高温合金综述2010_high_temperature_SMA

LEADING EDGE REVIEW

High temperature shape memory alloys

J.Ma1,I.Karaman*1,2and R.D.Noebe3

Shape memory alloys(SMAs)with high transformation temperatures can enable simplifications and improvements in operating efficiency of many mechanical components designed to operate at temperatures above100u C,potentially impacting the automotive,aerospace,manufacturing and energy exploration industries.A wide range of these SMAs exists and can be categorised in three groups based on their martensitic transformation temperatures:group I,transformation temperatures in the range of100–400u C;group II,in the range of400–700u C;and group III,above 700u C.In addition to the high transformation temperatures,potential high temperature shape memory alloys(HTSMAs)must also exhibit acceptable recoverable transformation strain levels, long term stability,resistance to plastic deformation and creep,and adequate environmental resistance.These criteria become increasingly more difficult to satisfy as their operating temperatures increase,due to greater involvement of thermally activated mechanisms in their thermomechanical responses.Moreover,poor workability,due to the ordered intermetallic structure of many HTSMA systems,and high material costs pose additional problems for the commercialisation of HTSMAs.In spite of these challenges,progress has been made through compositional control,alloying,and the application of various thermomechanical processing techniques to the point that several likely applications have been demonstrated in alloys such as Ti–Ni–Pd and Ti–Ni–Pt.In the present work,a comprehensive review of potential HTSMA systems are presented in terms of physical and thermomechanical properties,processing techniques, challenges and applications.

Keywords:High temperature shape memory alloys,Intermetallics,Thermomechanical processing,Shape memory effect,Superelasticity,Martensitic transformation

I.Introduction

Since the discovery of shape memory alloys(SMAs), much progress has been made both in the scienti?c understanding and application of these multifunctional materials.Owing to the unique behaviours of shape memory effect and superelasticity,SMAs have become a major materials class of choice in the biomedical industry and are beginning to permeate into other technological areas.However,the complexity of their governing microstructural mechanisms and physical behaviours have rendered sporadic commercial interest in these materials.Nevertheless,there is a recent revitalisation of interest in SMAs,driven primarily by the aerospace and automotive industries,for their potential to operate as solid state actuators.

Shape memory effect is a phenomenon whereby a deformed material could recover its predeformed shape after being heated.When this procedure is performed against some biasing force,the material is capable of doing work from its shape change.Superelasticity is an isothermal phenomenon where the material is able to recover high amounts of strain(up to more than20%in a few single crystalline alloys)triggered by mechanical stress.These two behaviours are the result of reversible martensitic transformation–a diffusionless solid state phase transformation mechanism that can be activated by temperature,stress and magnetic?eld.

Current practical uses for SMAs are,however,limited to temperatures below100u C.This is the transformation temperature limit of the two most commercially successful SMA systems:the near equiatomic Ni–Ti binary and Cu based ternary alloys.During thermo-mechanical processes required to produce stable shape memory or superelastic behaviour,the transformation temperatures are further reduced.1Naturally,such limitation hinders the utility of SMAs in high tempera-ture applications,and necessitates design modi?cations for SMA containing components in order to reduce operating temperatures to below100u C,or completely abandon their use.On the other hand,the unique properties of SMAs become even more bene?cial at high temperatures,since it is preferable to adopt single piece adaptive and multifunctional components over more complex multicomponent assemblies due to the higher

1Department of Mechanical Eng.,Texas A&M University,College Station,

TX778433123,USA

2Materials Science and Eng.Interdisciplinary Graduate Program,Texas

A&M University,College Station,TX778433003,USA

3NASA Glenn Research Center,MS49–3,Cleveland,OH44135,USA

*Corresponding author,email ikaraman@https://www.wendangku.net/doc/504773849.html,

?2010Institute of Materials,Minerals and Mining and ASM International

Published by Maney for the Institute and ASM International

DOI10.1179/095066010X12646898728363International Materials Reviews2010VOL55NO5257

likelihood of wear or damage and the greater weight and volume required by the latter.These issues have triggered several studies on possible SMAs with transformation temperatures above100u C.This class of materials is simply referred to as high temperature shape memory alloys(HTSMAs).As of now,despite intensive research efforts in recent years,HTSMAs have yet to be utilised commercially in appreciable amounts due to a number of unresolved issues.

Several recent reviews on HTSMAs are available,2–7 but the majority is restricted in scope to the basic metallurgical properties of reported materials and/or focus on only a few alloy systems.The present work seeks to provide a more comprehensive coverage of the possible alloy systems that display high temperature shape memory and superelastic behaviours,as well as processing techniques and the potential applications of HTSMAs.In addition,we intend to provide a resource for industry and facilitate the introduction of SMAs into

commercial high temperature applications.The primary target of this article is centred upon thermomechanical properties of SMAs,namely the transformation tem-peratures,shape memory and superelastic behaviours, and the bulk of the discussion on individual alloy systems will focus on the quanti?cation of these properties,processes that have been shown to improve them,and governing microstructural phenomena in their operation.Topics such as physics,thermodynamics and crystallographic theory of martensitic transforma-tion will not be addressed in detail.

First,a brief introductory discussion of SMAs is included for readers unfamiliar with these materials in Section II.Section II.1is designed to provide a basic understanding of SMAs for non-experts in this?eld.The section evolves around the stress–temperature phase diagram where transformation temperatures are plotted as a function of applied stress.Various phenomena related to SMAs,such as shape memory and superelastic behaviour,are described based on the deformation temperature relative to the transformation temperatures, and microstructural changes that take place during these behaviours.The origin of two way shape memory effect and processes that create it are also discussed. Following this,the focus is shifted toward important engineering properties of shape memory and superelastic behaviour,such as recoverable strain,irrecoverable strain,thermal and stress hysteresis,and work output. In Section II.2,primary factors that affect these properties are discussed.These topics include effects on shape memory and superelastic behaviour from conventional processing techniques–work hardening and precipitation hardening,the role of crystallographic texture,the effect of martensite/austenite structure,and variables unique to HTSMAs,such as oxidation and creep.In essence,this section addresses the question of how one may be able to improve shape memory and superelastic behaviour.

Section III provides detailed information on individual HTSMA systems based on the following temperature ranges:100–400u C,400–700u C and above700u C.These temperature ranges were chosen based on temperature ranges of potential applications.The critical character-istic transformation temperatures of the alloys will be used for their classi?cation,i.e.martensite?nish tem-perature M f will be used for alloys studied for shape memory effect and austenitic?nish temperature A f will be used for those studied for superelastic behaviour. Unconventional processing techniques such as rapid solidi?cation,physical vapour deposition and severe plastic deformation will also be discussed in subsections for each alloy system.Finally,some proposed applica-tions of HTSMAs will be summarised in Section IV and the present article will conclude by recapping some major problems and challenges facing the development and commercialisation of HTSMAs.

II.Basics of SMAs and issues at high temperatures

II.1.Brief introduction to SMAs

For readers less familiar with SMAs,a brief overview of these materials is provided here.Since the present article is focused primarily on the thermomechanical beha-viours of HTSMAs such as shape memory effect and superelasticity,a detailed description of these behaviors and the underlying microstructural mechanisms are reviewed.

One way shape memory effect and superelasticity are the most frequently utilised SMA behaviours in applica-tions.One way shape memory effect refers to the ability of an SMA deformed at a low temperature to recover the deformation when heated to a higher temperature. In other words,the material is able to memorise its undeformed shape(Fig.1).Superelasticity refers to the ability of SMAs to recover large amounts of stress induced inelastic deformation immediately upon unload-ing.Both behaviours are a consequence of the reversible martensitic transformation.

II.1?1.Mechanisms of shape memory effect and superelasticity

Martensitic transformation is a solid to solid phase transformation that occurs through a coordinated shear movement of atoms over very short(on the order of angstroms)distances where atoms retain close relation-ship with one another,as opposed to random long range diffusion of atoms.The high temperature phase, austenite,transforms to a low temperature phase, martensite,upon cooling.Because the crystal structure of austenite is different than that of martensite,it

is 1The one way shape memory effect:the initial SMA strip is deformed to the‘formed’state,but upon heating,the strip is able to return to its nearly undeformed shape.7 (Reproduced with permission from The Taylor&Francis Group)

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possible to introduce a macroscopic shape change that accompanies the transformation.

When martensite forms in austenite,the difference in their crystal structures generates large amounts of local strain.This strain is large enough so that it cannot be purely accommodated elastically.Instead in SMAs,which undergo reversible martensitic phase transforma-tion,the strain is accommodated by producing a twinned martensite structure (see Fig.2).When the higher symmetry austenite transforms to the lower symmetry martensite,it may do so in several ‘ways’,called martensite lattice correspondence variants.The number of such variants that can be formed is determined by the symmetry of martensite and austenite;for example,there are 12lattice correspondence variants of a monoclinic martensite to a cubic austenite.8In essence,each lattice correspondence variant is a ‘variation’of the martensite with a different orientation relationship to the austenite,but they are all energetically equivalent to one another under stress free conditions.By forming a structure of twin related lattice correspondence variants,the marten-site is able to accommodate a signi?cant portion of the strain associated with the change in crystal structure,as shown in Fig.2.These twin related lattice correspon-dence variants are collectively referred to as a habit plane variant,and several different habit plane variants can then be formed in such a way that together,they reduce the remaining strain of the transformation.This means that the transformation from austenite to martensite can be made to produce nearly no macroscopic shape change,and the resulting structure of the martensite phase that accomplishes this is then considered to be ‘self-accom-modated’,as seen in Fig.3.Under an external biasing stress,certain habit plane variants become energetically favoured and form/grow at the expense of others in a process known as martensite reorientation.In addition,the martensite may also detwin,where analogously,the lattice correspondence variant favoured under stress grows at the expense of others.Both martensite reorientation and detwinning results in a macroscopic shape change,and give rise to the shape memory behaviour and superelasticity.More details can be found in the literature regarding the nature of martensitic transformation,9–11structural description of twinning in martensite,12–18and self-accommodation.19–25For the sake of simplicity,detwinning and martensite reorienta-tion will be treated as the same mechanism in this introductory section.

The martensitic transformation can be induced both thermally and through the application of external stress.In other words,application of stress and reduction in temperature both act as driving forces for the austenite R martensite transformation.In fact,there is a linear relationship between the two.This relationship is derived from the thermodynamics relationships of phase transformation and is called the Clausius–Clapeyron relationship.Roughly,it states that d s /d T 5constant,and the transformation temperature is a straight line in the s 2T space seen in Fig.4.

The transformation process,however,is somewhat more complicated than that illustrated in Fig.4.In general,the transformation is not completed immediately at a single temperature,but occurs gradually over a range of temperature.The temperature during cooling at which the transformation from austenite to martensite,or the forward transformation,?rst begins is called the martensite start temperature,M s .The temperature at which the forward transformation reaches completion is called the martensite ?nish temperature,M f .Conversely,upon heating above the austenite start temperature,A s ,the martensite begins to transform back to austenite –the reverse transformation.The temperature at which reverse transformation is completed is the austenite ?nish tem-perature,A f .Each of these temperatures

approximately

2A simpli?ed illustration of the austenite and martensite

structures.In the absence of stress,austenite trans-forms to twinned martensite upon cooling in order to accommodate strain caused by a change in crystal structure.The twinned martensite is composed of mul-tiple (usually two)twin related lattice correspondence variants,labelled L 1and L 2in this ?gure.When stress is applied,the martensite may detwin,resulting in a single lattice correspondence variant structure and a net shape

change

3Process of self -accommodation in martensite;13L 1and

L 2are two different lattice correspondence variants.Under no stress,pairs of twin related lattice correspon-dence variants form a habit plane variant (H 1,H 2and H 3),shown in top ?gure,and several habit plane var-iants can then arrange themselves in such a way that results in no net shear,and approximate no volume change from the transformation,shown in the middle triangle.When external stress is applied,the degener-acy of the various habit plane variants and lattice cor-respondence variants are lifted,and the most favourable variant –the one most available to accom-modate the desired strain –is formed at the expense of others.Reproduced with permission from Springer Science and Business Media

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follows a Clausius–Clapeyron relationship,and by plotting all four of these temperatures on the same s 2T diagram,a shape memory phase diagram is obtained (Fig.5).However in reality,the slopes of the s 2T relationship of each transformation temperature are generally not the same,and may not even be a straight line due to the effect of microstructural features such as grain size,and microstructural mechanisms such as dislocation slip.Nevertheless,idealised versions are used here for simplicity and to convey the important parameters/mechanisms for clear introduction of thermomechanical behaviours of HTSMAs in the following sections.

The deformation response of SMAs depends on the temperature relative to the transformation temperatures (M s ,M f ,A s and A f )of the alloy.If the material is deformed below M f in a self-accommodated martensite structure (Fig.3),then the strain is accommodated by the growth of one variant favoured by the stress at the expense of others,as well as detwinning (Fig.6).Since all martensite variants are equally stable thermodynamically in the absence of external and internal stresses,the martensite stays in the reoriented and detwinned state,and the material remains in the deformed shape after unloading.When heated above A f after unloading,all martensite transforms back to austenite.When the austenite is once again cooled below M f ,the martensite will again form in a self-accommodated state,and all deformations from detwinning are recovered in the absence of plasticity;this is known as the shape memory effect (Fig.6).

If deformation takes place above A f where the alloy is fully austenitic,the material may deform by stress induced martensitic transformation and possibly det-winning of the transformed martensite.Upon unload-ing,the stress induced martensite is unstable at that temperature,and will completely transform back

to

5A s 2T phase diagram of SMAs undergoing martensitic

transformation.Above A f ,the specimen is fully in the austenite state,and below M f ,the material is fully mar-

tensitic

4The linear relationship between transformation tem-perature and applied stress:an increase in applied stress results in a corresponding increase in transfor-mation temperature

(a)(b)

6a demonstration of the shape memory effect using s 2T phase diagram.An initially twinned (self-accommodated)mar-tensite (state A)is deformed at a temperature below M f ,causing it to detwin (state B)and remain in the detwinned state after unloading (state C).This leads to an external shape change (shown in b ).Upon heating to above A f ,the detwinned martensite transforms fully back into austenite (state D),which again transforms to twinned (self-accommo-dated)martensite when cooled below M f ,restoring the initial shape;b demonstration of the shape memory effect on a s 2e diagram

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austenite and as a consequence,the deformation is recovered.This behaviour is known as superelasticity (Fig.7).

II.1?2.Characteristics of shape memory and superelastic behaviour

For any application of SMAs,there are several important properties other than transformation temperatures. Primarily,these properties can be condensed into the following characteristics:the magnitude(in terms of strain)of shape memory or superelastic behaviour and their reversibility.The former is represented by the recoverable strain e rec,and the latter by the irrecoverable strain e irr and recovery rate.The recoverable strain is the total amount of strain that can be recovered after a complete shape memory or superelastic cycle.It is important to note that e rec depends on how the characterisation experiment is carried out and the experimental conditions such as applied stress level and temperature.In terms of shape memory behaviour evaluated from stress free shape recovery experiments as seen in Fig.6,e rec is the sum of the elastic recovery from unloading and the strain recovery corresponding to detwinned martensite transforming?rst to austenite upon heating and then to twinned martensite upon cooling.For superelasticity(Fig.7),e rec is the total recovery upon unloading and contains a combination of elastic recovery and strain recovery upon reverse transformation from stress induced martensite to austenite.This topic will be addressed in further detail in the corresponding subsec-tions of Section II.1?4.

Clearly,e rec contains two components:elastic strain and shape strain.Shape strain may come from either detwinning/reorientation of martensite for the shape memory effect,e sme,or the transformation from austen-ite to martensite,as in superelasticity,e SE.Observed shape strains also depend on the experimental condi-tions,for example,whether the experiment is conducted in tension or compression,because of the anisotropic nature of SMAs(see Sections II.2?2and II.2?3).In this review article,we have chosen to use e rec as the main measurement of the magnitude of recovery,and its de?nition from different experiments is elaborated upon in Section III.1?4.

Reversibility of shape memory and superelastic behaviours is a measure of the magnitude of recovery relative to the magnitude of initial deformation.For example,if an SMA is deformed to5%strain at a temperature below M f,and recovers4?5%strain from shape memory effect after unloading and heating,there is a0?5%strain that is left over and not recovered.This 0?5%strain is then considered to be e irr and therefore permanent.The recovery rate is the ratio of e rec to the applied deformation,and would be90%in this example. Irreversibility in SMAs is normally considered to be generated by the creation and movement of dislocations, but can also be caused by stabilised martensite that does not transform to austenite even after heating above A f. While the former is truly irreversible,the retained martensite in the latter case may be recovered by heating to even higher temperatures.However,it is not easily possible to know the contribution of each to the total e irr without speci?cally testing for them.Such type of experiments have not in general been carried out in HTSMA studies,and the readers are advised to bear in mind that reported e irr most likely contains contribu-tions from both mechanisms,and possibly from addi-tional mechanisms that may play a role only at high temperatures.These additional mechanisms will be discussed in Section II.3?1.

During martensitic transformation,some of the driving force for the transformation is lost to the environment through non-reversible mechanisms.The magnitude of the associated energy loss or dissipation,is re?ected in the thermal(D T)or stress hysteresis(D s in Fig.7)of a full transformation cycle.Dissipation during the transformation can be due to several mecahnisms, including the creation of defects,emission of acoustic waves,and generation of heat due to internal friction at the phase interfaces.In actuator type applications,

(b)

(a)

7a demonstration of superelasticity using a s2T phase diagram.Initial austenite(state A)is deformed at temperatures above A f,and with suf?cient stress,becomes fully martensite in detwinned state(state B).When stress is removed upon unloading,the specimen returns to a fully austenite state and recovers all imposed deformation immediately (state C).b demonstration of superelasticity on a s2e diagram

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dissipation negatively impacts the ef?ciency of SMA devices.In thermally driven SMA devices,a large hysteresis means a larger heating and cooling range,and thus,higher energy cost and power demands are required for the operation of the device.Fortunately,hysteresis can be reduced through thermomechanical training (Section II.2?1),structural (Section II.2?2)and orientation/texture (Section II.2?3)manipulations,and controlling the thermoelastic nature of the transforma-tion through alloying (Section II.2?4).Furthermore,alloys with large hysteresis can be utilised for damping or energy absorption applications.

Finally,in actuation applications,work output is an important parameter.Work output can be simply de?ned as s 6e ,where e is e rec and s is the applied stress.The evaluation of work output is performed using constant stress thermal cycling experiments (also referred to as isobaric thermal cycling experiments).In these experi-ments,SMAs are cycled through their transformation temperatures under constant applied stress,so that the total e rec from each of these cycles is multiplied by the stress at that cycle to yield work output.This type of experiment is described in more detail in Section II.1?4.When individual HTSMA systems are discussed in Section III,their shape memory and superelastic behaviours will be reported in terms of the properties discussed in this subsection (e rec ,e irr ,recovery rate,D T and work output).This allows comparison to be made among the different systems.Unfortunately,for many alloy systems,several of these properties have not been reported in literature.

II.1?3.Other shape memory related behaviours

The shape memory effect described in previous sections is mainly one-way shape memory behaviour because only the austenite shape is ‘memorised’.The only shape that the martensite is capable of remembering is the same shape as the austenite due to self-accommodated structure of martensite variants.There is another shape memory behaviour where it is also possible to memorise a martensite shape different from the austenite shape.This behaviour is called the two way shape memory effect (TWSME):when the SMA is cooled from austenite to martensite,instead of adapting to a self-accommodated structure,some variants

of the martensite are favoured and the martensite adopts a shape different from that of the self-accommodated structure as seen in Fig.8.Two way shape memory effect is considered to be caused by internal stresses that develop in the SMA from plastic deformation in martensite,25–27superelastic cycles,28aging for precipitation under stress and/or under constraint.29–31The internal stresses gener-ated from these mechanisms are anisotropic,which may be created by directionally organised dislocations or retained martensite from prior thermomechanical training 25–28or by aligned coherent precipitates.29–31The symmetry and arrangement of point defects has also been suggested as a possible explanation for TWSME,32and this mechanism will be discussed in more detail in Section II.2?2.

In practice,TWSME is not used as commonly as the one-way shape memory effect.The reason behind this is that e rec from TWSME is generally smaller,12and because it depends on internal stress,two way shape memory strain tends to deteriorate at higher tempera-tures.Currently,reported TWSME in HTSMA systems are small and very unstable.For this reason,while TWSME is reported whenever available in this review article,it will not be subjected to the same attention and discussion as for the one way shape memory behaviour.Another unique behaviour of SMAs is called ‘rubber-like behaviour’.It is similar to superelasticity in austenite,where deformation is recovered immediately upon unloading.However,rubber-like behaviour occurs com-pletely in martensite.It is suggested that the behaviour stems from symmetry-conforming short-range order of point defects in martensite.32,33In HTSMA literature,rubber-like behaviour has not been reported.Thus,it will be excluded from the present article.

II.1?4.Evaluation of shape memory properties

Characterisation of shape memory and superelastic behaviour in SMAs require a different set of evaluation techniques than those used for ordinary engineering materials.In this section,some of these techniques and important material parameters and properties,that such techniques are designed to determine,will be summarised.II.1?4.a.Transformation temperatures

Transformation temperatures can be directly measured through many techniques including differential scanning calorimetry (DSC),electrical resistivity measurement as a function of temperature,and dilatometry.They can also be measured indirectly through extrapolation of transformation temperatures as a function of stress (similar to Fig.5)from a series of constant stress thermal cycling experiments.

In Fig.9,transformation temperatures are de?ned on a DSC plot.Two peaks are shown,marking the

forward

8Demonstration of one way and two way shape memory

effects.Whereas the martensite normally returns to a self-accommodated structure after cooling from auste-nite in the one way shape memory effect,the TWSME causes the martensite to adopt a more ‘single variant’con?guration.Owing to local oriented internal stresses or other reasons,certain habit plane variants become favoured,and the martensite changes shape upon cooling from

austenite

9Determination of transformation temperatures via DSC

measurements

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and reverse transformations.The forward transforma-tion is exothermic,and the reverse transformation endothermic.The temperatures at the peaks of the forward and reverse transformations are known as martensite peak temperature M p and austenitic peak temperature A p ,respectively.The difference between A p and M p (A p 2M p )is a measure of D T .

Another technique for measuring transformation temperatures is electrical resistivity measurement as a function of temperature.The temperature dependence of electrical resistivity is different for austenite and martensite due to their structural differences.Transformation temperatures are indicated by an abrupt change in resistivity versus temperature slope during heating and cooling.Other techniques such as dilato-metry can be used if there is a signi?cant volume change between martensite and austenite,but the volume change in thermoelastic martensitic transformations is generally very small.In situ X-ray,neutron and electron diffraction can also be used to detect transformation temperatures by observing the temperatures at which diffraction corresponding to martensite or austenite appears or disappears.The indirect measurement of

transformation temperatures via constant stress thermal cycling experiments is discussed in the next subsection.II.1?4.b.Shape memory properties

In most studies,shape memory behaviour is charac-terised by stress free recovery experiments where the specimen is deformed in martensite,unloaded,and then allowed to recover its shape upon heating under no external stress.This type of experiment is described in Fig.6,and the corresponding properties are de?ned in Fig.10.

In actuation applications,however,the shape memory behaviour is never used in this fashion.Instead,an external biasing force always exists on SMA actuators during thermal cycling.During transformation,the associated shape change causes SMA to push against the biasing force,thus doing mechanical work.For characterisation of shape memory behaviour under this condition,constant stress thermal cycling experiments,as shown in Fig.11,are used.A series of such thermal cycling experiments at various stress levels are usually conducted as shown in Fig.12a .The analysis of

the

10Shape memory properties from a one way shape

memory experiment.Detwinning/reorientation stress is denoted by s DT ,while irrecoverable strain,shape memory strain,and elastic strain are denoted by e irr ,e sme and e el ,

respectively

11Representative strain versus temperature response of

an SMA in constant stress thermal cycling experi-ments.Important shape memory characteristics are also shown,such as transformation temperatures,irrecoverable strain e irr ,total recovered strain e rec and transformation thermal hysteresis D T

(a)(b)

12Construction of stress versus temperature phase diagram b for an SMA using constant stress thermal cycling experi-ments in a .The lines for each transformation temperature in b can be extrapolated down to zero stress to determine the stress free transformation temperatures.

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strain versus temperature responses from these experi-ments and determination of transformation tempera-tures at each stress level lead to the construction of the stress–temperature phase diagrams in https://www.wendangku.net/doc/504773849.html,ing the stress–temperature relationships in Fig.12b,trans-formation temperatures at stress free condition can be extrapolated34.

Constant stress thermal cycling experiments are also used to determine the evolution of e rec,e irr,D T and work output of the material as a function of applied stress. Plotting e rec from each thermal cycle as a function of applied stress level generates a plot similar to Fig.13a.In this?gure,e rec increases initially with increasing stress, but eventually reaches a maximum before decreasing rapidly.At the same time,e irr is usually small up to a certain applied stress level,and then,increases mono-tonically with increasing stress.Multiplying the e rec at each stress level with the stress,a plot for the variation of work output with the stress(Fig.13b)is created.Similar to e rec,the work output also increases with increasing stress up to a certain level,reaches a maximum and then decreases with further increase in stress.Not surprisingly, the maximum work output corresponds closely to the maximum e rec.

II.1?4.c.Superelastic properties

Superelastic properties can be evaluated from the loading–unloading experiments shown in Fig.14at different temperatures.However,the temperature at which superelastic experiments are conducted relative to A f is critical.The larger the difference between the test temperature and A f,the greater the driving force will be required to initiate stress induced transformation which results in inferior superelasticity.Above a certain temperature,called M d,stress induced martensitic transformation becomes impossible because plastic deformation will occur?rst.It is,therefore,a good practice to conduct superelastic characterisation experi-ments at a consistent deformation temperature of A f z X,where X is a constant and A f depends on the alloy,in order to compare the superelastic properties of different alloys and the same alloys with different microstructures.

Important superelastic properties,as shown in Fig.14,are similar to those for the shape memory behaviour(Fig.10),namely e irr,e rec and s SIM.Reco-verable strain includes elastic recovery and recoverable shape change from the stress induced martensitic transformation and possibly also martensite detwinning.With increasing applied strain,both e rec and e irr tend to increase.Similar to the shape memory behaviour,e rec reaches a maximum at some strain level while e irr increases monotonically.

Not only can shape memory and superelastic experi-ments be performed through different experiment types, they can also be conducted under different stress states, such as in tension,compression,torsion or bending. Deformation modes in SMAs are highly orientation dependent,especially in martensite,so whether a speci-men is deformed in tension or compression can play a signi?cant role on the outcome of shape memory or superelastic properties.This is illustrated by the differences in the tensile and compressive behaviour of an SMA shown in Fig.15.The fundamental reason behind this has to do with the structures of and lattice correspondences between martensite and austenite,and crystallographic texture of the sample.This issue will be discussed in Section II.2?3.However,we mention this effect here because in many HTSMA studies,the evaluation of shape memory and superelastic effects are performed in different conditions without a uni?ed standard.Some researchers conduct experiments

in

(a)(b)

13a e rec,e irr and b work output as a function of applied stress.The curves are constructed using the data extracted from constant stress thermal cycling experiments in Fig.12

a

14Superelastic properties from a typical experiment:s SIM

denotes the critical stress for stress induced martensitic

transformation;e irr,e se and e el represent irrecoverable

strain,superelastic shape strain and elastic strain,

respectively.Total strain recovery,e rec,during supere-

lastic behavior is the sum of e se and e el

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tension,while others in compression or bending.As a result,reported shape memory properties of the exact same material often vary signi?cantly.In this article,the mode of deformation as well as the temperature at which deformation takes place will always be reported along-side the shape memory or superelastic properties.

II.2.Factors affecting shape memory and superelastic behaviours

Shape memory and superelastic behaviours are functions of a large number of factors that include both inherent microstructural and structural features of the alloy as well as external in?uences such as applied stress and temperature.The complexity causes dif?culty in char-acterisation and the comparisons between the properties of various materials since it is dif?cult to isolate the effects of a single factor on shape memory and superelastic behaviour.This problem has been encountered in a large number of published studies on HTSMAs,where a change in shape memory or superelastic behaviour is ascribed to a single factor,even though some others were not held constant at the same time.Such instances will be alluded to in discussions on alloy systems in Section III.A few of these factors will be discussed in this subsection.While the list is by no means exhaustive,the selected topics are considered,by the authors,to possess the greatest in?uence on shape memory and superelastic behaviour.These factors apply to all SMAs,but they are especially important for HTSMAs because the constraint of excess thermal energy at high temperatures causes a general deterioration of shape memory and superelastic properties,making their improvements through some factors less effective and increasing the importance of the ability to in?uence them through alternative methods.In Section II.3,we will address the topics that concern explicitly with issues at high temperatures,both in terms of shape memory and superelastic behaviour and metallic alloys in general.It is important to realise that not all of the issues in Sections II.2and II.3have been addressed by HTSMA researchers.In fact,a great deal is completely missing in certain alloy systems.The goal of Sections II.2and II.3is both to highlight these issues

relevant to HTSMAs and to note missing points in current research literature and topics that have yet to be studied.These topics are echoed whenever possible during discussions on individual alloy systems in Section III,and revisited in Section IV.

II.2?1.Plastic deformation and strengthening of SMAs

The factor that most greatly in?uences the level of irrecoverable strain e irr in SMAs is considered to be the ability of the material to resist plastic deformation,or its yield strength s y .During martensitic transformation,local internal stresses can often become several times higher than external the applied stresses,and the alloy may deform plastically even though the applied stress level does not exceed s y .For this reason,higher s y generally results in lower e irr and more stable shape memory and superelastic behaviours.Additionally,instead of considering s y alone,it is helpful to compare the difference between s y and the critical stress of the mechanism responsible for the shape memory s DT or superelastic behaviours s SIM ,as seen in Fig.16.

Shape memory behaviour occurs through reorienta-tion and/or detwinning of martensite.The associated critical stress for their onset is usually called the reorientation or detwinning stress.We will not differ-entiate between the two here,and will simply refer to the critical stress shown in Fig.10as s DT .Superelasticity is activated by stress induced martensitic transformation,and the corresponding critical stress is s SIM .In general,s y in both martensite and austenite decreases with increasing temperature (Fig.16),but this is not necessa-rily true for s DT and s SIM .Depending on the alloy system,s DT may increase or decrease with increasing temperature.On the other hand,s SIM always increases with temperature because austenite is stabilised by increasing deformation

temperature.

15The tension compression asymmetry in a polycrystal-line Ni rich Ni–Ti SMA.The superelastic behaviours in tension and compression are notably different in the same sample.35(Reproduced with permission from Springer Science and Business

Media)

16A model for the critical stresses of various deformation

modes as a function of temperature in SMAs.If the yield stresses are above reorientation/detwinning stress s DT and stress induced transformation stress s SIM ,one expects good reversibility and repeatability of shape memory and superelastic behaviours.However,all of these critical stresses depend on alloy type and compo-sition,microstructure and crystal orientations/textures.For many alloys,s y is very close to s DT and s SIM ,making reversible transformations exceedingly dif?cult.Above M d ,plasticity sets in and stress induced martensitic transformation is no longer possible

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Classical work hardening methods,such as cold working,effectively increase s y of ductile SMAs.Cold working improves the shape memory behaviour by increasing s y ,thereby decreasing e irr and allowing higher values of recovered transformation strain e rec to be reached.As a rule of thumb for SMAs,cold working in a phase will stabilise it,and for HTSMAs,the room temperature phase is martensite.Cold working at room temperature should therefore stabilise martensite and increase reverse transformation temperatures.In some alloy systems,hard precipitate phases can be produced by aging at relatively low temperatures around 300–500u C,and the combination of precipitation hardening and work hardening can further improve shape memory and superelastic properties.Unfortunately,a large number of prospective HTSMA systems are intermetal-lics with very limited ductility at room temperature,making them very dif?cult to process.Precipitates that can be used for strengthening usually cause further deterioration in ductility in these alloys.

It is also possible to improve shape memory and superelastic behaviour of SMAs through training.In this process,SMAs are thermally or mechanically cycled between the austenite and martensite a number of times.During training,some level of irrecoverable deformation takes place,but gradually saturates as the number of cycles increases.On one hand,the process acts some-what as work hardening,but more importantly,dislocations or remnant martensite are generated through the transformation process itself,and are therefore located at the ‘correct’places leading to favourable oriented internal stresses and strengthening.These dislocations and remnant martensite decrease the dissipation during transformation and discourage further dislocation creation upon subsequent cycles.Training is most commonly performed by temperature cycling under stress (shape memory training)or through stress–strain cycling at constant temperature (super-elastic training),both of which are demonstrated in Fig.17.Most SMAs can be trained,but desirable results often require a large number of training cycles.It is,however,dif?cult to train SMAs with limited ductility or poor fatigue resistance.

The aforementioned processes for increasing s y are very widely used because they are relatively easy to conduct,but in many HTSMA systems,the lack of ductility,phase instability and the high operating

temperatures prevents these conventional techniques from being carried out,or limits their effectiveness.

II.2?2.Structure of transforming phases in SMAs

Structural factors of austenite and martensite,such as their crystal symmetry and lattice parameters,are central to the understanding of shape memory and superelastic behaviour of SMAs.Structural parameters de?ne maximum capabilities of a transforming system,such as maximum transformation strain achievable.Although these maximum capabilities are often unreach-able in practice,they provide an upper limit for shape memory and superelastic properties.

II.2?2.a.Structural restrictions for martensitic transformation

Crystal structure of martensitically transforming phases dictates whether shape memory and superelastic beha-viour can occur at all between the two phases.According to Bhattacharya et al.,36,37transforming phases must have a group–subgroup relationship in order for shape memory behaviour to exist.Since martensite has a lower symmetry structure than austenite,the martensite must always be a subgroup of austenite.This explains why there are no thermoelastic martensitic transformations between a bcc austenite and an hcp martensite,since they are excluded by the theory.Another important role of structures of transforming phases is that they control the maximum possible transformation strain,e max tr .At the most basic level,the maximum shape strain possible between austenite and martensite is determined by the magnitude of the shear required to go from one structure to the other.As a rule of thumb,the greater the magnitude of this shear,the greater the maximum transformation strain.37Of course,the maximum transformation strain is rarely reached in practice,but it is generally observed that for a given austenite structure,the less symmetrical martensite phases (such as monoclinic B199martensite)produce,

the greater e max tr

than the martensite phases with higher symmetry (such as orthorhombic B19martensite).37In addition,a greater distortion of martensite from austenite often results in higher transformation strains.For example,transformation strain in the Ni–Mn–Ga system with tetragonal martensite increases when the c/a axis ratio of the martensite unit cell deviates further away from 1,because a c/a ratio far from 1signi?es a greater distortion from the cubic austenite.

38

17Thermomechanical training of SMAs through a constant stress thermal cycling (shape memory training)and b con-stant temperature superelastic cycling (superelastic training)

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II.2?2.b.Symmetry of point defects

Otsuka and Ren32,33proposed a symmetry conforming short range ordering(SC-SRO)process to explain the nature of martensite/austenite stabilisation,TWSME,as well as the‘rubber-like’pseudoelastic behaviour observed in some martensite.In this theory,the equilibrium arrangement of point defects is assumed to be identical to the underlying crystallographic symmetry of the respective phases.That is,equilibrium defects will adhere to an‘orthorhombic arrangement’if the crystal-lographic structure of martensite is orthorhombic.Since martensite and austenite possess different crystal sym-metries,the equilibrium arrangement of point defects will be different in each phase.Upon transformation from austenite to martensite,the defect arrangement of austenite is inherited,but this arrangement is not favourable in martensite,and the defects will attempt to rearrange themselves into a new equilibrium arrange-ment based on the martensite structure through short range diffusion processes.A similar rearrangement of defects will have to take place once again on the reverse transformation from martensite to austenite.

If defect structure is allowed to reach equilibrium in a

phase mimicking the symmetry of that phase,the phase would be stabilised;if,for example,this occurs in martensite,A s and A f would increase.Otsuka and Ren32,33concluded that such a phenomenon is present in all martensitic transformations and its kinetics can be predicted by considering the reduced homologous melting temperature(M s/T m),where T m is the melting temperature of the alloy.For a ratio of under0?2, kinetics(aging)is too slow for this phenomenon to produce any measurable changes,and for a ratio above 0?6it occurs almost instantaneously.For a reduced temperature between0?2and0?6,aging occurs gradually and is manifested as a time dependent change in transformation characteristics.

This type of martensitic aging is one of the prime contributors to time dependent instabilities in SMAs, particularly those that operate at high temperatures.The effect of martensitic aging on shape memory properties is re?ected in the change in transformation temperatures during prolonged exposure at high temperatures.Since it is a diffusion mechanism,the time scale at which this change in transformation temperatures occurs depends on the temperature at which the SMA is held.In the range of transformation temperatures of many HTSMAs,the reduced melting temperatures are in the range of0?2to0?6,and thus,martensitic aging occurs gradually.A gradual change in transformation tempera-tures occurring during the lifetime of an SMA makes its use dif?cult in applications requiring precise control of transformation temperatures.These issues will be discussed in greater detail in Section III.3?2.

II.2?2.c.Role of structure on transformation hysteresis Structure and lattice parameters of transforming phases also play a role in determining the energy dissipation during transformation,and thus,transformation ther-mal hysteresis,D T,and stress hysteresis,D s.Martensitic transformation between two phases can be described by a transformation stretch tensor which describes the lattice and orientation relationships between the auste-nite and martensite.For example,the transformation stretch tensor between a bcc austenite and one particular lattice correspondence variant of an orthorhombic martensite is39

U1~

b00

a z c

2

a{c

2

a{c

2

a z c

2

2

66

64

3

77

75

where b5b/a0,a5a/a0and c5c/a0(where a0is the lattice parameter of the austenite,and a,b,c are the lattice parameters of the martensite).According to Cui et al.,39

the energy dissipation,in the form of thermal hysteresis,

can be minimised by maximising the compatibility between martensite and austenite.This compatibility could be measured using the middle eigenvalue(the second largest/ smallest eigenvalue)of the transformation stretch tensor:

the closer the middle eigenvalue is to one,the better the compatibility between austenite and martensite.Figure18 shows experimental data from the Ni–Ti–X alloy systems,

and it certainly appears that D T reaches a minimum when

the middle eigenvalue approaches unity.

II.2?3.Effects of crystallographic orientation and texture:

single crystals and polycrystals

II.2?3.a.Single crystal orientation

e max

tr

of an SMA depends not only on the structures of the transforming phases,but also on the orientation relation-

ship between the crystal and the axis of the applied stress

along which the transformation strain is to be deter-mined.This is because a martensitic transformation can

be considered as a deformation mode,and activate only

along certain crystallographic directions on certain crystallographic planes,similar to deformation slip and twinning.Any directionally applied stress can be resolved

into a shear component and a normal component on the deformation plane,and these can be further decomposed

into deformation directions(according to the shear and dilatation components of the transformation).Based on

the orientation relationship between the direction of the applied stress and a particular habit plane variant, described by its unique pair of twinning plane and transformation shear direction,it is possible to calculate

a

18Variation of transformation thermal hysteresis,D T, with the middle eigenvalue of the transformation

stretch tensor in the Ni–Ti–X alloy systems.D T

appears to be minimised when the middle eigenvalue

approaches one.39(Reproduced with permission from

Macmillan Publishers and the Nature Publishing

Group)

Ma et al.High temperature shape memory alloys International Materials Reviews2010VOL55NO5267

‘resolved shear stress factor’(RSSF)for this variant.18The greater the RSSF factor for a variant,the more likely that variant will be activated/favoured.

In a single crystal,the RSSF of each habit plane variant can be found for a given applied stress state,and the s SIM and e max tr

of that single crystal is dictated by the variants most favoured by that particular stress state.18This is indeed the case,as seen in Fig.19,where s SIM of single crystals loaded along different orientations are indeed different.This effect is a consequence of a higher RSSF in those orientations with lower s SIM .Moreover,e rec is greater in orientations that have lower s SIM ,the reasons of which are described below.In addition,D s is also crystallographic orientation dependent as seen in the ?gure;however,the mechanisms responsible for such dependence will not be discussed here for the sake of brevity,more details can be found in work by Hamilton et al.40

Similarly in single crystals,it is possible to calculate the e max tr

expected along the known crystallographic direction of any uniaxially applied stress.One approach

for such calculations is the energy minimisation theory developed by Ball and James.23The calculations can be conducted with the assumptions of either full detwinning

of martensite or no detwinning,18while the real e max tr

is likely to be somewhere in between the values obtained from these two methods.These calculations can be performed for all orientations within the stereographic triangle for a cubic austenite,for instance,and plotted as transformation strain contours,if the lattice parameters and crystallographic structures of austenite and marten-site are known.18One such example is shown in Fig.20for a Ni rich NiTi SMA.41It is clear that single crystal orientations along which external stress is applied have a signi?cant effect on e max tr .In addition,the sense of loading (i.e.tension versus compression)has notable

effect on e max tr

as seen in the ?gure.This is because of the fact that martensitic transformation shear is unidirec-tional,similar to conventional deformation twinning shear observed in many metals and alloys with hcp and fcc structures.Likewise,the detwinning process in martensite variants is also unidirectional,ease of detwinning under tension may not necessarily indicate easy detwinning under compression.Thus,tension and compression loadings favour different martensite var-iants and degree of detwinning,and consequently,lead to remarkably different transformation strains.18,35,41,42II.2?3.b.Polycrystalline texture

It is reasonable to expect that orientation effects in single crystals would similarly be applicable to poly-crystals.This is indeed true,and shape memory and superelastic behaviours of polycrystalline SMAs is dependent on orientation distribution of the grains in the material,also known as texture.In polycrystalline SMAs,materials with strong texture in a particular orientation would be expected to have similar behaviour as a single crystal of that orientation.If the texture is nearly random,the behaviour of the polycrystal approximately becomes an average of the behaviour of single crystals of all orientations.Therefore,if shape memory and superelastic behaviours are known for single crystals of different orientations,then one

can

19Crystal orientation dependence of superelastic beha-viour in a series of single crystalline Ni rich Ni–Ti SMA samples under compression:s SIM ,e rec and Ds are all strongly affected by the single crystal orientation.35(Reproduced with permission from Springer Science and Business

Media)

20Maximum transformation strains as a function of the crystallographic direction of uniaxially applied stress in a Ni rich

NiTi SMA under a tension and b compression 41calculated using the energy minimisation theory.23Note the large dif-ference between the orientation with the highest e max tr (11?2%in tension),and that with the lowest e max tr (essentially 3?0%in tension).Reproduced with permission from Springer Science and Business Media

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attempt to process a polycrystalline SMA to achieve

strong texture close to the orientation exhibiting the best

sets of shape memory or superelastic properties.Texture

can be introduced by traditional cold working techni-

ques such as cold rolling,and the specimens can be cut

at certain angles to achieve the desired texture.

Nevertheless,it is dif?cult to control texture precisely

using these methods.Texture evolution under a given

stress state and strain level is dictated by the available

plastic deformation mechanisms limiting the achievable

textures using conventional processing techniques.

Recently developed severe plastic deformation techni-

ques such as equal channel angular extrusion(ECAE)

can expand the attainable texture space and provide

more precise texture control.43–45

The exact response of polycrystalline SMAs with

weak texture is,however,complicated due to the

required geometrical accommodation across the grain

boundaries.It is known that based on the Taylor

Criteria,a minimum of?ve independent deformation

systems are needed for general deformation to occur in

randomly textured polycrystals.Martensitic transforma-

tion is a deformation mode,and the number of

independent deformation systems is analogous to the

number of martensite lattice correspondence variants for

the transformation.37The number of lattice correspon-

dence variants,of course,depends on the structures of

martensite and austenite.Assuming a cubic austenite,

the transformation will not have the required?ve

martensite variants to satisfy the Taylor Criteria if

martensite has tetragonal or trigonal structure.On the

other hand,both orthorhombic and monoclinic mar-

tensites possess more than?ve variants if they are

transformed from a cubic austenite.For this reason,

completely recoverable martensitic transformation

would be extremely dif?cult to achieve in non-textured

polycrystalline SMAs with cubic austenite and tetra-

gonal or trigonal martensite.37More importantly,

repeated transformation cycles can lead to intergranular

fracture due to the strain incompatibility across grain

boundaries limiting cyclic deformation of this kind of

SMAs.37Published data on polycrystalline HTSMAs

are in accord with these observations in conventional

SMAs.Polycrystalline HTSMAs with cubic austenite to

tetragonal martensite transformation,such as Ni–Al,

Ni–Mn and Ni–Mn–Ga have shape memory behaviour

with very poor recoverability.

From experiments,large and highly reversible shape

memory and superelastic behaviours have been found in

n100m single crystal SMAs with tetragonal martensite, for example,the Co–Ni–Al/Ga SMAs discussed in

Section III.1?4.Thus it would appear that obtaining

the n100m texture in polycrystalline form of these alloys

would be the most important step for improving shape

memory behaviour.While exceptions exist,the potential

for good shape memory behaviour from tetragonal

martensite appears to be problematic in general.

II.2?4.Thermoelastic martensitic transformation Martensitic transformations can be classi?ed as thermo-elastic or non-thermoelastic.In thermoelastic transfor-mations,the interfacial boundary between martensite and austenite is very mobile,and interfacial strain between the two phases is converted into elastic lattice strain instead of being relieved through generation of defects such as dislocations.During the reverse transformation,the stored lattice strain simply causes

a reversion of austenite back to the original martensite.

On the other hand,non-thermoelastic transformation requires nucleation of austenite during the reverse transformation.Therefore,non-thermoelastic transfor-mation is in general not reversible,and SMAs are normally associated with thermoelastic martensitic transformation.However,some materials with non-thermoelastic transformations are somewhat reversible,

such as some cobalt and iron based alloys,and are also considered to be SMAs.In general,the thermoelastic nature of martensitic transformations is re?ected in the thermal hysteresis,D T.Non-thermoelastic transforma-

tions possess large D T up to several hundred degrees Celsius,while D T of thermoelastic transformations is typically less than100u C.In certain systems,the transformation can be made more thermoelastic through alloying,and is generally accompanied by a reduction in

D T and improvement in the reversibility of the shape memory behaviour.

II.3.Additional factors at high temperature

In addition to the factors mentioned in Section II.2, several others affecting SMAs uniquely in the high temperature regime are now considered.As many such factors are not exclusive to SMAs,only those directly impacting shape memory properties,such as transfor-mation temperatures and e rec,will be discussed in detail.

II.3?1.Mechanical behaviours at high temperatures

II.3?1.a.Effects of temperature on yield strength

Deformation behaviour of HTSMAs is complicated by

the availability of thermal energy at high temperatures.A common challenge for SMAs is to minimise e irr because of plastic deformation that occurs during phase transforma-

tion.This problem is exacerbated at high temperatures

due to the reduction of s y in both austenite and martensite.The impact of this reduction is practically

much more signi?cant on superelasticity than on the shape memory effect since the former requires the deformation temperature to be above A f(see Fig.16).As s SIM increases with temperature,the difference between s SIM

and s y of austenite or martensite quickly diminishes,and

slip becomes the dominant deformation mechanism.This

is one reason for the scarcity of high temperature superelastic alloys,even though many alloys are capable

of showing high temperature shape memory behaviour.

Figure16describes only the general trends in critical stresses for each deformation mechanism,but not their relative magnitudes.In materials that show shape memory or superelastic behaviour,s y is assumed to be above s DT and s SIM,respectively,but the relative magnitude of these stresses depends on deformation temperature,composition,and processing conditions of

the material,which dictate transformation temperatures. Detwinning and martensitic reorientation are diffusion-

less processes,so while s DT may decrease with increasing temperature as indicated,its dependence on temperature should be weaker than the temperature dependence of dislocation processes.In other words,as temperature increases,the decrease in s y should be much larger than

the decrease in s DT.As a result,HTSMA systems designed to operate at moderate to high temperature (above400u C)seldom show appreciable shape memory

and superelastic behaviour,and even in alloys that do,

full recovery at any applied strain level is rarely

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observed.Strategies for combating these problems were discussed in Section II.2.

II.3?1.b.Creep

The effects of creep are important for HTSMAs with very high transformation temperatures (above y 400u C).Although creep is probably not a vital deformation mechanism during the operation of HTSMA compo-nents with lower transformation temperatures,its effects are still important during high temperature processing and forming of the material.Creep manifests itself in many forms,such as diffusional creep,power law creep based on dislocation climb,and low temperature creep based on viscous effects of solute atmospheres.46A detailed description of the creep mechanisms will not be provided here,and interested readers are referred to Refs.46–48.

There are only a limited number of studies on the creep behaviour of SMAs,and most of them have focused on the binary Ni–Ti SMA.49–53A summary of the results,plotted for the creep rate as a function of applied stress and creep temperature,is shown in Fig.21.There is some incongruity in the results,as seen in the difference in the measured creep rates between the studies of Oppenheimer et al.,53Mukherjee,51Lexcellent,50and Kato 52at relatively high temperatures.Oppenheimer et al.53attributed these differences to four possible reasons:

(i)tension compression asymmetry:with the

exception of Oppenheimer et al.,who carried out experiments in compression,experiments in all other cited works were performed in tension (ii)differences in grain sizes of materials used in

each study

(iii)deviation of alloy composition from stoichio-metry

(iv)a possible change in creep mechanism over the

temperature range where the experiments were performed.

Possible effect of texture was also mentioned as an explanation for the discrepancies among different studies.Regardless,the exact reason for the discrepancies is not

yet known,and indicates the need for further detailed studies.Nevertheless,for operation of Ni–Ti SMAs below 500u C,creep does not appear to be an issue even at stress levels near 200MPa.

For creep responses of HTSMAS,however,the studies mentioned serve only as a starting point.Phase transformation often generates local stresses much higher than macroscopic stress,and it is anticipated that creep resistance will be worse when the SMA undergoes transformation cycles at high temperatures compared to static creep.If the high operating strain/stress levels are required for HTSMA components,creep would become a problem even at intermediate tempera-tures (y 500u C).It is necessary to directly test how phase transformation will affect creep resistance,as well as how creep affects shape memory properties.This type of information has not been reported in most HTSMA systems.

II.3?2.Microstructural instability

Transforming phases in many SMAs are non-equili-brium phases.Given suf?cient aging time at a high temperature,the equilibrium phase will often form.If the temperature at which precipitation of stable phases takes place is far above the operating temperature range of the alloy,controlled precipitation can be used to improve the properties of many SMAs.For example,strong precipitates are often used to improve the shape memory properties.In some brittle SMAs,ductile equilibrium phases are often created from aging in order to improve ductility of the alloy.

Unfortunately for HTSMAs,the formation of equili-brium phases may not be quite controllable.If the operating temperature range of the HTSMA is high enough for precipitation to occur at a suf?cient rate,then precipitation will cause a continuous change in the shape memory properties of the alloy due to composi-tion change of the matrix.These changes,such as in transformation temperatures,will eventually reach a degree such that the HTSMA component will no longer perform the designed function,i.e.the HTSMA effectively has a lifetime controlled by precipitation.For example,in nickel rich Ni–Ti binary SMAs,Ti 3Ni 4precipitation occurs during aging at temperatures above 300u C.These alloys are usually heat treated to a ‘peak aged’condition such that nanosized coherent precipi-tates provide the best hardening and thus,stability in the transformation behaviour of the alloy.However,if a similar Ni–Ti SMA is to be used at temperatures near 300u C,the precipitation process will continue during the operation of the alloy,the precipitates will grow larger in size and become incoherent,and the precipitation hardening effect will be diminished.The only effect remaining will be the compositional change of the matrix.This is a particular problem for certain Cu based and Ni based HTSMAs,which will be discussed in Sections III.1?3and III.1?6respectively.

Many HTSMA systems depend on work hardening or training to improve and stabilise their shape memory behaviour,but at a suf?ciently high temperature,dislocation recovery will become signi?cant enough that the effect of work hardening and training will be gradually lost over some period of time.Recovery is a thermal phenomenon,so the critical temperature at which the impact of recovery on the shape memory behaviour becomes unacceptable depends on

the

21Creep data for binary Ni–Ti;reproduced with permis-sion.‘Current Study’refers to the results of Oppenheimer et al.53Reproduced with permission from Elsevier

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expected time the SMA component will be exposed to the temperature of interest during its lifetime.The rate at which it occurs at a particular temperature depends on the melting point and microstructure of the alloy. Similarly,recrystallisation also has a negative effect on most SMAs.The temperature at which recrystallisation becomes a serious problem is usually high,but its onset renders strengthening by grain re?nement impossible. These two temperatures become an upper limit for alloys strengthened by conventional methods.

Finally,water quenching required after homogenisa-tion or solution heat treatments of many HTSMA systems cause some disordering of crystal structures and high levels of non-equilibrium vacancy concentrations. If self-diffusion is fast enough,the reordering process can drastically change the transformation temperatures, such as in the case of Cu based HTSMAs during repeated martensitic transformation as discussed in Section III.1?3.

II.3?3.Oxidation

Alloys exposed to the atmosphere at high temperatures will inevitably react with oxygen and other gaseous species.As a result,metal oxide products can form at the sample surface.If the reaction is not stopped and suf?cient amount of oxygen is available,which is usually the case in an uncontrolled atmosphere,then the alloy will continue to lose more of its mass to oxide products. Fortunately,oxide formation at the surface slows down the diffusion of metal ions and oxygen and controls the rate of oxidation.Therefore,in addition to temperature, oxidation rate is determined by the permeability of the oxide layer.Alloys designed to operate at high temperatures,such as the nickel based superalloys,are capable of forming a stable and impermeable oxide layer to prevent further oxidation once a critical oxide layer thickness is formed.

An oxidation problem unique to SMAs is that unlike high temperature structural metals,even a slight change in the matrix composition may become unacceptable for SMAs.Transformation temperatures in SMAs such as Ti–Ni and Ni–Al are very sensitive to composition,and often change more than100u C with each atomic percentage variation in composition.54–56Frequently,the oxide layer of an alloy is composed of mostly the oxide of a single element,thus reducing its concentration in the matrix.The oxidation effect of SMAs is most often studied on Ti–Ni alloys,and it was found that the sample surface consists of mostly titanium oxide,which results in a nickel rich layer immediately below the surface signi?cantly reducing local transformation temperatures.57–62Thickness of the oxide layer from aging increases sharply when oxidation temperature exceeds700–800u C,57–61,63such that the oxide layer formed from1h oxidation at500u C is about275nm thick,whereas the thickness of the layer formed at1h oxidation at700u C exceeds70m m.58

Further details on oxidation of Ti–Ni and other conventional SMAs are beyond the scope of this article; however,it is interesting to point out the reported effects of oxidation on their shape memory properties.Kim et al.60observed a decrease in all transformation temperatures in Ti49?6Ni50?6sheets with thickness of 0?8mm after oxidisation in air at900u C for1h when compared to the transformation temperatures of non-oxidised specimens.The decrease in M s was only y5u C, but A f and M f decreased by10–15u C,and A s by y25u C.The authors attributed this effect to a nickel rich layer

that forms immediately beneath the titanium oxide surface layer.On the other hand,Nam et al.58found an increase in both M s and A f as dry air oxidation temperature is increased from450to850u C while duration is held constant at10min in Ti51Ni49wire

with diameter of1?7mm.This increase is attributed

by the authors to the compressive stress exerted by the

oxide surface on the matrix beneath.This compressive stress is believed to be hydrostatic,so that it acts in

the normal direction to the habit planes of the martensitic transformation.Theoretically,a hydrostatic stress acts normal to the habit plane of martensites,

and will assist the transformation(increase transfor-mation temperatures)if the volume change from the austenite R martensite transformation carries the same

sign as the hydrostatic stress applied,where compressive stress is negative.64This is indeed the case for the Ti–Ni alloys.64,65Nam et al.suggest that the compressive stress is increased with an increase in the oxidation temperature, mostly likely because of the thickening of the oxide layer

and the change in its structure as oxidation temperature is increased.However,large hydrostatic pressure is needed

to signi?cantly impact transformation temperatures since

the volume change of martensitic transformation asso-ciated with thermoelastic SMAs is small.The stress at the matrix/oxide interface is not known,and it is not clear why compressive stress at the interface would be expected to increase when oxide layer thickens.More experimental

data is needed to determine whether the compressive stress

alone suf?ciently explains the increase in transformation temperature during oxidation.

While the results of the two reports appear to be

con?icting,this is not necessarily the case.The speci-

mens used in the study by Kim et al.60are already nickel

rich,so any additional decrease in the titanium concentration immediately causes immediate reduction

in transformation temperatures.54The specimens in the

study by Nam et al.58are titanium rich.In the titanium

rich region,transformation temperatures are not sensi-

tive to small changes in composition.As oxidation temperature is increased,the reduction of titanium content in the matrix right below the oxide layer becomes larger as titanium oxide is formed at the surface.However,the titanium rich composition of the

alloy means that a certain amount of titanium must be depleted?rst before the transformation temperatures

could decrease due to the compositional change.In the meantime,other mechanisms,such as possibly the compressive stress from the oxide layer,can cause an increase in transformation temperatures instead. Further evidence lies in the specimens oxidised at

1000u C for10min in the study by Nam et al.58Even though the transformation temperatures from DSC in

this condition were not reported directly,the composi-

tion of the matrix was changed to Ti48?2Ni51?8from the

initial composition of Ti51Ni49.It was also shown that

M s of the specimen oxidised at1000u C was lower as compared to that aged at850u C in a set of constant stress heating cooling curves reproduced in Fig.22.

The study by Nam et al.58also addressed the effects of oxidation on the shape memory behaviour of Ti51Ni49

alloy from thermal cycling experiments under constant tensile stress of60MPa.As oxidation temperature increased,both e rec and e irr remain constant up to an

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oxidation temperature of 550u C.However,if oxidation temperature is increased further,both e rec and e irr are reduced as shown in Fig.22.The authors did not provide an explanation for these observations.However,the reduction of e rec at high oxidation temperature should not necessarily be related to a deterioration of the shape memory effect.Reduction of both e rec and e irr at a constant stress level is often observed in a work hardened or precipitation strengthened material,because while the strengthening mechanism suppresses plastic deformation,it also poses a barrier to martensitic transformation.Therefore,a greater applied stress is needed to achieve the same level of e rec .Thus it is possible that if applied stress is increased beyond 60MPa,e rec would increase in speci-mens oxidised at higher temperatures.Also,since the oxide layer does not transform,the compressive stress imposed by it may act against the applied tensile stress.Therefore,the actual stress experienced by the matrix near the oxide matrix interface is smaller than the applied stress,resulting in a more self-accommodated martensite structure in this area and reducing e rec .

From these results,it would appear that the rate of oxidation is too low to be considered as an issue for Ni–Ti–X based HTSMAs designed to operate at lower temperatures (about 200–300u C),however,it becomes a more serious problem at higher temperature.Finally,the unique problem to SMAs is the transformation tem-perature sensitivity to compositional changes due to oxidation,which limits both the maximum temperature and exposure time of the SMA components.Shape change associated with the transformation of the matrix may also cause cracking in the brittle oxide layer,making the development of a stable oxide layer dif?cult even in alloys normally with good oxidation resistance.

III.Potential HTSMA systems

III.1.Alloys with characteristic transformation temperatures between 100and 400u C

Following the vast amount of research performed on SMAs operating near room temperature,most work

performed on HTSMAs has been within the 100–400u C temperature range for a number of reasons including:(i)the availability of numerous alloy systems with

transformation temperatures in this range

(ii)the similarity in processing of these materials to

the thermomechanical processing of current commercial alloys

(iii)the temperature range being comparatively low

and thus,the relative ease in experimentation.

At this temperature range,it may be expected that thermally activated processes would not have a sig-ni?cant effect on shape memory behaviour,but this has not necessarily been the case.Even so,in the short term,the best chance for developing a new family of commercial HTSMAs is expected to originate from the alloy systems described in this section.

III.1?1.Ti–Ni–(Pd,Pt)system

Interest in the Ti–Ni–Pd and Ti–Ni–Pt systems as potential HTSMAs was derived from three sets of studies:the comprehensive study of phase transforma-tions in binary B2titanium alloys,66the discovery of high transformation temperatures in the Ti–Pd and Ti–Pt binary systems by Donkersloot,67and the discovery of ternary alloying effects on the transformation temperatures of binary Ni–Ti SMAs by Eckelmeyer.68Based on the results from these studies,palladium and platinum were added to the Ti–Ni system in order to increase transformation temperatures,and nickel to the Ti–Pd and Ti–Pt systems to improve shape memory behaviour.

III.1?1.a.Ti–Ni–Pd alloys

Ti–Ni–Pd HTSMAs have received the most rigorous attention over the years.Initial focus was centred on improving their high temperature shape memory beha-viour,but more recently,the focus has shifted towards improving their work output,as well as dimensional and microstructural stability.In this system,transformation temperatures can be altered by replacing nickel with palladium.If the concentration of titanium is held constant at nearly 50at-%,the relationship between the transformation temperatures and relative concentration of nickel and palladium is parabolic,as shown in Fig.23.A minimum in transformation temperatures occurs at approximately 10at-%Pd,although the exact composition of this minimum is still subjected to debate.66,69,70In compositions with palladium concen-trations greater than the palladium concentration at this minimum,replacing nickel with palladium increases transformation temperatures by approximately 15u C/at-%.69,71–73On the other hand,if the concentration of palladium is lower than the composition at the minimum,replacing nickel with palladium actually lowers the transformation temperatures by 4u C/at-%.69This parabolic dependence of the transformation tem-peratures on composition stems from the change in the structure of martensite.On the higher palladium concentration side of the minimum,B2austenite trans-forms to B19orthorhombic martensite,and at the lower palladium concentration side,it transforms to B199monoclinic martensite or R phase.The composition at which the transformation temperatures are at a mini-mum corresponds to the point of the structure transi-tion.Because of the complete mutual miscibility of the Ti–Ni and Ti–Pd systems,the relationships

between

22Tensile thermal cycling curves under 60MPa for

Ti 51Ni 49SMA oxidised for 10min at various tempera-tures.At temperatures above 823K (550u C),both e rec and e irr are reduced.Reproduced with permission.58Reproduced with permission from Springer Science and Business Media

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transformation temperatures and composition hold over all ranges of palladium concentration.This enables the access to a continuous range of transformation tem-peratures from room temperature to over 500u C by adjusting the amount of palladium in the alloy.

Otsuka et al.74reported poor shape recovery of binary Ti 50Pd 50HTSMA,concluding that this was primarily caused by the low s y of both austenite and martensite.They proposed three possible solutions to improve its shape memory behaviour by strengthening the material through:

(i)solid solution hardening via ternary alloying (ii)thermomechanical processing

(iii)precipitation hardening or some combination of

these approaches.

As the following discussion demonstrates,these solu-tions have met with varying degrees of success.

While the main effect of nickel addition to Ti–Pd alloys is the reduction in transformation temperatures,it also has an indirect effect of improving shape memory behaviour by lowering the temperature range at which the alloy would be used.At a lower temperature,s y is at a higher level.Khachin 66observed full recovery of 4%applied strain in high temperature torsion experiments on Ti 50Ni 13Pd 37.However,Lindquist and Wayman 69studied the same alloy at room temperature under tension,and were only able to obtain 40%recovery of 6%applied strain.The reason for this discrepancy was not resolved,but it is likely due to the differences in the way the materials were processed and tested,which can have a remarkable effect on recovery rate in Ti–Ni–Pd alloys as shown in Fig.24.

Although the s y of Ti–Pd alloys can be raised indirectly through nickel addition,it is still too low for perfect shape recovery.Several researchers addressed the issue of low s y in Ti–Ni–Pd HTSMAs by further solid solution strength-ening.Suzuki et al.73and Yang et al.77added small amounts of boron to Ti 50Pd 30Ni 20and Ti 50?7Ni 22?3Pd 27alloys,respectively,but neither caused a notable reduc-tion in e irr with boron additions of up to 0?2at-%.Micro-metre sized Ti 2B particles were found along the grain boundaries,but they were too large and non-uniformly

distributed to possibly function as particles for precipitate hardening.73Instead,boron acted as a grain re?ner by reducing the grain growth rate in these alloys.73In identically solution treated or hot rolled specimens,0?12at-%boron reduced the grain size from y 40m m down to 10m m in the Ti 50?7Ni 22?3Pd 27HTSMA.77Presumably for the same reason,0?2at-%boron addition to Ti 50Ni 30Pd 20doubled the room temperature tensile elongation to failure from 8to 16%strain,73and increased the ultimate tensile strength from 460to 800MPa for a sample deformed in martensite at a temperature of 170u C.78It is not clear why the grain re?nement effect of boron did not improve the shape memory behaviour,since s y increases with grain size re?nement.In other studies on solid solution hardening,5at-%gold or platinum replacing palladium in Ti 50?5Ni 19?5Pd 30increased s y and mildly enhanced cyclic stability,but had little effect on total recoverable strain.78From these results,73,78it appears that increasing s y alone will not necessarily always result in improvements in shape memory behaviour.Another recent study similarly observed little impact on the shape memory recovery rate after alloying Ti 50?6Ni 19?4Pd 30with 1wt-%cerium,79no reason for this observation was given.

Although boron and cesium do not appear to improve the shape memory behaviour,a recent study by Atli et al.80showed that 0?5at-%scandium addition to Ti 50?5Ni 24?5Pd 25,replacing titanium,was more effective in this aspect.Although the scandium lowered transfor-mation temperatures by about 6–10u C,it was able to reduce e irr from constant stress thermal cycling experi-ments under tension by half at stress levels above 200MPa without adversely affecting e rec .This improve-ment was believed to be caused by the solution hardening effect of scandium,but the effect of scandium on the s y of martensite and austenite was not explicitly shown.Additionally,scandium addition also reduced D T and improved cyclic stability.After 10thermal cycles at 200MPa,the cumulative e irr was reduced by y 20%in the scandium containing specimen.

Golberg et al.71,75,81investigated the effects of thermomechanical treatments on the shape

memory

24Effect of processing on the recovery rate of Ti–Ni–Pd

alloys:all specimens were deformed in tension slightly below M f :170u C,75and 200u C.76Figure repro-duced with permission from The Taylor &Francis Group

7

23Martensite start temperature,M s ,as a function of the

palladium content in quasi-equiatomic Ti–Ni 502X Pd X alloys.Transformation temperatures increase linearly with the Pd content for alloys containing greater than approximately 10–15at-%.The trendline is a ?t through all data points

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behaviour of Ti–Ni–Pd HTSMAs.Cold rolling up to 24?5%reduction in thickness followed by annealing at 400u C for 1h proved effective in increasing s y levels of martensite in Ti 50Ni 20Pd 30at 170u C from y 200to 400MPa,such that up to 5?3%applied strain was fully recovered by heating above A f (y 250u C)after tensile deformation at 170u C.This was a signi?cant improve-ment over the solution treated alloy where only 2?5%strain could be fully recovered under the same testing conditions.Cold rolling and subsequent annealing,however,decreased the transformation temperatures by about 20–30u C as compared to that of the solution treated specimens,as expected from the observations on binary Ni–Ti SMAs.82On the other hand,while superelasticity exists in the cold rolled and annealed alloy tested above A f ,the strain recovery after unloading was incomplete.The authors concluded that the strength of the parent phase was still too low as compared to the high s SIM levels,and e irr was caused by plasticity.71

Investigating Ti x Pd 30Ni 702x (x 548?5to 51?0at-%)alloys encompassing both sides of the equiatomic composition,Shimizu et al.76realised (Fig.25)that with decreasing titanium content,transformation tempera-tures decrease only slightly in Ti rich compositions,but drops off dramatically on the Ni/Pd rich side,such that M s declines to room temperature in a Ti 48?5Pd 30Ni 21?5alloy.This is similar to the composition dependence of transformation temperatures in binary Ni–Ti SMAs,and is rationalised by the solubility of excess titanium or nickel atoms near the equiatomic composition.Solubility for extra titanium atoms in near equiatomic Ni–Ti is almost negligible,and although the solubility for extra nickel atoms is also small,it is possible to accommodate excess Ni concentrations under 1at-%in solution.Therefore,in titanium rich compositions,the extra titanium atoms tend to immediately form second phases,and do not affect the composition of the matrix and thus the transformation temperatures of the alloy.However,a small concentration of extra nickel atoms can dissolve in the matrix,changing its composition and transformation temperatures.Utilising precipitates in Ti rich compositions after proper heat treatments,Shimizu

et al.76(Fig.24)were able to demonstrate that

Ti 50?6Pd 30Ni 19?4

outperformed the equiatomic Ti 50Pd 30Ni 20alloys with a recovery rate of 90%versus 78%respectively,on samples deformed to 6%applied strain at 200u C under tension.The improvements were attributed to the homogeneous distribution of ?ne Ti 2(Ni,Pd)type precipitates formed during the annealing process.

Shirakawa et al .85and Nagasako et al.86studied the phase transformations in the (Ni,Pd)rich Ti–Ni–Pd alloys,and compensated for the decrease in transforma-tion temperature from the nickel rich matrix by increasing the palladium content.While the two way shape memory behaviour of a Ti 48Pd 31Ni 21alloy annealed at 400u C was qualitatively described,no quantitative data on either one way or two way shape memory behaviour were available.85Furthermore,it should also be possible to maintain reasonably high transformation temperatures in nickel rich compositions without increasing Pd content by applying appropriate heat treatments.The reported precipitates are nickel rich in nickel rich compositions,and were identi?ed to be Ti 2(Ni,Pd)3and Ti 3(Ni,Pd)4types,86so as these pre-cipitates grow,nickel content in the matrix should be reduced,thereby increasing transformation tempera-tures.Even though no published studies have explicitly demonstrated this in Ti–Ni–Pd alloys,this technique has been used in Ni–Ti–Hf (see Section III.1?2)alloys with some success.87

The combined effect of thermomechanical treatment and precipitate hardening through Ti 2(Ni,Pd)particles was investigated by Tian and Wu.83,88,89Tian et al.88claim that deformation in tension at room temperature followed by heating above A f produced full recovery of up to 7?2%strain and up to 95%recovery at 11%applied strain in a 9%cold rolled Ti 50?6Pd 30Ni 19?4alloy annealed for 1h at 400u C.However,these experiments were conducted without an extensometer,and from the published ?gures,the elastic modulus of martensite and austenite were calculated to be 9and 15GPa respectively.This indicates that the strain measurements are not quantitatively reliable.Nevertheless,there is useful information to be taken away from the work:when the specimen is deformed at 270u C (above A f ),it does not initially show superelasticity,but exhibits shape memory behaviour instead with a recovery rate of y 90%after reheating to y 400u C.If this deform heat process is repeated about ?ve times,a fully recoverable superelastic behaviour appears which is likely caused by the strengthening effect of training.88Owing to the uncertainty in the strain measurement,this superelastic behaviour cannot be quanti?ed.

Martensite reorientation,and thus the shape memory behaviour was absent in 21–29%cold rolled Ti 50Pd 30Ni 20alloy during subsequent deformation at 173u C when post-rolling annealing temperatures were below A s .81Moreover,shape memory behaviour dete-riorated if the post-rolling annealing temperatures were very high.Golberg et al.81concluded that annealing at a temperature below A s does not allow the preferentially oriented martensite inherited from the rolling process to reset to a self-accommodated martensitic morphology.Conversely,annealing far above A f initiated microstruc-tural recovery,destroying the work hardening effects of cold rolling.An ideal post-rolling annealing

temperature

25Dependence of the martensite start temperature,M s ,

on the Ti/(Ni,Pd)ratio.M s remains fairly constant on the Ti rich side of stoichiometry but drops off precipi-tously on the (Ni,Pd)rich side.The trendline is a ?t through all data points

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was proposed to be above A s but below the recrystallisa-tion temperature.81A follow-up study by Xu et al.72 examined the recovery and recrystallisation processes over a full range of cold rolled and annealed Ti50(Ni,Pd)50alloys with palladium contents ranging from0to50at-%.They observed that recovery started at450u C in Ti50Pd30Ni20alloy and recrystallisation began at550u C in Ti50Pd40Ni10alloy.Physically,these diffusion driven mechanisms reduce e rec by increasing grain size and reducing dislocation density from cold working.90If recrystallisation takes place in martensite, reverse transformation temperatures(A s and A f)91 increase due to the loss of internal twins in the martensite.This occurs because during phase transfor-mation,elastic energy is stored in the internal twins of martensite as elastic strain.Recrystallisation replaces these martensites with the‘new’strain free martensite, and the stored elastic energy is lost.Since stored energy is a driving force for the reverse transformation,its loss results in the increase in A s and A f.In particular, recovery and recrystallisation temperatures in the range

of450–600u C prevent effective thermomechanical treat-ment of alloys with high Pd contents and high transformation temperatures since recovery would occur within the operational temperature range of the alloy. Recovery should also cause thermomechanical training and TWSMEs in these HTSMAs to be almost impos-sible.This factor has severely limited the development of higher temperature Ti–Ni–Pd HTSMAs with Pd con-tents beyond about35at-%.

Similarly,thermal cycling can also affect the stability and properties of SMAs either by introducing a buildup of internal stresses or through their relaxation.For example,thermal cycling of Ti50?6Pd30Ni19?4under 200MPa increased transformation temperatures by up to16u C,92and stabilised after about30cycles.In contrast,thermal cycling under no stress reduced transformation temperatures by5–10u C,with alloys of higher Pd content exhibiting a greater shift in transfor-mation temperatures during cycling.91It was suggested that dislocations introduced during cycling suppressed the martensitic transformation during stress free thermal cycling,but the oriented internal stress?elds generated from isobaric thermal cycling assisted the external applied stress,increasing transformation temperatures.92 Another mechanism of thermal and time dependent instability in the Ti–Ni–Pd HTSMAs and many other HTSMAs in general is the martensite aging behaviour, whereby transformation temperatures increase with increasing aging time in the martensitic state.In alloys with less than30at-%Pd,this effect is very small; however at higher Pd concentrations,a two stage increase in transformation temperatures with increasing aging time was evident in cold worked and annealed samples.91Initially,a rapid increase in transformation temperatures with up to20h aging at240u C occurred, and was believed to be a consequence of the SC-SRO phenomenon of point defects,33while a more sluggish recovery process was responsible for the second stage. This behaviour is absent at lower Pd compositions because of lower aging temperatures and lower M s/T m ratio of near0?2in these alloys.This ratio was shown by Ren and Otsuka to be too low for martensitic aging by the SC-SRO mechanism to proceed.3In contrast,the M s/T m ratio is higher(0?3–0?4)in alloys with higher Pd concentrations,and allows for martensite aging to dominate the initial stages where transformation tem-peratures can rapidly increase.

While a signi?cant body of data exists concerning transformation temperatures and shape memory responses of Ti–Ni–Pd HTSMAs,the major physical property discussed up to this point has been the stress

free shape memory response seen in Fig.10.In reality,

few applications actually would use the alloy in that fashion.Instead,HTSMAs offer more potential as solid

state replacements for conventional actuation systems

such as solenoids,motors,or pneumatic and hydraulic systems and for use as the major motive force in adaptive and recon?gurable components for aircraft and

other related applications.Therefore,the main interest is

the shape memory response of the alloy,not under stress

free conditions,but acting against some kind of biasing force.Consequently,the most recent research on Ti–Ni–

Pd alloys has focused on the load biased or constant stress thermal cycling response.

The constant stress thermal cycling response of a typical SMA was shown Figs.11and12.The properties measured from this type of tests that are critical for applications are e irr and e rec at a given stress level. Another useful property is the work output(Fig.13b), which is the product of e rec and applied stress.It is desirable in actuator applications for HTSMA compo-

nents to be dimensionally stable(small e irr)while possessing a large‘stroke’(large e rec)under as high a stress level as possible.

Ti–Ni–Pd alloys containing less than y35at-%Pd

have acceptable work characteristics but tend to suffer

from rather high e irr during thermal cycling under load.

This irreversibility worsens with increasing applied stress78and increasing Pd concentration at constant stress.93Therefore,a signi?cant focus of the recent research on Ti–Ni–Pd alloys has been in pursuit of mechanisms that reduce e irr.In particular,Bigelow

et al.94have found that alloying with Pt or Au reduced

e irr considerably in the absence o

f thermomechanical trainin

g as summarised in Fig.26,but no signi?cant reduction of e irr from these alloying elements occurred in trained specimens.Thermomechanical training alone

also reduced e irr signi?cantly as seen in Fig.26.Not

only

26Effect of thermomechanical processing and alloying on e irr of Ti–Ni–Pd–X alloys94(Reproduced with per-

mission from the SPIE)

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are these approaches useful in reducing e irr in Ti–Ni–Pd alloys,they also have no detrimental effect on recover-able strain and thus relative work output for these materials.

The amount e irr for a given alloy and stress level increases as the maximum temperature to which the sample is thermally cycled increases,with all other considerations being constant,94thus placing another criterion on the upper temperature limit for these materials.However,certain alloying additions such as 5at-%Pt increases the maximum temperature capability of Ti–Ni–Pd alloys by y 30u C,94thus providing some measure of protection against overheating.95

Similar to conventional Ni–Ti SMAs,96,97the work output for Ti–Ni–Pd alloys peaks at some optimum stress level 84as shown in Fig.27a .This peak in work output is a result of competing factors.As the stress level increases,more martensite variants become favourably aligned and detwinned,forming more single variant morphology and accommodating larger strains.At the same time,the increasing stress level approaches s y of the alloy and causes irrecoverable deformation to occur instead of martensitic transformation,and negatively affects the recoverability of martensite.Consequently,these factors cause e rec to reach a maximum and then rapidly decrease with increasing stress as shown in Fig.27b for the Ti 50?5Ni 19?5Pd 30alloy.Therefore,even though the applied stress continues to increase,e rec diminishes and its product with applies stress,which is work output,reaches a maximum at a particular stress level as well.

The maximum work output for various Ti–Ni–Pd alloys as a function of the transformation temperatures range,such that the full range over which the transforma-tion occurs (M f to A f )for each alloy is clearly indicated,98is shown in https://www.wendangku.net/doc/504773849.html,rge and consistent work output of between 8and 11J cm 23is achieved for alloys with transformation temperatures between 100and 300u C.However,for alloys with transformation temperatures above this range,the work output drops off almost catastrophically so that alloys with transformation temperatures in the neighbourhood of 500u C are capable of essentially zero work output,placing another limit on the viability of HTSMAs beyond just the need for a high transformation temperature.This is a natural

consequence of the reduction of s y at higher temperatures and the onset of creep deformation.99

Most of the behaviours just described for the Ti–Ni–Pd alloys,such as generally poor recoverable strain capability and the sharp decline in work output beyond y 300u C,can be explained by comparing the deforma-tion characteristics of the individual martensite and austenite phases.To demonstrate this point,Bigelow et al.93determined s y of austenite,stress levels for martensite reorientation/detwinning,and s y of marten-site for a series of Ti–Ni–Pd alloys containing 15to 46at-%Pd which are shown in Fig.29.For each composition,one sample was loaded at a temperature of M f 250u C and then thermally recovered under stress-free conditions.This experiment resulted in a value for s DT of martensite,and s y of the martensite.Another sample was then loaded at A f z 50u C to determine s y of austenite.At around 37at-%Pd,the reorientation stress and s y of austenite cross.At lower Pd contents than this crossover composition,work is possible because an applied stress results predominantly in

martensite

27a speci?c work output for a Ti 50?5Ni 19?5Pd 30alloy as a function of applied stress loaded in both tension and compres-sion and b the corresponding transformation strain versus applied stress.84(Reproduced with permission from the

SPIE)

28Work output for a series of Ti–Ni–Pd and Ti–Ni–Pt

alloys as function of the transformation temperature range (M f to A f ).98Alloys containing ternary additions producing transformation temperatures greater than about 400u C show almost no usable work output.(Reproduced with permission from the SPIE)

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高温合金概述

1.1 高温合金 1.1.1 高温合金及其发展概况 高温合金是指以铁、钴、镍为基体，能在600℃以上温度，一定应力条件下适应不同环境短时或长时使用的金属材料。具有较高的高温强度、塑性，良好的抗氧化、抗热腐蚀性能，良好的热疲劳性能，断裂韧性，良好的组织稳定性和使用可靠性。高温合金为单一奥氏体组织，在各种温度下具有良好的组织稳定性和使用的可靠性，基于上述性能特点，且高温合金的合金化程度很高，故在英美称之为超合金（Superalloy）。 高温合金于20世纪40年代问世，最初就是为满足喷气发动机对材料的耐高温和高强度要求而研制的，高温合金的发展与航空发动机的进步密切相关，1939年英国Mond镍公司首先研究出Nimonic75，随后又研究出Nimonic80合金，并在1942年成功用作涡轮气发动机的叶片材料，此后该公司又在合金中加入硼、锆、钴、钼等合金元素，相继开发成功Nimonic80A、Nimonic90等合金，形成Nimonic合金系列。如今先进航空发动机中高温合金用量已超过50%。此外，在航天、核工程、能源动力、交通运输、石油化工、冶金等领域得到广泛的应用。高温合金在满足不同使用条件中得到发展，形成各种系列的合金，除传统的高温合金外，还开发出一批高温耐磨、高温耐蚀的合金。 高温合金是航空发动机、火箭发动机、燃气轮机等高温热端部件的不可代替的材料，由于其用途的重要性，对材料的质量控制与检测非常严格。高温合金的基本用途仍旧是飞行器的燃气轮发动机的高温部分，它要占先进的发动机重量的50％以上。然而，这些材料在高温下极好的性能已使其用途远远超出了这一行业。除了航空部件之外，规定将这些合金用于舰船、工业、陆地发电站以及汽车用途的涡轮发动机上。具体的发动机部件包括涡轮盘、叶片、压缩机轮、轴、燃烧室、后燃烧部件以及发动机螺栓。除了燃气发动机行业之外，高温合金还被选择用于火箭发动机、宇宙、石油化工、能源生产、内燃烧发动机、金属成形(热加工工模具)、热处理设备、核电反应堆和煤转换装置。

钛合金热处理

第十三章有色金属及合金 内容提要： 有色金属的产量和用量不如黑色金属多，但由于其具有许多优良的特性，如特殊的电、磁、热性能，耐蚀性能及高的比强度(强度与密度之比)等，已成为现代工业中不可缺少的金属材料。 1．铝及铝合金； 2．钛及钛合金； 3．铜及铜合金； 4．轴承合金。 基本要求： 掌握和了解各种有色金属的牌号、成分、性能和用途。 13.1铝及铝合金 13.1.1铅及铝合金的性能特点及分类编号 纯铝：纯铝具有银白色金属光泽，密度小(2.72 )，熔点低(660.4℃), 导电、导热性能优良。 耐大气腐蚀，易于加工成形。 具有面心立方晶格，无同素异构转变，无磁性。 1 铝合金及其特点 铝合金常加入的元素主要有Cu、Mn、Si、Mg、Zn等，此外还有Cr、Ni、Ti、Zr 等辅加元素。 ①比强度高（>>高强钢）。可用于轻结构件，尤其航空。 ②突出理化性能。导电、抗大气腐蚀。 ③良好加工性。高塑性、易冷成形；某些合金铸造性能好，宜作压铸件。 2 铝合金分类及分类编号 13.1.2铝合金的强化 1 形变强化 2沉淀强化 3 固溶强化和时效强化： 13.1.3变形铝合金 变形铝及铝合金牌号表示方法：根据国标规定，变形铝及铝合金可直接引用国际四位数字体系牌号或采用国标规定的四位字符牌号。GB 3190-82中的旧牌号仍可继续使用，表示方法为: ?防锈铝合金：LF+序号 ?硬铝合金： LY +序号 ?超硬铝合金：LC +序号 ?锻铝合金： LD +序号 常用变形铝合金 1 防锈铝合金：主要是Al-Mn和Al-Mg系合金。 Mn和Mg主要作用是提高抗蚀能力和塑性，并起固溶强化作用。 防锈铝合金锻造退火后组织为单相固溶体，抗蚀性、焊接性能好，易于变形加工，但切削性能差。不能进行热处理强化，常利用加工硬化提高其强度。常用的Al-Mn系合金有 LF21 （ 3A21 ），其抗蚀性和强度高于纯铝，用于制造油罐、油箱、管道、铆钉等需要弯曲、冲压加工的零件。常用的Al-Mg系合金有 LF5（ 5A05 ），其密度比纯铝小，强度比Al-Mn合金高，在航空工业中得到广泛应用，如制造管道、容器、铆钉及承受中等载荷的零件。

铜合金

牌号：白铜C7521prefix = o ns = "urn:schemas-microsoft-com:office:office" 标准：日本 C7521白铜： 以镍为主要添加元素的铜合金。纯铜加镍能显著提高强度、耐蚀性、硬度、电阻和热电性，因此白铜较其他铜合金的机械性能、物理性能都异常良好，延展性好、硬度高、色泽美观、耐腐蚀、富有深冲性能，被广泛使用于造船、石油化工、电器、仪表、医疗器械、日用品、工艺品等领域，并还是重要的电阻及热电偶合金。 C7521白铜分类： 普通白铜是铜和镍的合金﹔ 复杂白铜：加有锰、铁、锌、铝等元素的白铜合金称复杂白铜（即三元以上的白铜），包括铁白铜、锰白铜、锌白铜和铝白铜等。 ①铁白铜：铁白铜中铁的加入量不超过2％以防腐蚀开裂，其特点是强度高，抗腐蚀特别是抗流动海水腐蚀的能力可明显提高。 ②锰白铜：锰白铜具有低的电阻温度系数，可在较宽的温度范围内使用，耐腐蚀性好，还具有良好的加工性。 ③锌白铜：锌白铜具有优良的综合机械性能，耐腐蚀性优异、冷热加工成型性好，易切削，可制成线材、棒材和板材，用于制造仪器、仪表、医疗器械、日用品和通讯等领域的精密零件。 ④铝白铜：是以铜镍合金为基加入铝形成的合金。主要用于造船、电力、化工等工业部门中各种高强耐蚀件。 C7521白铜性能： 白铜是以镍为主要添加元素的铜基合金，呈银白色，有金属光泽，故名白铜。铜镍之间彼此可无限固溶形成连续固溶体，即不论彼此的比例多少，而恒为α--单相合金。当把镍熔入红铜里，含量超过16％以上时，产生的合金色泽就变得洁白如银，镍含量越高，颜色越白。白铜中镍的含量一般为25%。 C7521白铜应用： 产品广泛用于电器、电子、电力、汽车、通讯、五金等行业，如变压器铜带、引线框架材料带、射频电缆带、太阳能光伏铜带、高炉用铜冷却壁板、含银无氧铜板、电子接插件铜带、模具电极铜板、乐器铜板等。 C7521白铜化学成分： 牌号主要成份其他成份 日本Cu Ni Zn Fe Al Pb Mn C752164.5-66.516.5-19.5余量———— C7521白铜力学性能：

水平井钻井技术经验概述

第一章定向井（水平井）钻井技术概述 第一节定向井、水平井的基本概念 1．定向井丛式井发展简史 定向井钻井被（英）T．A．英格利期定义为：“使井筒按特定方向偏斜，钻遇地下预定目标的一门科学和艺术。”我国学者则定义为，定向井是按照预先设计的井斜角、方位角和井眼轴线形状进行钻进的井。定向井相对与直井而言它具有井斜方位角度而直井是井斜角为零的井，虽然实际所钻的直井它都有一定斜度但它仍然 石油管理局的河50丛式井组，该丛式井组长384米，宽115米，该丛式井平台共有钻定向井42口。 2．定向井的分类 按定向井的用途分类可以分为以下几种类型： 普通定向井 多目标定向井 定向井丛式定向井 救援定向井 水平井 多分枝井（多底井） 国外定向井发展简况

（表一）

10．井眼尺寸不受限制 11．可以测井及取芯 12．从一口直井可以钻多口水平分枝井 13．可实现有选择的完井方案 （4）．短曲率半径水平井的优缺点 优点缺点 1．井眼曲线段最短1．非常规的井下工具 2．侧钻容易２．非常规的完井方法 3．能够准确击中油层目标３．穿透油层段短（120—180米）4．从一口直井可以钻多口水平分枝井４．井眼尺寸受到限制

5．直井段与油层距离最小５．起下钻次数多 6．可用于浅油层６．要求使用顶部驱动系或动力水龙头 7．全井斜深最小７．井眼方位控制受到限制 8．不受地表条件的影响８．目前还不能进行电测 第三节定向井的基本术语解释 1）井深：指井口（转盘面）至测点的井 眼实际长度，人们常称为斜深。国外 称为测量深度（MeasureDepth）。 2）测深：测点的井深，是以测量装置 率是井斜角度（α）对井深（L?）的一阶导数。 dα Kα＝─── dL 井斜变化率的单位常以每100米度表示。 8）井深方位变化率：实际应用中简称方位变化率，?是指井斜方位角随井深变化的快慢程度，常用KΦ表示。计算公式如下： dΦ ＫΦ＝─── dL

第四章-钛合金的相变及热处理

第四章-钛合金的相变及热处理

第4章钛合金的相变及热处理 可以利用钛合金相变诱发的超塑性进行钛合金的固态焊接，接头强度接近基体强度。 4.1 同素异晶转变 1.高纯钛的β相变点为88 2.5℃，对成分十分敏感。在882.5℃发生同素异晶转变：α（密排六方）→β（体心立方），α相与β相完全符合布拉格的取向关系。 2.扫描电镜的取向成像附件技术（Orientation-Imaging Microscopy , OIM） 3.α/β界面相是一种真实存在的相，不稳定，在受热情况下发生明显变化，严重影响合金的力学性能。 4.纯钛的β→α转变的过程容易进行，相变是以扩散方式完成的，相变阻力和所需要的过冷度均很小。冷却速度大于每秒200℃时，以无扩散发生马氏体转变，试样表面出现浮凸，显微组织中出现针状α′。转变温度会随所含合金元素的性质和数量的不同而不同。 5.钛和钛合金的同素异晶转变具有下列特点： （1）新相和母相存在严格的取向关系 （2）由于β相中原子扩散系数大，钛合金的加热温度超过相变点后，β相长大倾向特别大，极易形成粗大晶粒。 （3）钛及钛合金在β相区加热造成的粗大晶粒，不像铁那样，利用同素异晶转变进行重结晶使晶粒细化。钛及钛合金只有经过适当的形变再结晶消除粗晶组织。 4.2 β相在冷却时的转变 冷却速度在410℃/s以上时，只发生马氏体转变；冷速在410～20℃/s时,发生块状转变；冷却继续降低，将以扩散型转变为主。 1.β相在快冷过程中的转变 钛合金自高温快速冷却时，视合金成分不同，β相可以转变成马氏体α′或α"、ω或过冷β等亚稳定相。 （1）马氏体相变 ①在快速冷却过程中，由于β相析出α相的过程来不及进行，但是β相的晶体结构，不易为冷却所抑制，仍然发生了改变。这种原始β相的成分未发生变化,但晶体结构发生了变化的过饱和固溶体是马氏体。 ②如果合金的溶度高，马氏体转变点M S降低至室温一下，β相将被冻结到室温，这种β相称过冷β相或残留β相。 ③若β相稳定元素含量少，转变阻力小，β相由体心立方晶格直接转变为密排六方晶格，这种具有六方晶格的过饱和固溶体称六方马氏体，以α′表示。 ④若β相稳定元素含量高，晶格转变阻力大，不能直接转变为六方晶格，只能转变为斜方晶格，这种具有斜方晶格的马氏体称斜方马氏体，以α′′表示。 ⑤马氏体相变是一个切变相变，在转变时，β相中的原子作集体的、有规律的进程迁移，迁移距离较大时形成六方α′相，迁移距离较小时形成斜方α′′相。 ⑥马氏体相变开始温度M S ；马氏体相变终了温度M f 。 ⑦钛合金中加入Al、Sn、Zr将扩大α相区，使β相变点升高；V、Mo、Mn、Fe、Cr、Cu、Si将缩小α相区（扩大β相区），使β相变点降低。 ⑧β相中原子扩散系数很大，钛合金的加热温度一旦超过β相变点，β相将快速长大成粗晶组织，即β脆性，故钛合金淬火的加热温度一般均低于其β相变点。

W85钨铜

W85钨铜 钨是理论上最好的金属电极材料。它的强度、密度、硬度都很高，熔点接近3400℃，因此在电火花和焊接加工过程中，钨电极实际损耗很小，但是纯钨作电极有两个困难：1.极难加工2.价格昂贵，所以利用纯铜的可塑性、高导电等优点，制成复合材料，就成了电极中的珍品--钨铜电极。 我司钨铜选用精细高纯钨粉、高纯铜粉，经静压成型、高温烧结、溶渗铜的一流工艺精制而成。针对硬质合金钨钢，高碳钢，淬火模具钢采用普通铜公电极损工大，精度低，加工慢的缺点，利用钨铜高导电，熔点高、热膨胀小的优点进行电火花放电加工，极大的改善了加工速度和精度。钨铜不仅导电性能好，软化温度高，而且耐电弧腐蚀，高耐磨性，使用寿命长，修正电极的频率低，在提高生产力效率的基础上还节省了加工工具的成本及维修费用，所以是理想的高级焊接材料。 W85钨铜的用途： ★高级电火花电极材料：针对钨钢（硬质合金）、高速钢、耐高温超硬合金制作的模具需电蚀加工时，普通电极损耗大，速度慢，而钨铜高的电腐蚀速度，低的损耗率，精确的电极形状，优良的加工性能，能保证被加工件的精确度大大提高。 ★高级焊接电极材料：综合了钨和铜的优点，耐高温、耐电弧烧蚀、抗熔焊和低截流、强度高、比重大、导电导热性好，易于切削加工，并具有发汗冷却和抗粘附等特性，经常用来做有一定耐磨性，抗高温的点焊、碰焊、对焊、凸焊电极。 ★电子领域的应用材料：钨铜具有强度高、导电导热性能好及热膨胀系数小等优点，所以作为一种新型的电子封装材料受到了电子工程师的青睐，被广泛的应用于功率电子器件，如整流管、晶闸管、功率模块、激光二极管、微波管等。在微电子器件中，如计算机CPU、DSP芯片等。钨铜在微波通讯、自动控制、电源转换、航空航天等领域发挥着重要的作用。目前，钨铜主要用在大功率微波管、大功率激光二极管以及某些大功率集成电路模块的热沉。★电触头材料：钨铜电触头在高压开关上已经使用多年，尤其以高压大电流断路器上使用量较大。如高压、无油、少油断路器、SF6断路器、隔离开关、重任务接触器等。 ★医疗设备和高比重材料：根据钨铜的特点，在医疗行业中用作防X射线和G射线的屏蔽材料。在民用工业中用作高比重合金配重，如手机振子、自动机械手表的重垂体、高尔夫球杆的杆体、飞镖等。 W85钨铜合金用途: 1.电阻焊电极： 综合了钨和铜的优点,耐高温、耐电弧烧蚀、强度高、比重大、导电、导热性好,易于切削加工,并具有发汗泠却等特性,由于具有钨的高硬度、高熔点、抗粘附的特点,经常用来做有一定耐磨性、抗高温的凸焊、对焊电极。 2.高压放电管电极： 高压真空放电管在工作时,触头材料会在零点几秒的时间内温度升高几千摄氏度,而钨铜的抗烧蚀性能、高韧性,良好的导电、导热性能给放电管稳定的工作提供必要的条件。 3、航天用高性能材料： 钨铜材料具有高密度、发汗冷却性能、高温强度高及耐冲刷烧蚀等性能，在航天工业中用作导弹、火箭弹的喷管喉衬，燃气舵的组件、空气舵、头罩及配重等。 4、真空触头材料： 触头材料必须有非常好的机械加工性能和抗热震性，由于接触和开断时打弧，触头材料会在零点几秒的时间内温度升高几千摄氏度。我公司生产的W-Cu触头材料由于其优异的物

钨铜

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高温合金

1、高温合金简介 (1) 2、高温合金的主要类别 (1) 2.1变形高温合金 (2) 2.1.1固溶强化型合金 (2) 2.1.2时效强化型合金 (2) 2.2铸造高温合金 (2) 2.3粉末冶金高温合金 (3) 2.4氧化物弥散强化(ODS)合金 (3) 2.5金属间化合物高温材料 (3) 3、高温合金的强化机理 (3) 3.1固溶强化 (3) 3.2沉淀强化及第二相强化 (4) 3.3晶界强化 (4) 3.4碳化物强化及质点弥散强化 (5) 4、常用高温合金的分类 (6) 4.1铁基超耐热合金 (6) 4.1.1铁基高温合金的成分和性能 (6) 4.2镍基超耐热合金 (6) 4.2.1镍基高温合金的组织特点 (6) 4.3钴基超耐热合金 (7) 4.3.1钴基高温合金的成分 (7) 4.3.2钴基高温的高温性能 (7) 5、高温合金的几种制造工艺 (7) 6、高温合金的应用 (8) 7、参考文献 (8)

1、高温合金简介 高温合金分为三类材料：760℃高温材料、1200℃高温材料和1500℃高温材料，抗拉强度800MPa。或者说是指在760--1500℃以上及一定应力条件下长期工作的高温金属材料，具有优异的高温强度，良好的抗氧化和抗热腐蚀性能，良好的疲劳性能、断裂韧性等综合性能，已成为军民用燃气涡轮发动机热端部件不可替代的关键材料。 按照现有的理论，760℃高温材料按基体元素主要可分为铁基高温合金、镍基高温合金和钴基高温合金。按制备工艺可分为变形高温合金、铸造高温合金和粉末冶金高温合金。按强化方式有固溶强化型、沉淀强化型、氧化物弥散强化型和纤维强化型等。高温合金主要用于制造航空、舰艇和工业用燃气轮机的涡轮叶片、导向叶片、涡轮盘、高压压气机盘和燃烧室等高温部件，还用于制造航天飞行器、火箭发动机、核反应堆、石油化工设备以及煤的转化等能源转换装置。 2、高温合金的主要类别 2.1变形高温合金 变形高温合金是指可以进行热、冷变形加工，工作温度范围-253～1320℃，具有良好的力学性能和综合的强、韧性指标，具有较高的抗氧化、抗腐蚀性能的一类合金。按其热处理工艺可分为固溶强化型合金和时效强化型合金。GH后第一位数字表示分类号即1、固溶强化型铁基合金2、时效硬化型铁基合金3、固溶强化型镍基合金4、钴基合金GH后，二，三，四位数字表示顺序号。 2.1.1固溶强化型合金 使用温度范围为900～1300℃，最高抗氧化温度达1320℃。例如GH128合金，室温拉伸强度为850MPa、屈服强度为350MPa；1000℃拉伸强度为140MPa、延伸率为85%,1000℃、30MPa 应力的持久寿命为200小时、延伸率40%。固溶合金一般用于制作航空、航天发动机燃烧室、机匣等部件。 2.1.2时效强化型合金 使用温度为-253～950℃，一般用于制作航空、航天发动机的涡轮盘与叶片等结构件。制作涡轮盘的合金工作温度为-253～700℃,要求具有良好的高低温强度和抗疲劳性能。例如：GH4169合金，在650℃的最高屈服强度达1000MPa；制作叶片的合金温度可达950℃，例如：GH220合金，950℃的拉伸强度为490MPa，940℃、200MPa的持久寿命大于40小时。变形高温合金主要为航天、航空、核能、石油民用工业提供结构锻件、饼材、环件、棒材、板材、管材、带材和丝材。 2.2铸造高温合金 铸造高温合金是指可以或只能用铸造方法成型零件的一类高温合金。其主要特点是： 1.具有更宽的成分范围由于可不必兼顾其变形加工性能，合金的设计可以集中考虑优化其使用性能。如对于镍基高温合金，可通过调整成分使γ’含量达60%或更高，从而在高达合金熔点85%的温度下，合金仍能保持优良性能。

第四章 钛合金的相变及热处理

第4章钛合金的相变及热处理 可以利用钛合金相变诱发的超塑性进行钛合金的固态焊接，接头强度接近基体强度。 4.1 同素异晶转变 1.高纯钛的β相变点为88 2.5℃，对成分十分敏感。在882.5℃发生同素异晶转变：α（密排六方）→β（体心立方），α相与β相完全符合布拉格的取向关系。 2.扫描电镜的取向成像附件技术（Orientation-Imaging Microscopy , OIM） 3.α/β界面相是一种真实存在的相，不稳定，在受热情况下发生明显变化，严重影响合金的力学性能。 4.纯钛的β→α转变的过程容易进行，相变是以扩散方式完成的，相变阻力和所需要的过冷度均很小。冷却速度大于每秒200℃时，以无扩散发生马氏体转变，试样表面出现浮凸，显微组织中出现针状α′。转变温度会随所含合金元素的性质和数量的不同而不同。 5.钛和钛合金的同素异晶转变具有下列特点： （1）新相和母相存在严格的取向关系 （2）由于β相中原子扩散系数大，钛合金的加热温度超过相变点后，β相长大倾向特别大，极易形成粗大晶粒。 （3）钛及钛合金在β相区加热造成的粗大晶粒，不像铁那样，利用同素异晶转变进行重结晶使晶粒细化。钛及钛合金只有经过适当的形变再结晶消除粗晶组织。 4.2 β相在冷却时的转变 冷却速度在410℃/s以上时，只发生马氏体转变；冷速在410～20℃/s时,发生块状转变；冷却继续降低，将以扩散型转变为主。 1.β相在快冷过程中的转变 钛合金自高温快速冷却时，视合金成分不同，β相可以转变成马氏体α′或α"、ω或过冷β等亚稳定相。 （1）马氏体相变 ①在快速冷却过程中，由于β相析出α相的过程来不及进行，但是β相的晶体结构，不易为冷却所抑制，仍然发生了改变。这种原始β相的成分未发生变化,但晶体结构发生了变化的过饱和固溶体是马氏体。 ②如果合金的溶度高，马氏体转变点M S降低至室温一下，β相将被冻结到室温，这种β相称过冷β相或残留β相。 ③若β相稳定元素含量少，转变阻力小，β相由体心立方晶格直接转变为密排六方晶格，这种具有六方晶格的过饱和固溶体称六方马氏体，以α′表示。 ④若β相稳定元素含量高，晶格转变阻力大，不能直接转变为六方晶格，只能转变为斜方晶格，这种具有斜方晶格的马氏体称斜方马氏体，以α′′表示。 ⑤马氏体相变是一个切变相变，在转变时，β相中的原子作集体的、有规律的进程迁移，迁移距离较大时形成六方α′相，迁移距离较小时形成斜方α′′相。 ⑥马氏体相变开始温度M S ；马氏体相变终了温度M f 。 ⑦钛合金中加入Al、Sn、Zr将扩大α相区，使β相变点升高；V、Mo、Mn、Fe、Cr、Cu、Si将缩小α相区（扩大β相区），使β相变点降低。 ⑧β相中原子扩散系数很大，钛合金的加热温度一旦超过β相变点，β相将快速长大成粗晶组织，即β脆性，故钛合金淬火的加热温度一般均低于其β相变点。 ⑨β相稳定元素含量越高，相变过程中晶格改组的阻力就越大，因而转变所需

钨铜主要用处有哪些

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钨铜都有哪些分类用途

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而且这些部件是在急剧温升的条件下(几秒钟升至工作温度)工作的，因此对部件产生激烈的热震破坏作用。对端头而言，当飞行器飞入太空再进入大气层时，由于速度快而受到粒子云的激烈摩擦产生高温和侵蚀。钨铜材料是能够满足上述要求的较好材料。 随着碳一碳(C—C)纤维复合材料的研制成功和发展，因它具有质轻和抗热震性好的优点，火箭喷管喉衬越来越多地用它来制造。但其抗烧蚀性远不如钨铜材料，对那些要求抗烧蚀性高的喷管喉衬、燃气舵和其他部件仍需用钨基材料制造。 3、破甲材料 钨铜材料可用作破甲材料，即一种所谓“药型罩”材料。用钨铜材料(常用W-30Cu 材料)制成杯形或漏斗形的罩，倒装于弹药简的前端，靠火药的温度和压力使罩变形成射流而穿甲。这种药型罩最早用紫铜制造，并大量应用。为了增大罩的单位质量从而提高破甲能力，后来研制单位质量比紫铜大的钨铡材料药型罩，在理想的情况下，它比紫铜罩的破甲能力提高30%左右。 钨铜用途： 钨铜合金综合铜和钨的优点，高强度/高比重/耐高温/耐电弧烧蚀/导电导热性能好/加工性能好，ANK钨铜合金采用高品质钨粉及无氧铜粉，应用等静压成型-（高温烧结）-渗铜，保证产品纯度及准确配比，组织细密，性能优异。本司铜钨系国内优质钨铜合金材料，极适合应用于高硬度材料及薄片电极放电加工，电加工产品表面光洁度高，精度高，损耗低，有效节约电极材料提高放电加工速度并改善模具精度。另可用作点焊/碰焊电极。钨铜与模具钢焊接成一体，在电极的使用上非常方便。 铜钨合金主要应用于：

镍基高温合金材料研究进展汇总-共7页

镍基高温合金材料研究进展 姓名：李义锋1 镍基高温合金材料概述 高温合金是指以铁、镍、钴为基，在高温环境下服役，并能承受严酷的机械应力及具有良好表面稳定性的一类合金[1]。高温合金一般具有高的室温和高温强度、良好的抗氧化性和抗热腐蚀性、优异的蠕变与疲劳抗力、良好的组织稳定性和使用的可靠性[2]。因此，高温合金既是航空、航天发动机高温部件的关键材料，又是舰船、能源、石油化工等工业领域不可缺少的重要材料，已成为衡量一个国家材料发展水平的重要标志之一。 在整个高温合金领域中，镍基高温合金占有特殊重要的地位。与铁基和钴基高温合金相比，镍基高温合金具有更高的高温强度和组织稳定性，广泛应用于制作航空喷气发动机和工业燃气轮机的热端部件。现代燃气涡轮发动机有50%以上质量的材料采用高温合金，其中镍基高温合金的用量在发动机材料中约占40%。镍基合金在中、高温度下具有优异综合性能，适合长时间在高温下工作，能够抗腐蚀和磨蚀，是最复杂的、在高温零部件中应用最广泛的、在所有超合金中许多冶金工作者最感兴趣的合金。镍基高温合金主要用于航空航天领域950-1050℃下工作的结构部件，如航空发动机的工作叶片、涡轮盘、燃烧室等。因此，研究镍基高温合金对于我国航天航空事业的发展具有重要意义。 镍基高温合金是以镍为基体(含量一般大于50 )、在650～1000℃范围内具有较高的强度和良好的抗氧化、抗燃气腐蚀能力的高温合金[2]。它是在Cr20Ni80合金基础上发展起来的，为了满足1000℃左右高温热强性(高温强度、蠕变抗力、高温疲劳强度)和气体介质中的抗氧化、抗腐蚀的要求，加入了大量的强化元素，如W、Mo、Ti、Al、Nb、Co等，以保证其优越的高温性能。除具有固溶强化作用，高温合金更依靠Al、Ti等与Ni形成金属问化合物γ′相(Ni3A1或Ni3Ti等)的析出强化和部分细小稳定MC、M23C6碳化物的晶内弥散强化以及B、Zr、Re等对晶界起净化、强化作用。添加Cr的目的是进一步提高高温合金抗氧化、抗高温腐蚀性能。镍基高温合金具有良好的综合性能，目前已被广泛地用于航空航天、汽车、通讯和电子工业部门。随着对镍基合金潜在性能的发掘，研究人员对其使用性能提出了更高的要求，国内外学者已开拓了针对镍基合金的新加工工艺如等温锻造、挤压变形、包套变形等。

钛合金及其热处理工艺简述样本

钛合金及其热解决工艺简述 宝鸡钛业股份有限公司：杨新林 摘要：本文对钛及其合金基本信息进行了简要简介，对钛几类固溶体划分进行了简述，对钛合金固态相变也进行了概述。重点概述了钛合金热解决类型及工艺，为之后生产实习中对钛合金热解决工艺结识提供指引。 核心词：钛合金，热解决 1 引言 钛在地壳中蕴藏量位于构造金属第四位,但其应用远比铜、铁、锡等金属滞后。钛合金中溶解少量氧、氮、碳、氢等杂质元素,使其产生脆性,从而妨碍了初期人们对钛合金开发和运用。直至二十世纪四五十年代,随着英、美及苏联等国钛合金熔炼技术改进和提高,钛合金应用才逐渐开展[5]。 纯钛熔点为1668℃,高于铁熔点。钛在固态下具备同素异构转变,在882.5℃以上为体心立方晶格β相,在882.5℃如下为密排六方晶格α相。钛 合金依照其退火后室温组织类型进行分类,退火组织为α相钛合金记为TAX,也 称为α型钛合金；退火组织为β相钛合金记为TBX,也称为β型钛合金；退火组织为α+β两相钛合金记为TCX,也称为α+β型钛合金,其中“X”为顺序号。国内当前钛合金牌号已超过50个,其中TA型26个,TB型8个以上,TC型15个以上[5]。 钛合金具备如下特点：

（1）与其她合金相比,钛合金屈强比很高,屈服强度与抗拉强度极为接近； （2）钛合金密度为4g/cm3,大概为钢一半,因而,它具备较高比强度； （3）钛合金耐腐蚀性能优良,在海水中其耐蚀性甚至比不锈钢还要好； （4）钛合金导热系数小,摩擦系数大,因而机械加工性不好； （5）在焊接时,钛合金焊缝金属和高热影响区容易被氧、氢、碳、氮等元素污染,使接头性能变坏。 在熔炼和各种加工过程完毕之后,为了消除材料中加工应力，达到使用规定性能水平，稳定零件尺寸以及去除热加工或化学解决过程中增长有害元素(例如氢)等,往往要通过热解决工艺来实现。钛合金热解决工艺大体可分为退火、固溶解决和时效解决三个类型。由于钛合金高化学活性,钛合金最后热解决普通在真空条件下进行。热解决是调节钛合金强度重要手段之一。 2 钛合金合金化特点 钛合金性能由Ti同合金元素间物理化学反映特点来决定，即由形成固溶体和化合物特性以及对α?β转变影响等来决定。而这些影响又与合金元素原子尺寸、电化学性质（在周期表中相对位置）、晶格类型和电子浓度等关于。但作为Ti合金与其他有色金属如Al、Cu、Ni 等比较，尚有其独有特点，如：（1）运用Tiα?β转变，通过合金化和热解决可以随意得到α、α+β和β相组织； （2）Ti是过渡族元素，有未填满d电子层，能同原子直径差位于±20％以内置换式元素形成高浓度固溶体；

钨铜合金-文献综述

目录 引言 (1) 一. 钨铜合金概况 (2) 1.1钨铜合金的性能及应用 (2) 1.2 钨铜合金的制备 (3) 1.2.1 熔渗法 (3) 1.2.2 活化液相烧结法 (5) 1.2.3金属注射成型（MIM） (7) 1.2.4 热压烧结法 (7) 1.2.5 超细混合粉末的直接烧结 (8) 二. 包覆粉及研究进展 (9) 2.1包覆粉的制备方法 (10) 2.1.1机械化学改性法 (10) 2.1.2溶胶-凝胶法 (11) 2.1.3 均相沉淀法 (11) 2. 1.4物理气相沉积法 (12) 2. 1.5化学镀法 (13) 三.钨铜板材的研究进展 (14) 3.1普通轧制 (14) 3.2金属粉末轧制 (14) 3.3其他制板技术 (15) 四.流延技术及应用 (16) 4.1.流延法 (16)

4.2.溶液流延法 (17) 参考文献 (19)

引言 钨铜合金由于自身的诸多优良特性，目前己广泛应用于大容量真空断路器和微电子领域。上世纪30年代中期，伦敦镭协会的Melennan和Smithells 最早进行了钨铜合金的研制。这类合金在国防、航空航天、电子信息和机械加工等领域中具有十分广泛的用途，在国民经济中占有重要的地位。钨基合金受到了世界各国的高度重视，已成为材料科学界较为活跃的研究领域之一。 钨具有高的熔点、高的密度、低的热膨胀系数和高的强度，铜具有很好的导热、导电性。由W和Cu组成的W-Cu合金兼具W和Cu的优点，即具有高的密度、良好的导热性和导电性、低的热膨胀系数。随着微电子信息技术的发展，电子器件的小型化和高功率化，器件的发热和散热是其必须面对的一个重要问题。W-Cu合金的高导热性可以满足大功率器件散热需要，尤为重要的是，其热膨胀系数(CTE)和导热导电性能可以通过调整材料的成分而加以设计，可以与微电子器件中不同半导体材料进行很好匹配连接，从而避免热应力所引起的热疲劳破坏。因此在大规模集成电路和大功率微波器件中，钨铜合金薄板作为电子封装基板、连接件、散热片和微电子壳体用材可以有效减少因散热不足和热膨胀系数差异导致的应力问题，延长电子元件的使用寿命，具有广阔的应用前景。 1