S.V. Raj and R.D. Noebe

Glenn Research Center, Cleveland, Ohio

Low Temperature Creep of Hot-Extruded 

Near-Stoichiometric NiTi Shape Memory Alloy 

Part I: Isothermal Creep

NASA/TM—2013-217888/PART1

June 2013
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S.V. Raj and R.D. Noebe

Glenn Research Center, Cleveland, Ohio

Low Temperature Creep of Hot-Extruded 

Near-Stoichiometric NiTi Shape Memory Alloy 

Part I: Isothermal Creep

NASA/TM—2013-217888/PART1

June 2013

National Aeronautics and

Space Administration

Glenn Research Center

Cleveland, Ohio 44135
Acknowledgments

The authors thank Mr. Darrell Gaydosh, Drs. Anita Garg, Bradley Lerch and Santo Padula for granting permission to cite or use 

their unpublished data. The authors also thank Prof. R. D. James, University of Minnesota, and Prof. P. Anderson, Ohio State 

University for discussions relating to some of the data included in this paper. Comments from Dr. Michael Nathal are greatly 

appreciated. This work was supported by the Subsonic Fixed Wing Project of NASA’s Fundamental Aeronautics Program.

Available from

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7115 Standard Drive

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Available electronically at http://www.sti.nasa.gov

Trade names and trademarks are used in this report for identification 

only. Their usage does not constitute an official endorsement, 

either expressed or implied, by the National Aeronautics and 

Space Administration.

This work was sponsored by the Fundamental Aeronautics Program 

at the NASA Glenn Research Center.

Level of Review: This material has been technically reviewed by technical management. 
NASA/TM—2013-217888/PART1 1 

Low Temperature Creep of Hot-Extruded 

Near-Stoichiometric NiTi Shape Memory Alloy 

Part I: Isothermal Creep

S.V. Raj and R.D. Noebe 

National Aeronautics and Space Administration

Glenn Research Center

Cleveland, Ohio 44135 

Abstract

This two-part paper is the first published report on the long term, low temperature creep of hotextruded near-stoichiometric NiTi. Constant load tensile creep tests were conducted on hot-extruded 

near-stoichiometric NiTi at 300, 373 and 473 K under initial applied stresses varying between 200 and 

350 MPa as long as 15 months. These temperatures corresponded to the martensitic, two-phase and 

austenitic phase regions, respectively. Normal primary creep lasting several months was observed under 

all conditions indicating dislocation activity. Although steady-state creep was not observed under these 

conditions, the estimated creep rates varied between 10–10 and 10–9 s–1. The creep behavior of the two 

phases showed significant differences. The martensitic phase exhibited a large strain on loading followed 

by a primary creep region accumulating a small amount of strain over a period of several months. The 

loading strain was attributed to the detwinning of the martensitic phase whereas the subsequent strain 

accumulation was attributed to dislocation glide-controlled creep. An “incubation period” was observed 

before the occurrence of detwinning. In contrast, the austenitic phase exhibited a relatively smaller 

loading strain followed by a primary creep region, where the creep strain continued to increase over 

several months. It is concluded that the creep of the austenitic phase occurs by a dislocation glidecontrolled creep mechanism as well as by the nucleation and growth of deformation twins. 

1.0 Introduction

Advanced aircraft (Ref. 1) and rotorcraft (Ref. 2) designs increasingly require an ability to move 

various parts of the vehicle to achieve certain design and performance objectives. For example, morphing 

aircraft with capabilities for real time, on-demand wing movements enable flexible design capabilities 

(Ref. 3). Similarly, aircraft equipped with chevron nozzles with an ability to move the chevrons in and out 

of the flow path of gases exhausting from the engines can reduce engine noise during take-off and landing 

while improving engine performance and efficiency during cruise (Refs. 4, 5, and 6). The incorporation 

of these advanced concepts in aircraft designs are expected to lower fuel consumption, reduce emissions, 

increase safety and diminish engine noise. The development of suitable actuator technologies is critical to 

enabling these adaptive structures. 

The emphasis of current actuator technology development is increasingly dependent on the use of 

lightweight “smart materials” solid state technology (Ref. 7). Among them, shape memory alloy (SMA) 

actuators undergo a dimensional change due to the martensite (M) to austenite (A) transformation on 

heating. Since these dimensional changes can be large, SMA actuators are very suitable for applications 

requiring large displacements and a high work output (Ref. 6). There is a large body of evidence which 

suggests that thermal cycling of many of these SMAs under an applied stress results in an offset strain at 

the end of each cycle, where the start and end points of the hysteresis loop are displaced by a finite 

amount of strain (Ref. 6). This offset strain increases cumulatively during each subsequent thermal cycle, 

which is likely to decrease the effectiveness of the actuator in actual applications. 
NASA/TM—2013-217888/PART1 2 

Since these thermal cycling tests are conducted under constant load conditions in the laboratory to 

simulate the actual operating conditions for the SMA actuators, the observance of these offset strains in 

binary and ternary NiTi alloys (Ref. 6) is intriguing. In this connection, some key questions arise: 

a) Does near-stoichiometric NiTi exhibit significant creep in the temperature ranges between 300 

and 473 K corresponding to the martensitic and austenitic phase fields, respectively? 

b) What is the low temperature creep design stress limit for near-stoichiometric NiTi?

c) What are the creep characteristics of the austenitic and martensitic phases?

d) Is there any evidence of dislocation activity in these two phases?

e) What is the long-term stability of NiTi under constant stress under isothermal and thermal cycling 

conditions? 

Determining answers to these questions is important for designing SMA actuators for long duration 

applications. Unfortunately, prior creep investigations on NiTi alloys were conducted above 743 K 

(0.47 Tm, where Tm is the absolute melting point) corresponding to the austenitic phase field and well 

outside the temperature range, where NiTi SMAs are likely to be used in aerospace applications 

(Refs. 8, 9, 10, 11, 12, 13, and 14). No prior creep data have been reported below 473 K (0.30 Tm) in 

either the austenitic or the martensitic phase fields. Additionally, there are no long term creep data on 

NiTi and it is difficult to assess the long term stability of these alloys for actuator applications if 

maintained under load for significant periods of time. The present investigation was undertaken to study 

the low temperature creep behavior of near-stoichiometric NiTi to seek answers to the above questions. 

The first part of the paper characterizes the isothermal creep of NiTi between 300 and 473 K, while the 

second part discusses the creep response of pre-crept NiTi when the temperature is cycled between 300 

and 473 K under constant stress conditions. 

2.0 Experimental Procedures

The single phase alloy used in this investigation had a nominal composition of Ni-45(wt.%)Ti 

approximately corresponding to the stoichiometric composition, Ni-50(at.%)Ti. Tensile specimens with 

threaded ends, and ridges precisely defining the gage section, were machined from hot-extruded rods 

(Figure 1). The grain size of the hot-extruded material varied between 20 and 60 μm (Ref. 15). The gage 

dimensions were measured using an optical traveling microscope, where five measurements were made 

along the gage length on each specimen. The average gage length was 20 ± 0.02 mm and the average 

gage diameter 3.2 ± 0.01 mm, where the errors represent 95 percent confidence values. Constant load 

tensile creep tests were conducted at absolute temperatures, T, of 300 (0.19 Tm), 373 (0.24 Tm) and 473 K 

(0.30 Tm) under applied initial stresses, σ, varying between 200 and 350 MPa using several constant load 

lever arm machines with lever arm ratios varying between 3:1 and 10:1. The specimen elongation was 

measured using extensometers attached to a pair of ridges machined on the specimen, where linear 

variable capacitance transducers measured the displacements of the extensometer heads. Three 

thermocouples were attached to the specimen to monitor the specimen temperature and the temperature 

gradient along the gage length. In order to ensure that the starting microstructures are similar in all the 

tests, all specimens were held at 523 K for 1 h before furnace cooling over a 20 to 25 h period to the test 

temperature under a low stress of 5 to 6 MPa. It should be noted that the starting microstructures of NiTi 

are expected to be martensitic, duplex and austenitic at 300, 373 and 473 K, respectively, at these low 

stresses (Ref. 16). The isothermal creep tests were conducted for up to about 15 months. Typically, each 

specimen was crept under a constant engineering stress for several hundred hours before conducting stress 

change tests. A few tests were also conducted on mild steel specimens to verify that the strain 

measurements were genuinely due to specimen elongation. 
NASA/TM—2013-217888/PART1 3 

Figure 1.—Tensile creep specimen showing gage dimensions and 

general design.

Figure 2.—Comparison of the creep strain, ε, versus time, t, 

plots for NiTi and mild steel at 300 and 473 K. 

3.0 Results and Discussion

3.1 Comparison of the Creep Behavior of NiTi With Mild Steel

Figure 2 compares the creep curves plotted as true creep strain, ε, against time, t, for specimens tested 

at 300 and 473 K under an initial applied stress of 300 MPa. As noted earlier, the initial microstructures 

of the specimens prior to loading were martensitic and austenitic, respectively, at these temperatures and 

applied stress. The creep curves for the steel specimens tested at 300 and 473 K tested under 300 MPa are 

also shown in Figure 2 for comparison. The steel specimens exhibit a very small loading strain followed 

by logarithmic creep. In contrast, NiTi exhibits a relatively larger loading strain at 300 and 473 K 

followed by primary creep lasting several months with no discernable steady-state creep behavior. These 

observations suggest that the external factors did not significantly influence the strain measurements and 

represent genuine material behavior. The creep characteristics of NiTi at different temperatures will be 

discussed in the later sections. 

3.2 Creep Limit 

The creep limits at 300 and 473 K were determined by increasing the applied stress from an initial 

value of 200 MPa. In the case of the martensitic phase, the specimen was crept for 2035.6 h under an 

applied stress of 200 MPa before increasing the stress to 300 MPa; subsequently, the stress was 

periodically increased by about 5 percent of the previous stress. The specimen deformed in the austenitic 

range was initially crept for over 1000 h under an applied stress of 200 MPa followed by periodic stress 

increases of about 10 to 13 percent of the previous stress. In all cases, the specimen was crept for several 

hundred hours after each stress change. 
NASA/TM—2013-217888/PART1 4 

Figure 3(a) and (b) show plots of ε versus time t, at 300 and 473 K, respectively, where the stress was 

periodically increased from 200 to 347.3 MPa. The broken lines depict the corresponding values of 

engineering stress; the temperature profile is also shown. The right hand axis gives the values of stress 

and temperature. Loading NiTi to 200 MPa at 300 K resulted in an increase in the creep strain to 

0.46 percent followed by a small primary creep region, where the creep strain increased to about 

0.53 percent in 2035.6 h corresponding to estimated creep rates, ε , less than 1.3×10–10 s–1 approaching 

the resolution limits of measurements. A stress increase, ∆σ+

, by 100 MPa to a new stress of 300 MPa 

resulted in a relatively large increase in strain by about 1.1 percent soon after the stress increase and 

further increased by an additional 0.41 percent in about 2000 h (Figure 4(a)). Interestingly, the variation 

in the creep strain, ∆ε+

, with increasing time, ∆t, after the stress increase exhibits a sigmoidal behavior for 

different values of ∆σ+

.

1

 The magnitude of ∆ε+ increases sharply with increasing ∆t after an “incubation 

period” of gradually increasing ∆ε+ with increasing ∆t. Significantly, the creep strain continued to 

increase more gradually for several weeks and months after the sharp increase in ∆ε+

. 

Figure 3.—Creep strain versus time plots for (a) NiTi martensite deformed at 300 K and (b) NiTi austenite deformed 

at 473 K showing the effect of stress changes on the nature of the creep curves. Insignificant creep was observed 

below 200 and 220 MPa for martensite and austenite, respectively. 

Figure 4.—The nature of the creep curves immediately after (a) a stress increase and (b) a stress decrease. The 

values of ∆ε and ∆t are defined with respect to the point at which the stress change was performed. The + and – 

superscripts indicate increase and decrease, respectively. The curves for ∆σ+ = 14.4, 15.9 and 17.0 MPa do not 

start at ∆t = 0 since the first data acquisition was acquired 10 s after the stress increase. 

 

1 The values of ∆ε and ∆t are defined with respect to the point at which the stress change was performed. The + or – 

superscripts indicate increase or decrease, respectively. It is noted that the curves for ∆σ+ = 14.4, 15.9 and 17.0 MPa 

do not start at ∆t = 0 since the first data acquisition was acquired 10 s after the stress increase. 
NASA/TM—2013-217888/PART1 5 

A close examination of Figure 3(a) and Figure 4(a) reveals that although the magnitudes of the stress 

increases above 300 MPa were approximately similar varying between 14 and 17 MPa, the corresponding 

increases in ∆ε+

 were not approximately equal. For example, ∆ε+ is much smaller when the stress is 

increased from 300 to about 314 MPa compared to that after a stress increase from about 314 to 330 MPa; 

the subsequent increase from about 330 to 347 MPa leads to an observed value of ∆ε+ somewhat 

intermediate (Figure 3(a) and Figure 4(a)). If the strain increase following the “incubation period” is 

attributed primarily to the reorientation and detwinning of the martensitic twin variants, then the present 

observations suggest that these mechanisms were sensitively dependent on the deformation 

microstructure present prior to the stress increase. In other words, there appears to be a close coupling 

between the deformation microstructure and the reorientation and detwinning of the martensitic twin 

variants at these low creep rates. Thus, it is suggested that the internal stress builds up during the 

“incubation period” as the alloy develops a dislocation microstructure as it creeps at a very low rate until 

the magnitude of this stress reaches a critical value when it triggers reorientation and detwinning of the 

martensitic twin variants. 

Negligible amount of strain recovery was observed on stepwise unloading from 347.3 to 314.4 MPa, 

and unloading to a minimum stress of 11.4 MPa results in an instantaneous strain recovery of 0.66 percent

and a total recovered strain of 0.82 percent after 3209.1 h following the stress decrease (Figure 4(b)). On 

decreasing the stress, ∆ε− decreases slightly before becoming essentially constant when ∆σ− ≤ 17 MPa. A 

large stress change of ∆σ− = 303 MPa, resulted in a large decrease in ∆ε−

 followed by a continuing almost 

linear decrease in ∆ε− with increasing ∆t. This observation of anelastic creep can be attributed to the back 

flow of dislocations after the removal of the applied stress as the material attempts to reduce the internal 

back stress accumulated during deformation. Using an experimentally determined value of dynamic 

Young’s modulus, E, of 72.2 GPa (Ref. 17), this recovered strain is close to the calculated elastic strain of 

0.43 percent. Thus, a significant amount of residual strain of about 2.4 percent was measured in the 

specimen after testing for about 15 months (Figure 3(a)). 

At 473 K corresponding to the austenitic phase, no measurable creep was observed below 220 MPa 

for creep times less than 1500 h but the creep strain increases dramatically above 220 MPa as the stress is 

increased by 23 MPa from 220 to 243 MPa (Figure 3(b)). However, the increases in creep strain were 

relatively small with subsequent increases in the applied stress varying between 24 and 38 MPa at each 

stress increase. The specimen continued to creep for several months at stresses exceeding 220 MPa. The 

observation of a very large increase in strain for a relatively small increase in stress from 220 to 243 MPa 

is interesting. This large increase in loading strain can be attributed to one or more factors: Dislocation 

nucleation and glide, stress-induced martensitic (SIM) transformation (Refs. 18, 19, 20, and 21) and 

deformation twinning of the B2 austenite (Refs. 20, 22, 23, 24, 25, 26, 27, and 28). Indeed, recent 

microstructural observations on a Ti-49.52 (at.%) Ni alloy crept at 773 K for 10 h and 0.75 h under 

stresses of 100 and 200 MPa, respectively, showed large amounts of deformation twins, “structural 

defects” at the interfaces of the twin plates and retained austenite in specimens furnace cooled to room 

temperature (Ref. 14). Interestingly, deformation twins were observed after a creep strain of 3 percent in a 

specimen crept at 773 K under a stress of 100 MPa corresponding to primary creep. 

Based on the observations shown in Figure 3(a) and (b), it is reasonable to conclude that NiTi is 

unlikely to exhibit significant isothermal creep below 200 MPa between 300 and 473 K for creep times 

less than 2000 and 1500 h, respectively. Since the estimated creep rates are comparable to the resolution 

limits of the equipment (~ 10–10 s–1), it is cautioned that measurable creep may occur in tests lasting 

several months or years.

3.3 Creep Characteristics of Martensitic NiTi at 300 K 

Figure 5(a) and (b) show the ε–t plots for NiTi deformed at 300 K under an initial applied stress of 

300 MPa corresponding to the martensitic phase. A very large strain increase of 2.6 percent was observed 

on loading followed by normal primary creep, where the creep rate rapidly decreases to values of about 

7×10–10 s–1 (Figure 5(c)). For comparison, the estimated elastic strain at this value of applied stress is 

0.42 percent using an experimental value of E = 72.2 GPa (Ref. 17). As will be demonstrated later, this 

additional contribution to the loading strain arises due to the reorientation and detwinning of the 
NASA/TM—2013-217888/PART1 6 

martensitic twin variants. It is important to note that the specimen continued to creep very slowly over 

several months, where the creep strain increased to 2.8 percent after 3168 h. The data in Figure 5(a) was 

well-represented by Garofalo’s logarithmic formulation (Ref. 29) so that (εM)creep (percent) = 3.0×10–2log 

(1+7.2×1086t), where t is in seconds. The coefficient of determination, Rd

2

, was observed to be 0.992. 

Figure 6(a) shows the creep curve for a NiTi specimen tested at 300 K under an initial applied stress 

of 350 MPa. The stress on the specimen was reduced to 5.2 MPa after it had crept for 2182 h and held 

at this lower stress for 1031 h before increasing the stress to 350 MPa. The stress was reduced once 

again to 5.2 MPa after 984.5 h, and the specimen was held at the lower stress for 676.2 h before 

increasing the stress to 350 MPa. The amount of decrease or increase in strain with each stress change 

varied between 0.62 and 0.67 percent. These values are 30 to 40 percent higher than the elastic strain, 

(ε)elastic ≈ 0.48 percent determined from the ratio, ∆σ/E. A small amount of anelastic creep was observed 

after the stress decrease, where the strain continued to decrease with increasing time at a very low rate2

for the entire time that the specimen was under 5.2. MPa (Figure 6(b)). On increasing the stress to 

350 MPa, the creep strain increases rapidly to the value prior to the stress decrease without any overshoot 

unlike observations in pure metals. The observations of primary and anelastic creep lasting several 

thousand hours are indicative of dislocation activity. It is reasonable to conclude that dislocation glidecontrolled creep is dominant in the martensitic phase at 300 K although its contribution to the total strain 

is quite small compared to the observed strain on loading. 

Figure 5.—Creep behavior of martensite at 300 K under an initial applied stress of 300 MPa. (a) Shape of the creep 

curve showing normal primary creep during the first hour after loading; (b) the full creep curve; (c) creep rate, ε̇, 

versus time, t, plot. 

 

2 It is cautioned that the asymptotic nature of the creep curve shown in Figure 6(b) after the stress decrease may be 

due to the result of the creep rates approaching the resolution limits of the strain measurement system. 
NASA/TM—2013-217888/PART1 7 

Figure 6.—Effect of stress changes on the creep behavior of NiTi at 300 K under an initial applied stress of 350 MPa. 

(a) Full creep curve; (b) magnified view of the creep curve after a stress decrease from 350 to 5.2 MPa showing 

anelastic creep; and (c) comparison of the creep data with the room temperature stress-strain curves for NiTi 

deformed under constant engineering strain rates varying between 5.0x10-6 and 5.0x10–4 s–1. The creep data 

represent the measured values of the creep strain observed at the new value of the applied stress after the stress 

increase.

It is evident from Figure 5(a) and Figure 6(a) that a substantial contribution to the total creep strain is 

from the initial strain observed on first loading the specimen. In order to verify if this relatively large 

loading strain is due to the reorientation and detwinning of favorably oriented martensitic twins, the 

magnitudes of the loading strain after a stress increase at 300 K were superimposed on the room 

temperature stress-strain curve for NiTi (Ref. 30). The present data are in excellent agreement with the 

“plateau region” of the conventional stress-strain curve (Figure 6(c)). Importantly, these observations are 

independent of the strain rate. This region is attributed to the reorientation and detwinning of favorably 

oriented martensite twins (Ref. 18 and 31). The plateau in the tensile stress-strain curves have been 

attributed to “domino detwinning”, where favorably-oriented martensitic twins detwin when the resolved 

stress reaches a critical value (Ref. 31). Such a process would require the existence of an “incubation 

period”, where the creep strain is either negligible or low before a sudden increase in creep strain with 

increasing time. Figure 4(a) is consistent with this expected behavior. 
NASA/TM—2013-217888/PART1 8 

Figure 7.—Creep curves for NiTi deformed at 473 K under an initial applied stress of 300 MPa. (a) Initial portion of 

the creep curve within the first 1 h; (b) entire creep curve showing the effects of stress changes; and (c) magnified 

view of the creep curve after a stress decrease from 300 to 5.7 MPa showing anelastic creep. 

3.4 Creep Characteristics of Austenitic NiTi at 473 K 

Figure 7(a) to (c) show the creep curves for the austenitic phase at 473 K under an initial applied 

stress of 300 MPa. The loading strain is 2.8 percent, which is similar to that observed for the martensitic 

phase (Figure 5(a)). The estimated value of (ε)elastic ≈ 0.35 percent, where E = 87.1 GPa (Ref. 17). 

Significant normal primary creep characteristic of pure metal behavior was observed, where the creep 

strain increased to 7.2 percent in 3838 h (Figure 7(b)). Since austenitic NiTi has a B2 crystal structure 

with only three slip systems, the observation of this large creep strain can attributed to contributions from 

dislocation glide, stress-induced martensitic transformation and deformation twinning of the B2 austenite. 

There is a considerable body of experimental and theoretical evidence in support of deformation twinning 

in B2 NiTi and its alloys (Refs. 20, 22, 23, 24, 25, 26, 27, and 28). Recently, Ahadi and Rezaei (Ref. 14) 

reported the observation of large amounts of deformation twins in Ti-49.52 (at.%)Ni specimens crept at 

773 K under a stress of 100 MPa for 10 h to a primary creep strain of 3 percent. These observations are 

consistent with the present results. On decreasing the stress to 5.7 MPa, the creep strain decreased to 

5.4 percent, where this recovered strain, (ε)recovery ≈ 6(ε)elastic. The strain continued to decrease with 

increasing time for as long as 502.3 h (~21 days) thereby indicative of the formation of a dislocation 

substructure in the alloy (Figure 7(b) and (c)). On increasing the stress back to 300 MPa, there is in an 

immediate increase in creep strain to 6.4 percent followed by a gradual increase to 7.4 percent after 

549.3 h. Normal creep was observed after the stress increase characteristic of a dislocation glidecontrolled process. The fact that the creep curve did not overshoot after the stress increase unlike 

observations in pure metals suggests that new dislocation sources were not significantly activated. Since 
NASA/TM—2013-217888/PART1 9 

Figure 8.—Creep curves for NiTi crept at 473 K under an initial applied stress of (a) 250 and (b) 350 MPa showing 

the effect of stress changes and the existence of anelastic creep after a stress decrease.

the (ε)recovery ≫ (ε)elastic, it is reasonable to assume that the austenite undergoes a stress-induced 

martensitic transformation on increasing the stress on the specimen (Refs. 18, 19, 20, and 21), and 

possibly during subsequent creep. In this case, the results from the stress decrease experiments suggest 

that (ε)SIM ≈ 5(ε)elastic, where (ε)SIM is the strain due to stress-induced martensitic transformation. 

Normal primary creep was also observed under an initial applied stress of 250 MPa with the 

measured instantaneous creep strain being 1.1 percent (Figure 8(a)). On decreasing the stress to 5.1 MPa, 

the creep strain decreased by 0.85 percent so that (ε)recovery ≈ 3(ε)elastic. In this case, measurable anelastic 

creep was observed lasting for 1152.2 h (~48 days) prior to increasing the stress to 250 MPa. Once again, 

normal primary creep is observed after the stress increase. Since (ε)recovery ≈ 3(ε)elastic, it is reasonable to 

assume that that (ε)SIM ≈ 2(ε)elastic since the amount of anelastic strain is much smaller (~0.04 percent).

Figure 8(b) shows the creep results for NiTi deformed at 473 K under an initial applied stress of 

350 MPa. Once again, anelastic creep lasting for over 47 days is observed on decreasing the stress from 

350 to 5.2 MPa. The magnitude of (ε)recovery ≈ 2(ε)elastic, so that (ε)SIM ≈ (ε)elastic. Based on the observations 

shown in Figure 7 and Figure 8, it is clear that (ε)elastic ≤ (ε)SIM ≤ 5(ε)elastic with no consistent variation with 

increasing applied stress. 

3.5 Creep Characteristics of Two-Phase NiTi at 373 K 

Figure 9(a) to (c) show the creep curves for NiTi at 373 K under an initial applied stress of 300 MPa 

for various time intervals. The strain after loading was about 0.8 percent followed by normal primary 

creep, where the estimated creep rate was about 3×10–6 s–1 (Figure 9(a)). After about 1200 s into the test, 

the strain increased rapidly from about 1.8 to 5.0 percent within 1200 s. Normal primary creep was 

observed after this strain increase at an estimated creep rate of about 2×10–8 s–1. A second rapid increase 

in strain was observed after a total creep time of 23.7 h, where the strain increased from about 5.4 to 

6.6 percent in 1 h (Figure 9(b)). Once again, normal creep was observed after this strain increase lasting 

several thousand hours without any further rapid increase in creep strain and the specimen continued to 

creep to a strain of about 7.4 percent after 3192.1 h at which point the test was stopped (Figure 9(c)). The 

sudden increase in strain during creep suggests microstructural instability presumably due to stressinduced martensitic transformation. Alternatively, any delayed detwinning of any martensite present in 

the microstructure after an “incubation period” (Figure 4(a)) can also result in a large increase in strain.

Under an initial applied stress of 250 MPa, the loading strain was observed to be 1.4 percent (Figure 

10), which is surprisingly twice as large as that observed under a creep stress of 300 MPa (Figure 9(a)). 

Normal primary creep was observed thereafter, where the creep strain increased to about 1.7 percent after 

about 2275 h. On reducing the stress to 5.4 MPa, the measured (ε)recovery ≈ 1.3 percent ≈ 4(ε)elastic was 
NASA/TM—2013-217888/PART1 10

close to the value of the instantaneous creep strain. This large recovery in the creep strain suggests that 

the sudden increase in the creep strain observed at 300 MPa (Figure 9(a) and (b)) is likely due to a SIM 

transformation. Thus, (ε)SIM ≈ 3(ε)elastic at 373 K under an applied engineering stress of 250 MPa. In 

contrast to observations on austenite (Figure 8), the observed anelastic creep strain was quite small. This 

is similar to the behavior of the martensitic phase (Figure 6). 

 

Figure 9.—Creep behavior of NiTi deformed at 373 K under an initial applied stress of 300 MPa. (a) and (b) show the 

relatively large increases in creep strain during the early stages of the test due to stress induced martensitic 

transformation, while (c) shows the entire creep curve. 

Figure 10.—Creep curve for NiTi deformed at 373 K under an 

initial applied stress of 250 MPa showing the effect of a 

stress reduction from 250 to 5.4 MPa.
NASA/TM—2013-217888/PART1 11

Figure 11.—Schematic showing the application 

of the Kelvin-Voigt model to adjacent austenitic 

(A) and martensitic (M) grains.

3.6 Strain Contributions

The present observations can be rationalized in a simple way using the Kelvin-Voigt model (Figure 

11). Figure 11 shows the idealized duplex austenitic and martensitic microstructures arranged to satisfy 

the Kelvin-Voigt model. In this case, the strain is identical in both phases and the stresses are partitioned 

between the two phases as:

εT = εA = εM (1a)

∑ ( ) σ+σ=σ n V AAT 1

V MM (1b) 

where εT, εA and εM are the total creep strain, and the strains in the austenitic and martensitic phases, 

respectively, σT, σA and σM are the applied stress, and the effective stresses in the austenitic and n-variants 

of the martensitic phases, respectively, and VA and VM are the volume fractions of the austenitic and 

martensitic phase, respectively. Based on the observations shown in Figure 5 to Figure 10, we can write 

εA = (εA)recovery + (εA)anelastic + (εA)creep + (εA)twin (2a) 

εM = (εM)recovery + (εM)anelastic + (εM)detwin + (εM)creep (2b) 

where (εA)recovery and (εM)recovery are the magnitudes of the recovered strain in the austenitic and martensitic 

phases, respectively, when the stress is decreased, (εA)anelastic and (εM)anelastic are the anelastic creep strains 

in the austenitic and martensitic phases, respectively, (εA)creep and (εM)creep are the plastic creep strains in 

the austenitic and martensitic phases, respectively, (εA)twin is the strain contribution from deformation 

twinning in the austenitic phase, and (εM)detwin is the strain due to the reorientation and detwinning of the 

martensitic phase. Since the measured recovery strain in the martensitic phase is within the measurement 

error of (ε)elastic (Fig. 6 (a)), it is reasonable to assume that (εM)recovery ≈ (ε)elastic. In contrast, (εA)recovery ≈

(εA)SIM + (ε)elastic ≈ 2 to 6(ε)elastic for the creep of the austenitic phase (Figure 7, Figure 8 and Figure 10). 

The contributions of (εA)anelastic and (εM)anelastic to the total creep strain are small (Figure 6 to Figure 10) , so 

that, for simplicity, they can be neglected from Equations (2a) and (2b). The creep strain in the martensite 

due to dislocation glide is given by (εM)creep = (αMbM)ʃ(ρMvM)dt, where αM is a constant, bM is the Burgers 

vector of the martensite, ρM is the dislocation density in the martensitic phase, and vM is the dislocation 
NASA/TM—2013-217888/PART1 12

velocity in the martensite. The dislocation glide in the martensite is given by vM ∝ exp (BMσM/RT), where 

BM is a constant and R is the universal gas constant.

However, it is important to note that the observations of measurable anelastic creep in NiTi especially 

in the austenitic phase suggests that back stresses develop during creep in the material. Additionally, the 

observation of primary creep in the austenitic regime (Figure 2, Figure 4, Figure 7 and Figure 8) suggests 

a significant increase in the dislocation density so that a generalized representation of the creep strain is 

provided by (εA)creep = (αAbA)ʃ(ρAvA)dt, where αA is a constant, bA is the Burgers vector of the austenite, 

ρA is the dislocation density in the austenitic phase, and vA is the dislocation velocity in the austenite. The 

lack of steady-state creep suggests that the rate of dislocation hardening is greater than the rate of 

softening. Dislocation climb is not expected to be significant at 473 K, so that the dominant dislocation 

creep mechanisms are expected to be obstacle-controlled dislocation glide and/or dislocation cross-slip 

mechanisms (Ref. 32). Both these mechanisms predict a stress-dependent activation energy so that vA ∝

exp (BAσA/RT), where BA is a constant. However, contributions from the nucleation and growth of 

deformation twins to (εA)creep are also likely to be important as indicated in Equation (2a). The precise 

mechanistic contribution of deformation twinning to the creep of NiTi at 473 K is unclear. 

3.7 Effect of Dislocation Activity on the Detwinning of the Martensite Twin Variants

The observation of sigmoidal creep behavior in the martensitic phase is significant since it suggests that 

detwinning of the martensitic twin variants is not instantaneous (Figure 4(a)). Following an “incubation 

period” consisting of a relatively insignificant increase in creep strain with increasing time, the subsequent 

sharp increase in the creep strain appears to be due to the reorientation and detwinning of the favorably 

oriented martensitic variants. However, the continued increase in creep strain after the strain rise with 

increasing time for as much as 2000 h suggests that most of the strain contribution is due to dislocation 

activity. A decrease in stress by 303 MPa results in an initial decrease in ∆ε ≈ 0.7 percent followed by a more 

gradual decrease in ∆ε with increasing time at an approximate creep rate of 1.5×10–10 s–1, where the elastic 

strain recovery is expected to be about 0.43 percent for ∆σ = 300 MPa similar to the experimental 

observations (Figure 4(b)). As noted earlier, a substantial amount of strain equal to about 2.4 percent was not 

recovered (Figure 1), which can be attributed to the detwinning of the martensitic variants as established by a 

comparison to the stress-strain behavior (Figure 6(c)). 

The additional gradual time-dependent decrease in ∆ε (Figure 4(b)) is similar to anelastic creep 

observed in pure metals, where the observation of anelastic creep after a decrease in stress was attributed 

to the development of long-range internal back stresses due to the formation of a dislocation substructure 

during creep (Ref. 33). In pure metals and class M solid solution alloys, the observed magnitudes of 

anelastic strains are much larger than those reported in the present investigation due to the fact that these 

materials easily form heterogeneous cell and subgrain dislocation substructures, which result in the 

development of large internal back stresses in the material (Refs. 33 and 34). Hasegawa et al. (Ref. 34) 

observed sigmoidal creep in a Cu-16 (at. percent)Al solid solution alloy, where the creep rate first 

increased and then decreased as the total dislocation density increased with increasing creep time before 

saturating to a constant value beyond the peak creep rate. Correspondingly, the measured values of the 

long-range internal stress increased from about 30 percent to about 80 percent of the applied stress with 

increasing time as the dislocation substructure changed from isolated dislocations to dislocation cells as 

the creep time increased. 
NASA/TM—2013-217888/PART1 13

Figure 12.—Plot of ∆ε+/∆t against ∆t derived from Figure 4(a) 

for ∆σ+ values of 15.9 and 100 MPa showing the 

occurrence of peaks corresponding to the detwinning of the 

martensitic variants. The subsequent decrease in ∆ε+

/∆t 

with increasing ∆t suggests a hardening of the creep curves 

due to an increase in the dislocation density. 

Figure 12 shows that ∆ε+

/∆t first increases and then decreases with increasing ∆t for ∆σ+ of 15.9 and 

100 MPa; the magnitudes of ∆ε+

/∆t were determined from the time-dependent variation of the two welldefined sigmoidal curves shown in Figure 4(a). It is noted that these observations are valid for applied 

stresses exceeding 300 MPa (Figure 3(a)). The peak values of ∆ε+

/∆t were estimated to be 2.8×10–7 and 

1.8×10–3 s–1 for 15.9 and 100 MPa, respectively. Significantly, there is a considerable decrease in the 

creep rate with increasing creep time after the occurrence of the maxima thereby indicating a hardening of 

the alloy due to the development of back stresses. Based on the observations of an “incubation period” in 

Figure 4(a) and anelastic creep in Figure 4(b), it is surmised that the detwinning of the martensitic 

variants is closely coupled with dislocation activity occurring within the twin plates or at the twin 

interfaces. The initial increase in ∆ε+

/∆t is primarily due to detwinning of the martensitic variants with 

some contributions from an increase in the dislocation density. The subsequent decrease in ∆ε+

/∆t after 

the peak values is mainly because of an increase in the long-range back stress due to the development of a 

dislocation substructure. These observations suggest that any attempt to model detwinning in this alloy 

must be coupled with a dislocation mechanism. Liu et al. (Ref. 35) observed a large number of 

dislocations networks at martensitic twin boundaries in NiTi deformed in tension. These observations are 

consistent with the present results.

3.8 Phenomenological Model for Detwinning 

The phenomenological model shown in Figure 13 attempts to rationalize the role of dislocations on 

martensitic detwinning during low temperature creep of NiTi under tensile loading. Martensitic plates 

with twin variants, 1 and 2, meet at the twin boundary AB, where the open and solid circles represent the 

Ni and Ti atoms, respectively. It should be noted that boundaries, such as AB, have been termed as “habit 

plane variants” (hpv) while the interfaces within the twin variants have been referred to as “lattice 

correspondent variants” (lcv) (Ref. 36). Although the movement of hpv and lcv interfaces have been 

theoretically distinguished as “reorientation” and “detwinning” (Ref. 36), such specific distinctions could 

not be inferred from the present creep data. Thus, the movement of all twin boundary interfaces are 

collectively referred to as detwinning in this present paper. 
NASA/TM—2013-217888/PART1 14

Figure 13.—Phenomenological model showing the possible role 

of dislocation pile-ups in martensitic variant 1, which results in 

the development of shear stresses in the boundary regions so 

that the boundary between variants 1 and 2 moves from AB to 

A1B1, and the corresponding detwinning of variant 2. 

Dislocation loops emitted from sources within the martensitic twin variant 1 will expand under the 

action of the applied stress and pile-up at boundary AB so that a pile-up stress builds up on the lead 

dislocation at AB and in twin variant 2.3 This stress build-up leads to a corresponding increase in the 

magnitudes of the shear stresses along AB and A1B1. When the magnitudes of the boundary shear stresses 

exceed a critical value, the pile-up loops will run into the boundary and it is anticipated that a shear wave 

will propagate in twin variant 2 due to instability brought about by a shuffling of the atoms from their 

present to their new positions along the shear direction, and a corresponding movement of the twin 

boundary from AB to A1B1 and beyond as the martensitic twin variant 2 undergoes detwinning. It is 

suggested that the observed “incubation time” (Figure 4(a)) is the time required for the build-up of the 

shear stresses along AB and A1B1. On the removal of the applied stress, any dislocation loops within 

variant 1 will contract thereby leading to anelastic creep. Alternatively, the shuffling of the atoms along 

A1B1 can involve short range diffusion in which case the process is likely to be thermally activated. 

4.0 Summary and Conclusions 

This paper presents the first published results on the effect of stress on the long term low temperature 

tensile creep behavior of nominally stoichiometric NiTi deformed at 300, 373 and 473 K under initial 

applied stresses between 200 and 350 MPa corresponding to the martensitic, two-phase and austenitic phase 

fields, respectively. The creep tests were conducted for several months to as long as 15 months. These 

observations revealed differences in the creep behavior of the two phases under constant stress and in stress 

change tests. The observed creep limit for the martensitic phase was determined to be 200 MPa at 300 K 

for creep times less than 2000 h. Similarly, the austenitic phase did not exhibit measurable creep below 

220 MPa at 473 K for creep times less than 1500 h. However, it is cautioned that the validity of these creep 

limits is unknown for test times exceeding 2000 h and lasting several months or years since the estimated 

creep rates at 200 MPa were comparable to the resolution limits of the equipment ( ε ~ 10–10 s–1).

The martensitic phase exhibits a very large strain increase on loading followed by normal primary 

creep above 200 MPa. It is demonstrated that the magnitudes of the loading strains observed after stress 

increases coincide with the “plateau region” of the room temperature stress-strain curve for NiTi 

generated at constant strain rate. It is noted that the plateau is attributed to the reorientation and 

detwinning of favorable martensitic variants (Refs. 6, 18, 19, and 20). Significantly, sigmoidal creep is 

 

3 Alternatively, the dislocations can be emitted from the twin interfaces.
NASA/TM—2013-217888/PART1 15

observed after a stress increase thereby suggesting the existence of an “incubation period” prior to the 

reorientation and the detwinning of the martensitic twins. A decrease in stress resulted in measurable 

anelastic creep lasting several thousand hours after an initial instantaneous strain decrease approximately 

equal to the elastic strain. The observation of normal primary creep and small amounts of anelastic creep 

is attributed to dislocation glide-controlled creep. A phenomenological dislocation creep model is 

proposed to rationalize the observations of an “incubation period” prior to detwinning of the twin variants 

in the martensite, where it is suggested that dislocation pile-up within one set of martensitic twin variants 

results in a build-up of shear stresses at its boundaries with other twin variants. Detwinning of the second 

twin variants occurs when the shear stresses exceed the critical stress required to cause the atoms to 

collectively shuffle to new positions thereby leading to the propagation of shear waves in these 

martensitic variants. The shuffling of these atoms can occur by short range diffusion. 

Measurable creep was observed in the austenitic phase above 220 MPa at 473 K. Significant normal 

primary creep lasting several months was observed during the creep of the austenite. The relatively large 

creep strains observed in the B2 austenitic phase suggest that both dislocation glide and deformation 

twinning contribute to creep. The magnitudes of instantaneous creep strains after a stress decrease varied 

between 3 and 5(ε)elastic depending on the value of the initial applied stress. It is suggested that a part of 

the austenitic creep strain is also due to stress-induced martensitic transformation of some of the austenite. 

Following a stress decrease, the creep strain continued to decrease for several days typical of anelastic 

creep and indicative of the formation of a dislocation substructure in the alloy. Normal primary creep was 

observed after a stress increase characteristic of a dislocation glide process. As expected, steady-state 

creep was not observed even after several months of testing. It is concluded that the creep of the austenitic 

phase occurs by a dislocation glide-controlled creep mechanism accommodated by the nucleation and 

growth of deformation twins. Normal primary creep was also observed at 373 K. However, sudden 

increases in the creep strain were observed within the first 1200 s after loading due to microstructural 

instability due to stress-induced martensitic transformation of the austenite present in the material. 

References

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Low Temperature Creep of Hot-Extruded Near-Stoichiometric NiTi Shape Memory Alloy 

Part I: Isothermal Creep

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14. ABSTRACT

This two-part paper is the first published report on the long term, low temperature creep of hot-extruded near-stoichiometric NiTi. Constant 

load tensile creep tests were conducted on hot-extruded near-stoichiometric NiTi at 300, 373 and 473 K under initial applied stresses 

varying between 200 and 350 MPa as long as 15 months. These temperatures corresponded to the martensitic, two-phase and austenitic 

phase regions, respectively. Normal primary creep lasting several months was observed under all conditions indicating dislocation activity. 

Although steady-state creep was not observed under these conditions, the estimated creep rates varied between 10-10 and 10-9 s-1. The creep 

behavior of the two phases showed significant differences. The martensitic phase exhibited a large strain on loading followed by a primary 

creep region accumulating a small amount of strain over a period of several months. The loading strain was attributed to the detwinning of 

the martensitic phase whereas the subsequent strain accumulation was attributed to dislocation glide-controlled creep. An “incubation 

period” was observed before the occurrence of detwinning. In contrast, the austenitic phase exhibited a relatively smaller loading strain 

followed by a primary creep region, where the creep strain continued to increase over several months. It is concluded that the creep of the 

austenitic phase occurs by a dislocation glide-controlled creep mechanism as well as by the nucleation and growth of deformation twins. 

15. SUBJECT TERMS

Low temperature creep; NiTi; Shape memory alloys; intermetallic alloys 

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