coatings Article Epitaxial Versus Polycrystalline Shape Memory Cu-Al-Ni Thin Films Doga Bilican 1 , Samer Kurdi 2 , Yi Zhu 2 , Pau Solsona 1 , Eva Pellicer 1 , Zoe H. Barber 2 , Alan Lindsay Greer 2 , Jordi Sort 1,3 and Jordina Fornell 1, * 1 Departament de Física, Facultat de Ciències, Universitat Autònoma de Barcelona, E-08193 Bellaterra (Cerdanyola del Vallès), Spain; bilican.doga@gmail.com (D.B.); pau.solsona@uab.cat (P.S.); Eva.pellicer@uab.cat (E.P.); jordi.sort@uab.cat (J.S.) 2 Department of Materials Science and Metallurgy, University of Cambridge, Cambridge CB3 0FS, UK; sk862@cam.ac.uk (S.K.); yz475@cam.ac.uk (Y.Z.); zb10@cam.ac.uk (Z.H.B.); alg13@cam.ac.uk (A.L.G.) 3 Institució Catalana de Recerca i Estudis Avançats (ICREA), Pg. Lluís Companys 23, E-08010 Barcelona, Spain * Correspondence: jordina.fornell@uab.cat; Tel.: +34-9358-114-01  Received: 9 April 2019; Accepted: 6 May 2019; Published: 8 May 2019  Abstract: In this work, two different approaches were followed to obtain Cu-Al-Ni thin films with shape memory potential. On the one hand, Cu-Ni/Al multilayers were grown by magnetron sputtering at room temperature. To promote diffusion and martensitic/austenitic phase transformation, the multilayers were subjected to subsequent heat treatment at 800 ◦ C and quenched in iced water. On the other hand, Cu, Al, and Ni were co-sputtered onto heated MgO (001) substrates held at 700 ◦ C. Energy-dispersive X-ray spectroscopy, X-ray diffraction, and transmission electron microscopy analyses were carried out to study the resulting microstructures. In the former method, with the aim of tuning the thin film’s composition, and, consequently, the martensitic transformation temperature, the sputtering time and applied power were adjusted. Accordingly, martensitic Cu-14Al-4Ni (wt.%) and Cu-13Al-5Ni (wt.%) thin films and austenitic Cu-12Al-7Ni (wt.%) thin films were obtained. In the latter, in situ heating during film growth led to austenitic Cu-12Al-7Ni (wt.%) thin films with a (200) textured growth as a result of the epitaxial relationship MgO(001)[100]/Cu-Al-Ni(001)[110]. Resistance versus temperature measurements were carried out to investigate the shape memory behavior of the austenitic Cu-12Al-7Ni (wt.%) thin films produced from the two approaches. While no signs of martensitic transformation were detected in the quenched multilayered thin films, a trend that might be indicative of thermal hysteresis was encountered for the epitaxially grown thin films. In the present work, the differences in the crystallographic structure and the shape memory behavior of the Cu-Al-Ni thin films obtained by the two different preparation approaches are discussed. Keywords: Cu-Al-Ni; shape memory alloys; thin film; sputtering; size effects 1. Introduction Shape memory alloys (SMA) exhibit displacive and reversible deformation behavior due to sensing thermodynamic and mechanical changes in their environment [1–7]. This plays a role in the development of components that can be cycled between two macroscopic shapes depending on temperature change. The fact that SMAs exhibit temperature-induced strain recovery makes them a type of advanced engineering material. Typical application fields for these materials are encountered in sensing–actuating systems in automotive, aerospace, robotics, and biomedical technologies [8]. Of the most functional shape memory alloys, Ni-Ti [9–11] Cu-Zn-Al [12], and Cu-Al-Ni [13,14] are some of the most widely used. Even though the Ni-Ti SMA system is widely studied and commercialized on account of its high percentage of shape recovery, Cu-based SMAs have become long-term proposed Coatings 2019, 9, 308; doi:10.3390/coatings9050308 www.mdpi.com/journal/coatings Coatings 2019, 9, 308 2 of 10 substitutions for the Ni-Ti system due to the fact that they have lower production costs in addition to having desirable properties such as a large superelastic effect, wide transformation temperature ranges, small hysteresis, and a high damping coefficient. Despite the fact that Fe-based SMAs, such as Fe-Mn and Fe-Mn-Si, also appear as substitute candidates for Ni-Ti due to their good workability and cost efficiency advantages, Fe-based SMAs undergo a large transformation hysteresis which limits their area for shape memory applications compared to Cu-based SMAs [15]. Regardless of the alloy composition, because of possessing large amounts of thermal capacitance, the application of shape memory behavior in bulk alloys is challenging. However, due to high actuation outputs per unit volume, the response time can be reduced substantially and the speed of operation may be increased sufficiently in shape memory thin films. As a consequence, these materials can be utilized in micro/nanorobotic platforms [16]. Rather than conventional elastic or plastic dislocation glide, the behavior of shape memory depends on displacive, diffusionless phase transformation, which takes place between a high temperature phase (austenite) and a low temperature one (martensite). As a non-equilibrium phase, the most commonly applied way to induce martensite in materials is to subject them to quenching after holding them at high temperature in order for the atoms to become locked into position before they reach their equilibrium states, with the lowest free energy [17]. Post-quenching is also convenient for introducing the shape memory effect in films grown at room temperature. Investigations on the synthesis of Fe-Pd [18], Ni-Ti [19], and Cu-Al-Ni [20] by sputter depositing and subsequent quenching have been reported in the literature. In thin film studies, epitaxy may be induced via growth on a single-crystal substrate, which dictates the crystal growth direction of the film. As a refractory, electrically insulating, and transparent substrate, MgO offers a flat interface for epitaxial growth. It shows a convenient lattice that matches many metals with face-centered cubic symmetry [21]. For thin film synthesis using in situ heating, epitaxial growth can be provided by techniques such as laser ablation [22], molecular beam epitaxy [23,24], and electron beam evaporation [25]. Since preferential orientation is favored, the enhancement of shape memory properties can be achieved. Among the sputtered shape memory alloys, Ni-Ti [26] and Ni2 MnGa [27,28] are the systems for which in-depth research on epitaxial growth is being carried out. Cu-Al-Ni SMAs are known for their good thermal and electrical conductivity, broad interval of transformation temperatures (between 70 and 470 K), good thermal stability of martensitic transformation, and large recoverable strains [29–31]. They typically have Ni and Al contents between 2–5 wt.% and 13–15 wt.%, respectively [32,33]. Cu-Al-Ni SMAs have some advantages over the widely used Ni-Ti SMAs, such as a lower melting temperature, which facilitates composition control [34]. The fact that they have better corrosion resistance and a lower cost compared to Ni-Ti, which is prone to oxidation, makes Cu-Al-Ni components favorable choices [35] for certain applications. Furthermore, it has been shown that achieving further enhancement in the microstructure and mechanical properties of Cu-Al-Ni SMAs is possible via the minor addition of cobalt to its composition [34]. The structural and mechanical properties of the Cu-Al-Ni system have been studied both for bulk alloys, which are produced by methods such as rapid solidification, casting [36], and powder metallurgy [37], and thin films, which are produced by methods including electron beam evaporation [38], thermal evaporation [39], and sputtering [40,41]. Here, the dependence of the crystallographic phase structure on the chemical composition of the Cu-Al-Ni shape memory system is reported for the first time for sputtered thin films. Different compositions were synthesized by adjusting the Al and Ni content. Additionally, the crystallographic differences due to the undertaken synthetic approach were investigated. Two methods were implemented to produce the thin films. Firstly, free-standing films with varying compositions were prepared by sputtering multilayers of Cu-Ni and Al. The as-grown films were conventionally annealed at 800 ◦ C and quenched in iced water, as would be done in a top-down approach, in order to obtain a set of samples with different chemical compositions: Cu-14Al-4Ni (wt.%), Cu-13Al-5Ni (wt.%) and Cu-12Al-7Ni (wt.%). Secondly, Cu-12Al-7Ni (wt.%) thin films were grown on MgO(001) substrates Coatings 2019, 9, 308 3 of 10 Coatings 2019, 9, 308 3 of 10 substrates held at 700 °C, taking advantage of the epitaxial relationship between MgO(001) and held at 700 ◦ C, taking advantage of the epitaxial relationship between MgO(001) and austenitic austenitic β-Cu-Al-Ni as a bottom-up approach. Unlike most of the previous studies for sputtered β-Cu-Al-Ni as a bottom-up approach. Unlike most of the previous studies for sputtered Cu-Al-Ni Cu-Al-Ni samples, where a Cu-Al-Ni target was used for synthesis [16,27,34], in this work all samples, where a Cu-Al-Ni target was used for synthesis [16,27,34], in this work all depositions used depositions used individual Cu, Al, and Ni targets (i.e., co-sputtering). Structural characterization individual Cu, Al, and Ni targets (i.e., co-sputtering). Structural characterization studies were carried studies were carried out to investigate the crystallographic properties of the films. The martensitic out to investigate the crystallographic properties of the films. The martensitic phase transformation phase transformation behavior of the samples was also investigated. behavior of the samples was also investigated. 2.2.Materials Materialsand andMethods Methods Cu, Cu,Al Aland andNiNitargets targetsof of 99.95% 99.95%purity were purity used were to grow used all samples. to grow all samples.In order to produce In order free- to produce standing Cu-Al-Ni films, Si substrates were covered with resin by spin coating before sputtering. free-standing Cu-Al-Ni films, Si substrates were covered with resin by spin coating before sputtering. The deposition of multilayers was carried out by DC sputtering from Cu and Al targets The deposition of multilayers was carried out by DC sputtering from Cu and Al targets and RF and RF sputtering from a Ni target. sputtering from Seven multilayers a Ni target. Seven consisting multilayersofconsisting three Al films deposited of three Al filmsbetween depositedCu-Ni films Cu-Ni between were prepared (Figure 1a). After the deposition, the samples were rinsed first in acetone films were prepared (Figure 1a). After the deposition, the samples were rinsed first in acetone and and later in ethanol to dissolve the resin and to separate the film from the substrate. In order to promote later in ethanol to dissolve the resin and to separate the film from the substrate. In order to promote diffusion todiffusion obtain a to homogeneous composition, obtain a homogeneous the films, wrapped composition, in Ta the films, foils, were wrapped in Tasubjected to annealing foils, were subjectedatto 800 °C for 60 min ◦in a sealed quartz tube and subsequently quenched in iced water. annealing at 800 C for 60 min in a sealed quartz tube and subsequently quenched in iced water. Sketchofofthe Figure1.1.Sketch Figure the(a) (a)prepared prepared Cu-Ni/Al Cu-Ni/Al multilayers multilayers before before quenching (forxxand quenching(for andyyvalues valuesrefer referto toTable Table1)1) and and(b)(b) Cu-Al-Ni Cu-Al-Nifilm onon film (001) oriented (001) MgO. oriented MgO. Table 1. Thickness values of the sputtered multilayers before quenching. The top and bottom Cu–Ni The top and bottom Cu-Ni layers had half the thickness of the Cu-Ni layers in between (Figure layers have half the thickness of the layers in between. 1a). The details of the thickness values for each composition are given in Table 1, while sputtering conditions are listed in Table 2. Layer Thickness (nm) Total Thin Film Thickness (nm) Composition Cu-Ni (x) Al (y) Cu-Al-Ni Table 1. Thickness values of the sputtered multilayers before quenching. The top and bottom Cu–Ni Cu-14Al-4Ni (wt.%) 230 58 864 layers have half the thickness of the layers in between. Cu-13Al-5Ni (wt.%) 242 63 915 Cu-12Al-7Ni (wt.%) 347 70 1250 Layer Thickness (nm) Total Thin Film Thickness (nm) Composition Cu-Ni (x) Al (y) Cu-Al-Ni The top and bottom Cu-14Al-4Ni Cu-Ni layers (wt.%) had half the58 230 thickness of the Cu-Ni864 layers in between (Figure 1a). The details of the thickness Cu-13Al-5Ni values for242 (wt.%) each composition 63 are given in Table 915 1, while sputtering conditions are listedCu-12Al-7Ni in Table 2. (wt.%) 347 70 1250 Table 2. The power (W) and time (min) conditions applied for the sputtering of samples Cu–14Al–4Ni Table 2. The power (W) and time (min) conditions applied for the sputtering of samples Cu–14Al– (wt.%), Cu–13Al–5Ni (wt.%), and Cu–12Al–7Ni (wt.%). 4Ni (wt.%), Cu–13Al–5Ni (wt.%), and Cu–12Al–7Ni (wt.%). Cu-Ni Cu-Ni Al Al Composition Composition PP(W) (W) t (min)t (min) P (W) tP(min) (W) t (min) Cu Cu NiNi Cu-14Al-4Ni Cu-14Al-4Ni (wt.%) (wt.%) 200 200 6060 10 10200 14 200 14 Cu-13Al-5Ni Cu-13Al-5Ni (wt.%) (wt.%) 200 200 75 75 9 9 200 15.5 200 15.5 Cu-12Al-7Ni (wt.%) 200 Cu-12Al-7Ni (wt.%) 200 8585 10 10200 200 17 17 The co-sputtered 650 nm thick Cu-12Al-7Ni (wt.%) thin film was deposited onto a 700 ◦ C heated MgO(001) substrate (Figure 1b) from individual Cu, Al, and Ni targets using powers of 22, 30, and 5 W, Coatings 2019, 9, 308 4 of 10 respectively. The deposition pressure was set to 0.65 Ar, and the deposition rate was approximately 4.3 nm/min. Compositional analyses were performed with a field emission scanning electron microscope (FE-SEM, Zeiss, Oberkochen, Germany) equipped with an energy dispersive X-ray spectroscopy (EDX) detector operated at 15 kV. To confirm the homogeneity of the film along its thickness, EDX analysis was also carried out on the film cross-section. Structural characterization was carried out by X-ray diffraction (XRD, PANalytical, Royston, UK) (θ/2θ diffraction with Cu Kα radiation) and transmission electron microscopy (TEM, JEOL 2011 200 KV, Peabody, MA, USA). For TEM observations, the free-standing films were prepared with a GATAN polishing ion device (Pleasanton, CA, USA) while the Cu-12Al-7Ni (wt.%) film grown on MgO was scratched from the substrate and milled with an agate mortar. A Brucker D8 theta/theta four circle diffractometer (Billerica, MA, USA): Ω, 2θ, X, Φ, graded mirror (GM; to give a nearly parallel beam) equipped with a scintillation counter detector was used to study the epitaxial films grown on MgO. Coating thickness measurements were made using the 3D optical surface metrology system, Leica DCM 3D (Leica Microsystems Inc., Buffalo Grove, IL, USA). In order to investigate martensitic transformation behavior, electrical resistance measurements were performed with a 2 K/min heating rate. 3. Results and Discussion The XRD patterns of films synthesized by both routes are shown in Figure 2a. Among the free-standing thin films produced by multilayer deposition, Cu-14Al-4Ni (wt.%) (Figure 2a) showed a mixture of β’ martensite and pure Cu. β’ martensite had a monoclinic structure (space group Cmcm), whereas Cu had a cubic Fm-3m lattice. As the content of Ni increased at the expense of Al, traces of β austenite (space group Fm-3m) along with β’ martensite were detected for Cu-13Al-5Ni (wt.%) (Figure 2b), and only β austenite was seen in the XRD pattern of Cu-12Al-7Ni (wt.%) (Figure 2c). In turn, the XRD pattern of the thin film co-sputtered at a high temperature on MgO mainly consisted of textured β austenite, but the Cu9 Al4 phase (space group P-43m) was also present (Figure 2d). Previous studies on Cu-Al-Ni bulk shape memory alloys have reported that a slight change in composition results in a shift in the transformation temperatures [33,41,42]. The studies carried out by Recarté et al. revealed that, for a fixed aluminum content of 13.2 wt.%, a decrease in nickel from 5 wt.% to 3.5 wt.% raised the austenite finish temperature from 10 to 80 ◦ C [33]. Agafonov et al. showed that Al content variation from 14.98 wt.% to 13.03 wt.% caused a change in the room temperature phase from austenite to martensite for samples synthesized by casting and quenching in water [41]. Similar trends were observed by Suresh and Ramamurty [42] where, at room temperature, a Cu-13.4Al-4Ni (wt.%) alloy was in the martensite state, but a Cu-14.1Al-4Ni (wt.%) alloy was austenitic. In our work, we observed that a change in Ni content followed the same trend in the transformation temperatures as the results reported by Recarté et al. [33], but the opposite tendency than that reported in [33,41] was observed when the aluminum content was modified. This may be because in our case, we were modifying Ni and Al content simultaneously, and, consequently, our transformation temperatures were influenced by both Al and Ni. The film grown on MgO(001) showed a Cu-12Al-7Ni (wt.%) composition. The epitaxial relationship MgO(001)[100]/Cu-Al-Ni(001)[110] was induced. The film grew with a 45◦ in-plane rotation on the cubic cell of the substrate. The lattice mismatch (f ) between the cubic lattices is shown in Equation (1): √ f = ( 2aMgO − aCuAlNi β )/aCuAlNi β (1) where aMgO = 4.212 Å, aCuAlNi β = 5.836 Å. The lattice mismatch between the film and the substrate was f = 2%, which was small enough to favor epitaxy. The epitaxial film, grown at 700 ◦ C, exhibited preferential (100) out-of-plane orientation. This showed up as strong (200) and (400) peaks in the XRD pattern (Figure 2d). Peaks belonging to the Cu9 Al4 phase (i.e., (200), (421), (332), and (550) planes) were also present. The formation of α-Cu and Coatings 2019, 9, 308 5 of 10 Cu9 Al4 phases is typically observed in Cu-Al-Ni systems [43,44]. The β phase in the Cu-Al-Ni system undergoes a eutectoid decomposition at ~840 K into α (Cu) and γ2 (Cu9 Al4 ), then, the stable phases at room temperature are Cu and Cu9 Al4 . However, if the alloy is quenched at a sufficiently high cooling rate from the β phase region to ambient temperature, the β phase may be retained or it may transform martensitically. However, if the cooling rate is not high enough, traces of Cu or Cu9 Al4 may be present in the alloy. Coatings 2019, 9, 308 5 of 10 Figure 2. XRD Figure 2. patterns of XRD patterns of the the sputtered sputtered Cu–Al–Ni Cu–Al–Ni thin thin films: films: (a) (a) θ/2θ θ/2θ scan scan for for Cu–14Al–4Ni (wt.%); Cu–14Al–4Ni (wt.%); (b) θ/2θscan (b) θ/2θ scanforfor Cu–13Al–5Ni Cu–13Al–5Ni (wt.%); (c) θ/2θ (wt.%); Cu–12Al–7Ni (c) θ/2θ Cu–12Al–7Ni(wt.%) films obtained (wt.%) by multilayer films obtained sputter by multilayer deposition and subsequent heat treatment; (d) θ/2θ scan for Cu–12Al–7Ni (wt.%) sputter deposition and subsequent heat treatment; (d) θ/2θ scan for Cu–12Al–7Ni (wt.%) film film prepared on MgO substrate at 700 ◦ C; and (e) φ scan for the (220) peak of the β–Cu–Al–Ni film deposited at 700 ◦ C prepared on MgO substrate at 700 °C; and (e) ϕ scan for the (220) peak of the β–Cu–Al–Ni film on MgO(001). deposited # peak at 700 °C onisMgO(001). unidentified. * peaks # peak originate from is unidentified. the sample * peaks holder. originate from the sample holder. The phi scan carried out for the (220) peak of this film confirmed that the film lattice was rotated The film grown on MgO(001) showed a Cu-12Al-7Ni (wt.%) composition. The epitaxial 45◦ in-plane, relative to the substrate (Figure 2e). relationship MgO(001)[100]/Cu-Al-Ni(001)[110] was induced. The film grew with a 45° in-plane For the free-standing thin films produced by multilayer deposition, further evidence that the rotation on the cubic cell of the substrate. The lattice mismatch (f) between the cubic lattices is shown compositional change from Cu-14Al-4Ni (wt.%) to Cu-12Al-7Ni (wt.%) resulted in a shift from in Equation (1): martensite to austenite was observed by TEM (Figure 3). Figure 3a,c shows TEM images of Cu-14Al-4Ni (wt.%) and Cu-12Al-7Ni (wt.%) thin 𝒇 = films, (√𝟐𝒂𝐌𝐠𝐎 − 𝒂𝐂𝐮𝐀𝐥𝐍𝐢 𝛃 )Typical respectively. 𝒂𝐂𝐮𝐀𝐥𝐍𝐢martensitic 𝛃 plates were present(1)in Cu-14Al-4Ni, while a regular polycrystalline structure was observed in the Cu-12Al-7Ni (wt.%) alloy. where aMgO = 4.212 Å, aCuAlNi β = 5.836 Å. The Selected Area Electron Diffraction (SAED) pattern from Cu-14Al-4Ni (wt.%) thin film (Figure 3b) The lattice mismatch between the film and the substrate was f = 2%, which was small enough to consisted of diffraction spots that belong to β’ martensite corresponding to the (040), (202), (0018), favor epitaxy. The epitaxial film, grown at 700 °C, exhibited preferential (100) out-of-plane (202), and (122) planes and to Face-Centered Cubic (FCC) Cu (i.e., (220), (311), (311), and (033) planes). orientation. This showed up as strong (200) and (400) peaks in the XRD pattern (Figure 2d). Peaks Conversely, the SAED pattern of the Cu-12Al-7Ni (wt.%) thin film (Figure 3d) consisted of diffraction belonging to the Cu9Al4 phase (i.e., (200), (421), (332), and (550) planes) were also present. The spots characteristic formation of α-Cu and of the Cu austenitic phase (β) corresponding to (111), (511), (711), (622), (422), and (533) 9Al4 phases is typically observed in Cu-Al-Ni systems [43,44]. The β phase planes. In Figure in the Cu-Al-Ni 4a, a undergoes system high resolution TEM image a eutectoid of the epitaxial decomposition at ~840 Cu-12Al-7Ni K into α (Cu) (wt.%) and γ2 thin film (Cu9Al 4), grown on MgO is shown. From the fast-Fourier transform (FFT) pattern then, the stable phases at room temperature are Cu and Cu9Al4. However, if the alloy is quenched atof the selected zone shown in Figure 4a, an a sufficiently highinterplanar cooling ratedistance from thevalue of 2.07 β phase Å (which region belongs to ambient to β austenite temperature, the(220)), β phaseas may well be as aretained distanceorofit2.93 mayÅtransform (which belongs to β austenite martensitically. (200)),ifwas However, theidentified cooling rate (Figure is not4b). high enough, traces In order to identify the transformation of Cu or Cu9Al4 may be present in the alloy. temperatures, electrical resistance measurements were carried Theoutphiasscan a function carried out offor temperature the (220) peak for of both this the filmquenched confirmedmultilayered sample that the film lattice was with the rotated composition 45° in-plane,Cu-12Al-7Ni relative to the (wt.%) and the substrate sample (Figure with the same composition grown on MgO. Martensitic 2e). transformation was not observed in the quenched For the free-standing thin films produced by multilayer multilayered Cu-12Al-7Ni deposition, further(wt.%) (Figure evidence that 5a), the whereas a trend that might be indicative of transformation hysteresis compositional change from Cu-14Al-4Ni (wt.%) to Cu-12Al-7Ni (wt.%) resulted in a shift from was found for the epitaxially grown Cu-12Al-7Ni martensite to austenite (wt.%) was sample observed (Figure by TEM 5b).(Figure The reason behind 3). Figure 3a,cthe fact that shows TEM a transformation images of Cu-14Al- was observed in the epitaxially grown (200) textured austenite thin film whereas 4Ni (wt.%) and Cu-12Al-7Ni (wt.%) thin films, respectively. Typical martensitic plates were present no change was seen in in the randomly while Cu-14Al-4Ni, oriented polycrystalline a regular one could polycrystalline be related structure to size effects. was observed in the Size effects, such Cu-12Al-7Ni as (wt.%) the volume of the material or structural components including precipitate alloy. The Selected Area Electron Diffraction (SAED) pattern from Cu-14Al-4Ni (wt.%) thin film particles and grains in polycrystals, have a huge influence on the martensitic phase transformation [37]. (Figure 3b) consisted of diffraction spots that belong to β’ martensite corresponding to the (040), (202),Decreasing grain size (0018), (202), and (122) planes and to Face-Centered Cubic (FCC) Cu (i.e., (220), (311), (311), and (033) planes). Conversely, the SAED pattern of the Cu-12Al-7Ni (wt.%) thin film (Figure 3d) consisted of diffraction spots characteristic of the austenitic phase (β) corresponding to (111), (511), (711), (622), (422), and (533) planes. In Figure 4a, a high resolution TEM image of the epitaxial Cu-12Al-7Ni (wt.%) thin film grown on MgO is shown. From the fast-Fourier transform (FFT) pattern of the selected zone Coatings 2019, 9, 308 6 of 10 and decreasing twin separation cause an increase in strain energy and twin interfacial energy which, in turn, increases the energy barrier [38]. For instance, Shi et al. [39] showed that for submicrometric NiTi particles (below 100 nm), the martensitic transformation was fully suppressed. Similarly, in our work, no transformation Coatings 2019, 9, 308 was observed in the randomly oriented polycrystalline austenitic Cu-12Al-7Ni 6 of 10 (wt.%) film 9, Coatings 2019, with 308 a crystallite size of 95 nm according to Scherrer’s formula. 6 of 10 Figure 3. (a) Transmission electron microscopy (TEM) image of Cu–14Al–4Ni (wt.%) thin film Figure 3.3. (a) Figure (a)Transmission electron Transmission microscopy electron (TEM) microscopy imageimage (TEM) of Cu–14Al–4Ni (wt.%) thin of Cu–14Al–4Ni film produced (wt.%) produced from multilayer deposition followed by annealing; (b) SAED pattern of (a,c) TEMthin film image of from multilayer deposition followed by annealing; (b) SAED pattern of (a,c) TEM image of Cu–12Al–7Ni produced from multilayer deposition followed by Cu–12Al–7Ni (wt.%) thin film; (d) SAED pattern of (c).annealing; (b) SAED pattern of (a,c) TEM image of (wt.%) thin film; (d) SAED pattern of (c). Cu–12Al–7Ni (wt.%) thin film; (d) SAED pattern of (c). 4. (a) HRTEM image of Cu–12Al–7Ni (wt.%) thin film grown on MgO; (b) FFT of the selected Figure 4. Figure Figure zone 4. (a) HRTEM (orange zone (orange image square) in square) in (a). of Cu–12Al–7Ni (wt.%) thin film grown on MgO; (b) FFT of the selected (a). zone (orange square) in (a). Additionally, In previous order to identify the studies demonstrate transformation that decreasing temperatures, filmresistance electrical thicknessmeasurements is also a reason for were In carried order the suppression to identify out as aoffunction the martensitic transformation temperatures, phase transformation of temperature for both the [45–48]. electrical Wan and quenched resistance Komvopoulos multilayered measurements sample[48]with were showedthe carried that for out as sputtereda function NiTi of films temperature with a for thickness both of lessthethan quenched 100 composition Cu-12Al-7Ni (wt.%) and the sample with the same composition grown on MgO.nm, multilayered no martensitic sample with transformation the composition was observed. Cu-12Al-7Ni In the case (wt.%) of sputteredand the sample polycrystalline with the Cu-Al-Ni same films, Martensitic transformation was not observed in the quenched multilayered Cu-12Al-7Ni (wt.%)composition transformation grown on MgO. temperatures Martensitic for films5a), (Figure transformation with whereas a trend was 2 µm thickness that not were might observed reported byinMoran be indicativetheofquenched multilayered and his co-workers transformation Cu-12Al-7Ni [31]. hysteresis To wasthe (wt.%) best for found of our the epitaxially grown Cu-12Al-7Ni (wt.%) sample (Figure 5b). The reason behind the fact thatthe (Figure knowledge,5a), whereas these are a trend the that sputtered might be Cu-Al-Ni indicative SMA of films transformation with the lowest hysteresis thickness was valuesfound for reported in a transformation was observed in the epitaxially grown (200) textured austenite thin film whereas noa epitaxially the literature.grown Torres Cu-12Al-7Ni et al. [20] also (wt.%) observed sample (Figure martensitic 5b). The transformation reason for 5 µmbehind thick the films. fact In that contrast, transformation no transformation change was seen was observed inwas theobserved randomlyin the epitaxially in oriented the grown 1.25 polycrystalline polycrystalline (200)µm textured one couldaustenite Cu-12Al-7Ni thin (wt.%) be related film effects. to sample size whereas preparedno Size change in this was work seen by in the randomly post-treatment of oriented the polycrystalline multilayers. It has one been could be indicated related effects, such as the volume of the material or structural components including precipitate particles by Chento size and effects. Schuh Size [49] effects, and suchinaspolycrystals, grains the volume of the amaterial have or structural huge influence on the components martensitic including precipitate particles phase transformation [37]. and grains in polycrystals, have a huge influence on the martensitic Decreasing grain size and decreasing twin separation cause an increase in strain energy and twin phase transformation [37]. Decreasing grain size and decreasing twin separation cause an increase in strain interfacial energy which, in turn, increases the energy barrier [38]. For instance, Shi et al. [39] showed energy and twin interfacial that energy which, in for submicrometric turn, NiTi increases particles the energy (below barrier 100 nm), the[38]. For instance, martensitic Shi et al. [39] transformation showed was fully that for submicrometric NiTi particles (below 100 nm), the martensitic suppressed. Similarly, in our work, no transformation was observed in the randomly oriented transformation was fully Coatings 2019, 9, 308 7 of 10 that as d(grain size)/D(sample thickness) decreases, the energy barrier for transformation increases. The interfacial energies depend on grain boundaries formed between martensite plates, the interfaces between Coatings austenite 2019, 9, 308 and martensite plates, and the twin interfaces within martensite plates [48]. 7 of 10 Resistance vs. temperature Figure 5. Resistance temperature measurement measurement of of Cu–12Al–7Ni Cu–12Al–7Ni (wt.%) films prepared by (a) heat treatment after sputtering and (b) in situ heating during growth. growth. The absence ofprevious Additionally, phase transformation in filmsthat studies demonstrate formed at low deposition decreasing film thicknesstemperatures couldfor is also a reason be attributed the to the small suppression grain sizesphase of martensitic causing a higher number transformation of grain [45–48]. Wan boundary interfaces and and Komvopoulos associated [48] showed excess free volume, and/or the thinner nature of the samples, as noted by Wan and Komvopoulos that for sputtered NiTi films with a thickness of less than 100 nm, no martensitic transformation was [48]. observed. In the case of sputtered polycrystalline Cu-Al-Ni films, transformation temperatures for 4. Conclusions films with 2 µm thickness were reported by Moran and his co-workers [31]. To the best of our In this work, knowledge, thesestructural characterization are the sputtered Cu-Al-Ni studies SMAwerefilmscarried out lowest with the to investigate thicknessthe values crystallographic reported properties of Cu-Al-Ni in the literature. Torresfilms et al.and [20]their also transformation temperatures. observed martensitic transformation for 5 µm thick films. In The no contrast, main conclusions from transformation this work in was observed arethe as polycrystalline follows: 1.25 µm Cu-12Al-7Ni (wt.%) sample prepared in this work by post-treatment of the multilayers. It has been indicated by Chen and Schuh • The microstructure of Cu-Al-Ni sputtered films is found to depend both on the alloy composition [49] that as d(grain size)/D(sample thickness) decreases, the energy barrier for transformation as well as the experimental procedure used to grow the films. increases. The interfacial energies depend on grain boundaries formed between martensite plates, the • A transition interfaces between from martensite austenite to austenite and martensite was observed plates, as the and the twin Ni content interfaces increased within and the martensite Al plates [48]. content decreased in samples prepared by multilayer sputtering followed by quenching. • Preferential The absencegrowth of phase along the (100) direction transformation in films was formedobserved at lowin β-austenite deposition Cu-12Al-7Ni temperatures (wt.%) could be grown on MgO at 700 ◦ C due to the epitaxial relationship MgO(001)[100]/Cu-Al-Ni(001)[110]. attributed to the small grain sizes causing a higher number of grain boundary interfaces and • Resistance associated change excess free with respect volume, to temperature, and/or the thinnersuggesting nature of martensitic the samples, transformation as noted by hysteresis, Wan and was observed Komvopoulos [48]. in the preferentially oriented austenitic Cu-12Al-7Ni (wt.%) film, whereas martensitic transformation was completely suppressed in the polycrystalline austenitic sample produced by multilayer sputtering with the same composition. 4. Conclusions In this work, structural characterization studies were carried out to investigate the Author Contributions: crystallographic Conceptualization, properties D.B. and of Cu-Al-Ni films andJ.F.; Methodology, their D.B.,temperatures. transformation P.S., S.K., Y.Z, and J.F.; Formal Analysis, D.B., S.K., and J.F.; Investigation, D.B., S.K., Z.H.B., and J.F.; Resources, D.B., S.K., Z.H.B., and J.F.; The mainD.B. Data Curation, conclusions from this work Draft and J.F.; Writing—Original are asPreparation, follows: D.B.; Writing—Review and Editing, D.B., S.K., Z.H.B., A.L.G., E.P., J.S., and J.F.; Supervision, E.P., J.S., and J.F.; Project Administration, E.P., J.S., and J.F.; Funding • The microstructure Acquisition, E.P. and J.S. of Cu-Al-Ni sputtered films is found to depend both on the alloy composition as well as the experimental procedure used to grow the films. Funding: This work was supported by the SELECTA (No. 642642) H2020-MSCA-ITN-2014 project. Partial financial •support A transition from government from the Spanish martensite (Project to austenite was observed asand MAT2017-86357-C3-1-R theassociated Ni content increased FEDER) and theand the Al Generalitat content (2017-SGR-292) de Catalunya decreased in samples prepared by multilayer sputtering followed by quenching. is acknowledged. • Preferential growth Acknowledgments: E.P. is along theto(100) grateful direction MINECO was“Ramon for the observed in β-austenite y Cajal” Cu-12Al-7Ni (wt.%) contract (RYC-2012-10839). J.F. acknowledges the “Juan de la Cierva” fellowship from MINECO (IJCI-2015-27030). grown on MgO at 700 °C due to the epitaxial relationship MgO(001)[100]/Cu-Al-Ni(001)[110]. •Conflicts Resistance change of Interest: with declare The authors respect to temperature, no conflict of interest. suggesting martensitic transformation hysteresis, was observed in the preferentially oriented austenitic Cu-12Al-7Ni (wt.%) film, whereas martensitic transformation was completely suppressed in the polycrystalline austenitic sample produced by multilayer sputtering with the same composition. Author Contributions: Conceptualization, D.B. and J.F.; Methodology, D.B., P.S., S.K., Y.Z, and J.F.; Formal Analysis, D.B., S.K., and J.F.; Investigation, D.B., S.K., Z.H.B., and J.F.; Resources, D.B., S.K., Z.H.B., and J.F.; Data Curation, D.B. and J.F.; Writing—Original Draft Preparation, D.B.; Writing—Review and Editing, D.B., S.K., Z.H.B., A.L.G., E.P., J.S., and J.F.; Supervision, E.P., J.S., and J.F.; Project Administration, E.P., J.S., and J.F.; Funding Acquisition, E.P. and J.S. Coatings 2019, 9, 308 8 of 10 References 1. Otsuka, K.; Wayman, C.M. Shape Memory Materials; Cambridge University Press: Cambridge, UK, 1999. 2. Ma, J.; Karaman, I. Expanding the repertoire of shape memory alloys. Science 2010, 327, 1468–1469. [CrossRef] 3. Cingolani, E.; Ahlers, M.; Van Humbeeck, J. Stabilization and two-way shape memory effect in Cu-Al-Ni single crystals. Met. Mater. Trans. A 1999, 30, 493–499. [CrossRef] 4. Ishibashi, M.; Tabata, N.; Suetake, T.; Omori, T.; Sutou, Y.; Kainuma, R.; Yamauchi, K.; Ishida, K. A simple method to treat an ingrowing toenail with a shape-memory alloy device. J. Dermatol. 2008, 19, 291–292. [CrossRef] [PubMed] 5. Seelecke, S.; Muller, I. Shape memory alloy actuators in smart structures: Modeling and simulation. Appl. Mech. Rev. 2004, 57, 23–46. [CrossRef] 6. Jani, J.M.; Leary, M.; Subic, A.; Gibson, M.A. A review of shape memory alloy research, applications and opportunities. Mater. Des. 2014, 56, 1078–1113. [CrossRef] 7. Montero-Ocampo, C.; Lopez, H.; Salinas Rodriguez, A. Effect of compressive straining on corrosion resistance of a shape memory Ni-Ti alloy in ringer’s solution. J. Biomed. Mater. Res. 1993, 32, 583–591. [CrossRef] 8. Bogue, R. Shape-memory materials: A review of technology and applications. Assem. Autom. 2009, 29, 214–219. [CrossRef] 9. Dehghanghadikolaei, A.; Ibrahim, H.; Amerinatanzi, A.; Hashemi, M.; Moghaddam, N.S.; Elahinia, M. Improving corrosion resistance of additively manufactured nickel-titanium biomedical devices. J. Mater. Sci. 2019, 54, 7333–7355. [CrossRef] 10. Eggeler, G.; Hornbogen, E.; Yawny, A.; Heckmann, A.; Wagner, M. Structural and functional fatigue of NiTi fhape memory alloys. Mater. Sci. Eng. A 2004, 378, 24–33. [CrossRef] 11. Ibrahim, H.; Jahadakbar, A.; Dehghan, A.; Moghaddam, N.S.; Amerinatanzi, A.; Elahinia, M. In vitro corrosion assessment of additively manufactured porous NiTi structures for bone fixation applications. Metals 2018, 8, 164. [CrossRef] 12. Perkins, J.; Sponholz, R.O. Stress-induced martensitic transformation cycling and two-way shape memory training in Cu-Zn-Al alloys. Met. Mater. Trans. B 1984, 15, 313–321. [CrossRef] 13. Ivanić, I.; Kožuh, S.; Grgurić, T.H.; Kosec, B.; Gojić, M. The influence of heat treatment on microstructure and phase transformation temperatures of Cu-Al-Ni shape memory alloys. Kem. Ind. 2019, 68, 111–118. [CrossRef] 14. Recarte, V.; Pérez-Landazábal, J.I.; Ibarra, A.; Nó, M.L.; San Juan, J. High temperature β phase decomposition process in a Cu-Al-Ni shape memory alloy. Mater. Sci. Eng. A 2004, 378, 238–242. [CrossRef] 15. Alaneme, K.K.; Okotete, E.A. Reconciling viability and cost-effective shape memory alloy options–A review of copper and iron based shape memory metallic systems. Eng. Sci. Technol. Int. J. 2016, 19, 1582–1592. [CrossRef] 16. Pan, Q.; Cho, C. The investigation of a shape memory alloy micro-damper for MEMS applications. Sensors 2007, 7, 1887–1900. [CrossRef] 17. Olson, G.B.; Cohen, M. A general mechanism of martensitic nucleation: Part I. General concepts and the FCC → HCP transformation. Metall. Mater. Trans. A 1976, 7, 1897–1904. 18. Sugimura, Y.; Cohen-Karni, I.; McCluskey, P.; Vlassak, J. Stress evolution in sputter-deposited Fe–Pd shape-memory thin films. J. Mater. Res. 2005, 20, 2279–2287. [CrossRef] 19. Chu, J.P.; Lai, Y.W.; Lin, T.N.; Wang, S.F. Deposition and characterization of TiNi-base thin films by sputtering. Mater. Sci. Eng. A 2000, 277, 11–17. [CrossRef] 20. Torres, C.E.; Condo, A.; Haberkorn, N.; Zelaya, E.; Schryvers, D.; Guimpel, J.; Lovey, F. Structures in textured Cu–Al–Ni shape memory thin films grown by sputtering. Mater. Charact. 2014, 96, 256–262. [CrossRef] 21. Tolstova, Y.; Omelchenko, S.T.; Shing, A.M.; Atwater, H.A. Heteroepitaxial growth of Pt and Au thin films on MgO single crystals by bias-assisted sputtering. Sci. Rep. 2016, 6, 23232. [CrossRef] 22. Gu, H.; You, L.; Leung, K.; Chung, C.; Chan, K.; Lai, J. Growth of TiNiHf shape memory alloy thin films by laser ablation of composite targets. Appl. Surf. Sci. 1998, 127, 579–583. [CrossRef] 23. Dong, J.W.; Chen, L.C.; Palmstro/m, C.J.; James, R.D.; McKernan, S. Molecular beam epitaxy growth of ferromagnetic single crystal (001) Ni2 MnGa on (001) GaAs. Appl. Phys. Lett. 1999, 75, 1443–1445. [CrossRef] Coatings 2019, 9, 308 9 of 10 24. Shih, T.C.; Xie, J.Q.; Dong, J.W.; Dong, X.Y.; Srivastava, S.; Adelmann, C.; McKernan, S.; James, R.D.; Palmstrøm, C.J. Epitaxial growth and characterization of single crystal ferromagnetic shape memory Co2 NiGa films. Ferroelectrics 2006, 342, 35–42. [CrossRef] 25. Kühnemund, L.; Edler, T.; Kock, I.; Seibt, M.; Mayr, S.G. Epitaxial growth and stress relaxation of vapor-deposited Fe-Pd magnetic shape memory films electron beam evaporation. New J. Phys. 2009, 11, 113054. [CrossRef] 26. Gisser, K.R.C.; Busch, J.D.; Johnson, A.D.; Ellis, A.B. Oriented nickel-titanium shape memory alloy films prepared by annealing during deposition. Appl. Phys. Lett. 1992, 61, 1632–1634. [CrossRef] 27. Jenkins, C.A.; Ramesh, R.; Huth, M.; Eichhorn, T.; Pörsch, P.; Elmers, H.J.; Jakob, G. Growth and magnetic control of twinning structure in thin films of Heusler shape memory compound Ni2 MnGa. Appl. Phys. Lett. 2008, 93, 234101. [CrossRef] 28. Niemann, R.; Backen, A.; Kauffmann-Weiss, S.; Behler, C.; Rößler, U.; Seiner, H.; Heczko, O.; Nielsch, K.; Schultz, L.; Fähler, S. Nucleation and growth of hierarchical martensite in epitaxial shape memory films. Acta Mater. 2017, 132, 327–334. [CrossRef] 29. Goryczka, T. Effect of wheel velocity on texture formation and shape memory in Cu-Al-Ni melt-spun ribbons. Arch. Metall. Mater. 2009, 54, 755–763. 30. Delaey, L. Diffusionless transformations. In Phase Tranformation in Materials; Kostorz, G., Ed.; Wiley-VCH: Weinheim, Germany, 2001. 31. Morán, M.; Condó, A.; Soldera, F.; Sirena, M.; Haberkorn, N. Martensitic transformation in freestanding and supported Cu–Al–Ni thin films obtained at low deposition temperatures. Mater. Lett. 2016, 184, 177–180. [CrossRef] 32. Niedbalski, S.; Durán, A.; Walczak, M.; Ramos-Grez, J.A. Laser-assisted synthesis of Cu-Al-Ni shape memory alloys: Effect of inert gas pressure and Ni content. Materials 2019, 12, 794. [CrossRef] 33. Recarte, V.; Pérez-Sáez, R.B.; Juan, J.S.; Bocanegra, E.H.; Nó, M.L. Influence of Al and Ni concentration on the Martensitic transformation in Cu-Al-Ni shape-memory alloys. Met. Mater. Trans. A 2002, 33, 2581–2591. [CrossRef] 34. Braga, F.D.O.; Matlakhov, A.N.; Matlakhova, L.A.; Monteiro, S.N.; De Araújo, C.J. Martensitic transformation under compression of a plasma processed polycrystalline shape memory CuAlNi Alloy. Mater. Res. 2017, 20, 1579–1592. [CrossRef] 35. Haidar, M.A.; Saud, S.N.; Hamzah, E. Microstructure, mechanical properties, and shape memory effect of annealed Cu-Al-Ni-xCo shape memory alloys. Metallogr. Microstruct. Anal. 2018, 7, 57–64. [CrossRef] 36. Saud, S.N.; Bakar, T.A.A.; Hamzah, E.; Ibrahim, M.K.; Bahador, A. Effect of quarterly element addition of cobalt on phase transformation Cu-Al-Ni shape memory alloys. Metall. Mater. Trans. A 2015, 46, 3528–3542. [CrossRef] 37. Sharma, M.; Vajpai, S.K.; Dube, R.K. Processing and characterization of Cu-Al-Ni shape memory alloy strips prepared via a novel powder metallurgy route. Metall. Mater. Trans. A 2010, 41, 2905–2913. [CrossRef] 38. Gómez-Cortés, J.; Juan, J.S.; Lopez, G.A.; Nó, M. Synthesis and characterization of Cu–Al–Ni shape memory alloy multilayer thin films. Thin Solid Films 2013, 544, 588–592. [CrossRef] 39. Canbay, C.A.; Tekatas, A.; Ozkul, I. Fabrication of Cu-Al-Ni shape memory thin film by thermal evaporation. Turk. J. Eng. 2017, 1, 27–32. 40. Minemura, T.; Andoh, H.; Kita, Y.; Ikuta, I. Shape memory effect and microstructures of sputter-deposited Cu-Al-Ni films. J. Mater. Sci. Lett. 1985, 4, 793–796. [CrossRef] 41. Agafonov, V.; Naudot, P.; Dubertret, A.; Dubois, B. Influence of the aluminium content on the appearance and stability of martensites in the Cu Al Ni system. Scr. Met. 1988, 22, 489–494. [CrossRef] 42. Suresh, N.; Ramamurty, U. Aging response and its effect on the functional properties of Cu–Al–Ni shape memory alloys. J. Alloy. Compd. 2008, 449, 113–118. [CrossRef] 43. Sharma, M.; Vajpai, S.K.; Dube, R.K.; Vajpai, S. Synthesis and properties of Cu–Al–Ni shape memory alloy strips prepared via hot densification rolling of powder preforms. Power Metall. 2011, 54, 620–627. [CrossRef] 44. Araujo, A.P.M.; Simões, J.B.; Araújo, C.J. Analysis of compositional modification of commercial aluminum bronzes to obtain functional shape memory properties. Mater. Res. 2017, 20, 331–341. [CrossRef] 45. Malygin, G.A. Nanoscopic size effects on martensitic transformations in shape memory alloys. Phys. Solid State 2008, 50, 1538–1543. [CrossRef] 46. Waitz, T.; Karnthaler, H. Martensitic transformation of NiTi nanocrystals embedded in an amorphous matrix. Acta Mater. 2004, 52, 5461–5469. [CrossRef] Coatings 2019, 9, 308 10 of 10 47. Shi, X.; Cui, L.; Jiang, D.; Yu, C.; Guo, F.; Yu, M.; Ren, Y.; Liu, Y. Grain size effect on the R-phase transformation of nanocrystalline NiTi shape memory alloys. J. Mater. Sci. 2014, 49, 4643–4647. [CrossRef] 48. Wan, D.; Komvopoulos, K. Thickness effect on thermally induced phase transformations in sputtered titanium-nickel shape-memory films. J. Mater. Res. 2005, 20, 1606–1612. [CrossRef] 49. Chen, Y.; Schuh, C.A. Size effects in shape memory alloy microwires. Acta Mater. 2011, 59, 537–553. [CrossRef] © 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).