Combinatorial development and discovery of
ternary and quaternary shape memory alloys
Dissertation
zur
Erlangung des Grades
Doktor-Ingenieur
der
Fakultt Maschinenbau
der Ruhr-Universitt Bochum
von
Robert Zarnetta
aus Karl-Marx-Stadt
(jetzt Chemnitz)
Bochum 2010
Dissertation eingereicht am: 23.06.2010
Tag der mndlichen Prfung: 22.07.2010
Erster Referent: Prof. Dr.-Ing. Alfred Ludwig
Zweiter Referent: Prof. Dr.-Ing. Gunther Eggeler
Fr meine Frau Christiane
und unsere Tochter Mariella Jocelin
The mere accumulation of facts, even an extremely extensive collection...does not constitute
scientific method; it provides neither a direction for further discoveries nor does it even deserve the
name of science in the higher sense of that word. The cathedral of science requires not only equip-
ment, but (needs) to indicate the pathway, by which the most fruitful new material might be
generated.
Dmitri Mendeleyev 19141
1 J. J. Tennenbaum, http://american_almanac.tripod.com/mendel94.htm
Executive Summary / Kurzfassung In der vorliegenden Arbeit wurden Methoden der kombinatorischen Materialforschung entwickelt und angewendet, um neue und/oder verbesserte ternre und quaternre Formgedchtnislegierungen zu erforschen. Fr die Herstellung sogenannter Dnnschicht-Materialbibliotheken wurde das Katodenzerstuben genutzt. Unter Verwendung von keilfrmigen Multilagenschichten konnten dabei komplette ternre Legierungssysteme oder groe Zusammensetzungsbereiche quaternrer Legierungssysteme abgeschieden werden. Zur Hochdurchsatz-Charakterisierung der Materialbibli-otheken wurden automatisierte Messverfahren fr die Bestimmung der Zusammensetzung (energiedispersive Rntgenanalyse), der Struktur (Rntgenbeugung) und des Phasenumwand-lungsverhaltens (temperaturabhngige Widerstandsmessung) eingesetzt. Fr die Charakterisierung des Formgedchtniseffektes dnner Schichten fand die sogenannte Biegebalken-Methode (tempe-raturabhngige nderung der Schichtspannung) Anwendung und wurde unter Nutzung mikrostruk-turierter Si-Biegebalken in Matrixanordnung zu einem parallelen Hochdurchsatz-Messverfahren weiterentwickelt. Mit den beschriebenen Methoden konnten in den Legierungssystemen Ti-Ni-Cu, Ti-Ni-Pd und Ti-Ni-Ag neue Formgedchtnislegierungen entdeckt und erstmals die vollstndigen Zusammensetzungsbereiche identifiziert werden, die eine reversible Phasenumwandlung zeigen. Das Ti-Ni-Ag System zeichnet sich dabei durch die Unlslichkeit von Silber in der Ti-Ni Phase aus, was zur Bildung von nanoskaligen Silberausscheidungen und zu herausragenden Formge-dchtniseigenschaften fhrt. Fr alle Legierungssysteme wurden die Zusammensetzungs-Struktur-Eigenschafts-Beziehungen bestimmt, die den Einfluss der Zusammensetzung auf die Phasenbil-dung und den Einfluss der verschiedenen Phasen auf die Phasenumwandlungseigenschaften, den Formgedchtniseffekt und die mechanischen Eigenschaften der Dnnschichten beschreiben. Diese ermglichten die Entwicklung verbesserter Formgedchtnis-Dnnschichten fr die Anwendung als Mikroaktoren.
Eine, der in den Dnnschichten neu entdeckten Legierungszusammensetzungen (Ti39Ni46Cu16) wurde als Massivmaterial hergestellt, um zu untersuchen, inwieweit die bertrag-barkeit der Ergebnisse von den Dnnschichten auf Massivmaterial gewhrleistet ist. Es wurden Hochskalierungseffekte entdeckt und der voraussagende Charakter der kombinatorischen Dnn-schichtexperimente fr die Entwicklung von Massivmaterialen besttigt. Basierend auf den Ergebnissen der ternren Legierungssysteme wurden im quaternren Legierungssystem Ti-Ni-Cu-Pd neue Legierungen gesucht, fr deren reversible Phasenumwand-lung eine verschwindend kleine thermische Hysterese theoretisch vorhergesagt worden war. Durch die gezielte Variation der Zusammensetzung und damit der Gitterparameter der Legierungen in den Dnnschicht-Materialbibliotheken konnten entsprechende Legierungen entdeckt und die Ergebnis-se anschlieend auf Massivmaterialien bertragen werden. Die quaternren Legierungen mit ver-schwindend kleiner thermischer Hysterese zeigten eine beispiellose funktionelle Stabilitt, die die der binren Ti-Ni und die der ternren Ti-Ni-Cu und Ti-Ni-Pd Legierungen deutlich bertreffen.
Contents 1 Introduction ............................................................................................................................... 1
1.1 Aims of the work .................................................................................................................. 3
2 Fundamentals ............................................................................................................................ 4
2.1 TiNi-based SMAs ................................................................................................................. 4 2.1.1 Phase transformation of TiNi-based SMAs .................................................................. 5 2.1.2 Thermal hysteresis of SMAs ......................................................................................... 8 2.1.3 Effect of third elements on transformation temperatures ............................................ 10
2.2 TiNi-based SMA thin films ................................................................................................ 12 2.2.1 SMA thin film deposition ............................................................................................ 12 2.2.2 SMA thin film stresses ................................................................................................ 15 2.2.3 Ageing effects of Ni-rich and Ti-rich Ti-Ni thin films ............................................... 17 2.2.4 Thickness effects of Ti-Ni thin films .......................................................................... 18 2.2.5 Functional fatigue of SMA thin films. ........................................................................ 19 2.2.6 Current and future developments ................................................................................ 20
2.3 Combinatorial materials science ........................................................................................ 21 2.3.1 Materials libraries ........................................................................................................ 21 2.3.2 High-throughput characterization ............................................................................... 22 2.3.3 Data analysis and visualization ................................................................................... 23 2.3.4 Current and future developments ................................................................................ 24
3 Methods .................................................................................................................................... 25
3.1 Combinatorial materials science methodology ............................................................... 25 3.1.1 Sputter deposition of SMA thin-film materials libraries ............................................. 25 3.1.2 Compositional analysis energy dispersive X-ray analysis (EDX) ........................... 30 3.1.3 Structural analysis X-ray diffraction methods (XRD) and temperature-dependent
X-ray diffraction (XRD(T)) ........................................................................................ 30 3.1.4 Phase transformation properties of thin films temperature-dependent resistance
measurements (R(T)) .................................................................................................. 31 3.1.5 Shape memory effect of SMA thin films cantilever deflection method .................. 33 3.1.6 Mechanical properties of substrate-attached thin films nanoindentation ................. 35
3.2 Preparation of bulk SMAs .................................................................................................. 36 3.2.1 Phase transformation properties of bulk SMAs differential scanning calorimeter
(DSC) and alternating current potential drop (ACPD) methods ................................. 36 3.2.2 Shape memory properties of bulk SMAs .................................................................... 37
3.3 Microstructural analysis Transmission electron microscopy (TEM) .............................. 37
4 Results and Discussion ............................................................................................................ 38
4.1 Phase transformation properties of Ti-Ni-Cu SMAs .......................................................... 38 4.1.1 The Ti-Ni-Cu system ................................................................................................... 38 4.1.2 Phase transformation characteristics of Ti-Ni-Cu thin films ...................................... 43 4.1.3 Structural properties of Ti-Ni-Cu thin films ............................................................... 47 4.1.4 Conclusions for Ti-Ni-Cu thin films fabricated using a multilayer approach ............ 53 4.1.5 Shape memory properties of Ti-Ni-Cu thin films - identification of optimized
compositions for high-frequency thin film microactuator applications ...................... 54 4.1.6 High-throughput measurements of mechanical properties for Ti-Ni-Cu thin films ... 59
4.2 Ti39Ni45Cu16 shape memory thin film and bulk alloys ....................................................... 67 4.2.1 Comparison of the phase transformation of thin film and bulk Ti39Ni45Cu16 ............. 67 4.2.2 Comparison of shape memory effects in thin film and bulk ....................................... 70 4.2.3 Microstructural investigation of bulk Ti39Ni45Cu16 ..................................................... 72 4.2.4 Transferability of thin film SMA compositions to bulk material ............................... 76
4.3 Phase transformation properties of Ti-Ni-Pd SMAs .......................................................... 77 4.3.1 Phase transformation characteristics for Ti-Ni-Pd thin films ..................................... 77 4.3.2 Structural properties of Ti-Ni-Pd thin films ................................................................ 81 4.3.3 Influence of precipitates on phase transformation temperatures and thermal hysteresis
of Ti-Ni-Pd thin films .................................................................................................. 82 4.3.4 Ti-Ni-Pd shape memory thin films in view of microactuator applications ................. 87
4.4 Identification of quaternary Ti-Ni-Cu-Pd SMAs with near zero thermal hysteresis and unprecedented functional stability ....................................................................................... 88
4.4.1 Geometric non-linear theory of martensite (GNLTM) ............................................... 89 4.4.2 Quaternary Ti-Ni-Cu-Pd thin film SMAs ................................................................... 89 4.4.3 Functional fatigue of quaternary Ti-Ni-Cu-Pd bulk SMAs ........................................ 94
4.5 Ti-Ni-Ag shape memory thin films .................................................................................... 98 4.5.1 Phase transformation characteristics for Ti-Ni-Ag thin films ..................................... 98 4.5.2 Structural characterization of Ti-Ni-Ag thin films .................................................... 101 4.5.3 Shape memory properties of Ti-Ni-Ag thin films ..................................................... 105 4.5.4 Effect of individual layer thicknesses of the multilayer Ti/Ni/Ag precursor thin films
on the phase transformation properties ..................................................................... 107 4.5.5 Internal stresses in Ti-Ni-Ag thin films proposed origin of negative hysteresis .... 110 4.5.6 Conclusions for Ti-Ni-Ag thin films fabricated using multilayer approach ............. 112
5 Conclusions and Outlook ..................................................................................................... 113
6 References .............................................................................................................................. 115
Publications ................................................................................................................................... 131
Acknowledgements ....................................................................................................................... 133
Curriculum Vitae .......................................................................................................................... 134
List of abbreviations & symbols
ACPD [-] alternating current potential drop Af [C] austenite finish temperature As [C] austenite start temperature at.% [-] atomic percent B2 [-] cubic austenite phase B19 [-] orthorhombic martensite phase B19 [-] monoclinic martensite phase CAW [-] cantilever array wafer CCS [-] continuous composition spread CTE [-] coefficient of thermal expansion DSC [-] differential scanning calorimetry [%] strain FIB [-] focused ion beam HAADF [-] high-angle annular dark field LN2 [-] liquid nitrogen Mf [C] monoclinic martensite finish temperature (B19) Ms [C] monoclinic martensite start temperature (B19) Of [C] orthorhombic martensite finish temperature (B19) Os [C] orthorhombic martensite start temperature (B19) SEM [-] scanning electron microscope SMA [-] shape memory alloy SME [-] shape memory effect STEM [-] scanning transmission electron microscope Rf [C] R-phase finish temperature Rs [C] R-phase start temperature RT [-] room temperature R(T) [-] temperature-dependent resistance [MPa] stress (T) [MPa] temperature-dependent thin film stress SAD [-] selected area diffraction TEM [-] transmission electron microscopy T [K] thermal hysteresis T [C] temperature TR [C] reference stress free temperature XRD [-] X-ray diffraction XRD(T) [-] temperature-dependent X-ray diffraction
1
1 Introduction
TiNi-Shape memory alloys (SMAs) are materials that after being deformed mechanically can re-
turn to their original shape upon heating. Thus, they seem to remember their original shape. By
doing so, these materials can induce large transformation stresses and strains, have the ability of
pseudoelasticity and show high damping capacity, good chemical resistance and biocompatibility.
Accordingly, shape memory alloys have attracted much attention as functional materials and for
medical applications and more recently for microelectromechanical system (MEMS).
Among different SMAs, Ti-Ni alloys exhibit the highest work output per unit volume,
which if employed as thin film microactuator exceeds that of other microactuation materials.2
Based on the large displacements and actuation forces possible, the low operation voltage and high
power density and in combination with the ease of creating friction free and non-vibrating move-
ment a number of thin film SMA applications have been developed and demonstrated. Fig. 1.1 and
Fig. 1.2a-c show micropumps47, -valves811, -grippers12,13, -wrappers14, -cages1, -mirrors15 and
Fig. 1.1 SMA thin film microactuators. (a) SEM picture of a microcage capturing a micro-polymer ball1; (b) sche-matic drawing of a SMA-actuated microgripper from Krulevitch et al.2; (c) stress-temperature behavior of Ti-Ni-Hf and Ti-Ni-Cu thin films sputtered on a (Mo+Nb) carrier. The nested hysteresis of the two SMAs allows switching the curved Ti-Ni-Hf/(Mo+Nb)/Ti-Ni-Cu composite upwards by the heat pulse T1T3T1 (1) and downwards by the heat pulse T1T4T1 (2); (d) and (e) pictures of the bi-stable actuator in downward and upward position, respectively.3
2
bi-stable actuators for tactile displays3. The compatibility with micro-fabrication techniques, i.e.
patterning using photolithography and selective chemical etching16 makes Ti-Ni-based thin films
attractive for MEMS applications. However, the limitations of SMA actuators, i.e. the low actua-
tion frequency still prevent a widespread use of SMA microactuators, despite the in this respect
positive scaling effects for thin films.
Additionally, the relatively high crystallization temperatures of Ti-Ni-based thin films limit the
number of suitable substrates. Only recently Ti-Ni-Cu SMA thin films were developed, that were
crystalline when deposited on heated polyimide substrates at 270 C, opening up a new field of
application.17
Next to micro-actuator applications, SMA thin films are of special interest in medical applica-
tions due to large obtainable strains, the constant stress level during pseudoelastic loading and their
good biocompatibility. In this respect, Ti-Ni microtubes are attractive materials for biomedical de-
vices, such as micro-catheters and micro-stents19, whereas superelastic TiNi-polymer-composites
could be used for novel applications in orthodontics and medical instrumentation.20,21 Fig. 1.2d-e
show SEM images of Ti-Ni microtubes fabricated by rotating-wire method, which allows to pro-
duce microtubes with uniform wall thickness and composition.19
Despite the considerable progress in the development of SMA and SMA thin films, no ground-
breaking actuation application has been commercialized, mainly due to unsolved material issues
that limit the temperature range of application, the actuation frequency and the repeatability (func-
tional stability) over a large number of cycles. Thus, a further improvement of SMA properties
based on an advanced understanding of the underlying composition-structure-property relation-
ships is needed to improve the material properties in order to facilitate the wide-spread use of
shape memory alloys for actuator applications.
Fig. 1.2 SMA thin film microactuators and microtubes for potential medical applications. (a) SEM image of a SMA actuated micromirror (top view)15, (b) microspring and (c) free-standing TiNiCu microtweezer structure18, (d) cross-section and (e) close-up SEM images of the fracture surface of Ti-Ni microtubes fabricated by rotating-wire method at 30 rpm. A potential starting material for thin film stents.19
3
1.1 Aims of the work
Within this work the aforementioned limitations of Ti-Ni SMAs will be addressed by investigating
ternary and quaternary SMA thin films using combinatorial materials science methods for the acce-
lerated development and discovery of improved SMAs and the underlying composition-structure-
property relationships. The following specific questions were aimed to be answered:
(I) Can ternary and quaternary materials libraries with well-defined composition gradients
suitable for investigating SMA properties be fabricated by sputter-deposition methods?
(II) Can temperature-dependent 4-point probe resistance and curvature measurements, as
well as nanoindentation be applied as high-throughput characterization tools for the charac-
terization of the phase transformation properties, the actuator response and mechanical
properties of SMA thin films, respectively?
(III) Do new, so far unknown SMA thin film compositions or composition-structure-
property relationships exist for ternary Ti-Ni-based alloys?
(IV) Do quaternary Ti-Ni-Cu-Pd compositions exist that show zero hysteresis, as post-
ulated by theory?
(V) Can composition-structure-property relationships determined for thin films be used to
predict bulk materials behavior?
Accordingly, the thesis is comprised of three main parts (I) introducing the fundamentals of SMA,
SMA thin films and combinatorial materials science, (II) the specific methods developed and used
in this work, followed by (III) the results and discussion. The results and discussion part is divided
into several sections, each related to a different ternary or quaternary alloy system, thin film or bulk
material, respectively. An overview and introduction to each alloy system will be given at the be-
ginning of the according section.
4
2 Fundamentals
2.1 TiNi-based SMAs
The shape memory effect (SME) is the ability of a material to recover its original shape after a de-
formation that extends well beyond the elastic regime. It appears in alloys showing reversible
thermoelastic martensitic transformations. The martensitic transformation is a diffusionless struc-
tural change associated with a large shear-like deformation. The martensite (low temperature
phase) is deformable and the deformation can be recovered perfectly by heating to the austenite
(high-temperature phase). The SME was first observed in the Au-Cd alloy system by Chang and
Read in 195122 and in 1963 the SME in Ti-Ni was discovered by Buehler et al..23
However, it was only at the start of the 1980s that the shape memory mechanism in the
Ti-Ni system was extensively investigated, due to difficulties with the fabrication of single crystals
and the complexity of the martensitic transformation.24 One of the difficulties being, the compara-
tively small equilibrium composition region for single phase TiNi at low temperatures (< 600 C:
~49.5 at.% Ni to ~50.5 at.% Ni), as shown in Fig. 2.1, which leads to the necessity of very precise
composition control. The comprehensive review on the physical metallurgy of Ti-Ni-based shape
memory alloys by Otsuka and Ren published in 2005 highlights the significant progress made in
understanding the martensitic transformations in this system in recent years.26
Fig. 2.1 Equilibrium phase diagram of the Ti-Ni system.25
5
2.1.1 Phase transformation of TiNi-based SMAs
The martensitic transformation is a first order diffusionless phase transformation. Thus, no change
in chemical composition occurs and a distinct phase boundary exists between the austenite phase
and the transformation product (martensite)28. The transformation starts upon cooling from the
high-temperature phase A (austenite) at the martensite-start-temperature (Ms). This temperature is
lower than the thermodynamic equilibrium temperature T0, since an additional driving force is
needed (GA,M) in order to compensate the nucleation energy of the martensite and the elastic
stress fields due to the formation of martensite in the austenite27 which results in an undercooling
(u), as illustrated in Fig. 2.2. The nucleation energy can be small or even negligible, if the auste-
nite and martensite phase are crystallographically compatible and thus the elastic stress field ener-
gy is minimized (see also 2.1.2).29
The mechanisms of the shape memory effect (temperature induced transformation) and superelas-
ticity (stress-induced transformation) are depicted in Fig. 2.3 schematically, using a two-
dimensional crystal lattice as a model. Fig. 2.3 shows that both superelasticity and the shape mem-
ory effect can occur in the same sample, and which occurs, depends upon the test temperature. The
austenite crystal structure (red) is completely transformed to the martensite (blue) upon cooling
below Mf (martensite finish temperature) and the martensite is characterized by the formation of
different variants (labeled A and B in Fig. 2.3), that have the same crystal structure but differ-
ent orientations (twinning related). Due to the self-accommodation mechanism of the martensite
variants the macroscopic shape of the material will not change.30 The variants in the thermally in-
duced martensite phase are twinning related and upon the application of an external load the twin-
ning planes are easily moved. The variants will reorient to accommodate the stress and the coales-
cence of various variants into a single variant will produce a transformation strain .31 The marten-
site phase at this temperature is stable, so the reoriented martensite phase is maintained. If the ma-
terial is heated above Af (austenite finish temperature), the reverse martensitic transformation will
be completed and the transformation strain will be recovered. This behavior is commonly termed
as the shape memory effect (SME).
Fig. 2.2 Schematic of the temperature dependence of the free enthalpy G. The enthalpy curves of the low-temperature phase A (austenite) and the high tempera-ture phase M (martensite) are shown. The phase transformation starts after an under-cooling (u) at the martensite start tem-perature Ms (after Hornbogen27).
6
For temperatures above Af, the austenite phase is stable, however on the application of an external
stress the martensite phase can form due to stress-induced martensitic transformation. The marten-
site phase in this case is unstable since the temperature is above Af, so upon the release of the ex-
ternal load the material will return to the austenite phase. During the deformation process the
stress-induced martensite can accommodate a transformation strain , and therefore it behaves as a
superelastic material. For the case of Ti-Ni the recovery strain can be as high as 10 %, which is
more than 10 times the elastic limit of normal alloys and metals.28
Three different transformation paths for the reversible thermoelastic martensitic transforma-
tion from the austenite (B2) to the monoclinic martensite (B19), to trigonal martensite (R-phase)
or orthorhombic martensite (B19) are known for Ti-Ni-based alloys, as shown in Fig. 2.4. For the
latter two, an additional successive transformation step to monoclinic martensite can be observed,
thus resulting in a two-step transformation.
Unique lattice correspondences between the austenite and martensite phases exist, such that the
reverse transformation path is restricted. Thus, the reversibility of the thermoelastic transformation
Fig. 2.4 Three martensitic transformation paths in Ti-Ni-based alloys. (Adapted from Otsuka and Ren26). Conti-nuous lines indicate a one-step transformation, the dashed line the possibility of a successive second transformation step, resulting in a two-step transformation.
Fig. 2.3 Mechanism of shape memory behavior. Shape Memory Effect (SME) left hand side, Superelasticity (SE) right hand side. The temperature scale is indicated at the top, Af is the finish temperature for the formation of austenite upon heating, Mf is the martensitic transformation finish temperature, at which the martensite phase has completely transformed upon cooling; is the transformation strain (see text for details). Arrows indicate a transformation or reo-rientation of martensite, due to heating (red), cooling (blue) or stress/load (black).
7
is guaranteed. The lattice correspondence of the B2/B19, B2/B19 and B2/R-phase are summarized
in Fig. 2.5. It is evident that the transformation strain is dependent on the crystal orientation. This is
seen most clearly for the B2R-phase transformation, where the R-phase forms by elongation
along any one of the [111] directions of the B2 structure, and thus the transformation strain along
the [111] direction is largest, whereas the transformation strain along [001] is nearly equal zero.32
In polycrystalline materials the orientation dependence is averaged, if no specific texture is present.
Due to the shear-like adeformation of the lattice during the reversible phase transformation, the
transformation temperatures of SMAs are influenced by stresses, and vice versa the plateau stress
during superelastic loading by temperature, according to the Clausius-Clapeyron equation:28,33,34
, (1) where is the uniaxial applied stress, S and H* are the transformation entropy and enthalpy per
unit volume, strain and T temperature26. The stress rate /T for B2B19 transformation in
Ti-Ni falls typically in a range between 4 20 MPa/K, whereas the R-phase shows a characteristi-
cally higher stress rates in the range of 30 70 MPa/K.35 In other words, the Rs temperature is less
sensitive to stress as compared to the Ms temperature, which in turn however can result in a change
of the transformation path from a two-step to a one-step transformation with increasing stress.36
The insensitivity of the Rs results from the significantly smaller lattice deformation of the R-phase
Fig. 2.5 Lattice correspondence of the martensites in Ti-Ni-based SMAs. (a) B2 austenite phase with a FCT cell delineated; (b) orthorhombic martensite B19, formed by shear/shuffle of the basal plane 110 B2 along 1 0 direc-tion; (c) monoclinic martensite B19, which is viewed as a B19 structure sheared by a non-basal shear 001 1 0 B2; (d) trigonal martensite R-phase formed by stretching the cubic austenite lattice along 111 diagonal direction. The axis a, b and c represent the principal axes in that deformation (from Otsuka and Ren26).
8
transformation, as compared to the B2B19 transformation. Fig. 2.6 shows typical stress-
temperature relations for free-standing thin films and bulk SMAs and depicts the increase of the
transformation temperatures with increasing stress.
However, for substrate-attached thin films, the Clausius-Clapeyron relationship is found to
be not applicable for describing the stress-dependence of the transformation temperatures during
reversible phase transformation or thermal hysteresis39,40, as discussed in more detail in 2.2.2.
2.1.2 Thermal hysteresis of SMAs
The reversibility of the structural phase transformation has a profound technological implication
for the application of SMAs and their fatigue life41 and thus understanding of the underlying me-
chanism govern reversibility, i.e. the thermal hysteresis of SMAs is essential. Currently, the most
widely known and commonly accepted explanations of hysteresis are the pinning of interfaces by
defects and thermal activation27,42. However, a close examination of the limited experimental data,
i.e. for alloys with different dislocation densities as a result of different quenching methods43 and
theoretical calculations of the energy barriers associated with nucleation (thermal activation) do not
seem to unambiguously support these ideas.29
More recently, a new theory on the origin of the reversibility of phase transformations has
emerged29,4447, suggesting that the growth of fully developed austenite/martensite interfaces is
responsible for determining the size of the hysteresis. Fig. 2.7 shows an optical micrograph and a
schematic of the microstructure of a twinned austenite/martensite interface, including an elastic
transition layer, which forms due to the crystallographic incompatibility of the two phases. An
energy barrier is associated with the developing austenite/martensite interface, due to stored elastic
Fig. 2.6 Stress-temperature relations for free-standing SMA thin films and bulk. (a) Comparison between the stress-transformation temperature relation for a solution treated, free-standing Ni-rich thin film and a Ti49.5Ni50.5 bulk sample37; (b) stresstransformation temperature relation of a free-standing Ti50.0Ni48.2Cu1.8 SMA thin film38.
9
energy and interfacial energy of the martensite twins, that needs to be overcome during the forward
and reverse transformation, thus giving rise to transformation hysteresis.29,41 Next to explaining the
fundamental cause of the transformation hysteresis the geometric non-linear theory of martensite
(GNLTM) predicts that the hysteresis can be drastically minimized by improving the crystallo-
graphic compatibility of the martensite and austenite phase.48,49
The theory specifies several conditions in order for a SMA to show near zero hysteresis (a
rigorous mathematical derivation can be found in the papers of Ball and James44,45,49). The first
condition, detU = 1, where U is the transformation stretch matrix, represents the condition of no
volume change. The second condition: 2 = 1, where 1 2 3 are the ordered eigenvalues (lat-
tice distortions) of the transformation stretch matrix U, represents the presence of an invariant plane between austenite and martensite, i.e. a perfectly coherent interface between both phases. At
such an interface, the energy contributions to the bulk energy through the usual elastic transition
layer or interfacial energy of fine arrays of twin bands is eliminated, thus leading to a decrease of
T. For the cubic (B2) to orthorhombic (B19) martensite transformation, the following six trans-
formation stretch matrices map the austenite lattice vectors to the martensite variant lattice vec-
tors45:
0 00 0
, 0 00 0
, 0 0 0
0 ,
0 0 0
0 ,
0 00
0 ,
0 00
0 , (2)
where / , /2 , /2 and a0 is the lattice parameter of the cubic austenite and a, b, c are the lattice parameters of orthorhombic martensite. U can be taken to be any one of these six matrices, since they all have the same eigenvalues. By measuring the lattice parameters of
the austenite and martensite phases the eigenvalues can be calculated. A clear correlation between
Fig. 2.7 Interface between austenite and martensite. (a) optical micrograph of an austenite/martensite interface in Cu69Al27.5Ni3.5 (from James and Zhang49; (b) schematic of the auste-nite/martensite interface. The different martensite variants are labeled A and B.
10
the middle eigenvalue 2 and was found for binary Ti-Ni50 and several ternary alloy systems,
i.e. Ti-Ni-X, X = Pt43, Pd41,43, Hf43, Au43, Cu41,43 in both thin film and bulk alloys as summarized
by Zhang et al.29, whereas the correlation between detU and was found to be weak.41 Addition-
ally, alloy compositions satisfying 2 = 1 were predicted in the quaternary alloy system
Ti-Ni-Cu-Pd and postulated to show close to zero .49
2.1.3 Effect of third elements on transformation temperatures
For the application of SMAs the transformation temperatures are next to the thermal hysteresis of
significant importance, since they determine the temperature range of application. The binary Ti-Ni
system shows a distinct composition dependence of the transformation temperatures for Ti-rich and
Ni-rich compositions. For the former, the transformation temperatures are constant, while for the
latter a strong decrease is observed with increasing Ni content, as shown in Fig. 2.8. Much effort
has been made to modify Ti-Ni shape memory alloys by adding various alloying elements to the
binary system, in order to reduce the strong composition dependence of the transformation temper-
atures and/or to increase them.
Fig. 2.9 summarizes the effect of third elements on transformation temperatures reported in
literature.26 Most alloying elements are found to lower the transformation temperatures (Fig. 2.9b).
Fig. 2.8 Influence of the Ni-concentration on the MS and T0 temperatures for binary Ti-Ni. In region I, MS and T0 (calculated equilibrium temperature; T0 = (Ms + Af) / 2) are constant. The concentration range of region I corresponds to the two-phase region, where Ti2Ni and TiNi are in equilibrium. Region II includes all chemical compositions which (after quenching) yield single phase alloys. In region II, T0 / MS decrease from 365/340 K at 50.0 at.% Ni to 227/211 K at 51.0 at.% Ni. Initially, this decrease is linear, with a slope of -83 C (at.% Ni)-1, followed by a stronger, non-linear decrease at higher Ni concentrations. From Frenzel et al..50
11
However, the addition of Pd, Pt, Au, Zr or Hf increases them (Fig. 2.9a) and thus Ti-Ni-X, X = Pd,
Pt, Au, Zr and Hf are considered as candidates for high-temperature SMAs.51
For future alloy design and the understanding of composition-structure-property relation-
ships the knowledge of the site preference of the alloying additions is of fundamental importance,
since the phase formation (e.g. precipitate phases) as well as the phase transformation characteris-
tics strongly depend on the substitution behavior. Experimental effort to determine the site occu-
pancy of alloying elements in Ti-Ni has been made by Nakata et al.52. They employed atom loca-
tion by channeling enhanced microanalysis (ALCHEMI) method to determine the site preference
of Cr, Mn, Fe, Co, Cu, Pd and found that Fe, Co and Pd preferentially substitute for Ni, while Mn,
Cr, and Cu seem to substitute for both Ti and Ni with similar preference. Nakata et al. summarized
Fig. 2.9 The effect of alloying ele-ments on the martensitic transforma-tion temperatures of Ti-Ni. (a) Pd, Pt and Hf in a wide alloying range; the circle indicates composition region covered in b; (b) Fe, Co, V, Mn, Au, Zr, Al, Pd, Pt, Hf and Cr in a narrow alloying range (low alloying level). From Otsuka and Ren.26 References for the individual sets of data can be found therein.
12
the following: (I) Fe, Co and Pd have strong preference for entering Ni-sites; (II) Sc has strong
preference for entering Ti-sites; (III) V, Cr, Mn, Cu and Au seem to have less preference for a par-
ticular site; their occupancy fractions are strongly affected by the way the alloying element is add-
ed. For example, if adding Cu in a way to give preference for the Ni-site, i.e. Ti50Ni50-xXx, it is con-
cluded that Cu occupies only Ni-sites; but if adding Cu giving preference for Ti-sites, i.e.
Ti50-xNi50Xx, there is only about 34% Cu entering into Ni-sites. This is quite different from the
well-accepted postulation that Cu only goes to Ni-sites.
Similar results are calculated by Sheng et al.53 and more recently for a wide range of ternary
additions by Bozzolo et al.5456 that strongly motivate the search for new and/or optimized shape
memory alloys in extended composition regions in ternary and quaternary alloy systems by combi-
natorial methods. Especially in the quaternary composition space, where alloy development using
conventional bulk preparation methods is costly and time-consuming and hence only a very limited
number of investigations are reported5763, the combinatorial thin film-based methods will be ad-
vantageous.
2.2 TiNi-based SMA thin films
Ti-Ni based thin films are the most frequently used thin film SMA materials, since they exhibit the
highest work output per unit volume, which if employed as thin film microactuator exceeds that of
other microactuation materials2, as mentioned above. Additionally, the positive scaling effect of
surface to volume enables fast cooling, thus significantly higher working frequencies can be rea-
lized, as compared to bulk SMA actuators. Several thin film deposition methods were used for the
fabrication of SMA thin films, including laser ablation64, ion beam deposition65, evaporation (mo-
lecular beam epitaxy)66 and sputter deposition67. The latter is most frequently applied and will be
discussed in more detail in the following section.
2.2.1 SMA thin film deposition
Sputter deposition of thin film SMAs was pioneered by Kim et al. in 198667 and in the 1990s sev-
eral groups demonstrated the feasibility of producing thin films exhibiting reversible phase trans-
formations using alloy targets.6872 In order to obtain transforming Ti-Ni films a precise control of
the composition is necessary. Fig. 2.10 shows a schematic of a magnetron sputtering setup and in-
dicates the processes involved. In order to avoid the characteristic loss of Ti, as the sputtering yield
of Ni is higher than for Ti, due to the different sputtering rates and angular distributions, the design
and composition of the sputtering target is essential.73
13
In previous work of Gyobu et al.74 and Quandt et al.72, the Ti deficiency was compensated by plac-
ing additional Ti pieces on top of the alloyed target or by using Ti-rich (54 at.% Ti) targets, respec-
tively. However, since in the case of magnetron sputtering from alloy targets the deposition rate is
additionally influenced by the depth of the sputter trench and thus changes with sputter duration2,
other sputter-deposition methods, namely co-deposition75 and multilayer deposition techniques7680
using elemental targets were developed, in order to control the thin film composition more precise-
ly. Moreover, the multilayer approach widely eliminates the influence of sputter power, sputtering
pressure and target-to-substrate distance usually found to strongly influence the resulting thin film
microstructure. Common to the described thin film deposition methods is the necessity of a subse-
quent annealing step. Since the Ti-Ni based alloys show a strong tendency to become amorphous
by sputter deposition at RT26, thin films sputtered from alloy targets or produced by co-deposition
need to be crystallized at elevated temperatures, as do multilayer thin films. Amorphization of the
alloys by sputtering is thereby an advantage for the applications of thin films, since the process of
amorphization and subsequent crystallization leads to small grain sizes, which are beneficial for the
mechanical properties.81 However, as-sputtered amorphous Ti-Ni thin films might contain excess
Ti or Ni atoms for Ti-rich of Ni-rich compositions, respectively, non-equilibrium phases may form
during the crystallization process, as will be discussed below.
Using heated substrates during the deposition of Ti-Ni based thin films the subsequent an-
nealing step can be avoided, since crystalline films can be obtained when deposition temperatures
above 250 C are used.17,65,8285
Fig. 2.10 Schematic of magnetron sput-tering. A magnetic field confines the orbits of the electrons to maintain an intense plasma and to increase the colli-sion rate with the Ar gas. The number of Ar ions created and impinging on the target, knocking out individual atoms, which are then deposited on the substrate, is enhanced. For sputter deposition a ultra-high vacuum system and high-grade Ar gas are beneficial, in order to prevent contamination of the thin films.
14
Fig. 2.11 shows a comparison between the crystallization behavior of an amorphous and a
multilayer Ti-Ni thin film (Fig. 2.11a) and depicts the alloying process of the latter
(Fig. 2.11b,e,g).80 The DSC curve during heating of the multilayer Ti-Ni thin film (Fig. 2.11a)
shows three exothermic peaks79, which correspond to (I) the formation of an amorphous phase at
each interface (interdiffusion of Ni into the Ti layer86, i.e. growth of the reaction layer (Fig. 2.11e),
(II) the formation of B2 phase (crystallization) and (III) the formation of precipitates by inter-
diffusion of Ti atoms into the Ti-Ni B2 from the residual Ti layers (Fig. 2.11g). The Ti2Ni precipi-
tates remain near the interface of the Ti-Ni B2 phase and no Ni4Ti3 precipitates were observed. In
contrast, for the as-sputtered amorphous Ti-Ni thin film only a single exothermic, crystallization
peak80 is observed (Fig. 2.11a) and the microstructure of Ti48.7Ni51.3 is known to consist of compa-
ratively larger grains with Ni4Ti3 precipitates.87
TEM images show the multilayer structure in the as-sputtered thin film (Fig. 2.11b), and the
Ti- (bright), Ni- (dark) and a thin reaction layer can be clearly seen. The corresponding selected
area diffraction patterns show the diffraction spots of Ti and Ni polycrystals (Fig. 2.11c,f,h) and
diffuse diffraction rings of the amorphous reaction layer in between (Fig. 2.11d).
Fig. 2.11 Alloying process of sputter-deposited Ti/Ni multilayer thin films. (a) DSC curves of the as-sputtered multilayer thin film and as-sputtered amorphous thin film79; (b), (e), (g) bright field TEM images and (c), (f), (h) the corresponding selected area diffraction pattern for as-sputtered Ti49.0Ni51.0, heated up to 640 K and 750 K multilayer thin films, respectively80; (d) nano-beam diffraction pattern of the interface layer between Ti and Ni layers.80
15
Additional insight into the alloy formation, i.e. amorphization reaction and recrystallization
behavior is provided by studies on Ti/Ni multilayers used for neutron optics components, such as
highly reflecting mirrors, polarizers, monochromators. Therein the thermal stability88, structural
and magnetic properties89,90 and the kinetics of alloy formation at the interface91 were investigated.
2.2.2 SMA thin film stresses
Stresses in SMA thin films can be intrinsic and extrinsic. The intrinsic stresses depend mainly on
the growth conditions (such as temperature, pressure, growth rates94,95) and thus play only a minor
role for SMA thin films in general and more specifically for the multilayer-deposition approach,
since a subsequent high-temperature annealing step is required for the crystallization of the films
leading to a relaxation of the intrinsic stresses.
Fig. 2.12 shows a schematic stress-temperature () curve of a constrained SMA thin film92
and an experimental () curve of a Ti-Ni-Cu thin film deposited from an alloy target onto an oxi-
dized Si substrate93, as determined using the cantilever deflection method (see 3.1.5). Upon anneal-
ing the compressive stresses of the as-deposited film initially increase and subsequently relax com-
pletely upon crystallization at high temperatures (Fig. 2.12b).82,93 Upon cooling, the extrinsic
stress, i.e. tensile stress rises linearly due to the different thermal expansion coefficients of SMA
film and substrate until, through the formation of martensite starting at Ms, the stress can be partial-
ly accommodated.
The slope of the linear part of the stress-temperature curve above Ms can be calculated us-
ing an estimate of the coefficient of thermal expansion (CTE) for TiNi (TiNi = 11.0 x 10-6/K) 96 and
of the Si substrate (Si = 2.5 x 10-6/K)97 and the following equation98 (the SiO2 diffusion barrier is
Fig. 2.12 SMA thin film stresses. (a) Schematic stress-temperature curve for transforming constrained SMA film. The evolution of the film microstructure and misfit is shown (austenite light grey; martensite black; substrate dark grey; reference stress-free state at TR) from Roytburd et al.92, (b) experimental stress-temperature curve of a 1.7 m Ti51.7Ni39.7Cu8.6 film on 180 m thick, thermally-oxidized Si (100) substrate after deposition. Heating and cooling rate was 5 K/min from Winzek and Quandt93.
16
neglected in this calculation):
/ / 1 , (3) where ETiNi is the Youngs modulus of TiNi austenite (~80 GPa)40; TiNi is the Poisson ratio of
Ti-Ni film (0.33)40. A slope of -1.03 MPa/K is calculated, which corresponds well to the experi-
mental data (see 4.1.6). According to (3), the absolute stress of a SMA thin film at the martensite
start temperature (Ms, Os, Rs) depends on the temperature interval between Ms (Os, Rs) and TR (ref-
erence stress free state, see Fig. 2.12a), on the thermal expansion coefficient as well as on the
Youngs modulus of the thin film. These values are likely to vary with composition in binary, ter-
nary or quaternary SMA thin films.
The effect of thermal stresses on the phase transformation properties of substrate-attached
(constrained) thin films was discussed in literature40,82,83,92,93,99,100, but up to now no consistent pic-
ture has emerged. Roytburd et al. calculated that the temperature interval of the two-phase equili-
brium is widened for constrained, single-crystal SMA thin films, which is supported by the expe-
rimental work of Liu and Huang38, but opposed by results reported by Winzek and Quandt93. The
latter find that the temperature interval of transformation remains constant, if constrained and free-
standing thin films are compared. Similarly, while the transformation temperatures are calculated
to shift to higher temperatures92 both, experimental verification93 and falsification38are reported in
literature.
With respect to the influence of thin film stresses on the thermal hysteresis, experimental
work indicates a significant reduction for constrained films38,93 as compared to free-standing thin
films.38,93 However, while the phase transformation properties of the constrained films were deter-
mined using the curvature method, the free-standing thin films were characterized using DSC mea-
surements. Thus, the observed differences may be due to different measurement methods and/or
different thin film thicknesses used.93 Ideally, the same measurement technique and the same sam-
ple dimensions (film thickness) should be used for comparison.
An additional controversy reported in literature is concerning a potential stress gradient in
the SMA thin films, which could influence the shape memory behavior. While Winzek101 calcu-
lates a homogenous stress distribution in substrate-attached thin films with a film to substrate
thickness ratio of < 0.1, both Grummon et al.40 and Liu et al.38 assume a stress gradient through the
film thickness based on the experimental observation that the intrinsic stresses in sputter-deposited
thin films decrease with increasing thickness100 and the assumption that a variation of shear stress
from a zero value at the free outer surface to a relative high value at the film/substrate interface
exists. However, the models based on such assumptions remain speculative, since no direct expe-
rimental observation of the stress gradients or the microstructural development during the phase
transformation in constrained films exists.
17
2.2.3 Ageing effects of Ni-rich and Ti-rich Ti-Ni thin films
Ageing treatments at elevated temperatures are an effective way of strengthening Ti-Ni thin films
by precipitate hardening, i.e. by Ni4Ti3 precipitates for Ni-rich compositions (Fig. 2.13a)36 or by
Ti2Ni precipitates for Ti-rich compositions102. Whereas the precipitation characteristics of Ni4Ti3
were found to be almost consistent with those reported in bulk samples36, non-equilibrium phases
were observed in Ti-rich thin films that are not observed in bulk alloys (Fig. 2.13b-d)102104. When
amorphous Ti51.8Ni48.2 thin films are crystallized at 500 C, the microstructure of the thin films
changes with annealing duration of 5 minutes, 1 hour and 10 hours in the following sequence: (I)
Guinier-Preston (GP) zones, (II) GP zones and Ti2Ni precipitates, (III) Ti2Ni precipitates, respec-
tively.102 The Ti2Ni precipitates within the TiNi grains have the same orientation as the TiNi ma-
trix, and the interface between both phases is partially coherent.
Fig. 2.13b-d show TEM images of a Ti-rich Ti-Ni thin film annealed at 500 C / 1 h. In
Fig. 2.13b the GP zones have a diameter ~17 nm, form along the 100 planes of the bcc (B2) phase, are completely coherent and have a disc / plate-like appearance.104,105 Additionally, spheri-
Fig. 2.13 Microstructure of aged Ti-Ni thin films. (a) TEM images of precipitates in Ti48.7Ni51.3 thin films aged at 500 C / 1 h after solution treatment at 700 C / 1 h36; (b) bright-field image of a Ti51.8Ni48.2 thin film annealed at 500 C / 1 h, (c) bright-field image taken in random orientation, and (d ) high-resolution TEM image in [100] orienta-tion.102 Precipitates and GP zones are indicated by solid arrows.
18
cal Ti2Ni precipitates can be observed and are clearly distinguishable, if the TEM image are taken
in random orientation, as shown in Fig. 2.13c. A high-resolution TEM image taken in [100]B2 di-
rection is presented in Fig. 2.13d and reveals the morphology of the disc-like GP zones and round
Ti2Ni precipitates in more detail. Owing to the existence of the GP zones, the critical stress for slip
in the parent phase is increased and, as a result, excellent shape memory properties are obtained for
Ti-rich Ti-Ni shape memory thin films.104 For annealing temperatures of 600 C and 700 C the GP
zones are lost and only Ti2Ni precipitates are observed, which tend to distribute at higher annealing
temperatures along the grain boundaries (the equivalent to the precipitation behavior reported in
bulk samples).106
Next to the observation of GP zones in Ti-rich Ti-Ni thin films, GP zones were also recent-
ly found in Ti-rich Ti-Ni-Cu thin films fabricated by co-deposition107109, whereas no GP zones
were found in annealed Ti-Ni multilayer thin films, where the precipitation behavior of Ti2Ni re-
sembles more closely the bulk-like behavior (Fig. 2.11a-d).
2.2.4 Thickness effects of Ti-Ni thin films
Another fundamental aspect for the application and investigation of SMA thin films is the critical
thickness needed to yield consistent properties, i.e. to be comparable to thicker films or even bulk
material. Several groups investigated the lower thickness boundary in sputtered Ti-Ni thin films for
shape memory application, i.e. the size effect on the martensitic transformation.110114 Fig. 2.14
shows TEM images of TiNi thin films with varying thicknesses and the grain size is indicated by
dashed lines. The surface oxide layer, as well as the affected zone (enriched in Ni) is highlighted in
Fig. 2.14e and the deduced strengthening mechanisms as proposed by Ishida and Sato are illu-
strated schematically in Fig. 2.14f.
Two kinds of resistances against deformations are considered: the constraints from neigh-
boring grains and from surface oxide layers. The former effect increases with increasing thickness,
whereas the latter increases with decreasing thickness.110 In addition, surface oxidation and the
formation of interdiffusion layers influence the transformation temperatures and the composition of
the transforming phase owing to the consumption of Ti.110 The constraints imposed by neighboring
grains was found to saturated, if the film thickness is greater than the average grain size.
From the above considerations two important conclusions can be drawn: (I) if the average
grain size and (II) the surface oxide layer thickness are decreased, the critical thickness boundary
can be significantly lowered. Thus, the multilayer deposition approach, yielding significantly
smaller grain sizes as compared to co-deposition or sputter deposition from alloy targets76,77 is
most suitable for investigating thin film shape memory properties, even for films with thicknesses
below 1 m, if the thin film samples are annealed under UHV conditions.
19
The final film thickness limit or grain size at which a reversible martensitic transformation is com-
pletely suppressed is found at < 100 nm111 for substrate-attached thin films and ~50 nm for nano-
crystallites in an amorphous matrix113.
2.2.5 Functional fatigue of SMA thin films.
Stability and fatigue of Ti-Ni thin films have always been concerns in the development of applica-
tions.2,87,93 The functional fatigue of Ti-Ni film is referred to the non-durability and deterioration of
the shape memory effect during cycling and results in changes of physical, mechanical and shape
memory properties115 due to irreversible processes, e.g. generation of dislocations116120 which
form during the martensitic phase transformation. Thus, the fatigue of thin films is influenced by
internal (alloy composition, lattice structure, precipitation, defects and film/substrate interface) and
additionally by external parameters (annealing treatment, applied stress, stress and strain rates,
heating/cooling rates). Consequently, from the scarce results published in literature no clear picture
Fig. 2.14 Thickness effect on shape memory behavior of Ti-Ni thin films. Cross-sectional TEM images of Ti50.0Ni50.0 thin films with thicknesses of (a) 7, (b) 5, (c) 2, (d) 1 m, respectively; (e) surface oxide layer and affected zone of Ti50.0Ni50.0 thin film after heat-treatment at 773 K for 300 s in a vacuum furnace (3 10-5 Pa) with infrared lamps (cross-sectional TEM image); (f) strengthening mechanisms of thin films.110
20
emerges regarding the different contributions of the aforementioned parameters on the fatigue of
SMA thin films.
However, an initial decrease in the recovery stress () is consistently reported, as shown in
Fig. 2.15a-b, which is stabilizing upon further cycling.2,87,93 The long term functional stability was
allegedly demonstrated by TiNi Alloy Companys laboratory, where a microvalve was successfully
tested for more than 50 million cycles (1 % deformation, 1 Hz).121
2.2.6 Current and future developments
The recent developments with respect to the fundamental understanding of SMA thin films and
their device applications were recently reviewed by Miyazaki et al.122. Future advances in the field
of SMA thin films will depend on scientific and technological progress with respect to improved
materials properties and improved manufacturing processes. From a thin film materials perspective
for actuator applications, higher recovery stresses, higher transformation temperatures, smaller
thermal hysteresis and a smaller temperature interval of transformation are desired in order to im-
prove the work output and frequency of operation. A clear opportunity to improve the shape mem-
ory effect lies thereby in targeting specific thin film textures using novel processing techniques in
order to optimize the attainable recoverable stress. Furthermore, ternary and quaternary alloying
additions are suitable for improving the phase transformation properties (thermal hysteresis, trans-
formation temperatures). From a manufacturing process point of view, the most challenging prob-
lems are related to keeping the exact stoichiometry of the thin films and to decrease the necessary
crystallization temperatures in order to decrease the thermal stresses. New thin film deposition
schemes based on multilayer deposition from elemental targets are promising in this respect.
Fig. 2.15 Fatigue of SMA thin films. (a) Ms and recovery stress (rec) as a function of thermal cycles for a Ti-Ni-Cu film. The strain during the test is ~0.2 %2; (b) recovery stress for Ti52.0Ni19.8Pd28.2 versus number of cycles N. The values originate from measurements with 3 K/min, the annealing cycles between were executed by pulses of electric current.93
21
2.3 Combinatorial materials science
Since 1995, when Xiang et al. coined the term combinatorial approach to materials discovery123,
numerous studies have demonstrated the applicability of the combinatorial methodology: synthesis
of materials libraries and their high-throughput characterization for the determination of phase dia-
grams124,125, for the discovery of new or optimized functional41,126,127, optical128 or catalyst mate-
rials129133 and polymers134. However, the principal idea of a multiple-sample concept for the
investigation of inorganic materials was already described by Kennedy et al.135 in 1965 or Hanak136
in 1970 and one of the earliest combinatorial approaches in material science can be traced back to
Thomas Edison and the year of 1878. At that time, Edison applied parallel and combinatorial me-
thods for the investigation of suitable filament materials for the incandescent lamp, as outlined in a
review by Schubert et al..134
2.3.1 Materials libraries
In general, two types of materials libraries can be used in the combinatorial approach to materials
discovery: discrete libraries sets of samples with individual compositions (Fig. 2.16a,b) or con-
tinuous libraries a single sample with a continuous compositional variation (Fig. 2.16c). While
the discrete approach allows rapid screening of large sets of different materials128, the latter is es-
pecially suited for the determination of composition-structure-property relationships.
The synthesis of continuous materials libraries of bulk materials, i.e. diffusion multiples,
and thin films so-called composition spreads were pioneered by Zhanpeng138 and Kennedy et
al.135, Miller et al.139 and Goldfarb et al.140, respectively.
Fig. 2.16 Material libraries, (a) A discrete 128-member BaCO3, Bi2O3, CaO, CuO, PbO, SrCO3, and Y2O3 library prior to sintering. Each site is 1 mm by 2 mm; the color of each is the natural color of reflected light from a white light source. BiSrCaCuO and YBaCuO superconducting films were identified.123 (b) Photograph of the as-deposited discrete quaternary library under ambient light used in the search for a blue photoluminescent composite material. The diversity of colors in the different sites stems from variations in film thicknesses and the optical indices of refraction.128 (c) Pho-tograph of an annealed quaternary continuous composition spread (CCS) for Si, Sn, Co and C.137
22
For the fabrication of thin film composition spreads, co-evaporation135, co-sputtering139 and more
recently, thin-film multilayer approaches140,141 were introduced. Co-deposition, e.g. the simultane-
ous sputtering of different materials, leads to an atomic mixing of the materials during deposition,
while the compositional variation (spread) is determined by the geometric arrangement of the de-
position sources and the deposition rates of the individual materials. Thus, the attainable composi-
tional variation is limited, as well as the variation of the composition gradients. In contrast, a
wedge-type multilayer approach, e.g. the sequential deposition of individual nanoscale, elemental
wedge-type layers requires a subsequent annealing step for the alloy formation; however, the com-
positional variation and the composition gradients can be adjusted in a broad range. Additionally,
the wedge-type multilayer approach can be extended in order to cover large continuous regions of
quaternary or pseudo-quaternary systems, as described in more detail in section 3.1.1. Due to the
sequential nature of the multilayer approach, the deposition is slower compared to the
co-deposition approach, thus the achievable film thicknesses are limited from a practical point of
view.
2.3.2 High-throughput characterization
For the high-throughput characterization of the materials libraries, various parallel and serial high-
throughput inspection methods are applied. In recent years, the characterization techniques for the
most fundamental properties, i.e. composition and structure energy or wavelength dispersive
X-ray analysis (EDX, WDX) and X-ray diffraction were improved significantly with respect to
their spatial resolution and measurement times, thus enabling their application to high-throughput
experimentation.
The development of more specialized characterization techniques for functional properties
of materials range from parallel optical methods for the characterization of hydrogen storage mate-
rials142, or photoluminescent materials128, to serial methods for the characterization of magnetic126,
thermoelectric143 or dielectric144 materials. For the latter, scanning techniques, i.e. scanning squid
microscopy126, a potential Seebeck microprobe143 or a scanning-tip microwave near-field micro-
scope144 were used, respectively. Furthermore, scanning characterization techniques have been
developed for the electrochemical screening of materials libraries, e.g. the scanning droplet cell145
or an alternating current scanning electrochemical microscope (AC-SECM).146
In general, the possibility to address specific locations on a materials library (x-y-z posi-
tioning system) combined with an automation of the measurement is fundamental for high-
throughput experimentation. However, equally important is the ability to design the materials li-
brary in order to match the composition gradients with the spatial resolution of the measurements.
Since the latter can vary significantly for different measurement methods, a technique capable of
23
controlling the composition gradients of the materials libraries is needed, as pointed out above and
outlined in more detail in section 3.1.1.
Additionally, to the scanning techniques, micro-electromechanical systems (MEMS) offer
powerful tools for the fabrication and processing of materials libraries as well as for accelerated
materials characterization on planar substrates such as Si wafers. MEMS can be used for parallel
materials processing, either as passive devices such as shadow mask structures, or as active devices
such as micro-hotplates.147,148 Microstructured wafers, which incorporate sensor or actuator struc-
tures such as electrodes149,150, cantilever arrays126,151,152 microtensile test devices153,154 or bulge test
structures155,156 can be used to probe materials properties in an efficient way. Furthermore, for the
characterization of thin film mechanical properties nanoindentation can be applied.157 However, the
limitations with respect to substrate effect and machine compliance when testing thin substrate-
attached materials libraries on the wafer scale need to be taken into account.158
Within this work, commercially available high-throughput characterization tools such as
automated EDX and XRD are used next to custom-made high-throughput test stands for the cha-
racterization of the temperature-dependent resistance and the stress change of thin-film materials
libraries. Nanoindentation at RT and elevated temperatures will also be used for the characteriza-
tion of the mechanical properties, as described in more detail in sections 3.1.2 to 3.1.6.
2.3.3 Data analysis and visualization
Data analysis and visualization capabilities are next to appropriate high-throughput characterization
techniques a fundamental requirement for the implementation of the combinatorial materials
science approach. Already in 1970, Hanak used computer data processing in tabular, graphical and
functional forms as part of his multiple-sample concept.136 However, he and his followers at that
time were lacking tools for data processing and analysis appropriate for hundreds of data sets. Only
later in the 1990s, computational power and commercially available software for the primary data
processing tasks: collecting characterization data from analysis instruments, automated visualiza-
tion of the raw data sets, and the ability to organize and share measurements results159, was widely
available and enabled the since-then ongoing success of the combinatorial approach.
Today, along with computer-based data acquisition and visualization using software, such
as LabVIEW, ORIGIN or MATLAB, respectively, also data mining techniques have been devel-
oped in order to facilitate the data analysis and will be used in this work. Two primary functions
are served by data mining techniques: pattern recognition and classification, both of which form
the foundations for understanding materials composition-structure-property relationships.160,161 As
an example, cluster analysis and principal component analysis (PCA) were recently implemented
within the XRDSuite software package162 in order to facilitate the arduous analysis of analyzing
24
large sets of XRD patterns, by sorting patterns by their similarity into discrete groups and subse-
quently deducing the representative basis X-ray patterns.163
2.3.4 Current and future developments
The recent developments and successes in the field of combinatorial solid-state chemistry of inor-
ganic material have been recently comprehensively reviewed by Zhao164 and Koinuma and Takeu-
chi165. Additionally, a number of books focusing on the synthesis of materials libraries166, their
high-throughput characterization167 or both168, and on data analysis and visualization169 were pub-
lished over the last years.
Future advances in the field of combinatorial materials science will depend on the devel-
opment of new characterization techniques for the high-throughput screening of materials libra-
ries.170 Thus, next to the development of sophisticated methods for the fabrication of materials li-
braries, the limitations of current high-throughput characterization tools need to be addressed.
Additional potential for the combinatorial approach lies in developing the ability to transfer
thin film composition spread results to bulk materials. While thin film composition spread tech-
niques were found to be especially useful for mapping composition-structure-property relationships
for thin films, equilibrium phase formation and structural properties (hardness, elasticity, creep)
were concluded to be more suitably investigated by bulk diffusion couples.171 Thus, new fields of
applications for the combinatorial thin film approach could emerge, if desired bulk properties could
be screened using thin film composition spreads and the results prove to be transferrable to the
bulk material.
Within this work, the latter concept will be pursued, next to the development of new sputter-
ing schemes for the fabrication of ternary and quaternary thin film composition spreads and the
development of new high-throughput characterization tools for the characterization of SMA thin
films, as outlined in the following section in more detail.
25
3 Methods
3.1 Combinatorial materials science methodology
The combinatorial materials science / high-throughput experimentation approach for the
development and discovery of SMA thin films, as well as other materials, is shown schematically
in Fig. 3.1. It comprises the deposition of thin-film materials libraries using special magnetron
sputter processes, the automated high-throughput characterization (screening) of these libraries,
and appropriate visualization and analysis of the data.
3.1.1 Sputter deposition of SMA thin-film materials libraries
The materials libraries, i.e. continuous composition spreads used in this work were fabricated by
magnetron sputtering using a dedicated ultra-high vacuum combinatorial sputtering system (CMS
600/400LIN, DCA, Finland) and 4-inch oxidized Si(100) wafers with 1.5 m thermal SiO2 as sub-
strates. The sputter system consists of a load-lock, distribution chamber, mask-storage chamber and
two sputtering chambers. The sputtering chamber used in these experiments is equipped with three
DC and three RF magnetron cathodes mounted on a moveable arm above a rotatable substrate
holder with integrated heater, as shown in Fig. 3.2. Additionally, four intermediate shutters placed
90 apart are located above the substrate, which can be moved independently and used to partially
shield the substrate or to create wedge-type films. The typical background pressure before sputter-
Fig. 3.1 Combinatorial materials science / high-throughput experimentation approach for the development of SMA thin films.
26
ing was lower than 2.0 x10-8 Torr. Elemental targets (purity: 99.99% or better) were used and pre-
sputtered in order to minimize contamination. The targets have a diameter of 4-inch to enable a
sufficient thickness uniformity of the deposited films on the 4-inch substrates. Deposition was done
with Ar (6N) at a pressure of 5 mTorr, a target-to-substrate distance of 87.5 mm and without inten-
tional heating. The wedge-shaped thickness profiles of the deposited material were created by one
of the intermediate shutters, which was set to shield the substrate and then slowly retracted during
the deposition (speed: 1 - 5 mm/s). Alternating wedge-type thin films of Ti, Ni, and X (X = Cu, Pd,
Ag) were consecutively deposited by switching the position of the targets above the substrate ac-
cordingly. In all cases the first layer was Ti in order to allow for good adhesion to the substrate and
the surface layer was the ternary element in order to prevent Ti surface oxidation during annealing.
Binary composition spreads were fabricated using opposing wedge-type films, realized by
rotating the substrate by 180 or the use of two opposing shutters. Due to the degrees of freedom
associated with the wedge-type multilayer approach, both compositional variation and gradients
can be adjusted in a broad range by selection of appropriate sputter powers and shutter speeds, re-
spectively. The depositions were controlled by a recipe with an internal loop repeating the deposi-
tion of each wedge-type film numerous times, in order to build up the chosen total film thickness.
Fabrication of pseudo-binary composition spreads, compromised of three elements, with
one kept constant, were deposited using a sequence of opposing wedge-type films of two elements
and intermediate layers of homogeneous thickness of the third element. The thickness variation in
Fig. 3.2 Ultra-high vacuum combinatorial magnetron sputter system for the fabrication of binary, ternary and qua-ternary thin-film materials libraries using wedge-type multilayer thin films.
27
Fig. 3.3 Schematic of the wedge-type multilayer approach for the fabrication of ternary materials libraries. (a) individual layers are rotated relative to each other by 120 and layer thicknesses of up to 10 30 nm were used (adapted from Goldfarb et al.140). (b) In order to cover the complete ternary system the wedge-type films are limited in length (70 mm). Black circle in b represents the contour of the 4-inch Si(100) substrate. (An animated movie of the sputter deposition scheme for the fabrication of ternary composition spreads is available at www.rub.de/wdm.)
the homogeneously deposited film was found to be less than 2 % within a radius of 2 cm around
the center of the substrate. Apertures were used to increase the homogeneity, but at a cost of signif-
icantly lower deposition rates.
Ternary composition spreads were realized by rotating the substrate by 120 each time be-
fore the deposition of the next material (Fig. 3.3a). For coverage of the complete ternary system,
the movement of the shutter was stopped at an intermediate position, and thus the length of the
wedge-type layer was intentionally limited. The shutter speed for all depositions was set to 1 mm/s.
By using a combination of such limited wedges, rotated by 120 relative to each other, the com-
plete ternary and binary systems can be covered in one experiment on a single substrate, as illu-
strated in Fig. 3.3b. For zoomed-in ternary composition spreads focused around a certain com-
position region, e.g. Ti50Ni50 wedge-type films extending over the full length were used. Thereby
the compositional variation and gradients can be adjusted by variation of sputter power and shutter
speed, respectively. Selection of appropriate deposition parameters was facilitated by calculating
the resulting compositions at selected positions prior to the deposition. The calculations were based
on the measured sputter rates of the individual elements and the deposition time at a given location
on the 4-inch substrate. The total deposition time is the sum of the exposure time related to the
movement of the shutter calculated using the location and the shutter speed and/or from a (sub-
sequent) static deposition.
28
Quaternary composition spreads were realized using the multilayer approach of alternating
wedge-type thin films. Sets of opposing wedges were used, with each set being a combination of a
single wedge and a sequence of three successive, opposing wedges, as illustrated in Fig. 3.4a. The
second set of opposing wedge-type thin films for the remaining two materials was rotated by 90
relative to the first set. For the single wedge-shaped thickness profile ( 30 mm around the sub-
strate center) one of the intermediate shutters was set to shield the substrate and then slowly re-
tracted (3 mm/s) during the deposition. The sequence of three successive wedges was realized by
updating the position of the paired shutters, so that one of them shielded the already covered areas,
whereas the other creates 20 mm wedges at 1 mm/s shutter speed.
Due to the approach taken, two of the three degrees of freedom for adjusting and control-
ling the composition range covered by the composition spread are fixed, namely shutter speed and
deposition time. Therefore, only the sputter power remains for controlling the composition range
covered within a deposition. The coverage realized using equal sputter rates of all four elements is
shown in Fig. 3.4b-c (calculated).137 The nine regions created on the sample (Fig. 3.4a) translate
each to a plane in quaternary composition space. The diagonal regions lie in the same plane: region
3 is contained in 5, which is contained in region 7. The seven distinct planes are almost parallel and
are separated by 7 at.% to 10 at.% in component A.
In Fig. 3.5 the measured compositional variations in a Ti-Ni-Fe-Au quaternary system are
shown, where all sputtering rates were chosen to be equal (Fig. 3.5a,b) or to have a specific ratio of
sputter powers, i.e. 2:2:1:1 for Ti, Ni, Fe, Au (Fig. 3.5c), respectively. Using equal sputter rates for
all elements, the coverage in the quaternary system can be maximized, however it will be concen-
trated in the center of the tetrahedron. A selection of a specific ratio of sputter powers can be used
to shift and confine the covered composition region in certain parts of the quaternary composition
Fig. 3.4 Scheme of the wedge-type multilayer approach for the fabrication of quaternary materials libraries. (a) Sets of opposing wedges were used, whereas each set was a combination of a single wedge and a sequence of three successive and opposing wedges (adapted from Chevrier and Dahn137). (b) Composition space covered within a qua-ternary alloy system (calculated). The two tetrahedrons are different views of the same data set.137
29
space, here onto the Ti-Ni axis. In order to focus on a single compositional plane out of the quater-
nary composition space, a combination of the wedge-type multilayers approach for the fabrication
of a ternary composition spread with intermediate homogenous layers of the fourth element was
used. An example of the achievable compositional variation is shown for Ti-Ni-Cu-Pd alloys in
section 4.4, Fig. 4.40.
In situ annealing of the multilayer composition spreads at temperatures of 500 C, 600 C
or 700 C for 1 h (heating rate 50 C/min, temperature stability 2 C) led to alloying of the ele-
mental multilayers via an amorphization reaction and complete recrystallization, as revealed by
XRD and TEM observations.7680
Fig. 3.6 shows images of an annealed binary, a complete ternary and a quaternary thin film
composition spread, that nicely compare with the schematic in Fig. 3.3 and Fig. 3.4. Thus, methods
for the fabrication of a broad range of materials libraries based on the sputter deposition of wedge-
type thin films were developed and demonstrated.
Fig. 3.6 Continuous thin film composition spreads. (a) binary composition spread the 301 point measurement grid is indicated (x, y = 4.5 mm), (b) complete Ti-Ni-Cu ternary composition spread, (c) Ti-Ni-Fe-Au quaternary com-position spread. White dashed lines in b and c indicate the ternary composition space (Fig. 3.3) and the 9 distinct re-gions in the quaternary composition spread (Fig. 3.4), repectively.
Fig. 3.5 Covered composition space within the quaternary alloy system Ti-Ni-Fe-Au. The two tetrahedra in (a) and (b) are different views of the same composition data determined by EDX, (sputtering rates all equal). In (a) the rotation of the tetrahedron is chosen to highlight the planes corresponding to the different regions of the composition spread (Fig. 3.4). (c) The relative sputtering rate ratio was 2:2:1:1 for Ti, Ni, Fe, Au, respectively.
30
3.1.2 Compositional analysis energy dispersive X-ray analysis (EDX)
The compositions of thin-film materials libraries and bulk samples were characterized by EDX
(Oxford INCA, Si:Li detector, LEO 1430 VP) using bulk standards, i.e. Ti50Ni40Cu10, Ti50Ni40Pd10,
for calibration of the instrument prior to each measurement. Thus, an accuracy of 0.5 at.% could
be achieved for the determination of thin film compositions using measuring times of 90 to 120 s
per measurement point. A typical measurement grid of 301 points with an x, y spacing of 4.5 mm
is shown in Fig. 3.6a. and at each point the EDX signal was integrated over an area of 400 m by
600 m. For the EDX characterization of bulk samples, measurement times of up to 1 h were used,
in order to increase the accuracy to 0.2 at.%.
3.1.3 Structural analysis X-ray diffraction methods (XRD) and temperature-dependent X-ray diffraction (XRD(T))
Structural properties of thin film and bulk samples were characterized by XRD at room tempera-
ture (RT) using a Bruker AXS D8 Discover (with GADDS, CuK radiation, spot size < 1 mm, in-
tegration time 600 s, area detector 2 range from ~26 to ~57), PANalytical XPert PRO MPD
(Pixel detector, CuK radiation, mono-capillary 0.8 mm, integration time 600 s, 2-range: ~30
to ~110) and PANalytical XPert PRO MRD (X'Celerator detector, CuK radiation, spot size
2 mm2, integration time 1800 s, 2-range: ~30 to ~110). The systems were calibrated using an
Al2O3 standard (Bruker) or Si standard (PANalytical), respectively. For the temperature-dependent
XRD measurements heating/cooling stages were used: Anton Paar TTK 450 (temperature range:
-100 C to 450 C, vacuum 0.75 mTorr), Anton Paar DHS 900 (temperature range: RT up to
900 C, Ar atmosphere, 150 Torr overpressure, flow 0.5 l/min) and a custom-designed Peltier stage
(temperature range: 5 C to 90 C, atmosphere). In order to monitor the reversible phase transfor-
Fig. 3.7 EDX analysis of Ti-Ni-Cu continuous thin film composition spreads. (a) distribution of Ni over the wafer, (b) Cu, (c) Ti. The arrows indicate the gradient direction of the wedge-type films from the thin to the thick end. A 301 point measurement grid (x, y = 4.5 mm) as shown in Fig. 3.6 was used for the automated compositional analysis.
31
mation upon heating and cooling by XRD, measurements were performed with a temperature step
size of 5 K.
Thin film lattice parameters were determined using synchrotron-based X-ray microdiffrac-
tion at the 2-BM beam line at the Advanced Photon Source at Argonne National Laboratory. For
this, the wafer was cut into 301 squares (4.5 x 4.5 mm) consistent with the analysis by EDX. Dif-
fraction measurements were performed at 110 C (i.e. austenitic state) and at -20 C (i.e. martensit-
ic state). Measurement times of 200 s per spot were used to obtain sufficient diffracted intensity for
a complete lattice parameter analysis of the thin film samples using an image-plate detector
(MAR 345). The beam size was focused to 15 x 15 m2 using a set of 30 Be compound refractive
lenses and the photon energy was set to 15 keV. The system was calibrated using a CeO2 standard
(National Institute of Standards and Technology). Lattice parameters were extracted from inte-
grated diffraction patterns using Fit2d software (http://www.esrf.eu/computing/scientific/FIT2D/).
Visualization and data management of the diffraction data for ternary alloy systems were
realized using the MATLAB-based XRDsuite software package.162 Implemented methods for the
identification of similar structural phases using cluster analysis163 and non-negative matrix factori-
zation172 were used to reduce the complexity of the data. For quaternary alloy systems the Quater-
naryViewer software package137 was used for visualization of structural data and functional prop-
erties.
For the identification of the detected phases, comparisons were made to databases of known
phases: inorganic crystal structures database (ICSD), Pauling Files binaries and Pearson's Crystal
database. Additionally, the CaRIne software package was used in order to calculate XRD pat-
terns of crystal structures reported in literature, but not recorded in the databases.
3.1.4 Phase transformation properties of thin films temperature-dependent resis-tance measurements (R(T))
The R(T) method for the characterization of the phase transformation of SMAs is based on a
change in the resistivity due to changes in the crystal lattice and number of internal interfaces (lat-
tice imperfections) and is well established for bulk173175 and thin film112,176 SMAs. DSC measure-
ments, commonly applied for the characterization of bulk materials, are likewise suitable for the
characterization of the phase transformation temperatures of thin films and are generally concluded
to agree well with the R(T) measurements.176 However, the electrical resistance measurements
were found to be more sensitive with respect to the appearance of the R-phase and the identifica-
tion of successive transformation steps.175 Additionally, the R(T) measurements can be conducted
locally, while a standard DSC measurement requires a certain amount of free-standing material (a
few mg), and thus is not suitable for localized screening of substrate-attached thin-film materials
32
libraries. However, recently introduced parallel nano-differential scanning calorimeter (PnDSC) a
micro-machined array of calorimetric cells177 could facilitate the use of calorimetry for the high-
throughput characterization of thin films, as demonstrated for Ti-Ni-Zr shape memory alloy thin
films178, and provide additional insight into crystallization kinetics and activation energies.179
R(T) measurements in the temperature range from -40 C to 250 C (heating/cooling rate
5 K/min) were made using a custom-designed, automated 4-point probe test stand (Fig. 3.8).180 In
general, two different measuring modes, i.e. screening and single mode measurements were per-
formed. For the screening of the thin film SMA materials libraries the 4-point probe was automati-
cally positioned to predefined loca
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