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 II: Effect of Thermal Cycling

NASA/TM—2013-217888/PART2

June 2013
NASA STI Program . . . in Profile

Since its founding, NASA has been dedicated to the 

advancement of aeronautics and space science. The 

NASA Scientific and Technical Information (STI) 

program plays a key part in helping NASA maintain 

this important role.

The NASA STI Program operates under the auspices 

of the Agency Chief Information Officer. It collects, 

organizes, provides for archiving, and disseminates 

NASA’s STI. The NASA STI program provides access 

to the NASA Aeronautics and Space Database and 

its public interface, the NASA Technical Reports 

Server, thus providing one of the largest collections 

of aeronautical and space science STI in the world. 

Results are published in both non-NASA channels 

and by NASA in the NASA STI Report Series, which 

includes the following report types:

• TECHNICAL PUBLICATION. Reports of 

completed research or a major significant phase 

of research that present the results of NASA 

programs and include extensive data or theoretical 

analysis. Includes compilations of significant 

scientific and technical data and information 

deemed to be of continuing reference value. 

NASA counterpart of peer-reviewed formal 

professional papers but has less stringent 

limitations on manuscript length and extent of 

graphic presentations.

• TECHNICAL MEMORANDUM. Scientific 

and technical findings that are preliminary or 

of specialized interest, e.g., quick release 

reports, working papers, and bibliographies that 

contain minimal annotation. Does not contain 

extensive analysis.

• CONTRACTOR REPORT. Scientific and 

technical findings by NASA-sponsored 

contractors and grantees.

• CONFERENCE PUBLICATION. Collected 

papers from scientific and technical 

conferences, symposia, seminars, or other 

meetings sponsored or cosponsored by NASA.

• SPECIAL PUBLICATION. Scientific, 

technical, or historical information from 

NASA programs, projects, and missions, often 

concerned with subjects having substantial 

public interest.

• TECHNICAL TRANSLATION. Englishlanguage translations of foreign scientific and 

technical material pertinent to NASA’s mission.

Specialized services also include creating custom 

thesauri, building customized databases, organizing 

and publishing research results.

For more information about the NASA STI 

program, see the following:

• Access the NASA STI program home page at 

http://www.sti.nasa.gov

• E-mail your question to help@sti.nasa.gov

• Fax your question to the NASA STI 

Information Desk at 443–757–5803

• Phone the NASA STI Information Desk at

443–757–5802

• Write to:

 STI Information Desk

 NASA Center for AeroSpace Information

 7115 Standard Drive

 Hanover, MD 21076–1320
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 II: Effect of Thermal Cycling

NASA/TM—2013-217888/PART2

June 2013

National Aeronautics and

Space Administration

Glenn Research Center

Cleveland, Ohio 44135
Acknowledgments

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

Available from

NASA Center for Aerospace Information

7115 Standard Drive

Hanover, MD 21076–1320

National Technical Information Service

5301 Shawnee Road

Alexandria, VA 22312

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

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/PART2 1 

Low Temperature Creep of Hot-Extruded 

Near-Stoichiometric NiTi Shape Memory Alloy

Part II: Effect of Thermal Cycling

S.V. Raj and R.D. Noebe 

National Aeronautics and Space Administration

Glenn Research Center

Cleveland, Ohio 44135 

Abstract

This paper is the first report on the effect prior low temperature creep on the thermal cycling behavior 

of NiTi. The isothermal low temperature creep behavior of near-stoichiometric NiTi between 300 and 

473 K was discussed in Part I. The effect of temperature cycling on its creep behavior is reported in the 

present paper (Part II). Temperature cycling tests were conducted between either 300 or 373 K and 473 K 

under a constant applied stress of either 250 or 350 MPa with hold times lasting at each temperature 

varying between 300 and 700 h. Each specimen was pre-crept either at 300 or 473 K for several months 

under an identical applied stress as that used in the subsequent thermal cycling tests. Irrespective of the 

initial pre-crept microstructures, the specimens exhibited a considerable increase in strain with each 

thermal cycle so that the total strain continued to build-up to 15 to 20 percent after only 5 cycles. Creep 

strains were immeasurably small during the hold periods. It is demonstrated that the strains in the 

austenite and martensite are linearly correlated. Interestingly, the differential irrecoverable strain, Δεirr, in 

the material measured in either phase decreases with increasing number of cycles, N, as Δεirr = 4.3(N)–0.5

similar to the well-known Manson-Coffin relation in low cycle fatigue. Both phases are shown to undergo 

strain hardening due to the development of residual stresses. Plots of true creep rate against absolute 

temperature showed distinct peaks and valleys during the cool-down and heat-up portions of the thermal 

cycles, respectively. Transformation temperatures determined from the creep data revealed that the 

austenitic start and finish temperatures were more sensitive to the pre-crept martensitic phase than to the 

pre-crept austenitic phase. The results are discussed in terms of a phenomenological model, where it is 

suggested that thermal cycling between the austenitic and martensitic phase temperatures or vice versa

results in the deformation of the austenite and a corresponding development of a back stress due to a 

significant increase in the dislocation density during thermal cycling. 

1.0 Introduction 

As noted in Part I (Ref. 1), NiTi shape memory alloys (SMA) are being considered for use as solid 

state actuators (Ref. 2). By design, the SMA actuators would be required to undergo thermal cycles under 

constant or variable loading conditions. However, thermal cycling tests on most SMAs under an applied 

stress result in irrecoverable strain at the end of each cycle so that the start and end points do not coincide 

thereby reducing the useful life of actuators in actual applications (Ref. 3). The occurrence of these 

irrecoverable strains has been attributed to transformation-induced plasticity (Refs. 4, 5, 6, 7, 8, 9, 10, 

and 11). These constant load thermal cycling tests have been conducted primarily using a servo-hydraulic 

machine so that long term creep effects have not been evaluated. In particular, the effect of prior long 

duration isothermal creep deformation lasting several months on subsequent constant load thermal cycling 

tests have not been investigated to date. Earlier constant load or constant stress creep studies on NiTi were 

conducted above 743 K corresponding to the austenitic phase field (Refs. 12, 13, 14, 15, 16, 17, and 18) 

and well above the application temperature range for NiTi actuators. Consequently, the long term 

isothermal creep behavior of NiTi between room temperature and 473 K under constant applied stresses 

varying between 200 and 350 MPa was investigated as part of this investigation and the results are 

reported in Part I (Ref. 1). Based on earlier studies on the deformation behavior of NiTi (Refs. 2, 3, 4, 
NASA/TM—2013-217888/PART2 2 

11, 19, 20, 21, 22, 23, 24, 25, and 26), the predominant creep microstructures in the specimens studied in 

Pt. I (Ref. 1) are expected to be primarily detwinned martensite (M) and “dislocated” austenite (A) 

consisting of a mixture of dislocations and deformation twins necessary to satisfy von Mises criterion 

(Ref. 27). Recent short-term constant load creep tests conducted on a Ti-rich NiTi at 773 K confirm the 

formation of an extensive dislocation and deformation twinned microstructure (Ref. 18). 

The objectives of the present paper are to report the effects of temperature cycling on the creep 

behavior of pre-crept NiTi under constant load conditions. Unlike previous investigations (Refs. 3, 4, 5, 6, 

7, 8, 9, and 10), where the specimens that were thermally cycled generally had an undeformed 

microstructure prior to testing, the specimens used in this investigation were first isothermally crept under 

constant load conditions for several thousand hours prior to thermally cycling them. 

2.0 Experimental Procedures 

Details of the alloy composition and testing procedures are described in Part I (Ref. 1). A total of 

three specimens were thermally cycled. The first two specimens were isothermally pre-crept either at 

300 K for 5205.7 h to a true creep strain, ε, of about 4.3 percent ((a) (Ref. 1)) or at 473 K for 4342.1 h to ε 

≈ 1.1 percent (Fig. 8(b) (Ref. 1)) under an applied stress,σ, of 350 MPa. As stated earlier, the expected 

creep microstructures at these values of ε are detwinned martensite and “dislocated” austenite at 300 and 

473 K, respectively. Subsequently, both specimens were thermally cycled between 300 and 473 K under 

an applied stress of 350 MPa. A third specimen was pre-crept at 473 K for about 5062 h to ε ≈ 1.6 percent

under an initial applied stress of 250 MPa (Fig. 8(a) (Ref. 1)) prior to thermally cycling between 373 and 

473 K while maintaining the applied stress. In all three cases, the specimen was held at each new 

temperature for 300 to 700 h before changing to the previous temperature in order to distinguish between 

the strain contributions from isothermal creep during the hold time from those due to thermal cycling.

3.0 Results 

3.1 Thermal Cycling Creep Curves 

Figure 1 shows the variations in the true creep strain and absolute temperature, T, with increasing 

creep time, t, as T varied between T1 = 473 K and T2 = 373 K, where T1 and T2 are the initial and final 

temperatures, respectively, for a specimen crept under an initial applied stress of 250 MPa for 5061.6 h 

(Fig. 8(a) in (Ref. 1)). Only the thermal cycling portions of the ε - t curve are shown in the figure for 

better clarity. The specimen was held at each temperature for varying hold times between 312 and 650 h. 

The two broken curves show the increase in total strain at either 373 or 473 K. An examination of Figure 

1 reveals that the total strain at the end of each cycle was higher than at the start of the cycle so that the 

strain in the material continued to ratchet up with each additional thermal cycle. Thus, the total strain 

increased from about 1.6 percent at the commencement of the first thermal cycle to about 5.4 percent at 

the end of the fifth thermal cycle; a maximum strain of 8.7 percent was observed after the fifth 

temperature decrease from 473 to 373 K. For comparison, the measured isothermal creep strain at 473 K 

increased from 1.1 percent on loading to 1.5 percent after 2709 h under an initial creep stress of 250 MPa 

(Refs. 1). This corresponds to an approximate creep rate of about 4×10-10 s-1, so that the relatively short 

duration hold times at each temperature are unlikely to significantly influence the measured strain in these 

thermal cycling tests. Thus, the observed thermal ratcheting can be attributed almost entirely to 

temperature cycling effects under the present loading conditions. 

Figure 2(a) shows the ε–t plot for variations of T between T1 = 300 and T2 = 473 K for the specimen 

pre-crept at 300 h for several months under an initial creep stress of 350 MPa to a ε ≈ 4.3 percent prior to 

the thermal cycling tests at the same stress (i.e., “detwinned” martensite); the same data are replotted as a 

ε–T plot for ease of comparison with data obtained from stress-strain curves (Refs. 2, 3, and 8) (Figure 

2(b)). At the start of the first cycle, ε decreased from 4.3 to 2.3 percent as the absolute temperature 
NASA/TM—2013-217888/PART2 3 

Figure 1.—Variation of the true creep strain and absolute 

temperature with creep time for a specimen pre-crept 

for 5061.6 h at 473 K and then thermally cycled 

between 473 and 373 K under an applied stress of 

250 MPa.

Figure 2.—(a) True creep strain versus creep time and (b) true creep strain versus absolute temperature for a 

specimen pre-crept at 300 K under an applied stress of 350 MPa and subsequently thermally cycled between 300 

and 473 K. 

increased from 300 to 473 K as the martensite transformed to austenite so that the decrease in the amount 

of creep strain was 2.0 percent. The specimen was held for about 505 h at 473 K prior to decreasing the 

temperature to 300 K, where the measured creep strain was 0.08 percent corresponding to an approximate 

creep rate of about 4.4×10–10 s–1. For comparison, the estimated isothermal creep rates at 300 K prior to 

thermal cycling were of the order of 1.5×10–10 s–1. The variation in strain was insignificant during this 

hold period. On decreasing the temperature from 473 K to room temperature, the magnitude of ε

increased to 8.7 percent as the austenite transformed to martensite. Significantly, the increase in the 

amount of true strain by 4.4 percent at the end of the cycle was much greater than the strain of 4.3 percent

at the beginning of the cycle thereby signifying a large irrecoverable strain residual in the specimen after 

the first thermal cycle. As noted above, this plastic strain is predominantly due to thermal ratcheting 

rather than creep under these conditions. Continued thermal cycling of the specimen resulted in a large 

strain build-up in the specimen so that ε ≈ 17.8 percent after N = 5 cycles, where N is the number of 

thermal cycles. The total strain in the austenite and the martensite continued to increase almost linearly as 

shown by the two broken lines in Figure 2(a). 
NASA/TM—2013-217888/PART2 4 

Figure 3.—(a) True creep strain versus creep time and (b) true creep strain versus absolute temperature for a 

specimen pre-crept at 473 K under an applied stress of 350 MPa and subsequently thermally cycled between 473 

and 300 K. 

Similarly, decreasing the temperature from T1 = 473 K to T2 = 300 K as in the case of the specimen 

pre-crept at 473 K resulted in an increase in ε from about 1.1 to 8.5 percent as the austenite transforms to 

martensite (Figure 3(a)); the corresponding ε–T plot is shown in Figure 3(b). The specimen was 

isothermally crept at 300 K for about 321 h at 300 K before increasing the temperature back to 473 K. 

The estimated creep rate during this hold time at 300 K was about 1.8×10–10 s–1 with a total strain 

accumulation of about 0.02 percent. For comparison, the isothermal creep rates at 473 K prior to thermal 

cycling were estimated to be about 1.2×10–10 s–1. A subsequent increase from 300 to 473 K resulted in a 

corresponding decrease in ε to 6.4 percent. The total true irrecoverable strain residual in the specimen at 

the end of the first thermal cycle was 5.3 percent. Once again, the total strain in the austenite and 

martensite continued to increase with each thermal cycle as indicated by the broken lines. At the end of 

N = 6 cycles, the creep strain measured in the martensitic phase was 19.6 percent. 

3.2 Effect of Thermal Cycling on Deformation and Transformation Strains

Several terms are defined in this paper using Figure 4(a) and (b), which are reproduced portions of the 

T–t and ε–t responses for the two thermally cycled specimens under an applied stress of 350 MPa (Figure 

2(a) and Figure 3(a)). The isothermal creep strains, εcreep, are represented by the broken lines. The range 

of each cycle, N, is also shown in the figures. It should be noted that Figure 4(a) and (b) are out-of-phase 

by half a cycle. First, εNA and εNM are the average creep strains in the austenitic and martensitic phases, 

respectively,

1 accumulated at the Nth thermal cycle and they are measured with respect to ε = 0, i.e., εNA

and εNM are absolute values of strain. They represent the average strain in the two phases at the Nth cycle 

due to thermal cycling between the two temperatures since the strain contributions during the hold periods 

were insignificant. Second, (εtr)NA and (εtr)NM are the transformation strains associated with either 

martensite to austenite or austenite to martensite transformations, respectively, in the Nth thermal cycle. 

Third, (∆εirr)NA and (∆εirr)NM are the differential irrecoverable austenitic and martensitic strains,

respectively, in the Nth thermal cycle. Fourth, ∆εA and ∆εM are the differential true strains with respect to 

the isothermal creep strain lines. In other words, they represent the cumulative strains due to thermal 

cycling. 

 

1 It is noted that εI (I = A or M), (εtr)I, and (∆εirr)I are identified with the major constituent, either austenite or 

martensite, in this paper for simplicity. This assumption is a reasonable first order approximation although it is noted 

that thermally cycling a specimen from one temperature to another could stabilize one or the other phase as a minor 

constituent in the microstructure to some extent. 
NASA/TM—2013-217888/PART2 5 

Figure 4.—Schematic of the strain- time plot for thermal cycles from 

(a) 473 to 300 K and (b) 300 to 473 K illustrating the various 

terminologies used in this paper. The broken horizontal line 

refers to the isothermal creep strain prior to thermal cycling. The 

suffixes refer to the strain in the phase (austenite (A) or 

martensite (M)) for the Nth cycle. 

An examination of Figure 4(a) for the “dislocated” pre-crept austenitic microstructure gives the 

following algebraic relationships for N = 1, 2, 3...N: 

(εtr)NA = εNM – εNA (1a) 

(εtr)NM = εNM – ε(N–1)A (1b) 

(∆εirr)NA = εNA – ε(N–1)A (1c) 

(∆εirr)NM = εNM – ε(N–1)M (1d) 

Similarly, for the “detwinned” pre-crept martensite corresponding to Figure 4(b), the relationships for 

N = 1, 2, 3 ....N are: 

(εtr)NA = ε(N–1)M – εΝA (2a) 
NASA/TM—2013-217888/PART2 6 

(εtr)NM = εNM – εNA (2b) 

(∆εirr)NA = εΝA – ε(N–1)A (2c) 

(∆εirr)NM = εΝM – ε(N–1)M (2d) 

The differential strain, ∆εA and ∆εM, are defined with respect to the isothermal creep line (Figure 4(a) and

(b)) as follows:

∆εA = εΝA – εcreep (3a) 

∆εM = εΝM – εcreep (3b) 

Figure 5 shows the variation of the differential strains, ∆εA and ∆εM with increasing N for the two 

specimens thermally cycled between 300 and 473 K under an initial applied stress of 350 MPa. Several 

observations can be made from Figure 5. First, the curves show a non-linear behavior so that the 

magnitudes of ∆εA and ∆εM decrease with increasing N thereby suggesting that both phases undergo strain 

hardening as the specimens are thermally cycled between the two temperatures due to the development of 

a back stress in the material. Second, the magnitudes of ∆εA and ∆εM for the “dislocated” and detwinned 

creep microstructures are nearly identical. This result is attributed to a fortuitous coincidence and no 

special significance is inferred from the observation. Third, the differences between ∆εA and ∆εM

observed in Figure 5 represent the transformation strains. Thus, (εtr)ΝA = ∆εΝM – ∆εΝA and (εtr)ΝM = ∆εΝM – 

∆ε(Ν−1)A for the “dislocated” austenite (Figure 4(a)), and ∆(εtr)NA = ∆ε(N–1)M – ∆εΝA and (εtr)NM = ∆εNM – 

∆εNA for the “detwinned” martensite (Figure 4(b)), respectively. 

Figure 5.—Variation of the differential strains in the 

austenite and martensite with the number of thermal 

cycles for specimens initially pre-crept at 300 K 

(circles) and 473 K (triangles) under an applied stress 

of 350 MPa.
NASA/TM—2013-217888/PART2 7 

Interestingly, εA and εM exhibit a linear correlation during thermal cycling under an initial applied 

stress of 350 MPa for pre-crept specimens except that both sets of data are laterally displaced with respect 

to each other (Figure 6). This lateral displacement can be attributed to the fact that the initial strain in the 

specimen pre-crept at 473 K was much higher than that in the specimen pre-crept at 300 K. The slopes of 

the regression lines are close to unity thereby suggesting that the strain in the austenitic phase in one-half 

cycle is in close correspondence with that in the martensitic phase in the other half cycle irrespective of 

whether the specimen was pre-crept at 300 or 473 K prior to the thermal cycling tests. However, the fact 

that the slopes of both plots are close to unity suggests that the strain evolution in both specimens is 

similar irrespective of the pre-crept conditions. Especially notable are the magnitudes of εA and εM, which 

are surprisingly large. It was postulated in Part I (Ref. 1) that the B2 austenitic phase with three 

independent slip systems exhibits significant creep due to a combination of dislocation glide and 

deformation twinning, which satisfy von Mises criterion (Ref. 27). A similar mechanism can account for 

the large values of εA observed in Figure 6 as the austenitic phase deforms during thermal cycling. 

Alternatively, the activation of several slip systems provides a rationale for these large strains (Ref. 11). 

This rationale would suggest that the deformation microstructure in the austenite influences the 

magnitude of εM both by influencing the nucleation of the martensitic plates as well as stabilizing the 

martensite, as the specimen is thermally cycled between 300 and 473 K. 

Figure 7 shows the variations of (εtr)A and (εtr)M with increasing values of N for the specimens precrept at 300 and 473 K. The effect of the initial creep microstructures on (εtr)A and (εtr)M are minimal so 

that the broken curves represent a reasonable approximation of the observed trend. The magnitudes of 

(εtr)A and (εtr)M decrease and increase, respectively, with increasing values of N until N = 3 after which 

both transformation strains approach a degree of constancy. In this case, (εtr)M decreases from a maximum 

value of about 7.5 percent at N =1 to a minimum value of about 4.8 percent at N = 6, while (εtr)A increases 

from a minimum value of about 1.9 percent at N =1 to a maximum value of about 3.7 percent at N = 5. 

Correspondingly, the ratio (εtr)M/(εtr)A decreases from about 3.5 at N = 1 to about 1.6 at N = 5 irrespective 

of the initial pre-crept microstructure. The fact that the austenite and the martensite undergo strain 

hardening due to the development of an effective back stress (Figure 5) and the transformation strains 

evolve to approximately constant values during thermal cycling indicates that the relative volume 

fractions of the austenite and the martensite are also influenced by the evolution of this back stress with 

continued thermal cycling. 

Figure 6.—Correlation between the cumulative strains 

in the austenitic and martensitic phases during 

thermal cycling in specimens pre-crept at 300 K (open 

circles) and 473 K (solid triangles) under an applied 

stress of 350 MPa.
NASA/TM—2013-217888/PART2 8 

Figure 7.—Variation of the transformation strains with 

increasing number of thermal cycles as the austenite 

transforms to martensite and vice versa during 

thermal cycling under an applied stress of 350 MPa. 

The data for the specimens pre-crept at 300 and 

473 K are represented by circles and triangles, 

respectively.

Figure 8.—Double logarithmic plot of the differential 

irreversible strains in the austenite and martensite 

against number of thermal cycles for specimens precrept at 300 K (circles) and 473 K (triangles). 

3.3 Effect of Pre-Crept Microstructures on Differential Irrecoverable Strains

Figure 8 shows a double logarithmic plot of (∆εirr)A and (∆εirr)M against N for the two specimens 

isothermally pre-crept under an applied stress of 350 MPa. Two important points may be noted from the 

figure. First, both (∆εirr)A and (∆εirr)M decrease linearly with increasing N with the data scattered 

uniformly about a single regression line drawn through all the data irrespective of whether the 

measurements of ∆εirr were determined for the austenite or martensite. Second, the data are well 

represented by the regression equation: 

∆εirr = (∆εirr)A = (∆εirr)M = 4.3 N–0.5 (Rd

2 = 0.838 ) (4) 
NASA/TM—2013-217888/PART2 9 

where Rd

2 is the coefficient of determination. The data are clustered along this regression line independent 

of the pre-crept microstructure. Significantly, Equation (4) is remarkably similar to the Manson-Coffin 

relation in low cycle fatigue (Ref. 27). Interestingly, the Manson-Coffin relation was observed during the 

low cycle mechanical fatigue of NiTi at room temperature, where the value of the exponent was reported 

to be about 0.2 (Ref. 28). Once again, Figure 8 suggests that the back stress increases in the material with 

increasing number of thermal cycles.

3.4 Creep Rate-Temperature Plots

Figure 9(a) to (d) show the variation of the creep rates, ε̇, with T for the specimens pre-crept at 300 

and 473 K during either the austenite to martensite forward transformation (Figure 9(a) and (b)) or the 

martensite to austenite reverse transformation (Figure 9(c) and (d)) of the thermal cycles, where the ε̇

values were evaluated by differentiating the ε–t curves shown in Figure 2(a) and Figure 3(a). The cool 

down curves for each half cycle during the forward transformation are superimposed on the same plot in 

Figure 9(a) and (b), whereas the corresponding heat-up curves during the reverse transformation are 

superimposed in Figure 9(c) and (d). Noting that the product σε̇ has units of power density, Figure 9 9(a) 

and (b) represent energy emission as the austenite transforms to martensite for the forward transformation 

while Figure 9(c) and (d) represent the energy absorption as the martensite transforms to austenite for the 

reverse transformation.

Figure 9.—True creep rate versus absolute temperature for specimens; (a) and (c) pre-crept at 300 K; (b) and (d) precrept at 473 K; (a) and (b) forward transformation; and (c) and (d) reverse transformation.
NASA/TM—2013-217888/PART2 10

A close examination of Figure 9(a) and (b) reveals that all the curves superimpose for the forward 

transformation irrespective of the pre-crept microstructures. Similarly, the curves nearly superimpose for 

the specimen pre-crept at 473 K for the reverse transformation (Figure 9(c)). In contrast, the ε̇–T curves 

for the specimen pre-crept at 300 K exhibited a somewhat more complex behavior for the reverse 

transformation (Figure 9(d)). In this case, the ε̇–T curves shifted to lower temperatures after the first and 

second thermal cycles before superimposing on each other from the third cycle onwards. Additionally, the 

slopes of the ε̇–T curves, as indicated by the magnitudes of ε̇, were steeper than the results shown in 

Figure 9(c). In other words, the kinetics of the martensite to austenite transformation occurs at a faster 

rate for the specimen pre-crept at 300 K (Figure 9(d)) than for that pre-crept at 473 K (Figure 9(c)). The 

precise reason for this difference in behavior between the two specimens during reverse transformation is 

unclear. 

A close examination of Figure 9(d) reveals that the tails of the transformation curves are increasingly 

skewed towards the lower temperatures with increasing number of thermal cycles. The curves are less 

skewed for the specimen pre-crept at 473 K (Figure 9(c)). These observations suggest that the reverse 

martensite to austenitic transformation commences at lower and lower temperatures with each thermal 

cycle presumably due to the build-up of the effective back stresses in the material. 

The absolute transformation temperatures, Ttr, where Ttr represents the austenitic start, As, austenitic 

finish, Af, martensitic start, Ms, and martensitic finish, Mf, temperatures, were determined from the ε-T 

curves (Figure 2(b) and Figure 3(b)). The rise and fall segments of these curves were extremely nonlinear, and therefore it was felt that Ttr determined by the intercept method (Ref. 8) was unreliable. Thus, 

the data were determined by the deviation method, where the magnitudes of Ttr were identified with the 

points at which the curves deviated from the extensions of the stable regions (Figure 10). 

Figure 11 shows the variation of Ttr with N for the two specimens with different pre-crept

microstructures. Several points may be noted from Figure 11. First, Ms and Mf are independent of the 

initial creep microstructure. Second, Ms increases slightly while Mf does not vary significantly with 

increasing N irrespective of the pre-crept microstructures. Third, the values of As decrease with increasing 

N and they are sensitive to the pre-crept microstructure such that the values of As are significantly larger 

for the specimen pre-crept at 300 K than that pre-crept at 473 K. In the latter case, the As for the specimen 

with an expected microstructure of dislocations and deformation twins is similar to the values of Ms after 

N = 3. Fourth, Af for the specimen pre-crept at 300 K decreases with increasing N and it is larger than that 

for the specimen pre-crept at 473 K for which Af is independent of N. Fifth, the transformation of the 

martensite to austenite during the reverse transformation cycle occurs over a larger temperature range for 

the specimen pre-crept at 473 K than for the one pre-crept at 300 K. 

Figure 10.—Illustration of the deviation method used to 

determine the transformation temperatures, As, Af, Ms and Mf. 
NASA/TM—2013-217888/PART2 11

Figure 11.—Effect of pre-crept microstructures on the 

variations of the transformation temperatures with the 

number of thermal cycles. 

4.0 Discussion

The observations reported in this paper have an important bearing on the design and applications of 

binary NiTi actuators. As shown in Figure 1 to Figure 3, thermally cycling NiTi between either 300 or 

373 and 473 K under an applied stress of either 250 or 350 MPa results in a ratcheting of the total strain 

with each cycle. The accumulated strains are extremely large and well above 15 percent after five thermal 

cycles. However, the effect of hold times varying between 300 and 700 h did not significantly contribute 

to the total accumulated strain presumably because the creep rates were very low varying between about 2 

and 5×10–10 s–1. Based on the isothermal creep results reported in Part I (Ref. 1), where the creep strain 

was significant after several months, it is anticipated that long hold times lasting several months and 

years, would contribute to some extent, although much less than the transformation strain, to the total 

accumulated strain during the thermal cycling tests. Although the effect of hold time on the total strain 

accumulated during thermal cycling tests must be evaluated more completely than reported in the present 

paper, it is reasonable to suggest that hold times effects are likely to be subtle due to the evolution of the 

deformation microstructures. As noted in Part I (Ref. 1), measurable primary creep was observed between 

300 and 473 K for stresses above 200 MPa although the creep of the austenitic phase was significantly 

higher than that of the martensitic phase. Thus, the effects of long hold times at 473 K lasting several 

months between each thermal cycle under a constant applied load are expected to be more significant than 

similar hold times at 300 K. 

In attempting to design SMA actuators, it is essential to understand the fundamentals of the creep 

behavior of the austenitic and martensitic phases. It was concluded in Part I (Ref. 1) that the creep of NiTi 

at 300 K results in the detwinning of the martensitic twin variants on loading followed by dislocation 

glide-controlled primary creep. However, it was hypothesized that primary creep of NiTi at 473 K 

resulting in significant strain accumulation is due to the development of dislocation and deformation twin 

microstructures in the austenitic phase. The formation of deformation twins is necessary to accommodate 

the limited slip systems in its B2 crystal structure. 

The effects of these two pre-crept microstructures on the subsequent thermal cycling behavior of NiTi 

reported in this paper can be summarized as follows: (a) The material undergoes strain hardening with 

increasing number of thermal cycles thereby indicating a build-up in residual stress (Figure 5). This 

increase in residual stress can be attributed to the development of dislocation microstructures during each 

thermal cycle. (b) Significantly, the creep strains in the austenitic and the martensitic phases are closely 

correlated so that the strain in one phase increases as the strain in the other phase increases (Figure 6). 
NASA/TM—2013-217888/PART2 12

Thus, the deformation microstructures formed in the austenite at 473K are not destroyed after it 

transforms to martensite when the specimen is cooled to 300 K. Instead, it appears that the deformation of 

one phase in the first half of the thermal cycle influences the deformation of the other phase in the second 

half of the thermal cycle irrespective of the initial creep microstructure. (c) The observed inverse linear 

correlation between ∆εirr and N shown in Figure 8, and described by Equation (2) similar to the MansonCoffin relation in low cycle fatigue (Ref. 27), is also consistent with the formation of a dislocation 

substructure and the development of an associated back stress. 

The observed variations in Ttr with increasing N shown in Figure 11 suggest that the As and Af

temperatures are acutely sensitive to the initial isothermal creep microstructure prior to thermal cycling 

unlike Ms and Mf. The deformation microstructures developed during isothermal creep of the austenitic 

phase at 473 K reduces the magnitudes of the As and Af transformation temperatures well below those for 

the specimen isothermally pre-crept at 300 K in the martensitic phase field. The decrease in Af with 

increasing N for both pre-crept microstructures suggests that it becomes progressively easier to complete

the reverse martensite-to-austenite transformation with each thermal cycle. The present observed trends in 

the transformation temperatures during thermo-mechanical cycling are not always consistent with earlier 

observations on the thermal cycling behavior of NiTi. For example, Grossman et al. (Ref. 29) reported 

that Ms and Mf increase, while As and Af were essentially constant, with increasing N during constant 

load thermomechanical cycling of a near-equiatomic binary NiTi. 

The observations shown in Figure 11 can be rationalized in terms of the time evolution of the ε–T plots 

(Figure 12). Prior to the commencement of the thermal cycling experiments corresponding to N = 0, the 

hysteresis of the ε-T plots is relatively large so that As is much greater than Ms as indicated by the two 

vertical broken lines in Figure 12. With continued thermal cycling, the As and Af decrease significantly 

while Ms increases slightly with increasing N for the specimen pre-crept at 300 K (Figure 11), thereby 

indicating a reduction in the width of the hysteresis with each thermal cycle as the reverse transformation 

curve moves to lower temperatures in Figure 12. The small increase in Ms to higher temperatures while Mf

remains unchanged indicates that the slope of the forward transformation (i.e., austenite to martensite) curve 

Figure 12.—Schematic showing the relative shifts in the 

positions of the transformation temperatures during thermal 

cycling with respect to the true strain versus absolute 

temperature plot.
NASA/TM—2013-217888/PART2 13

decreases in the ε-T plot with increasing N. The decrease in the As and a slight increase in the Ms with 

increasing N for the specimen pre-crept at 473 K also indicates a reduction in the hysteresis but since Af is 

independent of N, the slope of the reverse transformation curve in the ε-T plot decreases with increasing N.

The close correspondence between εA and εM (Figure 6) and the increase in ∆εA and ∆εM with increasing N 

(Figure 5) suggests that the ε-T plot moves upwards parallel to the creep strain axis in Figure 12. 

The slopes of the forward transformation curves observed in the ε–T plots are influenced by the 

stored elastic strain energy due to the formation of the martensite, whereas the slopes of the reverse 

transformation are influenced by the dissipation of this stored elastic energy (Refs. 30, 31, and 32). Thus, 

a large difference between the Ms and Mf transformation temperatures corresponds to a shallow slope and 

leads to a high stored elastic strain energy, whereas a small difference between these two transformation 

temperatures corresponds to a steeper slope and lower stored elastic energy. 

Similarly, the width of the hysteresis is indicative of the frictional resistance to the motion of the 

austenite-martensite interfaces and martensite-martensite interfaces with different crystallographic 

orientations (Refs. 30, 31, and 32). A narrow hysteresis suggests a decreased frictional resistance to 

interfacial motion and a low relaxation of the stored elastic energy during the martensite to austenite 

reverse transformation. However, a wide hysteresis indicates a large frictional resistance and a high 

relaxation of the stored elastic energy. The As and Af temperatures are expected to be lower for a ε-T plot 

with a narrow hysteresis than for one with a wider hysteresis. Thus, the observed decrease in As and Af

with increasing N (Figure 11) suggests that the hysteresis decreases for both pre-crept microstructures. As 

noted earlier, this corresponds to a decrease in the frictional resistance to interfacial motion and a lower 

rate of dissipation of the stored elastic energy (Refs. 30, 31, and 32). However, the hysteresis is narrower 

for the specimen pre-crept at 473 K than for the one pre-crept at 300 K. Since Ms increases with 

increasing N while Mf is constant for both pre-crept specimens (Figure 11), the stored energy in the 

martensite is high and less heating is required to initiate the reverse martensite-to-austenite transformation 

(Refs. 30, 31, and 32). 

There is now considerable microstructural evidence that thermal cycling conducted with and without 

an applied constant load results in a significant increase in the dislocation density with each thermal cycle 

(Refs. 5, 6, 10, 29, 32, 33, and 34). The measured dislocation densities increased from 1010 to 5×1014 m–2

after 100 thermal cycles under no external load conditions (Refs. 10 and 34). Several recent observations 

indicate that dislocation glide loops are emitted into the austenite in front of an advancing austenitemartensite interface during the forward transformation from austenite to martensite to accommodate the 

formation of the martensite (Refs. 6, 10, 11, 30, and 34). No doubt the generation of these dislocations 

can account for the observed irrecoverable strains (Figure 8). The presence of these dislocations results in 

a back stress on the austenite-martensite interface therefore requiring a larger driving force for martensitic 

transformation and a decrease in Ms with increasing N (Refs. 10 and 34). It has been suggested that the 

presence of retained martensite also contributes to the irrecoverable strain (Refs. 11 and 29) although 

retained martensite is not always observed (Ref. 7). In this case, the martensite is assumed to be stabilized 

by the dislocation structure so that Ms increases with increasing N (Ref. 29). Mechanical twins and 

pseudo-twins in the austenite have also been observed during thermomechanical cycling of NiTi (Ref. 4). 

Clearly, the plasticity of the austenite can be attributed to either the activation of at least five independent 

slip systems (Ref. 11) or a combination of deformation twins and dislocation slip to satisfy von Mises 

criteria (Ref. 27). 

In view of the above discussion, the areas under the ε̇-T plots are measures of either the rates of 

increase in the stored elastic strain energy during the forward transformation (Figure 9(a) and (b)) or the 

dissipation of the stored elastic strain energy during the reverse transformation (Figure 9(c) and (d)). An 

examination of Figure 9(a) and (b) reveals that the ε̇–T plots are fairly similar irrespective of the initial 

pre-crept microstructure thereby suggesting similar amounts of stored elastic energy in both cases. 

Remarkably, the curves are well-defined and overlap each other after each cycle with the peak strain rate 

values occurring between 364 and 368 K suggesting that the transformation of the austenite to martensite 

is unaffected by any microstructural variation caused by the martensite to austenite transformation in the 
NASA/TM—2013-217888/PART2 14

previous reverse cycle. In contrast, the pre-crept microstructures have a significant effect on the 

dissipation of the stored elastic energy during the reverse transformation. The ε̇–T plots for the specimen 

pre-crept at 473 K are well-defined and overlap after the first thermal cycle and the martensite transforms 

to austenite over a broad temperature range (Figure 9(c)). The troughs for the curves, which occur 

between about 410 and 430 K, have a relatively broad spread of about 20 K. Assuming that the initial precrept microstructure consisted of a mixture of dislocations and deformation twins in the B2 austenite, 

coupled with the fact that thermal cycling of NiTi generates additional dislocations in the austenitic phase 

(Refs. 6, 10, 11, 30, and 34), the present observations suggest that the wide temperature range over which 

the reverse transformation occurs is influenced by the development of the internal stresses in the material. 

Therefore, it is reasonable to suggest that these local internal stresses oppose the progress of the austenitemartensite interfaces in the material during the reverse transformation so that the martensite transforms 

sluggishly to austenite. Under local stress conditions favorable to stabilizing the martensite, it is unlikely 

that martensite will transform to austenite so that it is retained at 473 K.

In comparison, the ε̇–T curves are well-separated for each cycle with relatively sharp troughs for the 

specimen pre-crept at 300 K, where the trough decreases from 477 K for cycle no. 1 to 456 K for cycle 

no. 3 (Figure 9(d)). Interestingly, the spread in the curves increases from the first to the third cycle before 

becoming approximately constant while the martensite starts transforming to austenite at lower and lower 

temperatures with each thermal cycle. Since the pre-crept microstructure is expected to consist 

predominantly of single variant martensitic twins with a few dislocations generated during isothermal 

primary creep, the differences in the characteristics of the curves for cycles nos. 1 and 2 are revealing. It 

should be noted that since dislocations are generated during the forward transformation (Refs. 5, 6, 10, 

29, 30, 33, and 34), the microstructure of the specimen prior to cycle no. 2 is expected to consist of a 

significant number of dislocations. Thus, the observed broadening of the curves after cycle no. 1 can be 

attributed to a significant increase in the dislocation density and a corresponding increase in the internal 

stresses in the material similar to the case for the specimen pre-crept at 473 K. Thus, the observed 

differences in the nature of the plots shown in Figure 9(c) and (d) stem from the fact that the curves for 

the specimen pre-crept at 300 K are in the process of evolving towards those for the specimen pre-crept at 

473 K. Eventually, it is expected that the curves for both specimens would be similar after a sufficiently 

large number of thermal cycles. 

As noted earlier, the reverse transformation from martensite to austenite commences at lower and 

lower temperatures (Figure 9(c) and (d)) thereby resulting in a decrease in As with increasing N (Figure 

11). The precise reason for this is unclear at this time but several possibilities exist. First, it is possible 

that some of the austenite fails to transform as thermal cycling progresses. However, it is important to 

note that retained austenite has not been observed in the alloy used in the present investigation although 

recent results on the short-term creep of a Ti-rich NiTi deformed at 773 K revealed the presence of 

retained austenite (Ref. 18). Second, the skewness in the curves towards the lower temperatures could be 

attributed to the transformation of favorably oriented martensitic twins to twins in the austenite during the 

reverse transformation. Ii et al. (Ref. 21) observed a direct correlation between the �201�

B19′ twins in the 

deformed martensite and the {114}B2 twins in austenite during reverse transformation in in-situ 

microstructural experiments conducted on a Ti-50.6at.% Ni alloy. It is also worth noting that {114}B2

deformation twins were observed in a NiTi alloy after thermal cycling under a constant load (Ref. 4). 

Third, an increase in Ms with increasing N (Figure 11) has been attributed to the stability of the martensite 

(Ref. 29), so that it is possible that the local stress fields around some martensitic plates prevent them 

from transforming back to austenite during the reverse transformation so that the volume fraction of the 

retained martensite increases with increasing N. In that case, the amount of martensite available for

transformation back to austenite will correspondingly decrease so that Af would decrease with increasing 

N. Clearly, further studies are needed to elucidate these issues. 
NASA/TM—2013-217888/PART2 15

Figure 13.—Schematic showing the effect of thermal cycling on the evolution of the microstructure and back stresses, 

σb, for specimens pre-crept at: (a) 300 K and (b) 473 K. The austenite deforms by a combination of deformation 

twinning and dislocation slip, while the martensite deforms by detwinning. As the specimens are thermally cycled 

between 300 and 473 K, dislocations are generated in the austenite so that σb increases with increasing N until the 

microstructure stabilizes. The smooth curve represents the increase in σb with increasing N. Some of the 

martensite is retained at 473 K under the action of the local internal stresses. 

Figure 13(a) and (b) are schematic representations showing the evolution of the microstructure during 

thermal cycling for the two specimens with different pre-crept microstructures. In Figure 13(a), the 

specimen with the detwinned martensitic pre-crept microstructure is heated to the austenitic region above 

Af under an applied stress during the first half of the thermal cycle. The martensite transforms to 

austenite, where the deformation and transformation strains are accommodated by dislocation slip, and 

possibly deformation twinning, with a corresponding increase in the back stress, σb, in the specimen. It is 

expected that some of the martensite would be stabilized by local internal stresses, and therefore, these 

martensitic plates would not transform to austenite. Thus, the expected microstructure at 473 K in the 

austenite at the end of this half cycle consists of a combination of glide dislocations, deformation twins 

and untransformed martensitic plates. In the second half of the thermal cycle, the austenite transforms 

back to martensite but the deformation twins and dislocations remain in the matrix since there is no 

significant recovery at these temperatures. However, there is a distinct possibility that the {114}B2

deformation twins in the austenite transform to �201�

B19′ martensitic twins (Ref. 21). In addition, based 

on experimental evidence (Refs. 6, 10, 11, 30, and 34), it is expected that the dislocation density would 

increase substantially to accommodate the transformation strain. On reverse transformation in the first 

half of the subsequent cycles, it is suggested that the �201�

B19′ martensitic twins nucleate the {114}B2

deformation twins in the austenite so that these twins flip-flop every half thermal cycle. The back stress 

increases with increasing number of thermal cycles, so that the effective stress, σeff = (σ – σb), decreases 

in the matrix thereby resulting in a corresponding nonlinear increase in εA and εM (Figure 5), and a 
NASA/TM—2013-217888/PART2 16

decrease in (∆εirr)A and (∆εirr)M with increasing N (Figure 8). Assuming a Taylor-type hardening 

mechanism (Ref. 35), the effective stress is dependent on the dislocation density, ρd, and deformation 

twin density, ρt, 

σeff = (σ – σb) = σ0 + αMμAbA(ρd + ρt)

–0.5 (3) 

where σ0 is the lattice friction stress, M is the Taylor factor, α is a constant, μA is the shear modulus of the 

austenite and bA is the Burgers vector for the austenite. Eventually, when N = n, the microstructure 

stabilizes (Figure 13(a)) so that As and Af become independent of N. 

The response of the specimen pre-crept at 473 K results in a larger value of σb prior to the start of the 

thermal cycle experiments owing to the presence of dislocations and deformation twins generated during 

isothermal primary creep (Figure 13(b)). Cooling down to 300 K in the first half of the thermal cycle results 

in the transformation of the austenite to martensite while retaining the deformation twins and dislocations 

present in the pre-crept microstructure. From the second half of the first thermal cycle and beyond, the 

microstructure and σb develop in a manner similar to that described above for the specimen pre-crept at 300 

K (Figure 13(a)) except that the microstructure reaches a stable state at a lower value of N. 

The observations reported in this paper have important implications for actuator design and bring into 

question the stability of binary NiTi alloys during thermal cycling. As noted in Part I (Ref. 1), the 

objective of this investigation was to understand the origin of the irrecoverable strain, which arises as 

near-stoichiometric NiTi is thermally cycled under a constant applied stress (Figure 1 to Figure 3). It was 

demonstrated in Part I that NiTi undergoes significant primary creep at 473 K when the applied stress 

exceeds 220 MPa; the alloy creeps at a much lower rate at 300 K for applied stresses greater than 

200 MPa. However, as demonstrated in Part II, thermal cycling between 300 and 473 K under creep 

stresses identical to those used in the isothermal experiments results in a significant amount of thermal 

ratcheting to 15 to 20 percent. The fact that the creep strain was immeasurably small for hold times 

varying between 300 and 700 h at each temperature convincingly demonstrates that irrecoverable strain 

accumulation due to thermal cycling far exceeds that due to isothermal creep due to a significant increase 

in the dislocation density in the austenite (Refs. 5, 6, 10, 29, 30, 33, and 34). 

Clearly, the magnitude of the irrecoverable strain is determined by the deformation of the austenitic 

phase. Thus, strengthening of the austenitic phase against dislocation slip is expected to improve the long 

term stability of the actuator and increase its life. Traditionally, this process has involved thermally 

cycling NiTi over a large number of cycles in an attempt to develop a stable microstructure through a 

process called “training” (Ref. 36). However, the problem with this approach is that a stable dislocation 

substructure is generally attained only after several hundred thermal cycles, which is not only expensive 

but it is also likely to reduce the work output of the actuator. As demonstrated in this investigation, the 

dislocation substructure tends to stabilize within a few thermal cycles if the alloy is initially pre-crept in 

the austenitic phase region. Nevertheless, it must be recognized that the stability of the dislocation 

microstructure is limited to the stress and temperature conditions under which it was generated and 

changes in the operating conditions of the actuator can destabilize the dislocation microstructure. 

5.0 Summary and Conclusions

This paper is the first report on the effect of prior low temperature creep on the thermal cycling 

behavior of NiTi. Specimens isothermally pre-crept at 300 and 473 K for several thousand hours were 

thermally cycled between 300 or 373 and 473 K under applied stresses of either 250 or 350 MPa. The 

present observations revealed that the total strain significantly ratchets up with each thermal cycle 

irrespective of the initial creep microstructure with the strain contribution from isothermal creep during 

hold times at each temperature being negligible. It is demonstrated that the strains in the austenite and 

martensite decrease non-linearly with increasing number of thermal cycles thereby suggesting a hardening 

of the alloy due to the development of back stresses as dislocations are generated in the material during 

each thermal cycle. The differential irrecoverable strain, ∆εirr, varies with N as ∆εirr = 4.3 N-0.5 similar to 
NASA/TM—2013-217888/PART2 17

the Manson-Coffin relationship observed in low cycle and thermal fatigue. The magnitudes of As and Af

temperatures were observed to be sensitive to the isothermal pre-crept microstructure with these 

transformation temperatures being higher when the specimen was pre-crept in the martensitic phase field 

at 300 K compared to that pre-crept in the austenitic phase field at 473 K. Both of these transformation 

temperatures decrease with increasing N when the pre-crept microstructure consists of detwinned 

martensite. Similarly, As also decreased with increasing N, but Af was constant when the pre-crept 

austenitic microstructure likely consisted of dislocations and deformation twins. A phenomenological 

model is proposed to rationalize the observations reported in this paper, which suggests that the back 

stress increases with increasing N due to dislocation generation in the austenite. 

References

1. S. V. Raj, R. D. Noebe, submitted to Mater. Sci. A. 

2. J. Ma, I. Karaman, R. D. Noebe, Intern. Mater. Rev. 55 (2010) 257-315. 

3. T. Saburi, in: K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge University 

Press, Cambridge, U.K., 1998, pp. 44-96. 

4. E. Goo, T. Duerig, K. Melton, R. Sinclair, Acta Metall. 33 (1985) 1725-1733. 

5. S. Miyazki, Y. Igo, K. Otsuka, Acta Metall. 34 (1986) 2045-2051. 

6. T. Simon, A. Kröger, C. Somsen, A. Dlouhy, G. Eggler, Acta Material., 58 (2010) 1850-1860. 

7. R. Delville, B. Malard, J. Pilch, P. Sittner, D. Schryvers, Intern. J. Plastic. 27 (2011) 282-297. 

8. S. Padula, S. Qiu, D. Gaydosh, R. Noebe, G. Bigelow, A. Garg and R. Vaidyanathan, Metall. Mater. 

Trans. 43A (2012) 4610-4621. 

9. S. Belyaev, N. Resnina, A. Sibirev, J. Alloy Comp. 542 (2012) 37-42, 

http://dx.doi.org/10.1016/j.jallcom.2012.07.082. 

10. A. R. Pelton, G. H. Huang, P. Moine, R. Sinclair, Mater. Sci. Eng. A 532 (2012) 130-138. 

11. T. Ezaz, J. Wang, H. Sehitoglu, H. J. Maier, Acta Mater. 61 (2013) 67-78. 

12. A. K. Mukherjee, J. Appl. Phys. 19 (1968) 2201-2204. 

13. H. Kato, T. Yamamoto, S. Hashimoto, S. Miura, Mater. Trans. JIM 40 (1999) 343-350. 

14. E. Kobus, K. Neuking, G. Eggler, I. Wittkamp, Prakt. Metallogr. 39 (2002) 177-186. 

15. G. Eggeler, J. Khalil-Allafi, K. Neuking, A. Dlouhý, Z. Metallkd. 93 (2002) 654-660. 

16. C. Lexcelent, P. Robinet, J. Bernardini, D. L. Beke, P. Olier, Materialwiss. Werkstofftech, 36 (2005) 

509-512. 

17. S. M. Oppenheimer, A. R Yung, D. C. Dunand, Scripta Mater. 57 (2007) 377-380. 

18. A. Ahadi, E. Rezaei, J. Mater. Eng. Perform. 21 (2012) 1806-1812. 

19. K. Otsuka, X. Ren, Intermetall. 7 (1999) 511-528. 

20. H. Sehitoglu, R. Hamilton, D. Canadinc, X. Y. Zhang, K. Gall, I. Karaman, Y. Chumlyakov, H. J. 

Maier, Metall. Mater. Trans. 34A (2003) 5-13. 

21. S. Ii, K. Yamauchi, Y. Maruhashi, M. Nishida, Scripta Mater. 49 (2003) 723-727. 

22. Y. Liu, Z. Xie, in: P. L. Reece (Ed.), Progress in Smart Materials and Structures, Nova Science 

Publishers, Inc., Hauppauge, NY, 2007, pp. 29-65. 

23. J. X. Zhang, M. Sato, A. Ishida, Acta Mater. 54 (2006) 1185-1198. 

24. T. Ezaz, Ph.D. Thesis, University of Illinois, Urbana-Champaign, IL, 2011. 

25. T. Ezaz, H. Sehitoglu, Appl. Phys. Lett. 98 (2011) 241906. 

26. T. Ezaz, H. Sehitoglu, W. Abuzaid, H. J. Maier, Mater. Sci, Eng. A 585 (2012) 422-430. 

27. M. A. Meyers, K. K. Chawla, Mechanical Metallurgy: Principles and Applications, Prentice-Hall, NJ, 

1984, p. 697. 

28. K. N. Melton, O. Mercer, Acta Metall. 27 (1979) 137-144. 

29. CH. Grossman, J. Frenzel, V. Sampath, T. Depka, G. Eggeler, Metall. Mater. Trans. 40A (2009) 

2530-2544. 

30. L. Delaey, J. Ortin, J. Van Humbeeck, Proc. in Phase Transformation 87, The Institute of Metals, 

London, U.K., 1988, pp. 60-66. 
NASA/TM—2013-217888/PART2 18

31. J. Van Humbeeck, E. Aeronoudt, L. Delaey, Lu Li, H. Verguts, J. Ortin, Revue Phys. Appl., 23 

(1988) 595-601. 

32. R. F. Hamilton, H. Sehitoglu, Y. Chumlyakov, H. J. Maier, Acta Mater. 52 (2004) 3383-3402. 

33. T. Nishura, M. Nishida, Mater. Trans. 50 (2009) 1219-1224. 

34. A. R. Pelton, J. Mater. Eng. Perform. 20 (2011) 613-617. 

35. L. Zhu, H. Ruan, X. Li, M. Dao, H. Gao, Jian Lu, Acta Mater. 59 (2011) 5544–5557. 

36. J. Van Humbeeck, R. Stalmans, in: K. Otsuka and C.M. Wayman (Eds.) Shape Memory Materials, 

Cambridge University Press, Cambridge, U.K., 1998, pp. 149-183. 

REPORT DOCUMENTATION PAGE Form Approved 

OMB No. 0704-0188 

The public reporting burden for this collection of information is estimated to average 1 hour per response, including the time for reviewing instructions, searching existing data sources, gathering and maintaining the 

data needed, and completing and reviewing the collection of information. Send comments regarding this burden estimate or any other aspect of this collection of information, including suggestions for reducing this 

burden, to Department of Defense, Washington Headquarters Services, Directorate for Information Operations and Reports (0704-0188), 1215 Jefferson Davis Highway, Suite 1204, Arlington, VA 22202-4302. 

Respondents should be aware that notwithstanding any other provision of law, no person shall be subject to any penalty for failing to comply with a collection of information if it does not display a currently valid OMB 

control number. 

PLEASE DO NOT RETURN YOUR FORM TO THE ABOVE ADDRESS. 

1. REPORT DATE (DD-MM-YYYY)

01-06-2013 

2. REPORT TYPE

Technical Memorandum 

3. DATES COVERED (From - To)

4. TITLE AND SUBTITLE

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

Part II: Effect of Thermal Cycling

5a. CONTRACT NUMBER

5b. GRANT NUMBER

5c. PROGRAM ELEMENT NUMBER

6. AUTHOR(S)

Raj, S., V.; Noebe, R., D. 

5d. PROJECT NUMBER

5e. TASK NUMBER

5f. WORK UNIT NUMBER

WBS 561681.02.08.47.05.04 

7. PERFORMING ORGANIZATION NAME(S) AND ADDRESS(ES)

National Aeronautics and Space Administration 

John H. Glenn Research Center at Lewis Field 

Cleveland, Ohio 44135-3191 

8. PERFORMING ORGANIZATION

 REPORT NUMBER

E-18699-2 

9. SPONSORING/MONITORING AGENCY NAME(S) AND ADDRESS(ES)

National Aeronautics and Space Administration 

Washington, DC 20546-0001 

10. SPONSORING/MONITOR'S

 ACRONYM(S)

NASA 

11. SPONSORING/MONITORING

 REPORT NUMBER

NASA/TM-2013-217888-PART2 

12. DISTRIBUTION/AVAILABILITY STATEMENT

Unclassified-Unlimited 

Subject Category: 26 

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

This publication is available from the NASA Center for AeroSpace Information, 443-757-5802

13. SUPPLEMENTARY NOTES

Submitted to Materials Science & Engineering A, Elsevier. 

14. ABSTRACT

This paper is the first report on the effect prior low temperature creep on the thermal cycling behavior of NiTi. The isothermal low temperature creep behavior of nearstoichiometric NiTi between 300 and 473 K was discussed in Part I. The effect of temperature cycling on its creep behavior is reported in the present paper (Part II). 

Temperature cycling tests were conducted between either 300 or 373 K and 473 K under a constant applied stress of either 250 or 350 MPa with hold times lasting at each 

temperature varying between 300 and 700 h. Each specimen was pre-crept either at 300 or at 473 K for several months under an identical applied stress as that used in the 

subsequent thermal cycling tests. Irrespective of the initial pre-crept microstructures, the specimens exhibited a considerable increase in strain with each thermal cycle so that 

the total strain continued to build-up to 15 to 20 percent after only 5 cycles. Creep strains were immeasurably small during the hold periods. It is demonstrated that the strains 

in the austenite and martensite are linearly correlated. Interestingly, the differential irrecoverable strain, Δεirr, in the material measured in either phase decreases with increasing 

number of cycles, N, as Δεirr = 4.3(N)-0.5 similar to the well-known Manson-Coffin relation in low cycle fatigue. Both phases are shown to undergo strain hardening due to the 

development of residual stresses. Plots of true creep rate against absolute temperature showed distinct peaks and valleys during the cool-down and heat-up portions of the 

thermal cycles, respectively. Transformation temperatures determined from the creep data revealed that the austenitic start and finish temperatures were more sensitive to the 

pre-crept martensitic phase than to the pre-crept austenitic phase. The results are discussed in terms of a phenomenological model, where it is suggested that thermal cycling 

between the austenitic and martensitic phase temperatures or vice versa results in the deformation of the austenite and a corresponding development of a back stress due to a 

significant increase in the dislocation density during thermal cycling. 

15. SUBJECT TERMS

Low temperature creep; NiTi; Shape memory alloys; Thermo-mechanical cycling 

16. SECURITY CLASSIFICATION OF: 17. LIMITATION OF

 ABSTRACT

UU 

18. NUMBER

 OF 

 PAGES

26 

19a. NAME OF RESPONSIBLE PERSON

STI Help Desk (email:help@sti.nasa.gov) 

a. REPORT

U 

b. ABSTRACT

U 

c. THIS 

PAGE

U 

19b. TELEPHONE NUMBER (include area code)

443-757-5802 

Standard Form 298 (Rev. 8-98)

Prescribed by ANSI Std. Z39-18