases lie in a temperature regime much narrower than that between Tmin and Tmax, the DSC data are shown from 75 C to 75 C only. Outside this range, heat flows were constant in all measurements. To better illustrate the general evolution of the transformation behavior for the different specimens, cycles 1, 5, 10, and 20 (and for Tmax ¼ 250 C, also cycle 30) are highlighted by full lines in Fig. 2. The as-processed material initially exhibits a one-stage transformation with single peaks on cooling and heating representing the B2 / B190 , and the B190 / B2 transformations, respectively. In all three graphs, one can observe a shift towards lower temperatures of the PTTs
associated with the initial, one-stage, transformation. Moreover, for all experiments, the transformation behavior changes to a two-step transformation during cooling. The occurrence of the intermediate R-phase (associated with the first, smaller peak during cooling) is first evident in the 9th cycle for Tmax ¼ 150 C; it is observed from the 6th cycle onwards for Tmax ¼ 200 C, and already from the 3rd cycle onward for Tmax ¼ 250 C. At the highest Tmax (Fig. 2c), a two-stage reverse transformation can also be observed during heating from the 9th cycle onwards. For further analysis, the evolution of PTTs during the three thermal cycling experiments is presented in Fig. 3. Fig. 3aec show the PTTs associated with the martensitic forward transformations during cooling, and Fig. 3def show the corresponding PTTs during heating, as a function of the number of cycles. With increasing cycle number, stronger cooling is generally required to trigger the formation of B190 martensite. As a consequence, the R-phase transformation becomes a thermodynamically favorable intermediate path. Fig. 3aec also demonstrate that the R-phase peak temperatures are shifted towards higher temperatures with increasing cycle number. This process results in an increasing gap between the B2 / R and R / B190 peaks (i.e., a clear separation of both phase transformations). Most importantly, the PTT data sets in Fig. 3 illustrate that there is a considerable effect of Tmax on the functional properties: with increasing maximum temperatures, the martensite peaks (B2 / B190 and R / B190 during cooling, and B190 / B2 and B190 / R during heating) are more strongly shifted towards lower temperatures. This effect is stronger during the forward transformation. At higher Tmax, functional degradation occurs faster and is therefore more pronounced after similar cycle numbers. This is, for instance, evident from the first occurrence of the R-phase. When an R-phase peak can be distinguished for the first time, the Rp temperature is quite similar in all three experiments (Rp ¼ 13.5 1.5 C). However, the first occurrence of the R-phase is observed earlier in experiments with higher Tmax. Moreover, as functional evolution is more pronounced at higher Tmax, Rp-values change more strongly with increasing cycle
Fig. 2. DSC data during thermal cycling between Tmin ¼ 100 C and different maximum temperatures Tmax: (a) 150 C, (b) 200 C, and (c) 250 C. The heat flow curves have been truncated at 75 C and remain constant outside this temperature range.
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
1175
Fig. 3. Transformation start, peak and finish temperatures as a function of the number of cycles, (a)e(c) during cooling, (d)e(f) during heating.
number; after 20 thermal cycles, Rp-values of approx. 11, 7, and 12 C are observed for Tmax ¼ 150, 200, and 250 C, respectively. Similarly, the other PTTs decrease more strongly with increasing Tmax. At Tmax ¼ 250 C, the shift of the martensite peaks is so pronounced that it even allows for a well-defined two-step reverse transformation during heating. The general trends observed here (decreasing PTTs associated with the martensite peaks and occurrence of the R-phase) are in full agreement with the literature [10e12]. The commonly accepted view behind this special type of shape memory degradation is that the formation and accumulation of defects (dislocations) occurs due to a local misfit between the phases during the cyclic phase transformation. In a recent study, it was shown that the crystallographic compatibility can be improved by alloying copper (substituting nickel), which results in a better functional stability because less defects are introduced during cycling [9]. Generally, the increase in dislocation density with increasing number of cycles hinders the martensitic transition, and therefore, stronger cooling is required to trigger the B2 / B190 transformation. Moreover, the formation of the R-phase (which is associated with smaller transformation strains than the formation of B190 martensite) is less affected by an increased dislocation density. The shift of the R-phase peaks towards higher temperatures with increasing cycle number may indicate that the internal stresses due to micro-plasticity further promote the nucleation of the R-phase at higher temperatures.
While this classical view fully rationalizes the general evolution of the DSC data presented in Figs. 2 and 3, we emphasize that it cannot account for the most prominent result of our study, namely, the effect of Tmax on the transformation behavior of the Ni-rich specimens studied here. If functional degradation during thermal cycling was simply related to an increasing dislocation density, there should be little or no effect of Tmax on the cyclic functional behavior at all, as long as all experiments ensured that the specimen was fully transformed in every thermal cycle (this is clearly the case for all data presented here, see Fig. 2). Indeed, one would at the most expect an opposite effect, where recovery (i.e., a decreasing dislocation density) during heating to higher temperatures might result in less pronounced functional fatigue; but given the relatively low temperatures used in the present study, such an effect is quite unlikely. Instead, our findings point out that an additional, temperature-dependent mechanism contributes to functional degradation. In the remainder of this paper, we demonstrate, both indirectly and by direct TEM observations, that precipitation of the Ni4Ti3 phase in the temperature interval near Tmax results in the functional changes observed in our DSC data. 3.2. Effect of transformation-free cycling and ageing heat treatment Some important indirect information on the microstructural mechanism behind the newly observed effect of Tmax e and further
1176
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
evidence that an increasing dislocation density alone cannot account for it e is provided by our additional, “transformation-free” DSC experiment. In this experiment, thermal cycling was performed in the temperature range near Tmax ¼ 250 C (using the same heating/cooling rates of 10 K/min, and hold times of 3 min, as before), but completely avoiding the martensitic phase transformation by increasing Tmin to 30 C, i.e., above Af (see also Figs. 1 and 2). While the effect of ‘cyclic ageing’ experiments with varying temperatures in the classical ageing temperature regime (w670e750 K) was investigated in [39], a cyclic experiment at the low temperatures used here has to our best knowledge not been performed before. Transformation-free thermal cycling in the fully austenitic state was performed for 19 cycles in the DSC apparatus. The DSC data (not presented here) simply showed constant heat flows during heating and cooling and thus indicated that no martensitic transformation occurred. During the 20th thermal cycle, the transformation behavior was recorded by first cooling to below Mf and then heating to above Af. The corresponding DSC curve is shown in Fig. 4, where it is compared to the initial material behavior prior to thermal cycling. Although the specimen is not subjected to cyclic martensitic transformations in the transformation-free cycling experiment, its transformation behavior is considerably altered. The DSC peaks after transformation-free cycling are not as well defined, but similar to the material behavior after several regular transformation cycles (see Figs. 2 and 3), a two-step transformation with relatively broad peaks is observed, and the PTTs of the martensite peaks are shifted towards lower temperatures. Again, since no martensitic transition takes place during the first 19 transition-free thermal cycles, under classical understanding of dislocationcontrolled functional degradation one should expect the 20th thermal cycle to be equivalent to the 1st regular thermal cycle for Tmax ¼ 250 C (dashed curve in Fig. 4). In contrast, an additional thermal process obviously contributes to microstructural evolution in the temperature regime near higher Tmax-values. Considering that the alloy studied here is Ni-rich, a likely candidate mechanism is precipitation of Ni4Ti3 particles, in particular in the light of previous results [37] that show that nano-scale precipitates can be formed even at the relatively low temperatures similar to Tmax in our cyclic DSC experiments.
Fig. 4. After 19 transformation-free cycles (Tmax ¼ 250 C), the transformation behavior of the Ni-rich NiTi specimen clearly exhibits functional fatigue. The martensite peak during cooling and the austenite peak during heating are shifted towards lower temperatures, and a second peak indicates the occurrence of the Rphase during cooling.
To further corroborate this hypothesis, one additional specimen was placed into an evacuated Quartz tube, using Ti foil as getter material to further reduce oxidation, and annealed in a furnace at 240 C for 4.5 h, followed by water-quenching. The reasoning behind this choice of heat treatment parameters is as follows: Precipitation is diffusion-controlled and will hence be most active in the temperature range close to Tmax. In our cyclic experiments with Tmax ¼ 250 C, the specimen temperature is in the temperature range between 220 and 250 C for 9 min in every thermal cycle (3 min each during heating and cooling at 10 K/min, and 3 min during holding time at 250 C). The average temperature in this time interval is 240 C, and a total ageing time of 4.5 h corresponds to 30 cycles (at 9 min per cycle). The DSC curve of the aged specimen is compared to the transformation behavior after 20 and 30 thermal cycles with Tmax ¼ 250 C in Fig. 5. Similar to the transformation-free cycling measurement, the transformation behavior of the aged specimen differs considerably from the as-processed material behavior. As no cyclic transformation, and hence no micro-plasticity, occurred during ageing, the DSC curve in Fig. 5 clearly indicates that Ni4Ti3 precipitation occurs in the studied alloy (with a Ni-concentration of 50.7 at.%) at T ¼ 240 C, which agrees well with the timeetemperaturee transformation diagram published in [37]. The peaks associated with the two-stage transformation during cooling and the single peak during heating are well-defined and the overall DSC curve very closely resembles the cyclic data, Fig. 2c. Interestingly, the aged specimen’s PTTs correspond best to the 20th thermal DSC cycle, as opposed to the 30th thermal cycle that was initially targeted with the heat treatment. This deviation may well result from our very simple, non-weighted averaging procedure used to estimate appropriate ageing conditions. But more importantly, the data presented in Fig. 5 clearly demonstrates that both processes, classical functional fatigue due to an increasing dislocation density during cyclic forward and reverse transformations, and precipitation of Ni4Ti3, contribute to the functional degradation at Tmax near 250 C. 3.3. Precipitation during thermal cycling of NiTi The relatively modest differences between cyclic measurements with Tmax ¼ 150 and 200 C (Figs. 2 and 3) show that the
Fig. 5. The effect of a representative heat treatment (T ¼ 240 C, t ¼ 4.5 h) is very similar to that of pure thermal cycling (Tmax ¼ 250 C). The functional changes indicate that precipitation occurs even at the relatively low temperatures used in the thermal cycling experiments.
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
evolution of functional characteristics due to cyclic transformation and an increasing dislocation density is more dominant than precipitation at lower temperatures, while nucleation of Ni4Ti3 considerably affects functional degradation near Tmax ¼ 250 C. Incidentally, Tmax ¼ 250 C is about one third of the melting temperature of NiTi (Tmelt ¼ 1310 C, [40]), and diffusioncontrolled processes generally become relevant near this temperature range and above [41]. Therefore, our DSC results indirectly suggest the diffusion-controlled thermal formation of Ni4Ti3 particles during thermal cycling with elevated Tmax. Motivated by these indirect findings, a careful TEM analysis was performed on the as-processed material and on the thermally cycled DSC specimens. Fig. 6aec show selected area electron diffraction patterns of the [210] zone of B2 austenite for the three most interesting material states (as-processed, after 20 cycles with Tmax ¼ 200 C, and after 30 cycles with Tmax ¼ 250 C). Fig. 6d shows a schematic representation of the expected diffraction image for B2 austenite also containing R-phase and Ni4Ti3 particles. All three specimens considered here (Fig. 6aec) are clearly austenitic. In addition to the austenitic [210] reflections, small, diffuse intensities can also be noticed in Fig. 6b and c: In Fig. 6b (specimen after 20 thermal cycles with Tmax ¼ 200 C), one can also identify 1/3 <120> reflections of the R-phase [42]. These spots are shown as black rings in Fig. 6d. The presence of some R-phase is expected at room temperature after thermal cycling (see also Fig. 2b). The diffraction pattern after thermal cycling with Tmax ¼ 250 C (Fig. 6c) also indicates the
1177
presence of the R-phase. The corresponding peaks are more clearly visible than in Fig. 6b. Moreover, the diffraction image also exhibits 1/7 <321> reflections that clearly indicate the presence of the Ni4Ti3 phase (the corresponding reflections are schematically shown as grey circles in Fig. 6d). Bright field images of the Ni-rich precipitates could not to be captured under conventional usage of TEM. This is in line with the observations of Kim and Miyazaki that, under the thermal conditions prevalent here, Ni4Ti3 precipitates with diameters below 15 nm are formed [37]. Our findings confirm that, despite their small size, these particles significantly influence the functional properties of NiTi during thermal cycling. Considering furthermore the cyclic data in Figs. 2 and 3, it is obvious that nucleation of the nano-scale precipitates already occurs during the very first thermal cycles. While longer ageing times are necessary for growth of the precipitates, the driving force for nucleation is present from the very beginning, and the fast nucleation of Ni4Ti3 particles results in the stronger decrease of PTTs and the earlier observation of the R-phase at higher Tmax. This fast nucleation after very short times indicates that the cyclic accumulation of defects during thermal cycling may well provide an increase of the number of preferential nucleation sites, and therefore speed up Ni4Ti3 precipitation compared to conventional ageing experiments. This interpretation also means that the effect of Tmax on the transformation characteristics during thermal cycling is expected to be even more dominant in alloys with higher Ni-concentrations, such as 50.9 at.%, which is commonly used to obtain pseudoelastic behavior.
Fig. 6. TEM diffraction patterns in the [210] zone: (a) as-processed material, (b) after 20 thermal cycles with Tmax ¼ 200 C, (c) after 30 thermal cycles with Tmax ¼ 250 C, (d) Schematic representation of the diffraction pattern in the [210] zone. Black circles represent reflexes of the B2 structure, open circles are associated with the 1/3 <120> reflexes of the R-phase, and grey circles represent reflexes of type 1/7 <321> associated with Ni4Ti3.
1178
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179
4. Summary and conclusions Thermal cycling with different maximum temperatures (Tmax ¼ 150, 200 and 250 C) was carried out on a Ni-rich NiTi alloy (50.7 at.% Ni) in a differential scanning calorimeter. The evolution of phase transformation temperatures was analyzed for up to 30 cycles. In all cases, typical changes in phase transition temperatures (shifting of peaks towards lower temperatures) and a transition from single- to two-step transformations associated with the occurrence of the R-phase were observed. Most importantly, it was observed for the first time that the transformation characteristics also are a function of Tmax. Contrary to what one would expect from the classical studies on functional fatigue, functional degradation increases with increasing Tmax. Dislocation accumulation due to repeated cycling through the transformation range cannot account for this phenomenon. Indirect evidence from DSC measurements as well as complementary TEM analyses of thermally cycled specimens clearly demonstrate that precipitation of Ni4Ti3 particles occurs at increased Tmax and plays an important role for the transformation behavior in subsequent cycles. Traditionally, functional fatigue has been attributed to microplasticity during the repeated martensitic transformation, which, by definition, is diffusion-less. Here, we find that a rather substantial portion of the observed functional changes are in fact due to the diffusion-controlled mechanism of Ni4Ti3 precipitation even at (for heat treatments) relatively low maximum temperatures. The generation of dislocations, the precipitation of Ni4Ti3, and the interaction between these microstructural processes, have a combined effect on the functional properties of NiTi during thermal cycling. It is interesting to note that precipitation prior to thermal cycling results in an increase of functional stability, whereas precipitation during thermal cycling actually contributes to the effects typically associated with functional fatigue. While further work (for instance, comparing alloys with different Ni-concentrations, and high-resolution TEM to directly characterize the nano-scale precipitates) is necessary to obtain a more systematic understanding of this phenomenon, our first results presented here demonstrate that functional degradation during thermal cycling is more complex than previously thought. These observations must be taken into account for the interpretation (in terms of dislocation-based mechanisms, hysteresis, dissipation and fatigue) of fatigue data published previously in the literature. As thermal cycling experiments are also used to verify the functional stability of novel high-temperature shape memory alloys (for instance, ternary NiTi-based alloys with transformation temperatures above 120 C, [43,44]), similar effects may well become an important factor affecting functional fatigue of these materials. Acknowledgements This work was supported by the German Research Foundation (Deutsche Forschungsgemeinschaft, DFG) in the framework of the Collaborative Research Center SFB459 “Shape Memory Technology” and by funding through the Ruhr University’s Research Department “Integrity of small-scale systems/High Temperature Materials e IS3/ HTM”. MW gratefully acknowledges funding in the DFG’s Emmy Noether programme and by the North Rhine Westfalia Academy of Science. References [1] Heinen R, Hackl K, Windl W, Wagner MF-X. Microstructural evolution during multi-axial deformation of pseudoelastic NiTi studied by first principles based micromechanical modeling. Acta Mater 2009;57:3856e67.
[2] Wagner M, Richter J, Frenzel J, Grönemeyer D, Eggeler G. Design of a medical non-linear drilling device: the influence of twist and wear on the fatigue behaviour of NiTi wires subjected to bending rotation. Materialwiss Werkstofftech 2004;35:320e5. [3] Duerig T, Pelton A, Stöckel D. An overview of nitinol medical applications. Mater Sci Eng 1999;A273e275:149e60. [4] Pelton AR, Russell SM, DiCello J. The physical metallurgy of Nitinol for medical applications. JOM 2003;55:33e7. [5] Krone L, Mentz J, Bram M, Buchkremer HP, Stöver D, Wagner M, et al. 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 behaviour. Adv Eng Mater 2005;7:613e9. [6] Mentz J, Frenzel J, Wagner MF-X, Neuking K, Eggeler G, Buchkremer HP, et al. Powder metallurgical processing of NiTi shape memory alloys with elevated transformation temperatures. Mater Sci Eng 2008;A491:270e8. [7] Huang W. On the selection of shape memory alloys for actuators. Mater Des 2002;23:11e9. [8] Großmann C, Frenzel J, Sampath V, Depka T, Oppenkowski A, Somsen Ch, et al. Processing and property assessment of NiTi and NiTiCu shape memory actuator springs. Materialwiss Werkstofftech 2008;39:499e510. [9] Großmann C, Frenzel J, Sampath V, Depka T, Eggeler G. Elementary transformation and deformation processes and the cyclic stability of NiTi and NiTiCu shape memory spring actuators. Metall Mater Trans 2009; A40:2530e44. [10] Miyazaki Y, Igo Y, Otsuka K. Effect of thermal cycling on the transformation temperatures of TieNi alloys. Acta Metall 1986;34:2045e51. [11] Tadaki T, Nakata Y, Shimizu K. Thermal cycling effects in an aged Ni-rich TieNi shape memory alloy. Trans Japan Instit Metals 1987;28:883e90. [12] Matsumoto H. Transformation behaviour with thermal cycling in NiTi alloys. J Alloys Compd 2003;350:213e7. [13] Miyazaki S, Imai T, Igo Y, Otsuka K. Effectt of cyclic deformation on the pseudoelasticity characteristics of TieNi alloys. Met Trans 1986;A17: 115e20. [14] Gall K, Maier H. Cyclic deformation mechanisms in precipitated NiTi shape memory alloys. Acta Mater 2002;50:4643e57. [15] Hurley J, Ortega AM, Lechniak J, Gall K, Maier HJ. Structural evolution during the cycling of NiTi shape memory alloys. Z Metallkunde 2003;94:547e52. [16] Wagner MF-X, Nayan N, Ramamurty U. Healing of fatigue damage in NiTi shape memory alloys. J Phys D Appl Phys 2008;41:185408. [17] Morgan NB, Friend CM. A review of shape memory stability in NiTi alloys. J Physique IV 2001;11:325e32. [18] Scherngell H, Kneissl AC. Generation, development and degradation of the intrinsic two-way shape memory effect in different alloy systems. Acta Mater 2002;50:327e41. [19] Bhattacharya K, Conti S, Zanzotto G, Zimmer J. Crystal symmetry and the reversibility of martensitic transformations. Nature 2004;428:55e9. [20] Norfleet DM, Sarosi PM, Manchiraju S, Wagner MF-X, Uchic MD, Anderson PM, et al. Study of transformation-induced plasticity during pseudo-elastic deformation of NieTi microcrystals. Acta Mater 2009;57:3549e61. [21] Yawny A, Sade M, Eggeler G. Pseudoelastic cycling of ultra-fine-grained NiTi shape-memory wires. Z Metallkunde 2005;96:608e18. [22] Eggeler G, Hornbogen E, Yawny A, Heckmann A, Wagner M. Structural and functional fatigue of NiTi shape memory alloys. Mater Sci Eng 2004; A378:24e33. [23] Pfetzing J, Schaefer A, Somsen Ch, Wagner MF-X. Nanoindentation of pseudoelastic NiTi shape memory alloys: thermomechanical and microstructural aspects. Int J Mater Res 2009;100:936e42. [24] Wagner M, Sawaguchi TS, Kausträter G, Höffken D, Eggeler G. Structural fatigue of pseudoelastic NiTi shape memory wires. Mater Sci Eng 2004; A378:105e9. [25] Wagner MF-X, Eggeler G. New aspects of bending rotation fatigue in ultrafine-grained pseudo-elastic NiTi wires. Int J Mater Res 2006;97:1687e96. [26] Vaidyanathan R, Dunand DC, Ramamurty U. Fatigue crack-growth in shapememory NiTi and NiTieTiC composites. Mater Sci Eng 2000;A289:208e16. [27] Miyazaki S, Otsuka K. Development of shape memory alloys. ISIJ Int 1989;29:353e77. [28] Undisz A, Fink M, Rettenmayr M. Response of austenite finish temperature and phase transformation characteristics of thin medical-grade NieTi wire to short-time annealing. Scripta Mater 2008;59:979e82. [29] Jiang F, Liu Y, Yang H, Li L, Zheng Y. Effect of ageing treatment on the deformation behaviour of Ti-50.9 at.% Ni. Acta Mater 2009;57:4773e81. [30] Khalil-Allafi J, Dlouhy A, Eggeler G. Ni4Ti3-precipitation during aging of NiTi shape memory alloys and its influence on martensitic phase transformations. Acta Mater 2002;50:4255e74. [31] Dlouhy A, Khalil-Allafi J, Eggeler G. Multiple-step martensitic transformations in Ni-rich NiTi alloys e an in-situ transmission electron microscopy investigation. Phil Mag 2003;83:339e63. [32] Zheng Y, Jiang F, Li L, Yang H, Liu Y. Effect of ageing treatment on the transformation behaviour of Ti-50.9 at.% Ni alloy. Acta Mater 2008;56:736e45. [33] Wagner MF-X, Windl W. Elastic anisotropy of Ni4Ti3 from first principles. Scripta Mater 2009;60:207e10. [34] Tirry W, Schryvers D. High resolution TEM study of Ni4Ti3 precipitates in austenitic Ni51Ti49. Mater Sci 2004;A378:157e60. [35] Tirry W, Schryvers D. Quantitative determination of strain fields around Ni4Ti3 precipitates in NiTi. Acta Mater 2005;53:1041e9.
M.F.-X. Wagner et al. / Intermetallics 18 (2010) 1172e1179 [36] Tirry W, Schryvers D, Jorissen K, Lamoen D. Electron-diffraction structure refinement of Ni4Ti3 precipitates in Ni52Ti48. Acta Crystallogr 2006; B62:966e71. [37] Kim J, Miyazaki S. Effect of nano-scaled precipitates on shape memory behavior of Ti-50.9 at.%Ni alloy. Acta Mater 2005;53:4545e54. [38] Frenzel J, Zhang Z, Neuking K, Eggeler G. High quality vacuum induction melting of small quantities of NiTi shape memory alloys in graphite crucibles. J Alloys Compd 2004;385:214e23. [39] Jiang F, Li L, Zheng Y, Yang H, Liu Y. Cyclic ageing of Ti-50.8 at.% Ni alloy. Intermet 2007;16:394e8.
1179
[40] Murray J. NieTi (nickeletitanium). Binary alloy phase diagrams. Materials Park, OH, USA: ASM International; 1990. p. 2874e2876. [41] Nabarro F, de Villiers H. The physics of creep. London: Taylor & Francis; 1995. [42] Hara T, Ohba T, Okunishi E, Otsuka K. Structural study of R-phase in Ti50.23 at.%Ni and Ti-47.75 at.%Ni-1.50 at.%Fe alloys. Mater Trans JIM 1997;38:11e7. [43] Firstov G, van Humbeeck J, Koval Y. High-temperature shape memory alloys: some recent developments. Mater Sci Eng A 2004;378:2e10. [44] Firstov G, van Humbeeck J, Koval Y. High temperature shape memory alloys: problems and prospects. J Intell Mater Syst Struct 2006;17:1041e7.