NiTiHf-based shape memory alloys

H. E. Karaca*, E. Acar, H. Tobe and S. M. Saghaian

NiTiHf-based shape memory alloys have been receiving considerable attention for high

temperature, high strength and two-way shape memory applications since they could have

transformation temperatures above 100uC, shape memory effect under high stress (above500 MPa) and superelasticity above 100uC. Moreover, their shape memory properties can betailored by microstructural engineering. However, NiTiHf-based alloys have some drawbacks

such as low ductility and high slope in stress induced martensite transformation region. In order to

overcome these limitations, studies have been focused on microstructural engineering by aging,

alloying and processing. It has been revealed that microstructural control is crucial to govern the

shape memory properties (e.g. transformation temperatures, matrix strength, shape recovery

strain, twinning type, etc.) of NiTiHf-based alloys. A summary of the most recent improvements on

selected NiTiHf-based systems is presented to point out their significant shape memory

properties, effects of alloying, aging and microstructure of transforming phases and precipitates.

Keywords: NiTiHf, High temperature shape memory alloys, Microstructure control, Precipitation hardening, Superelasticity, Work output

This paper is part of a special issue on Smart Materials

IntroductionShape memory alloys (SMAs) are a unique class ofsmart materials with the ability of changing their shapesdepending on the applied temperature, stress and in spe-cial case of ferromagnetic alloys, magnetic field. Shapememory alloys can produce very high actuation strains,stresses and work outputs as they undergo reversiblemartensitic phase transformation.1 In addition to theirremarkable properties in actuation, vibration damping,noise reduction and sensing, they are compact, robust,lightweight, frictionless, quiet, environment-friendly (nohydraulic liquids), easy to inspect and have low after-market costs for inspection and maintenance.24 Shapememory alloys are playing a growing role in supplying keyactuation forces and sealing functions in oil and gas,automotive, aerospace and biomedical industries.24 Theability to remain elastic under very large deformationmakes SMAs potential candidates for superelastic devicesfor civil structures.5,6 Moreover, their superelasticity, goodcorrosion resistance, biological and magnetic resonancecompatibility and high bending resistance resulted in theiremployment as the biomedical devices in the orthodontic,orthopaedic, vascular, neurosurgical fields.7,8

Among the various SMA systems, NiTi alloys havegood dimensional stability, shape memory properties,ductility and workability. Currently, NiTi alloys are themost commercially viable SMAs and practically beingused in various medical and engineering applicationswhere the operating temperature is below 100uC.9 It hasbeen found that the transformation temperatures (TTs)

of NiTi can be adjusted by tailoring the stoichiometry orformation of precipitates.10,11 However, the TTs ofbinary NiTi cannot be increased above 120uC. Thedevelopment of a shape memory material with proper-ties similar to those of near equiatomic NiTi, but withhigher strength and TTs, especially above 100uC, isurgently needed for a broad range of applications in theaerospace, automotive and oil and gas industries to serveas compact actuators for flow and clearance controls,actuation tubes for rotors, moving or morphing surfacesas well as inlet/exhaust configurations, linear actuatorsand sealants.24

Ternary element addition to NiTi alloys is the mostpromising method to obtain commercially available hightemperature shape memory alloys (HTSMAs) in thenear future.12 Ternary element addition should not onlyincrease the TTs, but also help to maintain the goodmechanical and shape memory properties of NiTi alloys.It has been found that the addition of Hf, Zr, Pd, Pt andAu elements to NiTi increases its TTs.2,9 Among thoseelements, Pd, Pt and Au are very expensive and will limitthe use of their respective ternary alloys to some criticalapplications only (i.e. aerospace), while Zr is associatedwith high oxygen affinity.2,9,13 Among the potentialHTSMAs, due to its low cost, medium ductility and highwork output NiTiHf seems to be the most encouragingHTSMA for a wide range of applications in the critical100300uC temperature range.12

The TTs of NiTiHf alloys do not increase much up to10 at-% Hf content, however, at chemical concentrationshigher than 10 at-%, they tend to increase linearly up to525uC for 30 at-% Hf when Hf is added at the expense ofTi.1416 Transformation temperatures of NiTiHf alloysare not notably affected by a change in Ni composition aslong as the alloys are Ni-lean, but dropped steeply when

Department of Mechanical Engineering, University of Kentucky, Lexington,KY 40506, USA

*Corresponding author, email karacahaluk@uky.edu

2014 Institute of Materials, Minerals and MiningPublished by Maney on behalf of the InstituteReceived 1 November 2013; accepted 23 June 2014DOI 10.1179/1743284714Y.0000000598 Materials Science and Technology 2014 VOL 30 NO 13a1530

Ni content is increased beyond the equiatomic (50 at-%)composition, consistent with the behaviour of NiTialloys.1517

The main disadvantages of Ni-lean NiTiHf alloys aretheir large hysteresis (.50uC), poor ductility at roomtemperature, lack of cyclic stability due to the high stressfor the reorientation of martensite and detwinning, thelow strength for slip and poor formability.2,18 It shouldbe noted that Wojcik19 studied the possibility of thecommercialisation of the NiTiHf (Hf content less than10 at-%) alloys and showed that hot rolling can besuccessfully utilised to produce thin sheets. Anotherdrawback of the alloy is the absence of stress plateauduring phase transformation that results in the lack ofsuperelasticity. This behaviour has been attributed to thesimultaneous occurrence of stress induced martensite(SIM) and dislocation slip.20,21 To increase the strengthfor slip, NiTiHf alloys were severely deformed that resultedin increased recoverable transformation strain, decreasedirrecoverable strain levels and thermal hysteresis underconstant stress experiments, as well as improved cyclicstability.18 However, no superelasticity was observed dueto large hysteresis and low material strength.

Precipitation strengthening has been used to improvethe mechanical properties of NiTiHf alloys as a successfulmethod.Meng et al.22,23 revealed that it is possible to formprecipitates in Ni-rich NiTiHf alloys and TTs can beincreased drastically to temperatures above 100uC. Theyhave also reported that coherent precipitates increase thematrix strength and enhance the thermal stability.22 Ifthe chemical composition is slightly Ni-rich with high Hfcontent (15 to 20 at-% Hf), fine nanometer size precipi-tates which are face centred orthorhombic structure,simply referred to as the H-phase,24,25 are formed uponaging treatments. The formation of fine precipitatesprovides high resistance to dislocation motion resultingin exceptional strength and stability limiting residualstrain during transformation under isothermal and iso-baric conditions.2628

Quaternary alloying and precipitation strengtheninghave also been used to improve the overall behaviour ofNiTiHf polycrystalline and single crystal alloys. The shapememory properties of heat treated Ni45?3Ti29?7Hf20Pd5 (at-%) alloys in single crystalline and polycrystalline formshave been reported.2933 The replacement of 5% Pd withNi of Ni50?3Ti29?7Hf20 alloy resulted in a very high strengthalloy that has high damping capacity of 35 J cm23

in polycrystalline form and 44 J cm23 in [111] orientedsingle crystals.32-33 Transformation strain of 2% wasobserved in aged [111] oriented Ni45?3Ti29?7Hf20Pd5 singlecrystals under a compressive biasing stress of 1500 MPa.31

Moreover, perfect superelastic behaviour with recoverablestrain of 4?2% was observed in the solutionized conditioneven when compressive stress levels as high as 2?5 GPawere applied.32 However, it is also known thatNi45?3Ti29?7Hf20Pd5 alloys are brittle, since they generallyfail after limited plastic deformation in compression andduring phase transformation in tension in superelasticityexperiments.34

It has been considered that low workability is one ofthe main problems with NiTiHf alloys for practical use.Kim et al.35 reported that an addition of Nb to NiTiHfalloys caused the formation of a soft Nb-rich b phaseand improved the cold workability, although the TTsand plastic strain in thermal cycling experiments under

stress were decreased. Cu has been another alloyingelement to NiTiHf systems where, in general, it improvedthe glass forming ability and thermal stability of NiTiHfalloys while decreasing their TTs.36,37 NiTiHfCu alloyshave also demonstrated two-way shape memory effect.38

It has recently reported that Ni45?3Ti29?7Hf20Cu5 alloyshave the capability to recover compressive strains of 2%above 100uC and two-way shape memory strain of 0?8%above 80uC.39

Hsieh and Wu40 investigated the TTs and hardnessvalues of Ti50?52xNi49?5Zrx/2Hfx/2 (x5020 at-%)

40 andrevealed that TTs can be increased from 50 to 323uCwith increased Zr and Hf contents. Their shape memoryresponses under stress (e.g. constant stress thermal cycling,superelasticity) have not been reported yet.

This article reviews the effects of alloying, aging andprocessing on the shape memory properties and micro-structure of NiTiHf-based alloys. Special attention isgiven to recently developed Ni-rich NiTiHf-based alloys.

Transformation temperatures of NiTiHf-based shape memory alloysMany studies have been conducted in order to gain thefundamental understanding on how to change the TTsof SMAs.41 It is known that chemical compositionalteration is very effective to change the properties suchas TTs, transformation strain and matrix strength ofSMAs. Figure 1a shows the effects of Ni content on theMp (martensite peak temperature) of NixTi902xHf10.

16

It is clear that Mp is insensitive up to 50 at-% Ni andthen suddenly decreases to below 0uC with increased Nicontent.

Figure 1b shows the change in Mp as a function ofHf.1416,42 It is clear that Mp does not change up to 3%of Hf and then increases after 5%. Up to 10% Hf, theincrease ofMp is about 5uC/at-% Hf. As the Hf increasesbeyond 10%, there is an abrupt increase ofMp by almost20uC/at-% Hf in NiTiHf alloys and Mp reaches up to400uC for 25% Hf.Figure 2a shows the differential scanning calorimetry

responses of the Ni50?3Ti29?7Hf20 alloys after heat treat-ment at selected temperatures from 300 to 900uC for3 h.27 Initially, TTs slightly decreased compared to the asextruded (extruded at 900uC) material when aged at 300and 400uC. Then, TTs increased with heat treatmenttemperature up to 700uC and then TTs decreased. Themaximum Af (austenite finish temperature) was revealedto be 210uC in Ni50?3Ti29?7Hf20 alloys aged for 3 h at600uC. Figure 2b shows the change in TTs for Ni45?3Ti29?7Hf20Pd5 polycrystalline specimens aged for 3 h attemperatures between 400 and 900uC.33 The trend in TTswith heat treatment temperature was similar to thatof Fig. 2a. The maximum Af was about 150uC inNi45?3Ti29?7Hf20Pd5 after aging at 600uC for 3 h. Themain reason for the TTs change with aging in the bothalloys could be attributed to the change in the chemicalcomposition of matrix33 due to the formation of preci-pitates that will be discussed in details in the micro-structure part.

Zarinejad et al.41 revealed a practical relationshipbetween the chemical composition and TTs by consider-ing the number (ev/a) and concentration (cv) of valenceelectrons in NiTi-based alloys. The number of d and selectrons is accepted as the number of valence electrons

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Materials Science and Technology 2014 VOL 30 NO 13a 1531

for an atom in transition metals while the number ofvalence electrons is considered to be p and s electrons foran atom in non-transition metals.43 The number ofvalence electrons of alloys can be calculated with thefollowing equation44

ev

a~fAe

AvzfBe

BvzfCe

Cvz

::: (1)

where fA, fB and fC are atomic fractions of A, B and C

elements, eAv ,eBv ,e

Cv are the related valence electrons for

the elements in an alloy system.

The following equation can be used to determine theaverage concentration of valence electrons44

cv~ev

et~

fAeAvzfBe

BvzfCe

Cvz

:::

fAZAzfBZBzfCZCz:::(2)

where ZA, ZB and ZC are the atomic numbers ofelements A, B and C, respectively.

Figure 3 shows the relationships between the Ms(martensite start temperature) (orMp) and ev/a and cv inNiTiHf-based SMAs.30,35,37,40,41,44,45 It is clear that the

TTs do not have a clear trend with ev/a while theygenerally decrease with increasing cv. It is commonlyagreed that higher electron concentration results inhigher bulk (resistance to volume change) and shear(resistance to shape change) moduli.44,46 Thus, theconcentration of the electrons may affect the strengthof atomic bonds in metallic materials. In general, as theconcentration of valence electrons increases, the resis-tance to shear also increases. Thus, further energyprovided by undercooling is necessary for the martensi-tic transformation resulting in decreased TTs.

As stated above, even though there are some guidelinesin predicting the TTs of NiTiHf-based alloys, therelationship between the nominal chemical compositionand TTs is not completely established since there aremany other factors that may alter TTs such as precipita-tion and grain size effects.37,4345 For instance, if theprecipitates are fine and interparticle distances are small,nucleation of martensite could be more difficult andrequire additional undercooling, resulting in decreasedTTs. Transformation temperatures are also sensitive to

1 Mp temperature as a function a Ni and b Hf contents in NiTiHf alloys1416,42 (chemical compositions are in at-%)

2 Transformation temperatures of a Ni50?3Ti29?7Hf20 and b Ni45?3Ti29?7Hf20Pd5 alloys after heat treatment of 3 h at selected

temperatures27,33

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local chemical composition changes due to formation ofprecipitates. Moreover, it is known that internal stressfields increase the TTs in SMAs.37,4345

Crystal structure and microstructure ofNiTiHf-based alloysThe shape memory and superelastic properties of NiTi-based alloys are significantly influenced by the mic-rostructure, such as the precipitation size, interparticledistance and martensite morphology. In order to obtaingood shape memory and superelastic properties, it isimportant to strengthen the matrix to prevent the intro-duction of dislocations during the martensitic transforma-tion. One of the well-known procedures to improve thestrength of the matrix is the precipitation hardening.Aging of NiTiHf alloys produces several precipitates thataffect the martensite morphology that will be discussed indetails.

Crystal structure of NiTiHf alloysIn general, the crystal structures of austenite andmartensite phases in NiTiHf alloys are cubic (B2) andmonoclinic (B199), respectively, which are similar to thosein NiTi binary alloys. Zarinejad et al.47 investigatedthe effect of Hf on the lattice parameters of the B199martensite in NiTiHf alloys. The lattice parameters a, b, cand b angle of the martensite are plotted in Fig. 4 asa function of Hf content for Ni(1002x)/2Ti(1002x)/2Hfx,Ni502xTi50Hfx and Ni50Ti502xHfx (x5520 at-%) alloys.The addition of Hf increased all the lattice parameters forthe Ni(1002x)/2Ti(1002x)/2Hfx and Ni502xTi50Hfx alloys. Onthe other hand, when Ni is constant, the increase in Hf inthe Ni50Ti502xHfx alloy increased a, c and b but decreasedb. Potapov et al.45 also observed a similar dependence oflattice parameters on the Hf content for Ni49?8Ti50?22xHfx(x5825 at-%) alloys where the increase in Hf while Niwas kept constant to 49?8% slightly decreased the latticeparameter b, while it increased a, c and b of B199martensite. It was also reported that the addition of Hfincreased the lattice parameter of B2 austenite.45 Thevolume change during transformation was smaller than0?5% which was similar to that in NiTi binary alloys(y0?3% or less).48,49 It should be noted that in somestudies, NiTiHf alloys with more than 15 at-% Hf in

Ni48?5(Ti51?52xHfx)50 and between 20 and 30 at-% Hf of

Ni50(Ti502xHfx)51 were reported to have orthorhombic

B19 martensite.

Precipitation characteristics and their effects onmartensite morphologyNi-lean NiTiHf-based alloys

Konig et al.52 fabricated NiTiHf thin films with a widecomposition range by magnetron sputtering method andinvestigated their TTs, precipitate structure and thermalcycling properties. Multilayer thin films (individual layersy15 nm thick) were sputtered from elemental targetsand annealed at 550uC for 1 h in order to transform theirmultilayer structure into alloys. Figure 5 depicts thecomposition regions in which different precipitates areformed.52 The relative intensity of one characteristic X-ray diffraction peak belonging to the phase of interest wasplotted colour-coded within a section of the NiTiHfternary phase diagram. Four different precipitates, i.e.HfNi(Ti), Ti2Ni(Hf), Hf2Ni(Ti) and Laves phase, wereconfirmed in Ni-lean composition regions. They con-cluded that the observation of reversible phase transfor-mation was limited by the formation of Ti2Ni(Hf),HfNi(Ti) and/or Hf2Ni(Ti) precipitates. These precipi-tates restricted the transforming region to compositionswith Ni contents abovey40 at-% and Hf contents belowy30 at-%.The Ti2Ni(Hf) precipitates have also been observed by

many other researchers in Ni-lean NiTiHf alloys.17,36,5355

It has been reported that the volume fraction of theTi2Ni(Hf) precipitates decreased with increasing the Nicontent, although the Ti2Ni(Hf) precipitates were stillobserved in slightly Ni-rich compositions.17,23 Fine Ti2Ni(Hf) precipitates strengthen the matrix and improveshape memory and superelastic properties of NiTiHf-based alloys.36,53 The effects of aging temperature and timeon the formation of Ti2Ni(Hf) precipitates were investi-gated by Meng et al. in Ni49Ti36Hf15

53 and Ni44Ti36Hf15Cu5

36 alloys. The size of the precipitates increased withincreasing aging temperature and time. Figure 6a and bshows the bright-field transmission electron microscopy(TEM) images of the Ni44Ti36Hf15Cu5 ribbons annealedat 500 and 700uC for 1 h, respectively.36 According to theselected area diffraction (SAD) pattern (Fig. 6c) takenfrom the specimen annealed at 500uC, the precipitate was

3 Ms as a function of a ev/a and b cv in NiTiHf-based shape memory alloys30,35,37,40,41,44,45

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Materials Science and Technology 2014 VOL 30 NO 13a 1533

confirmed to be Ti2Ni(Hf). The diameter of the precipi-tates was estimated to be 2040 nm when the annealingtemperature was 500uC. After annealing at 700uC, thesize of the precipitates increased to y150 nm. The fineprecipitates formed in the ribbon after annealing at 500uCstrengthened the matrix and prohibited plastic deforma-tion, which resulted in a perfect superelastic shape recoveryafter deformation to 3?5% strain. On the other hand, theribbon annealed at 700uC showed an incomplete shaperecovery due to the lower strength of the matrix with largeprecipitates.

It is important to note that the size of the Ti2Ni(Hf)precipitates are very effective to control the martensitemorphology. It was found that (001)B199 compoundtwins were dominant when the material containedhomogeneously distributed Ti2Ni(Hf) precipitates with2040 nm in diameter (Fig. 6a). Similar martensitemorphology has been observed in a Ti-rich NiTi thinfilm with a homogeneous distribution of fine Ti2Niprecipitates.56 When the annealing temperature was700uC, {011}B199 type I twins became dominant andthe martensite variants showed mainly spear-like andmosaic-like morphologies as shown in Fig. 6b. Marten-site domains with (001)B199 compound twins were alsoobserved around the coarse Ti2Ni(Hf) precipitates. Thespear-like and mosaic-like morphologies have beenreported as typical morphologies of the martensite inHf-added NiTi alloys.57,58

Ni-rich NiTiHf-based alloys

Meng et al.23,59 have reported that Ni4(Ti, Hf)3 pre-cipitates were formed in Ni-rich NiTiHf alloys similar tothe Ni4Ti3 precipitation in NiTi binary alloys. However,recently, it has been reported that a new precipitate whichhas a more complicated structure than that of Ni4(Ti,Hf)3forms in Ni-rich NiTiHf alloys24,25,60 and improvestheir shape memory and superelastic properties due toprecipitation strengthening.26,27,33 Initially, Han et al.61

reported a precipitate with a face-centred orthorhombiclattice with a space group of F 2/d 2/d 2/d in an agedNi48?5Ti36?5Hf15. There are six different variants in thisorthorhombic precipitate with habit planes of (100)P//{001}B2 and long axes of [001]P//,-110.B2. However, theydid not provide an atomic structure model for theobserved precipitate.

Recently, Yang et al.25 proposed an atomic structuremodel which contains of 192 atoms in an orthorhombicunit cell for the observed precipitate in Ni-rich NiTiHfalloys. The orthorhombic precipitate phase was namedas H-phase and Fig. 7a shows the unit cell of thisprecipitate.25 In order to refine the structure model, abinitio density functional theory calculations have alsobeen performed to relax the structure model.24,25 Selectedarea diffraction patterns obtained from a single large H-phase precipitate in a Ni52Ti28Hf20 alloy are shown inFig. 7bd.25 All the SAD patterns revealed the orienta-tion dependence between the precipitate and austenite B2

4 Lattice parameters a a, b b, c c and d b of B199 martensite as a function of Hf in NiTiHf alloys47

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phase (the diffraction spots are indexed according to theaustenite phase). There were additional reflections at 1/3positions along ,110.B2

* in reciprocal space as shownby arrows, which was a characteristic of the H-phase. Thecomposition of the proposed H-phase was Ni50Ti16?7Hf33?3, whereas it has been indicated by energy dispersivespectroscopy analysis that the Ni content of the H-phaseprecipitate was always slightly richer than that of thenominal composition of Ni-rich NiTiHf alloys in contrastto the proposed Ni content of 50 at-% .24,25,62 Thereforethe formation of H-phase precipitates depleted Ni fromthe matrix and increased TTs as shown in Fig. 2. Yanget al.25 observed anti-site defects within the precipi-tate which may slightly change the composition of theprecipitate, and proposed that the H-phase did not have

a unique composition. The effects of the alloy composi-tion on the H-phase precipitation were investigated bySantamarta et al.24 They concluded that the H-phaseprecipitates grew faster in alloys with higher Ni contentsince the precipitates were richer in Ni content comparedto the nominal composition of the alloys. Similarly, for afixed Ni content, the growth of the H-phase became fasterwhen the Hf content was increased.

The control of the size and interparticle distance of H-phase precipitates is important to obtain good shapememory and superelastic responses. It has been reportedthat the aging temperature and time significantly affectedthe size and interparticle distance of the precipitatesformed inNi-richNiTiHf-based alloys.23,24,33 Figure 8acillustrates the representative microstructure of Ni50?3Ti29?7

5 Composition regions in which different precipitate phases exist. The relative intensity of an X-ray diffraction peak for

each phase is plotted colour-coded within a section of the ternary NiTiHf diagram for a HfNi(Ti), b Ti2Ni(Hf), c

Hf2Ni(Ti), and d Laves phase (colour code: red5high; green5medium; blue5low intensity)52 Figure 5 will be repro-

duced to be mono on the printed version

6 a typical bright-eld image of martensite in Ni44Ti36Hf15Cu5 ribbon annealed at 500uC for 1 h and SAD pattern takenfrom region W, electron beam//[1-10]M,T; b typical martensite structure in the ribbon annealed at 700uC for 1 h and theSAD pattern taken from region D, electron beam//[2-11]M1,M2//[ -2 -11]M3; c SAD pattern obtained from Ti2Ni(Hf) type preci-pitates formed in ribbon annealed at 500uC for 1 h, electron beam//[110]Ti2Ni(Hf)

36

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Materials Science and Technology 2014 VOL 30 NO 13a 1535

7 a unit cell of unrelaxed orthorhombic model of the H-phase.25 SAD patterns of the b [001]B2, c [1-11]B2, d [1-10]B2 and e[110]B2 zone axes obtained from a single large particle in a Ni52Ti28Hf20 alloy.

24 The small arrows and circles mark the

additional reections arising from the precipitate

8 Bright-eld images of the Ni50?3Ti29?7Hf20 alloy a extruded at 900uC, b aged at 550uC for 3 h and c aged at 650uC for3 h.27 Bright-eld images of the Ni45?3Ti29?7Hf20Pd5 alloy aged at d 550uC and e 650uC for 3 h.

33 Inset in d is the enlar-

gement of area D. The SAD patterns shown in d and e were taken from the area D and E, respectively

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Hf20 alloys in as extruded and aged conditions.27 The

bright-field image of the as extruded Ni50?3Ti29?7Hf20alloy is shown in Fig. 8a. Precipitate formation was notconfirmed in the as extruded condition. Figure 8b and cshows TEMmicrographs of the extruded Ni50?3Ti29?7Hf20alloy aged at 550 and 650uC for 3 h, respectively. Fine andcoherent H-phase precipitates were formed in the 550uCaged specimen. When the aging temperature increasedfrom 550 to 650uC, the precipitate size increased fromabout 20 to 4060 nm. The interparticle distance alsoincreased after aging at 650uC for 3 h compared with the550uC for 3 h case. Figure 8d and e shows the H-phaseprecipitates and B199 martensite in slightly (NizPd)-richNi45?3Ti29?7Hf20Pd5 alloys in aged conditions.

33 The sizeof the spindle-shaped H-phase was increased when theaging temperature was increased from 550 to 650uC.The fine and coherent H-phase precipitates in the Ni50?3Ti29?7Hf20 and Ni45?3Ti29?7Hf20Pd5 alloys aged at 550uCimproved the shape memory and superelastic propertiesdue to precipitation strengthening. However, the alloysaged at 650uC exhibited relatively poor shape memoryand superelastic properties due to large precipitate sizes.

The martensite morphology in Ni-rich NiTiHf-basedalloys is affected by the size and interparticle distance ofH-phase precipitates. The martensite variants in the asextruded Ni50?3Ti29?7Hf20 alloy show spear-like mor-phology and high density of twins can be seen inside themartensite plates (Fig. 8a). Han et al.57,58 have reportedtwo types of martensite morphologies; spear-like andmosaic-like in NiTiHf alloys and they also revealed thateach martensite lath is consisted of (001)B199 compoundtwins. If the precipitates were small and interparticledistance was short, the growing martensite plates canabsorb all the precipitates during growth as it can beseen in the 550uC aged Ni50?3Ti29?7Hf20 alloys (Fig. 8b).The large martensite plates were related by the {011}B199type I twinning mode, which was confirmed by the SADpattern shown in Fig. 8b taken at the interface of theplates. It should be noted that no internal twins wereobserved in the large martensite plates in the 550uC agedspecimen. On the other hand, when the precipitates werebig and interparticle distance was large, martensiteplates can be formed between the precipitates and thethickness of the plates was controlled by the interparticledistance of the precipitates (Figs. 8ce). In Ni45?3Ti29?7Hf20Pd5, the SAD patterns were taken from the area Dfor the 550uC aged specimen (Fig. 8d) and from the areaE for the 650uC aged specimen (Fig. 8e). It was revealed

that the main twinning mode observed in the martensitewas (001)B199 compound twin in both aging conditions.It was suggested that the internal twinning type was notaffected by the size of the H-phase precipitates if themartensite plates are formed between the precipitates.

Addition of Nb and Pd to NiTiHf alloys

In NiTiHf-based alloys, the lattice invariant shear (LIS)of the martensitic transformation depends on the alloycomposition. The (001)B199 compound twins have beenfrequently observed in martensite plates and consideredas the LIS in NiTiHf alloys.57,58 However, recently, it wasfound that the ,011.B199 type II twin was the LIS in a(NizPd)-rich Ni45?3Ti39?7Hf10Pd5 alloy which was homo-genised at 900uC followed by furnace cooling.30 TheNi45?3Ti39?7Hf10Pd5 alloy exhibited less hardening duringtransformation compared to a Ni45?3Ti29?7Hf20Pd5 alloywhich has (001)B199 compound twins.

Figure 9a shows a bright-field TEM image for theNi45?3Ti39?7Hf10Pd5 alloy

30 which consisted of two phases,B2 austenite and B199martensite at room temperature. Inthe SAD pattern taken from the austenite phase (Fig. 9b),there were diffuse streaks along the ,110.B2

* directionsin reciprocal space. The diffuse streaks could be attributedto the formation of very small precipitates during the slowfurnace cooling process from the homogenisation tem-perature. Sandu et al.63 also observed similar diffusestreaks in an aged Ni-rich NiTiZr alloy. The SAD patterntaken from the martensite phase (Fig. 9c) indicated thatthe internal twins formed in the martensite variants werethe ,011.B199 type II twins. Compared to the (001)B199compound twin, lower density of twins is found when theLIS is the,011.B199 type II twin. It is noted that the LISin NiTi binary alloys is known as the ,011.B199 type IItwin and the (001)B199 compound twin has been observedin NiTi alloys as a deformation twin.64 The (001)B199compound twin has been also found in nanocrystallineNiTi alloys65 and in aged Ni-rich NiTi alloys with fineNi4Ti3 precipitates.

66 These results suggested that the LISin NiTiHf-based alloys depends on the alloy compositionand the size and interparticle distance of precipitates.

Kim et al.35 reported that addition of Nb to NiTiHfalloys causes the formation of a soft Nb-rich b phase andimproves the cold workability. The stability of shapememory properties is improved by the precipitation of theb phase, although the shape recovery strain decreases bythe addition of Nb. Figure 10 shows the back-scatteredscanning electron images of (Ni49?5Ti35?5Hf15)Nb alloys.

35

9 a bright-eld image of Ni45?3Ti39?7Hf10Pd5 alloy homogenised at 900uC followed by furnace cooling, b SAD pattern takenfrom B2 austenite phase and c SAD pattern taken from martensite phase indicating B199 monoclinic structure30

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In Fig. 10a, the Ti2Ni type precipitate can be seen in theNi49?5Ti35?5Hf15 ternary alloy with a slightly dark contrast.The b phase, which appears white on the images, wasobserved even after 1% Nb addition (Fig. 10b), indicatingthat the solubility limit of Nb in the matrix was less than1%. The amount of the b phase increased with increasingNb content. When 15% Nb was added, it exhibited a fullylamellar microstructure as shown in Fig. 10c, which is a

characteristic of eutectic solidification. This fine lamellarstructure strengthened the matrix and prohibited plasticdeformation during transformation.

Morphologies of reoriented martensite andstress induced martensiteAcar et al.67 have reported the morphology of thereoriented martensite in a Ni45?3Ti34?7Hf15Pd5 alloy.

10 Back-scattered scanning electron images of a Ni49?5Ti35?5Hf15, b (Ni49?5Ti35?5Hf15)Nb1 and c (Ni49?5Ti35?5Hf15)Nb15 alloys35

11 Bright-eld image (TEM) of a as homogenised Ni45?3Ti34?7Hf15Pd5 and b 8% deformed alloy with corresponding SAD

pattern.67 Bright-eld image of Ni49Ti36Hf15 c deformed to 8% at 250uC and d deformed to 16% at 250uC.21 The SAD

pattern shown in c was taken from area II

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1538 Materials Science and Technology 2014 VOL 30 NO 13a

Transmission electron microscopy observation was car-ried on a homogenised sample after 8% compressivedeformation at 15uC (below martensite finish tempera-ture, Mf). Figure 11a and b shows the TEM micrographsobtained from the as homogenised and deformed samples,respectively. There are fine twins in the martensite platesin the as homogenised sample (Fig. 11a). In the deformedsample (Fig. 11b), thicker martensite plates were formedby the reorientation of martensite variants as compared tothe as homogenised sample. The thick martensite platesare considered to be favourable martensite variants understress. The inset in Fig. 11b is an SAD pattern taken fromthe interface between the martensite plates A and B. It wasrevealed that the fine twins in the martensite plates are(001)B199 compound twins and the boundary between theplates A and B is close to the {111}B199 type I twin plane(so called {111}B199-type boundary).

36 It is consideredthat the {111}B199-type boundary can move under stresswithout significant detwinning of fine (001)B199 compoundtwins in martensite plates.

Dalle et al.68 investigated the morphology of reorientedmartensite of annealed (800uC for 1 h) Ni49?8Ti42?2Hf8after 10% tensile deformation. They observed finer(001)B199 compound twins in the deformed material com-pared to the as annealed material with self-accommo-dated martensite. They suggested that the detwinning ofthe (001)B199 compound twins is difficult and proposedthat, instead of the detwinning, a supplementary (001)B199mechanical twinning could take place during deformationby a mechanism of the repetition of the dislocation slip onthe (001)B199 plane.

Meng et al.20,21 investigated the morphologies of theSIM inNi49Ti36Hf15 which were solution treated at 1000uCfor 1 h and deformed in tension at 250uC. Figure 11cshows the typical morphology of the preferentiallyoriented SIM variants and the SAD pattern taken fromthe area II for the 8% deformed Ni49Ti36Hf15.

21 (001)B199compound twins were mainly observed in the SIM plates.The SAD pattern revealed that the SIM plates were twin-related with {011}B199 type I mode, which was similar tothe thermally transformed martensite.57,58 The preferen-tially oriented SIM variants were disappeared and severalmartensite variants were intersected into each other after

deformation. Figure 11d shows the variant-crashed/var-iant-intersected morphology after deformation of 16%.The interfaces of the martensite variants are blurred in thevariant-crashed/variant-intersected morphology. They notedthat the stress induced martensitic transformation anddislocation slip occurred simultaneously during loadingand suggested that the introduction of dislocations in-creases the martensite variants with the variant-crashed/variant-intersected morphology.

Mechanical behaviour of NiTiHf-basedshape memory alloysThe relatively high degree of brittleness or poor cyclicstability in NiTiHf alloys are the main obstacles fortheir commercial high temperature applications. It hasbeen observed that ductility of NiTiHf alloys could beimproved by deformation at higher temperatures in Ni-lean NiTiHf alloys. Ni49Ti36Hf15 alloys failed after 7% ofbending deformation at room temperature while they didnot fracture until 30% tensile strain at 260uC.42,69

Material properties of NiTiHf-based alloys can becontrolled by aging at different temperatures and time asillustrated in Fig. 12. Meng et al.53 illustrated that yieldstrength of Ni49Ti36Hf15 can be adjusted by aging at700uC while ductility was constant as shown in Fig. 12a.The strength of matrix was improved after 20 h butfurther increase in aging time decreased the strength ofalloy which can be related to the size, interparticledistance and volume fraction of Ti2Ni(Hf) precipitates.Figure 12b shows the hardness (HV) of Ni50?3Ti29?7Hf20

and Ni45?3Ti29?7Hf20Pd5 alloys as a function of agingtemperature for 3 h aging. The increase in the hardness inthe both alloys can be attributed to formation of nano-size coherent precipitates that minimises the dislocationmotion. The decrease in hardness at high aging tempera-tures in the both alloy systems can be linked to formationof semi-/non-coherent precipitates and larger interparticledistance due to over-aging and thus the lack of precipita-tion strengthening which was also demonstrated in Fig. 8.It is also clear that Ni45?3Ti29?7Hf20Pd5 is harder in naturewhen is compared to Ni50?3Ti29?7Hf20.Initially, SMA properties of Ni-lean NiTiHf alloys

were mainly investigated due to low TTs of Ni-rich

12 a effect of aging time on yield strength and elongation of Ni49Ti36Hf1553 and b hardness values of Ni50?3Ti29?7Hf20 and

Ni45?3Ti29?7Hf20Pd5 alloys as a function of aging temperature

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Materials Science and Technology 2014 VOL 30 NO 13a 1539

NiTiHf alloys.23,59 Shape memory effect with 3% recover-able strain or 80% recovery of 6% applied strain isobserved in compression and bending while 80% reco-very of 2?5% applied tensile strain is observed inNi49Ti36Hf15.

20,69 The poor shape memory effect isattributed to high stress (y500 MPa) for martensitereorientation and high slope in SIM transformation region(no plateau region observed) confirmed by tensile experi-ments.69,70 Although no superelasticity is observed in Ni-lean NiTiHf alloys,21,23 0?88% strain for two-way shapememory effect has been observed.71 Unstable cyclicbehaviour is a major problem in Ni-lean NiTiHf alloyswhere it has been observed that TTs were decreased by40uC during stress free thermal cycling of Ni49Ti41Hf10after 20 cycles.14

In Ni-rich NiTiHf alloys, almost perfect dimensionalstability with 3% strain under a compressive stress of500 MPa was observed as illustrated in Fig. 13a.26 It canbe seen from Fig. 13b that the stressstrain curve ofsolutionised Ni-lean Ni49Ti36Hf15 at temperature aboveAf showed high slope in SIM transformation region anddeformation was not fully recovered upon unloadingwhich was similar to that in cold worked TiNi,72 TiPd73

and as extruded or overaged Ni-rich NiTiHf27 alloys.Aging can improve the shape memory and materialproperties of Ni-rich NiTiHf alloys. Figure 13c showsthe superelasticity responses of as extruded and agedNi50?3Ti29?7Hf20 alloys.

27 Perfect superelastic behaviourwith 4% recoverable strain was revealed at 240uC afteraging at 550uC for 3 h in Ni50?3Ti29?7Hf20. The improve-ment in superelastic behaviour with aging can be attributedto the presence of coherent and fine H-phase precipitates(as discussed in the microstructure part and shown inFig. 8), which strengthen the matrix. Poor superelasticresponse after aging at 650uC for 3 h can be attributed toloss of the coherency of the coarsened precipitates. It isworth to note that beside the fully recoverable strain, Ni-rich NiTiHf exhibited high yield strength at hightemperature and the ClausiusClapeyron (CC) slopeswere between 7 and 13 MPa uC21. It should also be notedthat almost fully recoverable strain with small amount ofplastic deformation under 1000 MPa with low plastic andperfect superelastic behaviour was obtained in Ni-richNi50?3Ti29?7Hf20 along the [111] orientation.

28

Another method to improve the shape memoryproperties of NiTiHf has been the quaternary alloying.Recently, several studies29,3133 have revealed the effectsof Pd addition on the mechanical properties of Ni50?3Ti29?7Hf20. It was shown that perfect superelastic curves(with negligible plastic strain) at stress levels as high as2 GPa were possible for aged polycrystalline Ni45?3Ti29?7Hf20Pd5

33 for a temperature window of 50130uC asillustrated in Fig. 14a. However, full strain recovery wasnot observed in over aged materials due to the formationof large precipitates as shown in Fig. 8 and high degree ofhardening was observed during transformation. Whenthe Hf content was decreased to 10% in Ni45?3Ti39?7Hf10Pd5, high slope in SIM transformation region wasnot observed that can be attributed to the formation ofdifferent type and density of twinning as discussed in themicrostructure section (see Fig. 9). It was mentionedbefore that the ,011.B199 type II twin was the LIS ina (NizPd)-rich Ni45?3Ti39?7Hf10Pd5 alloy

30 in contrastto the (001)B199 compound twin in Ni45?3Ti29?7Hf20Pd5alloys.33 Type II twins could be held responsible for thelack of hardening in the transformation region of Ni45?3Ti39?7Hf10Pd5 alloys containing 10% Hf compared toNi45?3Ti29?7Hf20Pd5 alloys that has (001)B199 compoundtwins. The growth ofmartensite variants with thin (001)B199compound twins is more difficult in contrast to,011.B199

13 a thermal cycling under stress response of Ni-rich Ni50?3Ti29?7Hf20,26 b superelastic behaviour of hot rolled Ni-lean

Ni49Ti36Hf1520 and c superelastic behaviour of Ni-rich Ni50?3Ti29?7Hf20

27

14 Stressstrain responses of a Ni45?3Ti29?7Hf20Pd533 and

b Ni45?3Ti39?7Hf10Pd530 at temperatures above Af

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1540 Materials Science and Technology 2014 VOL 30 NO 13a

type II twins. Thus, the required energy to complete theSIM transformation increased and resulted in high slopein SIM transformation region.Another alloying addition to NiTiHf alloys has been

Cu where it generally improved the two-way shapememory effect and thermal stability of NiTiHf alloyswhile decreased their TTs.3638 It has been reported thatNi45?3Ti29?7Hf20Cu5 can recover compressive strains ofy2?2% under 700 MPa at temperature above 100uC andcan produce 0?8% two-way shape memory strain attemperature above 80uC.39 Perfect superelasticity responsewas not observed in Ni45?3Ti29?7Hf20Cu5 due to its high CC slope (about 1425 MPa uC21) and work hardeningcoefficient in addition to low yield stress for plasticdeformation.39

It can be seen from Fig. 15 that the Ni-rich Ni50?3Ti29?7Hf20 alloys had higher TTs and strength than the Ni-leanNi49?5Ti35?5Hf15 alloys. This can be attributed to thedifference in precipitates types where H-phase particleswere observed in Ni-rich and Ti2Ni(Hf) precipitates wereformed in Ni-lean NiTiHf after aging. NiTiHfNb hashigher transformation strain but lower strength than

Ni-rich NiTiHf. Also, as the Nb content was increasedthe TTs were decreased and the shape memory propertiesof (Ni49?5Ti30?5Hf15)Nb5 alloys became more stable as theplastic strain was decreased due to the fine lamellarstructure that strengthen the matrix as shown in Fig. 10.Moreover, the shape recovery ratio was increased as Nbcontent increased. Addition of Pd to Ni50?3Ti29?7Hf20decreased the TTs and transformation strain while itimproved the strength of the alloy.

Cycling instability is a major concern of the NiTiHfalloys for high temperature applications. Figure 16 showsthe thermal cycling under 200 MPa experiments of Ni49?8Ti42?2Hf8 in homogenised and equal channel angularextruded at 650uC conditions.18 It is clear that severeplastic deformation improved the thermal cyclic stabilityand decreased the thermal hysteresis of Ni-lean Ni49?8Ti42?2Hf8 alloys. Moreover recoverable strain increasesand irrecoverable strain decreases. Formation of preci-pitates influences the cyclic degradation resistance inNiTiHf27 where small coherent precipitates generallyimprove the thermal cyclic stability while larger precipi-tates do not affect the stability.

Work output, damping capacity andpotential applications of NiTiHf-basedalloysShape memory alloy based actuators can be employed aslight weight and energy efficient alternatives of hydraulicor pneumatic systems2 in automotive, aerospace anddown-hole energy exploration industries. Alloys withhigher work output values can be used to decrease therequired weight or size of actuators.

The maximum work output levels of various NiTiHf-based SMAs as a function of their average operatingtemperature range are shown in Fig. 17a. Work outputcan be calculated as the mathematical multiplication ofreversible transformation strain and applied stress inconstant-stress thermal cycling experiments. NiTi alloyshave work output densities of about 1218 J cm23,74 whileNiTiPd andNiTiPt alloys have work output capabilities of69 and 13 J cm23 respectively75 at temperatures above150uC. The work output of Ni-rich NiTiHf polycrystallinealloys was found to be 1820 J cm23.27 Ni45?3Ti29?7Hf20Cu5 alloys can generate work outputs of around 1415 J cm23 while NiTiHfNb alloys have work output levels

15 Thermal cycling under constant compressive stress of

500 MPa results for NiTiHf-based SMAs27,33,35

16 Straintemperature response of Ni49?8Ti42?2Hf8 under 200 MPa a homogenised and b equal channel angular extruded

at 650uC using route 2C under 200 MPa18

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Materials Science and Technology 2014 VOL 30 NO 13a 1541

of 1718 J cm23 above 100uC and 150uC, respectively.35,39

On the other hand, Ni45?3Ti29?7Hf20Pd5 alloys can generatehigher work outputs of 3235 J cm23 (up to 120uC)compared to other NiTiHf-based SMAs, while uppertemperature capability is somewhat limited compared tothe above mentioned NiTiHf-based alloys.33

Figure 17b shows the damping capacities/absorbedenergies of NiTi-based alloys as a function of transfor-mation stress. Damping capacities can be calculatedfrom the area between the loading and unloading curvesin a superelastic cycle and can be explained as the abilityto repeatedly dissipate unwanted energy from a system.SMA based mechanisms could be employed under highnumber of cycles in real practical applications. Thus,stability of a superelastic curve is essential for dampingapplications. A HTSMA that could absorb large energywill be very appealing for high temperature dampingdevices. In addition to the high work output, NiTiHf-based SMAs have high damping capacities. They couldbe employed in aircraft engines as a damper for acousticenergy and construction for countering seismic move-ments in impact damping devices.The damping capacity of Ni45?3Ti29?7Hf20Pd5 alloys is

3034 J cm23 stemming from its outstanding mechanicalhysteresis (around 900 MPa) and good superelastic strainof 4%.33 In related systems, the damping capacity is 16,1820, 38, and 54 J cm23 for NiTi, NiTiHf, NiTiNb andNbTi/NiTi nanocomposites, respectively.7678 The Ni45?3Ti29?7Hf20Pd5 alloy has similar damping capability toNiTiNb alloys that are often used in coupling applica-tions. However, it should be noted that Ni45?3Ti29?7Hf20Pd5 alloys have the ability to operate at much higherstresses (y2 GPa) than the NiTiNb systems. Dampingcapacities of NiTiHfCu and NiTiHfNb were not com-pared since full recoverable superelastic cycles have notbeen reported in literature.

ConclusionsFrom the present review of NiTiHf-based alloys, it is clearthat NiTiHf-based alloys are attractive candidates forhigh temperature, high strength and damping applications.Their TTs and strength can be adjusted by heat treatments.They could show perfect superelasticity above 100uC andshape memory effect under high stress levels. However,some of their drawbacks such as low ductility, high slope in

SIM transformation region and low cyclic stability are stillremained to be improved.

It has been shown that microstructural control bycomposition alteration and aging is essential in tailoringshape memory and mechanical properties (e.g. TTs,strain, hysteresis and strength) in NiTiHf-based alloys.Based on the composition and precipitation character-istics (e.g. precipitate size and interparticle distance), themain microstructural features such as twin type, marten-site morphology can be adjusted that would affect theshape memory and mechanical properties. In Ni-leanNiTiHf-based alloys, the size of the Ti2Ni(Hf) precipi-tates are effective to control the martensite morphology.(001)B199 compound twins are dominant when Ti2Ni(Hf)precipitates are small (about 2040 nm) and homoge-neously distributed while {011}B199 type I twins becomedominant with increasing the size of Ti2Ni(Hf) precipi-tates. The martensite morphology in Ni-rich NiTiHf-based alloys is affected by the size and interparticledistance of H-phase precipitates. When the precipitatesare small and interparticle distance is short, martensiteplates can absorb the precipitates during their growth andthey are mainly twin-related with the {011}B199 type Imode. On the other hand, when the precipitates are bigand interparticle distance is large, martensite plates canbe formed between the precipitates. The thickness of theplates is governed by the interparticle distance of theprecipitates. The formation of fine H-phase precipitatessignificantly improves the shape memory and superelasticproperties due to precipitation strengthening.

Pd addition decreases the TTs of NiTiHf alloys while thematrix strength was increased by solid solution strengthen-ing. Ni45?3Ti29?7Hf20Pd5 alloys can show perfect superelasticresponse under extremely high compressive stress levels of2 GPa with negligible plastic deformation. NiTiHf(Pd)alloys have high work outputs and damping capacitiesreaching up to 3035 J cm23 owing to their good strain,high strength and large mechanical hysteresis. Nb additionto NiTiHf alloys improves the cold workability and thestability of shape memory properties while decreases theshape recovery strain by the precipitation of the b phase.Detailed studies are needed to gain the fundamental

understanding on processingcompositionmicrostruc-tureproperty relationships and reveal the true potentialof NiTiHf-based alloys. Currently, they are the most

17 Comparisons of a work outputs and b damping capacities for typical NiTi-based SMAs

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1542 Materials Science and Technology 2014 VOL 30 NO 13a

promising alloys for high temperature and strengthapplications.

Acknowledgement

This work was supported by the NASA EPSCORprogram under grant no. NNX11AQ31A.

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