Shape Memory Ceramics SMC

Viscoelastic

Martensitic

Ferroelectric

Ferromagnetic

SMC Viscoelastic

• Mica based glass-ceramics

• Shape recovery up to 0.5% of deformation

• Process similar to shape memory polymers:

• plastic deformation at high temperature

• cooling down to room temperature to “freeze” the shape

• heating up to high temperature to recover the original shape

• Heterostructure composed by:

• 40-60% volumetric of mica

• glass matrix for the rest

Deformation recovery: mica glass-ceramic

above 823 K, however, readily leads to the formationof Ti

!Ni, Ti

"Ni

#and Ti

$$Ni

$#precipitates

[57, 61, 95]. As a result, the shape-memory behaviourof the thin films is very sensitive to the annealingconditions. Taking into account that the SMA filmsmay combine with dissimilar materials to form hybridcomposites, the crystallization temperatures should beas low as possible. Jardine et al. [57] reported that thecrystallization kinetics at 723K were very slow, fullcrystallization occurring after a 7.5 h anneal. However,the crystallization temperature can be further loweredat 673K through cold-working via rolling [57, 83].

2.1.2.6. Residual stress and adhesion. The residualstresses in the films will affect the mechanical behav-iour and the static and dynamical compatibility withthe substrates or dissimilar materials. Generally, thethin films deposited at low working gas pressure showlarge compressive intrinsic stresses. The biaxial resid-ual stress, !

!, in the film is given by the well-known

Stoney equation [61]

!!"1

6E"

1!""

t!"t!

1#

(1)

Upon heating, the stresses increase in magnitude withincreasing temperature and the rate of the stressevolution, d!/d¹ is determined by [116]

d!d¹" E

!1!"

($!!$

") (2)

where E is the film modulus, $ the coefficient of ther-mal expansion (CTE), " the film Poisson’s ratio, t thethickness, # the substrate radius of curvature, and thesubscripts s and f refer to substrate and film, respec-tively. For the TiNi film deposited on (1 0 0) Si, the ratewas estimated to be about !1.5MPaK%$ [116]. Thecompressive stress tends to zero when approachingthe crystallization temperature and then becomesa tensile stress of 50—100 MPa as a result of the filmdensification. The evolution of the high tensile stressmay lead to cracking and delamination when anneal-ing the amorphous SMA films on Si [89, 116]. How-ever, the stress can be relaxed on cooling by thestress-induced martensitic transformation, and onheating a recovery stress will be generated, formingthe basis for reversible cyclic actuation of the SMA/Sicomposites applicable to MEMS [57, 61, 89, 116, 117].

The deposited thin films can exhibit both perfectshape-memory effect and pseudoelasticity, and thethermomechanical properties are comparable to thoseof their bulk counterpart [61, 88, 91—94, 96]. A morespecific review of the sputtering process, characteriza-tion of the microstructures and phase transformations,and the thermomechanical properties of the depositedSMA thin films will be presented in a separate paper[118].

2.2. Shape-memory ceramics2.2.1. Visocelastic shape-memory ceramicsSome mica glass-ceramics exhibit clear shape-memoryphenomenon, i.e., a nearly perfect recovery of up to

Figure 1 Torsional strain recovery of a mica glass-ceramic asa function of time. The material was deformed in axial compressionat 773K before cooling under load to room temperature, and thenreheated to different temperatures [120].

0.5% prestrain, after high-temperature plastic defor-mation, cooling under load to room temperature, andthen reheating [119—124]. These materials typicallyhave a heterostructure of a volume fraction of between0.4 and 0.6 mica as the principal crystalline phasedispersed in a continuous glassy phase. Unlike shape-memory alloys, the shape-memory phenomenon in theheterostructure arises from the elastic energy intro-duced into the rigid matrix driving a viscous plasticstrain reversal in the dispersed crystalline element onreheating. At temperatures above 573K, the mica canbe deformed plastically by basal slip and the plasticstrain in the crystalline constituent accommodatedelastically by the surrounding rigid glass. Because theglide through dislocations in mica is not possible atlow temperature, the deformation of the material atelevated temperatures will be retained even if the loadis removed, after cooling under load to ambient tem-peratures. The elastic strain energy stored in the glassyphase thus will provide a driving force for recoveringthe original shape. If the deformed mica is reheated toa high temperature at which the stored elastic strainenergy is sufficient to activate the dislocation glide, thephase mixture will reverse the original plastic defor-mation. Fig. 1 shows the torsional strain recovery ofa mica glass-ceramic as a function of time at severaltemperatures. The sample was deformed in axial com-pression at 773K before cooling under load to roomtemperature, and a strain recovery of 99% was ob-served on prolonged annealing at 1073K [120]. Theviscoelastic shape-memory phenomenon is not limitedto mica glass-ceramics or glass-ceramics. %-spodu-mene glass-ceramics and 2ZnO—B

!O

"glass-ceramics,

and a variety of sintered ceramics that contain verylittle glass phase, including mica (KMg

"AlSi

"O

$&F!),

silicone nitride (Si"N

#), silicone carbide (SiC), zirconia

(ZrO!) and alumina (Al

!O

"), also exhibit the shape-

memory behaviour [121—124]. However, the recover-able strain of the ceramics is much smaller (about0.1%). Corresponding to the shape-recovery process,stress relaxation was observed. The activation ener-gies for the shape-recovery process in the glass-ceram-ics were found to be much lower than those for high-temperature creep of the same mica glass-ceramics,

3747

Deformed in axial compression at 773K

Mica based glass-ceramic

• Plastic deformation by slip of the basal plane of mica over 573 K

• The glass matrix follow the deformation elastically -> residual stress generation

• At room-/low temperature the deformation is retained

• Increasing the T up to the point the mica basal plane can slip under the residual stress from the glass matrix

• Recovery depends on several factors: max deformation, temperature, deformation rate, heating time

Others viscoelastic SMC

• Not only mica glass-ceramics show shape memory effects but also:

• β-spomudene or 2ZnO-B2O3 glass-ceramics

• ceramics containing small percentage of glass like: mica, silicon nitride, carbon nitride, zirconia and allumina

• The deformation recovery in ceramics is a lot smaller (around 0.1%)

• On observe a relaxation of residual stresses in the process

• The recovery energy is much lower than in mica based glass-ceramics

SMC martensitic

• The process is quite similar to the SMA

• Martensitic transformation:

• Ceramics based on zirconia or partially stabilized zirconia

• Martensitic transformation and re-alignement of ferro-elastic domains:

• Typical of some ionic materials like Pb3(PO4)2 & LnNbO4 (Ln=La, Nd)

• Superconductors like V-Si, Zr-Hf-V, Y-Ba-Cu-O, Bi-(Pb)-Sr-Ca-Cu-O & Ti-Ba-Ca-Cu-O

• All show both shape memory effects and pseudo-elastic deformation

Zirconia-stabilizer phase diagram

• Stabilizers:

• Y2O3

• CeO2

• MgO

• CaO

• ......

• TZP, CZP, Ce-TZP......

Stabilized zirconia

thereby differentiating the two phenomena. Becausethe materials are visco-elastic and thermally stimu-lated, the shape-recovery percentage and shape-recov-ery force show a strong dependence on the prestrain,pre-deformation temperature, pre-deformation rate,reheating temperature and holding time [121, 123].

2.2.2. Martensitic shape-memory ceramicsSome inorganic or ceramic compounds undergo mar-tensitic or displacive transformations which can beeither stress or thermally activated, often resulting intransformation plasticity or transformation toughen-ing. As a matter of fact, transformation toughening viathe martensitic transformation is one of the mosteffective ways of improving the reliability and struc-tural integrity of engineering ceramics, and has led tothe wide recognition of the technological importanceof the transformations in ceramics [125]. If the trans-formations in some ceramics are thermoelastic orferroelastic, reversible strain or shape recovery, that is,pseudoelasticity and shape-memory effect can be ex-pected. In certain ZrO

!-containing ceramics, the

transformation between the tetragonal structure (t-ZrO

!) at intermediate temperatures and the mono-

clinic structure (m-ZrO!) at low temperatures, can

occur thermoelastically [125]. The magnesia-partiallystabilized zirconia (Mg-PSZ) [119, 126] and theceria-stabilized, tetragonal zirconia polycrystals(CeO

!—TZP) [119, 127—129] have been reported to

show shape-memory behaviour. In the CeO!—TZP,

the t—m transformation occurs at a characteristic tem-perature M

!on cooling, which depends on the grain

size. The t—m transformation can also be stress-in-duced at temperatures above M

!, resulting in seem-

ingly ‘‘plastic’’ deformation. On subsequent reheating,the m—t transformation occurs and the strain can berecovered. Fig. 2 shows the shape-memory effect ob-served in a ZrO

!—12mol%CeO

!ceramic. In the ce-

ramic, total strains up to 1.5%—2% in uniaxialcompression can be almost fully recovered, but signifi-cant grain-boundary microcracking and other irre-versible damage occur which lead to a degradation inshape recovery and premature fracture. However, a nearfull shape recovery of axial compressive plastic strains ofup to 4.5% can be achieved while the tendency ofmicrocracking be suppressed in the samples deformed inhydraulic compression [119, 127]. More displaciveor martensitic-like transformations and potentialshape-memory materials can be found in a varietyof structural ceramics [130, 131]. In addition, someionic materials such as Pb

"(PO

#)!

and LnNbO#

(Ln"La, Nd), crystals [132—136], and some super-conductors such as V—Si and Zr—Hf—V [137],Y—Ba—Cu—O, Bi—(Pb)—Sr—Ca—Cu—O and Ti—Ba—Ca—Cu—O [138—141], also exhibit remarkable pseudoelas-ticity and shape-memory effect, due to first-order mar-tensitic transformations or the rearrangements of theferroelastic domains in the materials. Using theceramics, some new shape-memory device can bedesigned for high temperature application whereordinary shape-memory alloys are not applicable.However, the technological application of the

Figure 2 Axial stress—strain curve for ZrO!-12mol.%CeO2 under

uniaxial compression at room temperature, together with temper-ature—strain curve showing strain recovery on subsequent heating[129].

shape-memory capacity is limited by the small magni-tudes of recoverable strains and the tendency of theceramics to microcracking.

2.2.3. Ferroelectric shape-memory ceramicsIn pervoskite-type oxides, the crystallite domains mayexist in a variety of states such as cubic, tetragonal,rhombohedral or orthorhombic, which may be eitherparaelectric, ferroelectric or antiferroelectric, depend-ing upon the exact composition, as well as externalconditions such as temperature, stress and electricfield. The phase transitions between the differentstructures, such as the paraelectric—ferroelectric(PE—FE) transition and the antiferroelectric—ferroelec-tric (AFE—FE) transition, may be accompanied bya considerable strain in the ceramics. In particular, ithas been observed that the orthorhombic antifer-roelectric (AFE) to tetragonal ferroelectric (FE)transition in some ceramics usually generate largestrains [119, 142—148]. Very importantly, in suitablecompositions, the AFE—FE transition can be inducedby application of a sufficiently large electric field. Thephase transition, which is caused by the switching orreorientation of the polarized domains, is accom-panied by a lattice distortion leading to a linear‘‘digital’’ displacement and a net volume expansion[146, 147]. Because the total strain is comprised of (i)a spontaneous strain which occurs due to the phasetransformation, and (ii) a strain associated with do-main alignment on poling, the induced strain is muchlarger than that resulting from the converse piezoelec-tric effect in the conventional piezoelectrics. Totalstrains of up to 0.6% have been reported tooccur during the transition [144, 146]. When theelectric field is removed, conventional electroceramicswill return to their original state. This is a typical

3748

Zirconia: shape memory mechanism

Figure 8 Schematic illustration of the deformation processes in the ferroelastic, ferromagnetic, and ferroelectric materials. (a) Stress-inducedmartensitic transformation by twinning and reorientation of martensite domains by detwinning, (b) magnetic field-induced transformationand reorientation of magnetic domains, (c) electric field-induced AFE—FE transformation and polarization of FE.

rhombohedral or tetragonal structure [175, 212]. Asmentioned above, the compositions of the ferroelectricshape-memory ceramics are usually so selected thatthe ceramics have an antiferroelectric structure butthey are close to the ferroelectric phase boundary, asshown in Fig. 10. In this region, the antiferroelectriccan be transformed into the ferroelectric at high elec-tric fields. Both the metastable AFE and FE phasesmay coexist, rendering the forward switchingAFE—FE under increasing field and the backwardswitching FE—AFE under decreasing field. The phasestability of the phases can be predicted based ona thermodynamic model [145]. The transformationscan also be approached within the ‘‘constrained the-ory’’ established by James and Wuttig for ferromag-netics [207]. Because antiferroelectrics do not displayany macroscopic polarization and very little strain isachieved when applied at low electric fields, theAFE—FE transformation and the subsequent poling of

the ferroelectric domains will cause mechanicalstrains, as illustrated in Fig. 8c. Suppose the magni-tude of the sublattice polarization remains essentiallyunchanged during the transformation, then, accordingto the two-sublattice model [146, 150], the spontan-eous strain associated with field-induced paraelectric—ferroelectric (PE-FE) transition is described by

!!"#

"QP$%

(9)

while the spontaneous strain in the antiferroelectricstate, !

&#', and the strain in the ferroelectric state

induced from the antiferroelectric by applying a field,!#'

, can be expressed, respectively, as

!&#'

"Q(1!!)P$&

(10)

!#'

"Q (1#!)P$#

(11)

where P&

and P#

are two-sublattice based polariza-tions, P

!is the field-induced polarization, Q and ! are

3755

Ce stabilized zirconia

Figure 6 Typical stress—strain curves at various temperatures forshape-memory materials. (a) A Ti—49.8 at % Ni shape-memory alloy[13]; (b) a CeO

!-stabilized tetragonal zirconia polycrystal (Ce-TZP)

shape-memory ceramic [127] (—) axial, (- - -) radial. In both thematerials the critical stress to induce the transformation showsa remarkable increase with increasing temperature.

transformed phase), and can be experimentally deter-mined or theoretically estimated from the constitutiverelations [182—186]. Fig. 7 shows the measured recov-ery stress as a function of temperature for a Ti—Nialloy for different prestrains. Through thermo-mechanical training, both one-way SME and two-waySME can be achieved. These characteristics ofSMAs can be applied for actuation, shape or position

Figure 7 Recovery stress as a function of temperature and prestrainfor a binary Ti—Ni shape-memory alloy [20].

control, impact or creep resistance and energyconversion.

From a thermodynamic point of view, pressure, liketemperature, is an independent variable that canchange the free energy and thus phase state of a mater-ial. It is well known that the Gibbs’ chemical freeenergy, G, is defined as

G"H!¹S (6)

where ¹ is temperature, H and S are the enthalpy andentropy of the system, respectively. Based on the firstand second laws of thermodynamics, it can be easilyderived that

dG"»dP!Sd¹ (7)

where P is pressure and » is volume. It is evident thatthe free energy can be altered by varying either pres-sure or temperature. At a given temperature, the in-crease in pressure will increase the free energy of thesystem. As a consequence, the system will tend totransform into other phases that have a lower freeenergy under the pressure, or, to decrease its volumethrough contraction, which will change the electronicstructures and in turn the physical properties, andthermodynamic state of the material. Pressure-in-duced phase transitions have been observed in a verywide range of materials. Of particular interest are thereversible pressure-induced transitions, which mayimplicate shape-memory effect. More recently, a mem-ory glass exhibiting structural reversibility has beendiscovered by Kruger and Jeanloz [189]. At 300K,both polycrystalline and single-crystal AlPO

"be-

rlinite became amorphous when the pressure was in-creased above 15GPa. On reducing the pressurebelow 5GPa, they reconverted from the glassy state tothe crystalline state with the same orientation as theoriginal crystal, implying a memory capacity of thecrystallographic structure. Similar phase transitionshave been observed in GaAsO

"and other quartz-like

structure [190, 191].The effects of hydrostatic pressure on martensitic

transformations in some shape-memory alloys, in-cluding Cu—Al—Ni, Ti—Ni, Fe—Pt, Fe—Ni, Fe—Ni—C,Fe—Ni—Mn and Fe—Ni—Co—Ti alloys have been

3753

Figure 6 Typical stress—strain curves at various temperatures forshape-memory materials. (a) A Ti—49.8 at % Ni shape-memory alloy[13]; (b) a CeO

!-stabilized tetragonal zirconia polycrystal (Ce-TZP)

shape-memory ceramic [127] (—) axial, (- - -) radial. In both thematerials the critical stress to induce the transformation showsa remarkable increase with increasing temperature.

transformed phase), and can be experimentally deter-mined or theoretically estimated from the constitutiverelations [182—186]. Fig. 7 shows the measured recov-ery stress as a function of temperature for a Ti—Nialloy for different prestrains. Through thermo-mechanical training, both one-way SME and two-waySME can be achieved. These characteristics ofSMAs can be applied for actuation, shape or position

Figure 7 Recovery stress as a function of temperature and prestrainfor a binary Ti—Ni shape-memory alloy [20].

control, impact or creep resistance and energyconversion.

From a thermodynamic point of view, pressure, liketemperature, is an independent variable that canchange the free energy and thus phase state of a mater-ial. It is well known that the Gibbs’ chemical freeenergy, G, is defined as

G"H!¹S (6)

where ¹ is temperature, H and S are the enthalpy andentropy of the system, respectively. Based on the firstand second laws of thermodynamics, it can be easilyderived that

dG"»dP!Sd¹ (7)

where P is pressure and » is volume. It is evident thatthe free energy can be altered by varying either pres-sure or temperature. At a given temperature, the in-crease in pressure will increase the free energy of thesystem. As a consequence, the system will tend totransform into other phases that have a lower freeenergy under the pressure, or, to decrease its volumethrough contraction, which will change the electronicstructures and in turn the physical properties, andthermodynamic state of the material. Pressure-in-duced phase transitions have been observed in a verywide range of materials. Of particular interest are thereversible pressure-induced transitions, which mayimplicate shape-memory effect. More recently, a mem-ory glass exhibiting structural reversibility has beendiscovered by Kruger and Jeanloz [189]. At 300K,both polycrystalline and single-crystal AlPO

"be-

rlinite became amorphous when the pressure was in-creased above 15GPa. On reducing the pressurebelow 5GPa, they reconverted from the glassy state tothe crystalline state with the same orientation as theoriginal crystal, implying a memory capacity of thecrystallographic structure. Similar phase transitionshave been observed in GaAsO

"and other quartz-like

structure [190, 191].The effects of hydrostatic pressure on martensitic

transformations in some shape-memory alloys, in-cluding Cu—Al—Ni, Ti—Ni, Fe—Pt, Fe—Ni, Fe—Ni—C,Fe—Ni—Mn and Fe—Ni—Co—Ti alloys have been

3753

NiTinol

Ferroelectrics

Figure 8 Schematic illustration of the deformation processes in the ferroelastic, ferromagnetic, and ferroelectric materials. (a) Stress-inducedmartensitic transformation by twinning and reorientation of martensite domains by detwinning, (b) magnetic field-induced transformationand reorientation of magnetic domains, (c) electric field-induced AFE—FE transformation and polarization of FE.

rhombohedral or tetragonal structure [175, 212]. Asmentioned above, the compositions of the ferroelectricshape-memory ceramics are usually so selected thatthe ceramics have an antiferroelectric structure butthey are close to the ferroelectric phase boundary, asshown in Fig. 10. In this region, the antiferroelectriccan be transformed into the ferroelectric at high elec-tric fields. Both the metastable AFE and FE phasesmay coexist, rendering the forward switchingAFE—FE under increasing field and the backwardswitching FE—AFE under decreasing field. The phasestability of the phases can be predicted based ona thermodynamic model [145]. The transformationscan also be approached within the ‘‘constrained the-ory’’ established by James and Wuttig for ferromag-netics [207]. Because antiferroelectrics do not displayany macroscopic polarization and very little strain isachieved when applied at low electric fields, theAFE—FE transformation and the subsequent poling of

the ferroelectric domains will cause mechanicalstrains, as illustrated in Fig. 8c. Suppose the magni-tude of the sublattice polarization remains essentiallyunchanged during the transformation, then, accordingto the two-sublattice model [146, 150], the spontan-eous strain associated with field-induced paraelectric—ferroelectric (PE-FE) transition is described by

!!"#

"QP$%

(9)

while the spontaneous strain in the antiferroelectricstate, !

&#', and the strain in the ferroelectric state

induced from the antiferroelectric by applying a field,!#'

, can be expressed, respectively, as

!&#'

"Q(1!!)P$&

(10)

!#'

"Q (1#!)P$#

(11)

where P&

and P#

are two-sublattice based polariza-tions, P

!is the field-induced polarization, Q and ! are

3755

• PLZT-(Pb, La)(Zr, Ti)O3, PZST-Pb(Zr,Sn,Sn,Ti)O3, PLSnZT-(Pb, La)(Zr, Sn, Ti)O3, (Pb, Nb)(Zr, Sn, Ti)O3

• (Sr,Ba)Nb2O6

• Hexagonal manganites RMnO3 (R=Ho, Y)

PZTs phase transformations

Figure 5 Phase diagrams of three prototypes of smart materialsshowing some common features. (a) Piezoelectric Pb(Zr, Ti)O

![212], (b) magnetostrictive (Tb, Dy) Fe

"[176], and (c) ferroelas-

tic Cu—xZn—10 at% Al shape-memory alloy (only a portion ofmetastable phase diagram of interest is shown, estimated froma modified Bragg—Williams model). The arrows indicate the mor-photropic phase boundaries and the compositions of the materialsin practical use.

proper stresses are applied at temperatures near theparent—martensitic phase boundary. Experimentally,it has been verified that both uniaxial tensile andcompressive stresses increase the martensitic trans-formation temperature, M

!, without regard to the

sense of external stresses and the kind of alloy systems[179—181]. The critical stress, (!"), to induce the mar-tensitic transformation at a given temperature isguided by the well-known Clausius—Clapeyonequation

d!"

d¹" !H»¹

#!#

(4)

where !H is the transformation enthalpy, !# thetransformation strain, » the molar volume and ¹

#is

the temperature at which the parent and martensitephases are in equilibrium at zero stress. Generally, thevalue of !" increases linearly with increasing temper-ature. Fig. 6 shows the stress—strain curves at differenttemperatures for a Ti—Ni SMA and Ce-TZP SMM.More systematic descriptions of the stress—strain—tem-perature space can be demonstrated by the well-estab-lished constitutive relations [182—186]. The mostsimple form of the one-dimensional SMA constitutiveequation is given by

!!!#"D(#!#

#)#"(¹!¹

#)##($!$

#) (5)

where D is Young’s modulus, " is the thermoelastictensor, # is the transformation tensor, and !, # and¹ are the state variables for stress, strain and temper-ature, respectively. $(!, ¹) is the martensitic fraction.The stress-induced martensitic transformations inmost SMAs (excluding some iron-based alloys) arereversible when unloading, giving rise to a mechanicalshape memory: this effect is usually called pseudo-elasticity. More importantly, the recoverable trans-formation strain for most polycrystalline SMAs can beas much as up to 8%, for some single-crystallineSMAs may exceed 10% [13, 14, 18], resulting ina very high elastic energy storage capacity.

The martensite phase consists of highly twinnedpolydomains (martensite variants). There are 24 pos-sible kinds of martensite variants in SMAs. The crys-tallographic characteristics can be well described bythe phenomenological crystallographic theory of mar-tensitic transformations [13]. When subjecting thematerial in the matensite state to an applied stress, thevariants undergo a self-accommodating pattern ofshear-induced shrinkage, growth and reorientation bydetwinning, resulting in twin-induced inelastic strainswhich can be described by some simple relations, asestablished by Roytburd [187, 188]. Unloading fromthe inelastic state may leave a residual strain, but itcan be recovered by heating the material to the parentphase. This is the so-called shape-memory effect(SME). The mechanism of the shape-memory effectwas overviewed in more detail by Miyazaki and Ot-suka [13]. Besides the unique behaviour, heat-inducedformation of the parent phase generates a consider-able recovery stress, which can be used for mechanicalactuation. The recovery stress is a function of pre-strain and temperature (i.e. the volume fraction of

3752

Ferroelectrics & anti-ferroelectrics

Figure 3 Comparison of the longitudinal strains for ferroelectricand antiferroelectric materials. (a) Spontaneous strain due to polar-ization in a ferroelectric, (b) field-induced transformation strainfrom AFE to FE state [148].

ferroelectric behaviour. Some ceramics, however, aremetastable in either of the ferroelectric and antifer-roelectric states at zero field, and they will remain inthe ferroelectric state as the field is removed. To returnto their original states, it can be achieved either rap-idly by reversing the polarity of the applied field, orslowly by heating to effect the reverse FE—AFE trans-formation, thus giving rise to a shape-memory behav-iour similar to those observed in shape-memory alloys.Fig. 3 shows a comparison of the filed-induced strainsfor ferroelectric and antiferroelectric materials. Theshape-memory effect has been observed in the perov-skite-type oxides (Pb, La) (Zr, Ti)O

!(PZSTs) [119, 142,

143], Pb(Zr, Sn, Sn, Ti)O!

(PZSTs) [119, 142, 144,145], (Pb, La) (Zr, Sn, Ti)O

!(PLSnZTs) and (Pb,Nb)

(Zr, Sn, Ti) O!

[146—150], (Sr, Ba) Nb"O

#[151] and

the hexagonal manganites RMnO!

(R"Ho, Y)[152]. Although the shape-memory ceramics havelower strain levels than shape-memory alloys, theyhave some clear advantages and may be more suitablefor certain applications. For instance, because an elec-tric field can be readily changed at much higher ratesthan temperature, the shape-memory ceramics may beactuated at higher bandwidths, with the maximumresponse speed of only a microsecond [142]. A proto-type adaptive structure using the shape-memory cer-amics has been demonstrated by Ghandi and Hagood[147, 148].

2.2.4. Ferromagnetic shape-memoryceramics

Some transition metal oxides undergo paramag-netic—ferromagnetic, paramagnetic—antiferromagnetictransformation, or orbital order—disorder transitionsand the reversible transformations are also accom-panied by recoverable lattice distortions. In the tetra-gonal manganite spinels Mn

!(Zn, Cd)

$%!Mn

"O

&[153, 154] and in the non-stoichiometric orthoman-ganites RMnO

!"!(R"Nd, Sm, Eu, Gd, Tb, Dy) [155],

the orbital ordered and disordered phases coexistin a wide temperature interval, and short-rangeferromagnetic (or antiferromagnetic) ordering orJahn—Teller phase transitions may take place,resulting in a shape-memory effect. Because, mostof the manganites are antiferromagnets and theirNeel temperatures are very low, spontaneousmagnetization of the compounds is only achievableat very low temperatures. As a result, the effects ofmagnetic field on the transformation and ferromag-netic shape-memory effect in the compounds are lessinvestigated.

2.3. Shape-memory polymers and gelsShape-memory polymers (SMPs) were introduced in1984 in Japan, and have since gained much attentionin Japan and in the United States [156, 157]. It is wellknown that polymers exist in a rubbery state at highertemperatures and in a glassy state at lower temper-atures. Owing to the lower rubbery modulus, poly-mers can be subjected to a very large deformation athigher temperatures. Because the glassy modulus is atleast two orders of magnitude times the rubberymodulus, the stored elastic stress is not large enoughto drive the reverse deformation in the glassy state asthe load is removed. As a result, the deformation canbe frozen in the glassy state on quenching or aftercooling under load to the lower temperatures. Usu-ally, ordinary polymers cannot completely restoretheir residual inelastic deformation upon reheating tothe rubbery state. In contrast, shape-memory poly-mers can recover almost all the residual deformation.These polymers typically consist of two phases, name-ly fixed phase and reversible phase. Amongst them, thepolynorborene, the trans-isopolyprene and the styrene—butadiene copolymer were the first few polymers re-ported to exhibit shape-memory effect. However, thecommercial application of the early-developed shape-memory polymers were much limited because of theirundesirable properties, namely a narrow range of glasstransition temperature, ¹

!, and poor processibility

[157—159].More recently, segmented polyurethane thermo-

plastic polymers have received intensive attention[156—160]. They are usually polymerized fromdiphenylmethane diisocyanate (OCN—R—NCO) andpolycaprolactone diols (HO—R!—OH), with butanediol(HO—R!!—OH) as chain extenders. The polyurethaneare thus basically multiblock copolymers consisting ofalternating sequences of soft segments with molecular

3749

Figure 3 Comparison of the longitudinal strains for ferroelectricand antiferroelectric materials. (a) Spontaneous strain due to polar-ization in a ferroelectric, (b) field-induced transformation strainfrom AFE to FE state [148].

ferroelectric behaviour. Some ceramics, however, aremetastable in either of the ferroelectric and antifer-roelectric states at zero field, and they will remain inthe ferroelectric state as the field is removed. To returnto their original states, it can be achieved either rap-idly by reversing the polarity of the applied field, orslowly by heating to effect the reverse FE—AFE trans-formation, thus giving rise to a shape-memory behav-iour similar to those observed in shape-memory alloys.Fig. 3 shows a comparison of the filed-induced strainsfor ferroelectric and antiferroelectric materials. Theshape-memory effect has been observed in the perov-skite-type oxides (Pb, La) (Zr, Ti)O

!(PZSTs) [119, 142,

143], Pb(Zr, Sn, Sn, Ti)O!

(PZSTs) [119, 142, 144,145], (Pb, La) (Zr, Sn, Ti)O

!(PLSnZTs) and (Pb,Nb)

(Zr, Sn, Ti) O!

[146—150], (Sr, Ba) Nb"O

#[151] and

the hexagonal manganites RMnO!

(R"Ho, Y)[152]. Although the shape-memory ceramics havelower strain levels than shape-memory alloys, theyhave some clear advantages and may be more suitablefor certain applications. For instance, because an elec-tric field can be readily changed at much higher ratesthan temperature, the shape-memory ceramics may beactuated at higher bandwidths, with the maximumresponse speed of only a microsecond [142]. A proto-type adaptive structure using the shape-memory cer-amics has been demonstrated by Ghandi and Hagood[147, 148].

2.2.4. Ferromagnetic shape-memoryceramics

Some transition metal oxides undergo paramag-netic—ferromagnetic, paramagnetic—antiferromagnetictransformation, or orbital order—disorder transitionsand the reversible transformations are also accom-panied by recoverable lattice distortions. In the tetra-gonal manganite spinels Mn

!(Zn, Cd)

$%!Mn

"O

&[153, 154] and in the non-stoichiometric orthoman-ganites RMnO

!"!(R"Nd, Sm, Eu, Gd, Tb, Dy) [155],

the orbital ordered and disordered phases coexistin a wide temperature interval, and short-rangeferromagnetic (or antiferromagnetic) ordering orJahn—Teller phase transitions may take place,resulting in a shape-memory effect. Because, mostof the manganites are antiferromagnets and theirNeel temperatures are very low, spontaneousmagnetization of the compounds is only achievableat very low temperatures. As a result, the effects ofmagnetic field on the transformation and ferromag-netic shape-memory effect in the compounds are lessinvestigated.

2.3. Shape-memory polymers and gelsShape-memory polymers (SMPs) were introduced in1984 in Japan, and have since gained much attentionin Japan and in the United States [156, 157]. It is wellknown that polymers exist in a rubbery state at highertemperatures and in a glassy state at lower temper-atures. Owing to the lower rubbery modulus, poly-mers can be subjected to a very large deformation athigher temperatures. Because the glassy modulus is atleast two orders of magnitude times the rubberymodulus, the stored elastic stress is not large enoughto drive the reverse deformation in the glassy state asthe load is removed. As a result, the deformation canbe frozen in the glassy state on quenching or aftercooling under load to the lower temperatures. Usu-ally, ordinary polymers cannot completely restoretheir residual inelastic deformation upon reheating tothe rubbery state. In contrast, shape-memory poly-mers can recover almost all the residual deformation.These polymers typically consist of two phases, name-ly fixed phase and reversible phase. Amongst them, thepolynorborene, the trans-isopolyprene and the styrene—butadiene copolymer were the first few polymers re-ported to exhibit shape-memory effect. However, thecommercial application of the early-developed shape-memory polymers were much limited because of theirundesirable properties, namely a narrow range of glasstransition temperature, ¹

!, and poor processibility

[157—159].More recently, segmented polyurethane thermo-

plastic polymers have received intensive attention[156—160]. They are usually polymerized fromdiphenylmethane diisocyanate (OCN—R—NCO) andpolycaprolactone diols (HO—R!—OH), with butanediol(HO—R!!—OH) as chain extenders. The polyurethaneare thus basically multiblock copolymers consisting ofalternating sequences of soft segments with molecular

3749

FerroelectricFerroelectric-antiferroelectric

metastable

In the case of antiferroelectric metastability on observe a ferroelectric or antiferroelectric phase at zero field depending on the history. A deformation is associated. With the

temperature the stable phase is recovered and the deformation released.

Pb0.99Nb0.02((Zr0.6Sn0.4)0.94Ti0.06)0.98O3

• Elastic deformation induced by the electric field at different temperatures

Ferromagnetic SMC

Figure 8 Schematic illustration of the deformation processes in the ferroelastic, ferromagnetic, and ferroelectric materials. (a) Stress-inducedmartensitic transformation by twinning and reorientation of martensite domains by detwinning, (b) magnetic field-induced transformationand reorientation of magnetic domains, (c) electric field-induced AFE—FE transformation and polarization of FE.

rhombohedral or tetragonal structure [175, 212]. Asmentioned above, the compositions of the ferroelectricshape-memory ceramics are usually so selected thatthe ceramics have an antiferroelectric structure butthey are close to the ferroelectric phase boundary, asshown in Fig. 10. In this region, the antiferroelectriccan be transformed into the ferroelectric at high elec-tric fields. Both the metastable AFE and FE phasesmay coexist, rendering the forward switchingAFE—FE under increasing field and the backwardswitching FE—AFE under decreasing field. The phasestability of the phases can be predicted based ona thermodynamic model [145]. The transformationscan also be approached within the ‘‘constrained the-ory’’ established by James and Wuttig for ferromag-netics [207]. Because antiferroelectrics do not displayany macroscopic polarization and very little strain isachieved when applied at low electric fields, theAFE—FE transformation and the subsequent poling of

the ferroelectric domains will cause mechanicalstrains, as illustrated in Fig. 8c. Suppose the magni-tude of the sublattice polarization remains essentiallyunchanged during the transformation, then, accordingto the two-sublattice model [146, 150], the spontan-eous strain associated with field-induced paraelectric—ferroelectric (PE-FE) transition is described by

!!"#

"QP$%

(9)

while the spontaneous strain in the antiferroelectricstate, !

&#', and the strain in the ferroelectric state

induced from the antiferroelectric by applying a field,!#'

, can be expressed, respectively, as

!&#'

"Q(1!!)P$&

(10)

!#'

"Q (1#!)P$#

(11)

where P&

and P#

are two-sublattice based polariza-tions, P

!is the field-induced polarization, Q and ! are

3755

• Tetragonal manganites (spinels: Mnx(Zn,Cd)1-xMn2O4)

• Ortomanganites non stoichiometric: (RMnO3+x, R=Nd, Sm, Eu, Gd, Tb, Dy)

Comparison SMA-Piezo-Magnetici

TABLE II Comparison of characteristics of shape-memory alloys, piezoelectric ceramics and magnetostrictive materials as actuationmaterials

Properties Shape-memory alloy Piezoelectric Magnetostrictive(Ti—Ni) (PZT) (Terfenol-D)

Compressive stress (MPa) !800 60 700Tensile strength (MPa) 800—1000 30—55 28—35Young’s modulus (GPa) 50—90 (P) 60—90 (Y! )! 25—35 (Y" )"

10—35 (M) !110 (Y# )# 50—55 (Y$ )$Maximum strain !0.1 !0.001 !0.01Frequency (Hz) 0—100 1—20 000 1—10 000Coupling coefficient ! 0.75 0.75Efficiency (%) 3—5 50 80Energy density (kJm%&) 300—600 !1.0 14—25

!modulus for constant electric field#modulus for constant electric displacement"modulus for constant magnetizing field$modulus for constant induction field

conventional piezoelectric or electrostrictive ceramicshave a superior dynamic response but their displace-ments are quite small and most of them are verybrittle. Combining SMAs with piezoelectric or mag-netostrictive materials, field-activated smart com-posites can be designed, which may generate a largerdisplacement than conventional piezoelectric ceramicsor magnetostrictive materials and have an improveddynamic response as compared to monolithic SMAs.More recently, some pioneers have explored the tech-nical feasibility of smart thin-film heterostructures bydepositing the SMA thin films on piezoelectric ormagnetostrictive substrates. However, the complexityof the fabrication processing and the interface bondingand dynamic coupling of dissimilar components re-main tough issues for the composites [213].

The alternative is to improve the monolithic shape-memory materials by employing new processing tech-niques or to design a new generation of shape-memorymaterials. The development of deposited thin-filmshape-memory alloys, as we described above, is one ofthe efforts directed to this objective. Also worth men-tion are the recently developed porous shape-memoryalloys [214, 215]. Bulk Ti—Ni alloys with different por-osity, exhibiting superelasticity and shape-memoryeffect, have been successfully manufactured via thepowder metallurgical route. The porous SMAs arevery desirable for some biomedical applications be-cause the alloys have good biocompatibility and theirporous structure favours in-growth of living tissuesand firm fixation. Naturally, it reminds us of bone— a typical biomimetic model. Bone is also porous;moreover, it exhibits pyroelectricity and piezoelectric-ity, and maintains the skeletal homeostasis and min-eral homeostasis for the body [1]. After the model,biomimetic artificial bone materials based on theporous SMAs and other advanced materials may bedeveloped. For instance, microballoons or microtubescoated by some functional material layers can beconstructed in the porous SMAs which may providea suitable substrate or skeleton to grow heterostruc-tures with certain intelligence.

In principle, the deformation of the polydomains inthe ferromagnetic and ferroelectric materials by ap-plying external fields can be controlled just the same

way as the stress-induced deformation of the marten-sites in ferroelastic SMAs. The next challengingobjective, therefore, is to explore new potentially com-mercial materials wherein the martensitic-like trans-formations and the reorientation of the domains canbe induced by magnetic fields or electric fields atambient temperatures. The design concepts and strat-egies for finding new ferromagnetic and ferroelectricshape-memory materials have been proposed [142,205—211]. In this aspect, the remarked common fea-tures shared by several smart material systems, andthe successful development story of the giant mag-netostrictive materials Terfenol-D [176, 216—218]may offer some clues or inspirations.

AcknowledgementsThe authors thank Prof. J. Van Humbeeck of K. U.Leuven and Mr J. Cederstrom, Scandinavian MemoryMetals, AB, for critical reading of the manuscript,Professor R. D. James, University of Minnesota, andProfessor D. S. Grummon, Michigan State University,for providing their papers prior to publication. One ofthe authors (Z G W) wishes to acknowledge the RoyalInstitute of Technology (KTH) for offer of a fellowshipto work at KTH.

References1. M. V. GANDHI and B. S. THOMPSON, ‘‘Smart Materials

and Structures’’ (Chapman & Hall, New York, 1992).2. T. TAKAGIL, J. Intelligent Mater. System Struct. 1 (1990)

149.3. Idem, in ‘‘Proceedings of the 3rd International Conference on

Intelligent Materials’’, edited by P. F. Gobin and J. Tatibouet(Technomic, Lancaster, 1996) p. 2.

4. C. A. JAEGER and C. A. ROGERS, in ‘‘Proceedings of theARO Workshop on Smart Materials, Structures and Math-ematical Issues’’, edited by C. A. Rogers (ARO, Washington,1988) p. 14.

5. S. NOZAKI and K. TAKAHASHI , in ‘‘Proceedings of the2nd International Conference on Intelligent Materials’’,edited by C. A. Rogers and G. G. Wallace (Technomic,Lancaster, 1994) p. 1230.

6. G. BECK and P. F. GOBIN, in ‘‘Proceedings of the 1stInternational Conference on Intelligent Materials’’, edited byT. Takagi, K. Takahashi, M. Aizawa and S. Miyata (Tech-nomic, Lancaster, 1992) p. 9.

3758

Comparison different types of actuatorsConfronto tipi di attuatori

IBM SYSTEMS JOURNAL, VOL 35, NOS 3&4, 1996 FLETCHER 633

been discovered but are still in the developmentalstage.18

Piezoresistive polymers and polymer composites.Another promising class of materials for force trans-duction consists of electronic polymer materials. Thefield of polymer science continues to advance tremen-dously, and there now exist various types of polymersthat behave to some degree like the more commonmetal- or ceramic-based smart materials.19 However,since most smart polymers are still in early develop-ment, an alternate approach has been to use polymersfor sensing or measuring force indirectly. These mate-rials are grouped separately from smart materialsbecause the mechanism employed is not reciprocal(e.g., applying an electrical stimulus will not producea force output). Rather than performing force trans-duction through direct conversion of mechanicalenergy to electrical energy, as in PVDF, certain types ofpolymers can be used to modulate an electric currentvia a force-dependent electrical resistance. Polymersensors can be used in compression as well as exten-sion and can function over a greater range of strainsthan conventional metal strain gauges.20

The two most common physical mechanisms em-ployed in resistive polymer force sensors are piezo-resistance and geometric deformation. The deforma-tion type of force sensor is comprised of a conductingpolymer strip or foam (such as polypyrole) and reliessimply on the deformation of the polymer to vary theconducting cross section, thus changing the resis-tance. Piezoresistive sensors, in contrast, generallyconsist of a semi-insulating polymer matrix contain-

ing some type of conducting particulates, such asgraphite. Such compounds are macroscopicallypiezoresistive but also exhibit noise and hysteresisdue to the percolative nature of the conduction mecha-nism. Although some mildly piezoresistive polymershave been synthesized, practical force sensors com-prised of a homogeneous piezoresistive polymer arenot yet commercially available.

Rare-earth permanent magnets. The discovery ofnew rare-earth magnetic materials, such as NdFeB,has led to the development of compact electromag-netic actuators, which represent perhaps the mostattractive technology for force actuators requiringlarge displacements, as a possible substitute forhydraulic actuators.21 Using a geometry similar to thatof solenoid actuators, linear electromagnetic actuatorsmade from permanent magnets can produce forces ashigh as 2000 foot-pounds over a 24-inch stroke.Cleaner and less noisy than hydraulic actuators, com-pact powerful permanent magnets driven by magneticfields may also be employed as large-displacementdynamic elements to transmit sound or vibration.

Other materials. Some of the materials mentioned inthis paper will most likely be superseded by newclasses of force transduction materials that are cur-rently being synthesized or are yet to be discovered.In some cases, new process technologies will enablenew materials to be processed, such as the powdermetallurgy and rapid solidification process thatenabled the fabrication of amorphous metals. In othercases, the advent of powerful computers and theincreasing ability to engineer materials on the molec-

Table 1 Comparison of active materials

PZT5H

Max strain (! 10–6)

Max stress (MPa)

Density (kg/m3)

Modulus (GPa)

Efficiency (%)

Bandwidth (Hz)

300

19

7500

62

56

1000

Energy density (KJ/m3) 2.9

600

72

7500

120

75

1000

22

3000

180

~7500

~60

?

?

?

200

~1

1780

3

2

1000

~1

1800

90

9250

40

40

100

19

70000

190

6450

78

>3

3

>10

20000

180

~1500

5

>30

>1

>1

400000

0.3

~1300

<0.1

30

.1

0.4

400000

0.3

1037

.06

>35

4

0.8

PMN PLZT PVDF Terfinol-D SMA(NiTi)

ConductingPolymer

PolymerHydro-gel

HumanMuscle