Electrochimica Acta 51 (2006) 4250–4254

Nanostructured Fe–Pd thin films for thermoelastic shape memoryalloys—electrochemical preparation and characterization

Feng Wang a,∗, Sayaka Doi a, Kaori Hosoiri a, Hirohisa Yoshida a,Toshio Kuzushima b, Masao Sasadaira b, Torhu Watanabe a

a Department of Applied Chemistry, Graduate School of Engineering, Tokyo Metropolitan University,1-1 Minami-ohsawa, Hachioji-shi, Tokyo 192-0397, Japan

b Electroplating Engineers of Japan Ltd., 5-50 Shinmachi, Hiratsuka, Kanagawa 254-0076, Japan

Received 4 August 2005; received in revised form 28 November 2005; accepted 30 November 2005Available online 26 January 2006

Abstract

Nanostructured Fe–Pd thin films with about 30 at.% Pd have been successfully synthesized on the substrates of Pt buffer layer/Cr seed layer/Siby an electrodeposition process from a plating bath containing ammonium tartrate, citric acid and ammonia solution as complexing agents. Resultscfutfi©

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learly show that the as-deposited films with body-centered cubic structure were transformed into face-centered cubic structure by heating at 900 ◦Cor 45 min followed by quenching into iced water. The in situ X-ray diffraction analysis results indicate that the quenched film with 29.8 at.% Pdndergoes a reversible thermoelastic austenite-to-martensite transformation with a narrow temperature hysteresis and a martensite transformationemperature of about −30 ◦C. The present study demonstrates the effectiveness of electrodeposition for synthesizing nanostructured Fe–Pd thinlms for the application of low-temperature-type thermoelastic shape memory alloys.2005 Elsevier Ltd. All rights reserved.

eywords: Iron–palladium; Thin films; Electrodeposition; Martensitic phase transformation

. Introduction

Shape memory alloys (SMAs) as the intelligent/smart mate-ials have attracted considerable attention for their applicationsf sensors and microactuators, which can be controlled either bylternating temperature in a certain range or by alternating inten-ity of an external magnetic field at constant temperature [1]. Thehape memory effect in the thermoelastic SMAs is related to theustenite-to-martensite reversible thermoelastic transformationhat is driven by a thermal field. The bulk thermoelastic SMAs,owever, are not suitable for the rapid actuation because theesponse speed of actuators is significantly limited by the heatonduction of materials [2]. One of way to get rid of this dis-dvantage is to fabricate the SMAs nanostructured thin films. Itas been proven that the actuation response speed in thin films

∗ Corresponding author at: Department of Chemistry and Material Engineer-ng, Faculty of Engineering, Shinshu University, 4-17-1 Wakasato, Nagano-shi,agano 380-8553, Japan. Tel.: +81 26 269 5670; fax: +81 26 269 5667.

E-mail address: wangf@gipwc.shinshu-u.ac.jp (F. Wang).

can be improved from bulk form by two orders of magnitudes[3].

Fe–Pd alloy films with about 30 at.% Pd are considered tobe highly promising SMAs due to their strain reversibility witha narrow hysteresis and high ductility of the alloy films [4].Numerous methods, such as magnetron sputtering [5], electro-magnetic melt-spinning [6], and arc-melting methods [7–9] havebeen employed to prepare Fe–Pd thin films. Although an elec-trodeposition process has been proven to be a feasible andeffective electrochemical means for synthesizing various func-tional nanostructured materials [10], however, there is only afew research works on Fe–Pd thin films [11] by the electrode-position method because of the great difficulty in preparingstable and effective deposition electrolytes. In particular, to thebest of our knowledge, there has no report on the thermoelasticbehavior of Fe–Pd thin films prepared by the electrodepositiontechnique.

In the present study, we describe an effective and reliablemethod based on the electrodeposition method that is suit-able for the preparation of nanostructured Fe–Pd thin films.Firstly, we have developed a novel plating bath and appropriate

013-4686/$ – see front matter © 2005 Elsevier Ltd. All rights reserved.oi:10.1016/j.electacta.2005.11.044

F. Wang et al. / Electrochimica Acta 51 (2006) 4250–4254 4251

operating parameters for fabricating the Fe–Pd thin films withabout 30 at.% Pd. Subsequently, by X-ray diffractometer (XRD),scanning electron microscope (SEM) and high-resolution trans-mission microscope (HR-TEM), the microstructures and thethermoelastic behaviors of the films have been investigated indetail.

2. Experimental

The plating bath consisted of 0.1875 mol l−1 FeSO4·7H2Oand 0.0125 mol l−1 PdCl2·6H2O as the source of metal ions(Fe2+/Pd2+ = 15:1). In addition to the metal salts, 0.375 mol l−1

(NH4)2C4H4O6, 0.0625 mol l−1 C6H8O7·H2O and 0.30 mol l−1

NH3·H2O were employed as a complex agents, while 0.6 mol l−1

NH4Cl was added as the supporting electrolyte. The pH of plat-ing bath was finally adjusted to 8.0 with dilute NH3·H2O. TheFe–Pd thin films were electrodeposited on Pt buffer layers with30 nm thickness, which were grown on Cr-seeded Si substrates.The buffer/seed layers were deposited by the sputtering method.The anode material was a titanium grid coated with platinum.The current density was ranged from 8.0 to 30 mA cm−2, and allthe deposition processes were performed at 25 ◦C under the totalelectric charge of 10 C cm−2. The plating cell was a beaker of200 cm3 with magnetic stirrer agitation. Finally, the as-depositedfilms were sealed in a quartz tube, and heated at 900 ◦C for45 min followed by quenching into iced water.

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Fig. 1. Voltammetric curves for deposition of Fe and Pd: (a) 0.1875 mol l−1

Fe2+; (b) 0.1875 mol l−1 Fe2+ + 0.375 mol l−1 (NH4)2C4H4O6 + NH3·H2O;(c) 0.0125 mol l−1 Pd2+; (d) 0.0125 mol l−1 Pd2+ + 0.0625 mol l−1

C6H8O7·H2O + NH3·H2O; (e) (b) + (d). Supporting electrolyte: 0.6 mol l−1

NH4Cl. Scan rate: 2 mV s−1.

voltammetric curves for Fe–Pd electrodeposition at the temper-ature of 25 ◦C. The deposition potentials of Fe2+ and Pd2+ werecompared with those of the solutions without complexing agent.As can be seen, the deposition potentials of Fe2+ and Pd2+ areshifted toward more negative potential by the complex formationwith ammonium tartrate, citric acid monohydrate and ammoniasolution, and the potential shift of Pd2+ by complex formation issignificantly larger than that of Fe2+. Deposition of the Pd2+ ismarkedly inhibited by the formation of [Pd(NH3)4]2+ complexions between Pd2+ and NH3·H2O since the complex ions cannotbe reduced as easily as the free Pd2+. It is evident that the additionof complexing agents in the plating baths causes the depositionpotentials of Fe2+ and Pd2+ shift to close with each other, thuspromoting the co-deposition of Fe2+ and Pd2+. Moreover, it isobserved that the deposition of Pd2+ in plating bath contain-ing complexing agents starts at about −0.65 V versus Ag/AgCl,while that of Fe2+ starts at about −1.1 V versus Ag/AgCl. Theco-deposition of Fe2+ and Pd2+ starts at about −1.1 V versusAg/AgCl. This indicates that for the co-deposition of Fe2+ andPd2+under galvanostatic condition, the cathodic current densityshould be controlled to be larger than 5.0 mA cm−2. The voltam-metric measurement results demonstrate the effectiveness of theaddition of complexing agents in the plating baths for achiev-ing the co-deposition of iron and palladium from the aqueoussolutions.

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Polarization behaviors of the plating baths were measuredsing an automatic polarization system (Hz-3000, Hokutoenko) with an Ag/AgCl electrode as the reference electrode.ll curves were measured at the scanning rate of 2 mV s−1.he chemical compositions of thin films were measured by

he energy dispersive X-ray analysis (EDX) (JED-2001, JEOL)erformed in a JOEL JEM-6100 scanning electron microscope.he structures of the deposited films were determined by X-rayiffractometer (XRD) (M21X TXJ-FO88, MAC Science) andhigh-resolution transmission microscope (HR-TEM) (JEM-

000FE, JEOL). To characterize the thermoelastic behaviorsf the quenched Fe–Pd films, in situ observation of crystaltructures was performed on the thermal cycling between roomemperature (RT) and −160 ◦C using the X-ray diffraction tech-ique. Magnetization measurements were performed with aibrating sample magneto-meter (VSM-5-18, Toei).

. Results and discussion

.1. Potentiodynamic studies of Fe–Pd electrodeposition

In Fe–Pd, electrodeposition system, it is difficult to obtaine–Pd thin films from the plating baths only containing sim-le metal ions since the standard electrode potentials of ironnd palladium differ by approximately 1.4 V. In order to bringhe potentials of iron and palladium close enough to achievehe co-deposition, ammonium tartrate, citric acid monohydratend ammonia solution were used to form complexes with metalons in the plating baths. All the plating baths prepared in theresent study showed superior electrochemical stability, therebyllowing the preparation of Fe–Pd thin films. Fig. 1 shows

.2. Preparation and characterization of Fe–Pd thin films

Fig. 2 shows the current density-dependence of Pd contentn the Fe–Pd films. A decrease of the Pd content in deposit waschieved by an increase of cathodic current density. This sug-

4252 F. Wang et al. / Electrochimica Acta 51 (2006) 4250–4254

Fig. 2. Influence of cathodic current density on the Fe and Pd contents in theFe–Pd thin films.

gests that the composition of Fe–Pd thin films can be controlledby adjusting the cathodic current density, showing a controllableelectrochemical feature.

Fig. 3 shows the representative HR-TEM image of as-deposited Fe–31.3 at.% Pd thin film and the correspondingselected area diffraction pattern (SADP). The TEM image showsa fine-grained structure in the film, and the average grain size is

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Fig. 4. X-ray diffraction patterns and SEM images of (a) as-deposited and (b)quenched Fe–29.8 at.% Pd film on the Pt/Cr/Si substrate. The as-deposited filmsare heat-treated at 900 ◦C for 45 min followed by quenching into iced water.

estimated to be around 15 nm, showing a nanostructured feature.The SADP clearly reveals that the as-deposited film consists ofrandomly oriented fine grains with typical body-centered cubicstructure (bcc).

According to their thermal equilibrium phase diagram [12],the Fe–Pd thin films with about 30 at.% Pd exist as face-centeredcubic structure (fcc, austenite) above approximately 770 ◦C.Thus it can be considered that the as-deposited films with bccsolid solution structures can be transformed into fcc austeniteby rapid quenching to avoid the separation of �-Fe and orderedFePd. Fig. 4 illustrates the typical X-ray diffraction patterns andSEM images of (a) as-deposited and (b) quenched Fe–29.8 at.%Pd thin films on the Pt/Cr/Si substrates (heating at 900 ◦C for45 min followed by quenching into iced water). It should bementioned that three peaks at 2θ = 33◦, 62◦ and 69◦ are due tothe Si substrate, while one peak at 2θ = 39◦ in the as-depositedfilm is due to the Pt buffer layer. Two peaks corresponding tothe (1 1 0) and (2 1 1) planes of bcc (�-Fe, Pd) solid solution areclearly observed in the diffraction pattern of the as-depositedfilm, while four peaks corresponding to the (1 1 1), (2 0 0), (2 2 0)and (3 1 1) of fcc (�-Fe, Pd) solid solution are observed in the

ig. 3. HR-TEM image of as-deposited Fe–31.3 at.% Pd film and the corre-ponding selected area diffraction pattern (SADP).

quenched film. These suggest that the rapid quenching convertsthe as-deposited films from bcc structure to fcc structure. More-

F. Wang et al. / Electrochimica Acta 51 (2006) 4250–4254 4253

over, it is worthy to note that the diffraction peak of Pt bufferlayer occurring in the as-deposited film completely disappearsin the quenched films. This can be ascribed to the formationof Fe–Pd–Pt solid solution through the atoms diffusion of Pttoward the Fe–Pd films under heating at 900 ◦C for 45 min. Theaddition of Pt lowers the martensite transformation temperatureand strengthens the films due to the solution hardening [13].The SEM images present that the rapid quenching roughens thesurface morphology of thin films probably associated with thechange of phase structure.

3.3. Thermoelastic behavior of Fe–Pd thin films

In order to observe the thermoelastic behavior of quenchedFe–Pd films, in situ observation of XRD spectral changes withthe change of temperature was performed. The measurementwas carried out in a diffraction angle from 46◦ to 52◦ whilekeeping the sample temperature constant at an interval of 20 ◦Cin thermal cycling between RT and −160 ◦C. Fig. 5 shows thechanges of (2 0 0)fcc and (2 0 0)fct (fct: face-centered tetragonal,martensite) peaks of quenched Fe–29.8 at.% Pd film at varioustemperatures during (a) cooling and (b) heating process. As pre-sented in Fig. 5(a), the peak of (2 0 0)fcc is observed at 2θ = 49.1◦at RT. This peak decreases with the decrease of temperature anddisappears at −40 ◦C, while (2 0 0)fct peak at 2θ = 47.3◦ appeara ◦TfpatpXt

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ment between cooling and heating shows a good reversibility ofthe process.

From the XRD results (Fig. 5), the variations of the latticeconstants of Fe–Pd film are calculated according to Bragg’s for-mula. Fig. 6 shows the changes of (1 0 0)fcc and (1 0 0)fct latticeconstants of Fe–Pd films as a function of the temperature ofthermal field. At the temperature from RT to −20 ◦C in thecooling process, the lattice constant of (1 0 0)fcc keeps as a con-stant of about 0.370 nm, showing a independence correlationto the cooling temperature. As the temperature changes from−40 to −160 ◦C, a new fct phase appears instead of fcc phase.The lattice constant of (1 0 0)fct gradually increases reachingas about 0.384 nm with lowering the temperature. The reversephenomenon is observed in the heating process, that is, the lat-tice constant of (1 0 0)fct gradually decreases with increasingthe temperature from −160 to −40 ◦C and completely disap-pears at the temperature between −20 ◦C to RT, meanwhile,the fcc phase appear as the new phase, and the lattice con-stant of (1 0 0)fcc keeps as a constant independent of the heatingtemperature. The changes of lattice constant as a function oftemperature both in the cooling and heating process indicate anarrow temperature hysteresis of austenite-to-martensite trans-formation. Accordingly, the hysteresis curves are approximatelyfitted as the broken lines in Fig. 6. Based the fitted curves, we canapproximately estimate a narrow temperature hysteresis of thefcc–fct transformation of about 10 ◦C and the fcc-to-fct trans-flciaFw

F((o

s second phase and gradually increases by cooling to −160 C.hen it can be considered that an austenite-to-martensite trans-

ormation takes place at −20 to −40 ◦C during the coolingrocess. The reverse behavior of the phase transformation islso observed during the heating process in Fig. 5(b), in whichhe fct structure returns back to the fcc structure during heatingrocess from −160 ◦C to RT. Moreover, it is apparent that the-ray diffraction patterns of cooling process are almost similar

o those of heating process. This suggests that the cyclic treat-

ig. 5. Changes of (2 0 0)fcc and (2 0 0)fct peaks of quenched Fe–29.8 at.% Pdlm at different temperature during (a) cooling and (b) heating.

ormation temperature TM of about −30 ◦C. The increase of theattice constant of (1 0 0)fct during the cooling process may beaused not only by the quantitative increase of martensite phasen the films, but also by the expansion of a-axis and the shrink-ge of c-axis, which are schematically illustrated in the inset ofig. 4. Martensite in Fe–Pd thin films has a tetragonal structureith the c/a ratio of 0.94–0.96 [14]. Further cooling may cause a

ig. 6. Variations of the lattice constants of (1 0 0)fct (cooling and heating) and1 0 0)fcc (cooling and heating): �, (1 0 0)fct (cooling); ♦, (1 0 0)fct (heating); �,1 0 0)fcc (cooling); �, (1 0 0)fcc (heating). Inset shows the schematic illustrationf austenite and martensite structures.

4254 F. Wang et al. / Electrochimica Acta 51 (2006) 4250–4254

Fig. 7. Magnetization curves of quenched Fe–29.8 at.% Pd film measured at RTand −160 ◦C.

non-thermoelastic fct–bct (bct: body-centered tetragonal) trans-formation proceeding in the Fe–Pd films, thus deteriorating theshape memory properties [15]. However, no bct martensite phaseis observed in the XRD patterns in the present study during wholecooling process, indicating that bct martensite phase would notbe thermally induced by cooling even to −160 ◦C.

3.4. Magnetic properties of Fe–Pd thin films

Magnetization curves of Fe–Pd films measured at RT (fcc)and −160 ◦C (fct) are shown in Fig. 7. The slopes of initialparts of the magnetization curves are similar for both fcc and fctphases, showing a similar initial susceptibility. The same sus-ceptibility in the martensite state reflects a possibility that themagnetic domain wall movement is as easy as in the austenitestate [16]. The coercivity of fct phase is 7.38 kA m−1, whichis larger than that of fcc phase (4.61 kA m−1), showing thatboth austenite and martensite phases of Fe–29.8 at.% Pd filmsare ferromagnetic. Considering that martensite phase is morestrongly magnetized under a magnetic field than austenite phase,

the Fe–Pd thin films may exhibit magnetic field-induced actu-ation in the martensite phase [17]. This implies the possibilitythat the Fe–Pd thin films prepared by an electrodeposition pro-cess also can be used as the ferromagnetic shape memory alloys(FSMAs).

4. Conclusions

We have successfully synthesized the nanostructuredFe–29.8 at.%Pd thin films by an electrodeposition process.The in situ X-ray diffraction analysis results indicate that thequenched film with 29.8 at.% Pd undergoes a thermoelasticaustenite-to-martensite transformation with a narrow tempera-ture hysteresis of about 10 ◦C, and the martensite transformationtemperature is about −30 ◦C. The present study has proven thatthe electrodeposition technique is available to the fabricationof Fe–Pd thin films for the application of low-temperature-typeshape memory alloys.

References

[1] D. Vokoun, T. Goryczka, C.T. Hu, J. Alloys Compd. 372 (2004) 168.[2] K. Tsuchiya, T. Nojiri, H. Ohtsuka, M. Umemoto, Mater. Trans. 44

(2003) 2502.[3] S. Inoue, K. Inoue, K. Koterazawa, K. Mizuuchi, Mater. Sci. Eng. A

339 (2003) 34.[4] H.P. Xu, H. Heinrich, J.M.K. Wiezorek, Intermetallics 11 (2003) 969.

[[[

[[[[

[

[5] S. Inoue, K. Inoue, S. Fujita, K. Koterazawa, Mater. Trans. 44 (2003)304.

[6] T. Kubota, T. Okazaki, Y. Furuya, T. Watanabe, J. Magn. Magn. Mater.239 (2002) 513.

[7] C.T. Hu, T. Goryczka, D. Vokoun, Scripta Metall. 50 (2004) 542.[8] T. Yamamoto, M. Taya, Y. Sutou, Y.C. Liang, T. Wada, L. Sorensen,

Acta Mater. 52 (2004) 5091.[9] T. Sakamoto, T. Fukuda, T. Kakeshita, T. Tekeuchi, K. Kishio, Mater.

Trans. 44 (2003) 2498.10] V. Fleury, W.A. Watters, L. Allam, T. Devers, Nature 416 (2002) 719.11] S. Doi, F. Wang, K. Hosoiri, T. Watanabe, Mater. Trans. 44 (2003) 652.12] T.B. Massalski, Binary Alloy Phase Diagram, The Material Information

Society, 1990, p. 1751.13] T. Wada, T. Tagawa, M. Taya, Scripta Mater. 48 (2003) 211.14] R. Oshima, Scripta Metall. 15 (1981) 833.15] D. Vokoun, C.T. Hu, V. Kafka, J. Magn. Magn. Mater. 264 (2003) 174.16] T. Wada, Y.C. Liang, H. Kato, T. Tagawa, M. Taya, T. Mori, Mater. Sci.

Eng. A 361 (2003) 82.17] H. Morito, A. Fujita, K. Fukamichi, R. Kainuma, K. Ishida, Appl. Phys.

Lett. 81 (2002) 1659.