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New Au–Cu–Al thin film shape memory alloys with tunable functional properties and high thermal stability
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ScienceDirect Acta Materialia 85 (2015) 378–386 www.elsevier.com/locate/actamat
New Au–Cu–Al thin film shape memory alloys with tunable functional properties and high thermal stability ⇑,à
Pio John S. Buenconsejo
and Alfred Ludwig
⇑
Institute for Materials, Ruhr-Universita¨t Bochum, 44801 Bochum, Germany Received 13 August 2014; revised 20 November 2014; accepted 21 November 2014
Abstract—An Au–Cu–Al thin film materials library prepared by combinatorial sputter-deposition was characterized by high-throughput experimentation in order to identify and assess new shape memory alloys (SMAs) in this alloy system. Automated resistance measurements during thermal cycling between 20 and 250 °C revealed a wide composition range that undergoes reversible phase transformations with martensite transformation start temperatures, reverse transformation finish temperatures and transformation hysteresis ranging from 15 to 149 °C, 5 to 185 °C and 8 to 60 K, respectively. High-throughput X-ray diffraction analysis of the materials library confirmed that the phase-transforming compositions can be attributed to the existence of the b-AuCuAl parent phase and its martensite product. The formation of large amount of phases based on face-centered cubic (Au–Cu), Al–Cu and Al–Au is responsible for limiting the range of phase-transforming compositions. Selected alloys in this system show excellent thermal cyclic stability of the phase transformation. The functional properties of these alloys, combined with the inherent properties of Au-based alloys, i.e. aesthetic value, oxidation and corrosion resistance, makes them attractive as smart materials for a wide range of applications, including applications as SMAs for elevated temperatures in harsh environment. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Combinatorial material science; Thin films; Au-based alloys; Shape memory alloys; HTSMA
1. Introduction Au-based alloys are widely used in various applications due to their inherent properties. For example, Au is the noblest and most chemically inert of the known elements [1], so alloys with a high Au content can be used in harsh environments (oxidizing and corrosive). Au is also known to have excellent catalytic properties [2] making it useful for energy generation or conversion. It is also highly sought after for its aesthetic value [3], making it popular in jewelry and ornamental alloys. Pure gold is biocompatible and, when combined with Cu and Ag as an alloy, has been reported to promote the corrosion of Cu and Ag, thus inducing antibacterial function [4]. Combining the excellent inherent properties of Au-based alloys with shape memory properties opens up new opportunities for personal, healthcare, home, energy and transportation applications. “Spangold” is an 18-carat gold jewelry alloy (Au7Cu5Al4 or Au43.75Cu31.25Al25 in at.%) that undergoes a reversible martensitic transformation (martensite transformation start temperature (Ms) = 22 to 30 °C, reverse transformation
⇑ Corresponding
authors; e-mail addresses: [email protected]; [email protected]; [email protected] à Present address: Facility for Analysis Characterisation Testing and Simulation (FACTS), School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore.
finish temperature (Af) = 77 to 90 °C) and exhibits shape memory properties [5–11]. The reversible martensitic transformation is attributed to the b-AuCuAl parent phase, which has a body-centered cubic (bcc) structure [8] which undergoes structural transformation to a martensite phase. Depending on the composition and temperature, the latter may be indexed with orthorhombic or monoclinic unit cells [9]; however, to a first approximation, it may also be considered as a body-centered tetragonal (bct) structure [10]. Here, for simplicity, the bct martensite structure will be used for the purpose of identifying the characteristic X-ray diffraction peaks. Previous studies revealed that the parent phase shows atomic ordering (B2, DO3, L21) that is dependent on the thermal history of the alloys, and it was revealed that L21-type ordering was critical for the alloy to exhibit shape memory behavior [8,11]. Functional properties resulting from its reversible martensitic transformation [12] makes these alloys attractive for use as actuators, superelastic materials and elasto-caloric materials. Examining the published phase diagram of the system Au–Cu–Al reveals that the stability region of the b-AuCuAl phase extends beyond the “Spangold” composition [13–15]. Accordingly, some compositions with reversible phase transformation besides “Spangold” have already been identified [16,17]. However, most of the studies made on this system were on bulk alloys of a limited compositional range, and focused mainly on the “Spangold” composition
http://dx.doi.org/10.1016/j.actamat.2014.11.035 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
P.J.S. Buenconsejo, A. Ludwig / Acta Materialia 85 (2015) 378–386
[5–11,13–28]. The compositional variations reported in the literature are limited to ternary variations (Au30Cu49Al21, with Ms about 350 °C, and Au29Cu47Al24, with Ms about 300 °C) [16], and quaternary additions of Ir and Co over a limited range (Ms < 100 °C) [17]. Moreover, very few studies have been done on melt-spun ribbons [29] and thin films [30]. Although b-AuCuAl exists over a wide range in the phase diagram, no systematic investigation of compositional variation has been reported. The relationship between composition and phase transformation behavior is still unknown. The main purpose of this paper is to explore the Au–Cu–Al ternary system in a large compositional region for potential new thin film shape memory alloys (SMAs) and to establish functional phase diagrams showing the correlations between composition, processing, structure and functional properties by using the thin film combinatorial materials science approach. This approach is an effective way to systematically explore a large compositional region, and it has been successfully applied to explore multi-component systems for the development of new thin film SMAs [31–38]. Although the results of the present study are based on thin films, thin film composition “hits” such as compositions with special and optimized functional properties, discovered by thin film combinatorial study have been successfully up-scaled to bulk scale [36]. Thus the results of this study could also be used as a guide for further development of bulk Au–Cu–Al SMAs. 2. Experimental details The Au–Cu–Al thin film materials library was prepared by using a wedge-type multilayer sputter deposition technique from elemental targets (>99.9% purity). A detailed description of thin film materials library synthesis using this method can be found in Ref. [31]. The substrate was a 4-inch-diameter oxidized Si wafer, with 1.5 lm SiO2 acting as a diffusion and reaction barrier. The surface of the substrate was patterned with photo-resist crosses, which were removed by lifting them off after deposition. The steps created were used for thickness measurements. Each element was deposited sequentially as a thickness wedge. The thickness gradient directions were oriented 120° away from each other, creating a compositional variation over the surface of the substrate. The precursor (Au/Cu/Al)30 multilayer film was annealed in UHV (>1 10 8 torr) at 500 °C for 1 h in order to induce interdiffusion, alloying and phase formation. After annealing, the materials library was left in the chamber to cool to room temperature. The thicknesses measured across the materials library using a profilometer varied between 438 and 870 nm, corresponding to an Au-poor region and an Au-rich region, respectively. The measurement regions of 4.5 mm 4.5 mm were laid over the materials library. Composition analysis was performed by automated energy-dispersive X-ray analysis (EDX) using INCA-EDX software and a scanning electron microscope (JEOL JSM 5800LV) equipped with an Oxford INCA X-act EDX detector (INCA 250 system). X-ray diffraction (XRD) patterns were collected from each measurement region using a fully automated XRD system (PANalytical X’pert PRO system with Cu Ka radiation and an X-ray probe size of 5 mm 3 mm) equipped with an x–y–z stage. Data visualization of the combinatorial data set was carried out using XRDsuite software [39]. Phase
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transformations were characterized by temperature-dependent electrical resistance (R(T)) measurements using a fully automated four-point probe set-up [40]. The temperature cycle range was limited to 20 to 250 °C (heating/cooling rate at 5 K min 1), so phase transformations occurring beyond these limits could not be measured. After the highthroughput screening process, the materials library was diced to separate each measurement region. This was necessary to carry out further experimentation where it is not possible to accommodate the complete materials library. The phase-transforming compositions were selected and were characterized further by temperature-dependent X-ray diffraction (XRD(T)), using a heating–cooling attachment (Anton Paar 450 + LNC) for the XRD system. Selected samples having different transformation behavior were examined further in order to identify their thermal cyclic stability by R(T) measurements.
3. Results and discussion 3.1. Phase transformation properties The Au–Cu–Al materials library was screened for its phase-transforming properties by cyclic R(T) measurements in the range of 20 to 250 °C. Fig. 1 shows typical non-linear R(T) curves confirmed in the materials library, which are indicating reversible phase transformations. Compositions exhibiting a reversible phase transformation have sharp inflection points on heating and cooling, corresponding to the transformation temperatures. The martensite transformation start (Ms) and reverse transformation finish (Af) temperatures were determined by the tangential method. For non-transforming compositions only a linear trend was confirmed. A strong variation of transformation behavior for different compositions is obvious in the R(T) curves, such as high or low transformation temperatures and wide or narrow hysteresis. The compositions showing a reversible transformation as indicated by the R(T) measurements were investigated by XRD(T) to identify the type of structural transformation occurring. Two representative XRD(T) results of
Fig. 1. Typical R(T) curves of selected compositions showing phase transformations. The transformation points, such as Ms and Af are indicated by arrows. DT is defined here as the difference between Ms and Af. The curves are offset vertically for clarity.
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phase-transforming compositions are shown in Fig. 2. The parent phase is indexed as b-AuCuAl [8] and the martensite phase is indexed as bct-AuCuAl [10]. The composition Au57Cu32Al11 (Fig. 2a) was thermally cycled between 50 and 250 °C, and exhibited high transformation temperatures (Ms = 130.6 °C and Af = 163.6 °C). XRD patterns were collected at 25 °C intervals on heating and cooling. It is confirmed that the (2 2 0) peak of b-AuCuAl (parent phase) appeared for temperatures >150 °C, while the (2 0 2) peak of bct-AuCuAl (martensite) completely disappeared at 175 °C. On cooling, the (2 0 2) peak of bctAuCuAl started to appear from 125 °C, while the (2 2 0) of b-AuCuAl disappeared completely at 100 °C. For the composition Au49Cu30Al21 (Fig. 2b), the same trend was confirmed but the structural transformation occurred at lower temperatures (Ms = 1.9 °C and Af = 19.9 °C). The structural transformation temperatures in XRD(T) corresponds very well with the transformation temperatures as determined by R(T) measurement. Fig. 3 summarizes the phase transformation properties confirmed in the Au–Cu–Al materials library: values of Ms, Af and DT are plotted in a ternary composition diagram. Ms varies from 15 to 149 °C, Af varies from 5 to 183 °C and DT varies from 8 to 66 K. The variation of Ms is strongly influenced by the Al content: alloys with <15 at.% Al exhibit high transformation temperatures. A similar trend could be observed for Af temperatures with respect to the Al content. Compositions with Al >15 at.% and Au <50 at.% have higher Af temperatures compared to compositions with Au >50 at.%. Thus there is a slight dependence of DT (=Af Ms) with respect to the Au content. Narrow hysteresis values were confirmed for alloys with Al content close to 15 at.%, whereas wide hysteresis values were measured for alloys with Au <50 at.% and Al
>15 at.%. The complete phase-transforming composition region as determined within the limits of this study ( 20 °C to 250 °C) is confined within the range of 35–70 at.% Au and 9–23 at.% Al, with the balance Cu. It is interesting to note that the Ms temperature changes strongly with Al content. The AuCuAl alloy is considered to be a b electron compound [8]. Therefore its stability is influenced by the valence electron per atom ratio (e/a). The valence electron numbers for Au and Cu are both equal to 1, but that of Al is 3. Thus the Al content changes the e/a ratio of the alloy. The Ms and Af temperatures for all phase-transforming compositions are plotted as a function of their e/a ratio in Fig. 4. The figure clearly shows an interesting trend whereby at e/a < 1.35 the transformation temperatures increases almost linearly with decreasing e/a, while at e/a > 1.35 the transformation temperatures reach a plateau. The scatter in the data plots may be explained by the existence of mixed phases causing a shift in the nominal composition of the b-AuCuAl phase. For alloys with a higher concentration of valence electrons, the b-AuCuAl phase is stabilized, thus Ms decreases with increasing e/a. The range where it reaches a plateau may indicate the critical concentration of valence electrons in the alloy. The composition range of potential SMAs confirmed in this study is wider than previously reported [5–11,16,17]. Interestingly, the transformation properties can be tuned by adjusting the composition, making these alloys suitable for a variety of potential applications, like devices operating at body temperature or device components operating at elevated temperatures (>100 °C). The transformation hysteresis (DT) value is also tunable between 8 and 66 K, where compositions with a narrow DT are suitable for high-speed microactuators while compositions with a wide DT are excellent for superelastic applications.
Fig. 2. XRD(T) profiles of the compositions (a) Au57Cu32Al11 and (b) Au49Cu30Al21 undergoing reversible structural transformations on heating and cooling between a b-AuCuAl parent phase and a bct-AuCuAl martensite phase.
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Fig. 3. Phase transformation properties measured on the Au–Cu–Al materials library, which was annealed at 500 °C for 1 h. Color-coded plots of (a) Ms, (b) Af and (c) DT over composition visualize the compositional dependence of the phase transformation properties. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 4. Transformation temperatures (Ms, Af) as a function of the e/a ratio (valence electron per atom ratio).
3.2. Phase constitution A high-throughput XRD analysis was performed on the Au–Cu–Al materials library in order to investigate the underlying relationship between functional properties and phase constitution. The results are visualized as colorcoded peak intensity profiles of the predominant phases according to their characteristic Bragg diffraction peaks. In all of the peak intensity profiles it should be noted that it is difficult to draw out exact phase boundaries due to:
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phase mixtures, similarity of diffraction peaks of different phases, peak shifts due to alloying and “shoulder effects” as influenced by neighboring peaks. For this reason, the peak intensity profiles in Fig. 5 are sometimes attributed to more than one phase, but they are differentiated according to their composition. For example in Fig. 5a the high peak intensity for compositions with <40 at.% Au is due to (1 1 1) Cu3Au. This peak intensity profile is also due to (2 1 1) AuAl and (2 2 0) b-AuCuAl for compositions having Au contents between 35 and 50 at.%. Nevertheless, the results clearly show the existence of predominant phases in certain composition ranges, and it is sufficient in this high-throughput study to make qualitative composition– structure–property relationships. For reference, the phasetransforming composition boundary as determined by R(T) screening is represented by a black dashed boundary line plotted on each peak intensity profile. The phase transforming phase is the b-AuCuAl (Fig. 5a) parent phase, which undergoes martensitic transformation to a bct-AuCuAl (Fig. 5b) martensite product. In bulk alloys the b-AuCuAl phase was reported to undergo order–disorder transformation from a disordered bcc (A2) at high temperatures to B2, L21 or DO3, depending on the thermal treatment history, and it has been suggested that L21 ordering was necessary for the appearance of martensitic transformation [8,11]. However, it was difficult to determine the type of ordering from the XRD patterns due to preferred orientation occurring in thin films and overlapping with neighbor diffraction peaks. To identify this phase regardless of the ordering state, the peak intensity profile for the (2 2 0) b-AuCuAl phase at 2h = 41.3° is selected since it is a diffraction peak that is common to them all and is the most prominent. This phase is predominant in the upper half of the phase-transforming composition range. This is because the compositions within this region have Af below or near room temperature (Fig. 3b). The martensite phase is indexed as a bct-AuCuAl phase, so the transformation from parent to martensite will split the (2 2 0) b-AuCuAl peak into (2 0 2) and (2 2 0) bctAuCuAl peaks. The (2 0 2) bct-AuCuAl peak was selected since it was most prominent (Fig. 5b), and it is predominant in the lower part of the phase-transforming composition range. This composition range shows the martensite phase since the Ms > room temperature (Fig. 3a). The composition range of the b/bct-AuCuAl phase found in this study did not coincide exactly with the phase diagram [13–15]. Possible reasons for this are as follows: (i) compositions measured by the EDX method have an error of ±1–2 at.%, causing a slight shift in the actual composition of the alloys; or (ii) the formation of phases by annealing a multilayer thin film precursor undergoes a different phase formation mechanism and the resulting phases do not necessarily correspond to the same phases formed by bulk metallurgical methods. Within the phase-transforming composition range, the b/bct-AuCuAl phase coexists to a large extent with secondary phases. For example, in Fig. 2, the XRD pattern shows the coexistence of several phases for the compositions Au57Cu32Al11 and Au49Cu30Al21. Outside the phasetransforming composition boundary, non-transforming phases are predominant, such as the intermetallic compounds of the Au–Al and Cu–Al systems, and Au–Cu solid solution. These were identified using an XRD database [41]. In the Au–Al part of the system, the intermetallic phases Au2Al (Fig. 5b, e and g), AuAl (Fig. 5a and c), AuAl2
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(Fig. 5d), Au4Al (Fig. 5e) and Au8Al3 (Fig. 5f and g) were confirmed. These phases formed on the upper and right side of the phase-transforming composition boundary. In the Cu–Al part of the system, the intermetallic phase Cu4Al (Fig. 5i) was confirmed, and this formed on the left side of the phase-transforming composition boundary. In the Au–Cu part of the system, face-centered cubic (fcc)-AuCu solid solution phases (Fig. 5a, b, e, f, g, h and i) were confirmed. They form on the lower side of the phasetransforming composition boundary. The fcc-AuCu phase is represented with different peak intensity plots due to a significant peak shift. The large difference in the atomic radii of Au and Cu causes a significant peak shift due to change in lattice constant with composition [42]. For example, the (1 1 1) fcc peak varies from about 39° at the Au-rich side to about 42° at the Cu-rich side. Au–Cu compositions at Au = 70, 50 and 33 at.% also undergo ordering, but in this study the ordering of the binary phases was not investigated further. In conjunction with R(T) screening and the XRD mapping results, it is concluded that the phase-transforming
composition range is limited by the formation of a large amount of non-transforming binary intermetallics and solid solutions. 3.3. Alloying behavior of b-AuCuAl The alloying behavior of b-AuCuAl is investigated in more detail by determining the influence of composition on the lattice constants. The XRD patterns collected at room temperature show either the b-AuCuAl phase or the bct-AuCuAl phase, depending on the Ms temperature. The phase-transforming compositions were selected after dicing the materials library. XRD patterns were recorded at 250 °C, at which temperature all the samples are in the b-AuCuAl phase (>Af). A preferred orientation was observed for all samples, therefore the lattice constant of the b-AuCuAl phase was estimated using only a single peak. The preferred orientation of most samples was (2 2 0), so the lattice constant was evaluated using this peak. However, on two samples the (2 2 0) peak did not appear due to a very strong (2 2 2) peak, so this peak was used
Fig. 5. Phase constitution map of phases confirmed in the Au–Cu–Al materials library. Peak intensities at characteristic 2h positions were plotted with respect to composition. The peak intensity color codes are as follows: red = high, green = medium, blue = low. The color-coded scheme identifies the most prominent phase with respect to composition. The black dashed line encircles the phase transforming region as identified with R(T). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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instead. Fig. 6 shows the plot of b-AuCuAl lattice constants (ab) measured at 250 °C with respect to composition. The lattice constants at 250 °C were all mostly higher than the reported values of 0.616–0.619 nm, which were measured at lower temperatures and for different compositions [8]. The increased values are due to the thermal expansion of the unit cell and a higher Au content. The plotted results show the compositional dependence of ab only, since the temperature was fixed at 250 °C. As a function of Au content (Fig. 6a), ab follows a decreasing trend for Au <50 at.% and plateaus for Au >50 at.%. As a function of Cu content (Fig. 6b), ab follows a decreasing trend for Cu >30 at.% while it plateaus for Cu <30 at.%. However, there are no clear trends confirmed with respect to Al content (Fig. 6c). The inflection point (Fig. 6a and b), where the trend changes from plateau to decreasing when Au <50 at.% and Cu >30 at.%, indicates a nominal composition of Au50Cu30Al20 for the b-AuCuAl phase. To observe a more meaningful relationship, ab was plotted as a function of the Au/Cu ratio (Fig. 6d). A clear trend is confirmed with respect to the Au/Cu ratio: ab decreases for Au/Cu <2 and it saturates to about 0.625 nm for Au/Cu >2. In a solid solution phase, like b-AuCuAl, changes in lattice constants are strongly dependent on the solubility of constituent elements. Thus the observation that the lattice constants reach a plateau for >50 at.% Au could be interpreted as the maximum amount of Au soluble in the b-AuCuAl phase. Moreover, the decreasing trend of ab for <50 at.% Au means that the Au site in the unit cell is substituted with a smaller atom. Comparing the metallic radius [43] of the constituent elements, Au (144 pm) and
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Al (143 pm) are similar, but significantly larger than Cu (128 pm). It is therefore reasonable that ab decreases with increasing substitution of Cu for Au. This is also consistent with the reported site occupancy in the crystal structure of b-AuCuAl phase [8], where the Au-rich site (C site) has the following site occupancy: Au = 0.653, Cu = 0.328, Al = 0.019. This means that Cu substitution for Au at the Au-rich site is more likely than Al substitution for Au. Furthermore, the ratio of Au/Cu site occupancy is 1.99, which coincides perfectly with the inflection point in Fig. 6d. 3.4. Visual appearance of the Au–Cu–Al materials library The aesthetic value of Au-based alloys is characterized by their color, so a photograph of the Au–Cu–Al thin film materials library annealed at 500 °C after complete R(T) screening was taken (Fig. 7a). The materials library was placed on a white paper with markers. The photograph was taken under neutral color temperature in order to be able to do a qualitatively correct color interpretation. The measurement region where phase transformations were confirmed is within the dashed black boundary. A clear color variation is observable over the materials library indicating a composition gradient. This is a well-known phenomenon in Au-based jewelry and in this alloy system, where color changes as a result of changes in the composition and phase constitution [3]. The dark blue area in the Cu-rich region could be attributed to Cu-oxide that formed on the surface as a result of repeated thermal cycling (up to 250 °C) during R(T) screening. The yellow-gold color is dominant for regions close to the Au-rich side. The color
Fig. 6. Lattice constants (ab) of b-AuCuAl measured at 250 °C plotted as a function of (a) Au content, (b) Cu content, (c) Al content and (d) Au/Cu ratio.
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Fig. 7. (a) Photograph of the Au–Cu–Al materials library after the R(T) screening step. Crosses mark the corners of the measurement regions (samples). Within the dashed boundary is the region where phase transformations were confirmed. (b) The corresponding plot of Ms temperature over the physical coordinates of the materials library.
becomes less yellow and more white when going towards the Al-rich side. In this direction a region of reddish yellow color is apparent, which falls partly within the phasetransforming composition range. When compared with the Ms temperature (Fig. 7b), the regions with high transformation temperatures exhibit a yellow-gold color, which is due to the higher Au–Cu content. On the other hand, the regions with lower transformation temperatures have a wider variety of colors, such as yellow-gold, light yellow-gold, whitish gold and reddish yellow-gold. These alloys could be used for various ornamental and jewelry applications. Functions derived from reversible martensitic transformation, such as shape memory effect, superelasticity and elasto-caloric effects, could be incorporated to provide new applications, such as smart ornaments and jewelry. 3.5. Thermal cyclic stability One of the most important considerations for SMA application is the cyclic stability of its functional property. During the R(T) screening process the whole materials library experiences about 52 thermal cycles. For screening purposes, the in-house R(T) test stand is equipped with a probe head having five sets of four-point electrical resistance probes, where each set can measure areas with 4.5 mm separation. It can thus measure five compositions simultaneously. The materials library is fixed by vacuum on the
hot/cold plate, so all compositions are simultaneously subjected to thermal cycling. This is one limitation of the R(T) screening test stand. To investigate further the stability of the observed reversible phase transformations, the materials library was diced to separate each measurement region corresponding to one composition (4.5 mm 4.5 mm size). Selected compositions were individually subjected to further thermal cycles using a probe head with a single set of fourpoint electrical resistance probes. Fig. 8 summarizes the R(T) thermal cycle test made after the R(T) screening. Raw data are given in Supp. Fig. 1. The selected compositions are Au55Cu34Al11, which has a high transformation temperature, and Au44Cu43Al13, which has a low transformation temperature. The difference between the first and fortieth cycles for Au55Cu34Al11 is minimal, but the hysteresis has narrowed slightly. Compared to the R(T) curve measured during R(T) screening, the difference in the transformation curve is likewise minimal. For Au44Cu43Al13, the first to third cycles exhibited a noisy signal (Supp. Fig. 1) around the transition point, so the fourth cycle was used for comparison instead. The change in the R(T) curves was minimal for the fourth and fortieth cycles, and likewise when compared to the R(T) curve measured during R(T) screening. The transformation temperatures with respect to thermal cycles for both compositions are compared in Fig. 8c. It can be clearly seen that for both compositions the transformation temperatures did not significantly change with thermal
Fig. 8. Thermal cycling test results of the selected compositions (a) Au55Cu34Al11 and (b) Au44Cu43Al13. The R(T) curve measured during the R(T) screening test is also shown as a reference. (c) The transformation temperatures (Ms and Af) are plotted with respect to thermal cycles. The measured values obtained during the R(T) screening are plotted as the zeroth cycle.
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cycles. Furthermore, it should be noted that during the R(T) screening all samples experienced several thermal cycles. For comparison, the transformation temperatures measured during R(T) screening are plotted as the zeroth cycle. The very small difference (Ms decreased by 4.8 °C for Au55Cu34Al11 and 4.1 °C for Au44Cu43Al13) between the zeroth cycle and the succeeding 40 cycles suggests that the phase transformation behavior of these alloys shows excellent thermal stability. The origin of this excellent thermal stability could not be explained within the scope of this study, but it will be a topic for future investigations on this new SMA system.
4. Conclusions The Au–Cu–Al system annealed at 500 °C for 1 h was systematically investigated by the thin film combinatorial materials science approach. As a result, very promising new SMAs with tunable functional properties and very stable functional properties were identified and characterized. This includes SMAs that are functional at room temperature and in a hot environment (>100 °C). The following conclusions are drawn from this study: 1. Reversible phase transformation over a wide range of compositions (35–70 at.% Au, 9–23 at.% Al and Cu balance) were identified by R(T) measurements. The phase transformation temperatures and hysteresis values span the following ranges: Ms = 15 to 149 °C, Af = 5 to 183 °C and DT = 8 to 66 K. The variation of these properties is strongly composition dependent. This means that it is possible to tune the desired property for a wide variety of potential applications. 2. The phase constitution across the materials library was investigated by high-throughput XRD analysis. The results revealed the correlation of phase constitution to composition and functional properties. Reversible phase transformation is attributed to the presence of b-AuCuAl and its martensite product. The formation of a large amount of non-transforming phases – Au– Al, Cu–Al and Au–Cu binary phases – was found to limit the range of compositions that show reversible phase transformations. 3. The lattice constants (ab) of the b-AuCuAl phase at 250 °C were found to vary with Au and Cu content. For compositions with >50 at.% Au and <30 at.% Cu, the lattice constant does not significantly change. However, for compositions with <50 at.% Au and >30 at.% Cu, the lattice constant decreases. This behavior is due to preferential substitution of Cu for Au on the Au-rich site of the b-AuCuAl unit cell with composition change. 4. The visual appearance of thin films within the transforming region varies with composition and shows different functional properties. The compositions that show high transformation temperatures are yellow-gold in color due to high Au and low Al contents. The compositions that show a low transformation temperature have various colors, such as light yellow-gold, reddish yellow-gold and whitish gold. 5. The thermal cyclic stability of two selected compositions with high (Au55Cu34Al11) and low (Au44Cu43Al13) transformation temperatures were investigated by R(T) thermal cycling test up to 40 thermal cycles. Both compositions show minimal change in their transformation
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temperatures during thermal cycles, suggesting that they have excellent thermal cyclic stability. Acknowledgements The authors would like to express their gratitude for support from the German Research Foundation (DFG) within the funding for the research unit FOR 1766 (High temperature shape memory alloys).
Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/ j.actamat.2014.11.035. References [1] B. Hammer, J.K. Norskov, Nature 376 (1995) 238–240. [2] M.B. Cortie, E. van der Lingen, Mater. Forum 26 (2002) 1– 14. [3] C. Cretu, E. van der Lingen, Gold Bull. 32 (4) (1999) 115–126. [4] D.C. Wright, R.F. Gallant, L. Spangberg, J. Biomed. Mater. Res. 16 (1982) 509–517. [5] I.M. Wolff, M.B. Cortie, Gold Bull. 27 (1994) 44–54. [6] M. Cortie, I. Wolff, F. Levey, S. Taylor, R. Watt, R. Pretorius, T. Biggs, J. Hurly, Gold Technol. 14 (1994) 30–36. [7] F.C. Levey, M.B. Cortie, L.A. Cornish, Metall. Mater. Trans. A 31 (2000) 1917–1923. [8] M.B. Cortie, F.C. Levey, Intermetallics 8 (2000) 793–804. [9] M.M. Elcombe, C.S. Kealley, V.K. Bhatia, G.J. Thorogood, D.J. Carter, M. Avdeev, M.B. Cortie, Acta Mater. 79 (2014) 234–240. [10] F.C. Levey, M.B. Cortie, Mater. Sci. Eng. A 303 (2001) 1–10. [11] M. Jin, J. Liu, X. Jin, Intermetallics 18 (2010) 846–850. [12] K. Otsuka, C.M. Wayman (Eds.), Shape Memory Materials, Cambridge University Press, Cambridge, 1998. [13] F.C. Levey, M.B. Cortie, L.A. Cornish, Metall. Mater. Trans. A 33 (2002) 987–993 (Isothermal section 500C). [14] F.C. Levey, M.B. Cortie, L.A. Cornish, J. Alloys Compd. 354 (2003) 171–180. [15] V.K. Bhatia, C.S. Kealley, R. Wuhrer, K.S. Wallwork, M.B. Cortie, J. Alloys Compd. 488 (2009) 100–107. [16] V.K. Bhatia, C.S. Kealley, A. Dowd, M.B. Cortie, in: 33rd Annual Condensed Matter and Materials Meeting, 4-6 February 2009, Australian Institute of Physics, Wagga Wagga, NSW, Australia, 2009, p. F8. [17] L. Fumagalli, S. Besseghini, F. Passaretti, G. Airoldi, J. Alloys Compd. 433 (2007) 332–337. [18] S. Urbano, A. Manca, S. Besseghini, G. Airoldi, Scr. Mater. 52 (2005) 317–321. [19] V.K. Bhatia, C.S. Kealley, M.J. Prior, M.B. Cortie, Acta Mater. 59 (2011) 2193–2200. [20] M.B. Cortie, F.C. Levey, Intermetallics 10 (2002) 23–31. [21] Y. Gu, M. Jin, X. Jin, Intermetallics 17 (2009) 704–707. [22] G. Fiore, L. Battezzati, Intermetallics 19 (2011) 1978–1982. [23] M.B. Cortie, C.S. Kealley, V. Bhatia, G.J. Thorogood, M.M. Elcombe, M. Avdeev, J. Alloys Compd. 509 (2011) 3502– 3508. [24] X.J. Jin, J.Y. Liu, M.J. Jin, J. Alloys Compd. 577 (2013) S155–S158. [25] M.J. Jin, Y.J. Gu, X. Jin, Mater. Charact. 60 (2009) 1395– 1399. [26] M.J. Witcomb, R.V. Finnie, L.A. Cornish, J. Mater. Sci. Lett. 16 (1997) 674–677. [27] J.Y. Liu, Z.Q. Liu, X.J. Jin, Philos. Mag. 94 (2014) 56–72. [28] F.C. Levey, M.B. Cortie, L.A. Cornish, Scr. Mater. 47 (2002) 95–100.
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[29] L. Battezzati, G. Fiore, M. Massazza, J. Alloys Compd. 434– 435 (2007) 262–267. [30] V. Bhatia, G. Thorogood, A. Dowd, M.B. Cortie, MRS Proc. 1295 (2011) 33–38. [31] A. Ludwig, R. Zarnetta, S. Hamann, A. Savan, S. Thienhaus, Int. J. Mater. Res. 99 (2008) 1144–1149. [32] J. Cui, Y. Chu, O.O. Famodu, Y. Furuya, J. HatttrickSimpers, R.D. James, A. Ludwig, S. Thienhaus, M. Wuttig, Z. Zhang, I. Takeuchi, Nat. Mater. 5 (2006) 286–290. [33] D. Ko¨nig, R. Zarnetta, A. Savan, H. Brunken, A. Ludwig, Acta Mater. 59 (2011) 3267–3275. [34] R. Zarnetta, P.J.S. Buenconsejo, A. Savan, S. Thienhaus, A. Ludwig, Intermetallics 26 (2012) 98–109. [35] P.J.S. Buenconsejo, R. Zarnetta, D. Ko¨nig, A. Savan, S. Thienhaus, A. Ludwig, Adv. Funct. Mater. 21 (2011) 113– 118. [36] R. Zarnetta, R. Takahash, V. Srivastava, M.L. Young, A. Savan, Y. Furuya, S. Thienhaus, B. Maaß, M. Rahim, J. Frenzel, H. Brunken, Y.S. Chu, R.D. James, I. Takeuchi, G. Eggeler, A. Ludwig, Adv. Funct. Mater. 20 (2010) 1917–1923.
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ScienceDirect Acta Materialia 85 (2015) 378–386 www.elsevier.com/locate/actamat
New Au–Cu–Al thin film shape memory alloys with tunable functional properties and high thermal stability ⇑,à
Pio John S. Buenconsejo
and Alfred Ludwig
⇑
Institute for Materials, Ruhr-Universita¨t Bochum, 44801 Bochum, Germany Received 13 August 2014; revised 20 November 2014; accepted 21 November 2014
Abstract—An Au–Cu–Al thin film materials library prepared by combinatorial sputter-deposition was characterized by high-throughput experimentation in order to identify and assess new shape memory alloys (SMAs) in this alloy system. Automated resistance measurements during thermal cycling between 20 and 250 °C revealed a wide composition range that undergoes reversible phase transformations with martensite transformation start temperatures, reverse transformation finish temperatures and transformation hysteresis ranging from 15 to 149 °C, 5 to 185 °C and 8 to 60 K, respectively. High-throughput X-ray diffraction analysis of the materials library confirmed that the phase-transforming compositions can be attributed to the existence of the b-AuCuAl parent phase and its martensite product. The formation of large amount of phases based on face-centered cubic (Au–Cu), Al–Cu and Al–Au is responsible for limiting the range of phase-transforming compositions. Selected alloys in this system show excellent thermal cyclic stability of the phase transformation. The functional properties of these alloys, combined with the inherent properties of Au-based alloys, i.e. aesthetic value, oxidation and corrosion resistance, makes them attractive as smart materials for a wide range of applications, including applications as SMAs for elevated temperatures in harsh environment. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Combinatorial material science; Thin films; Au-based alloys; Shape memory alloys; HTSMA
1. Introduction Au-based alloys are widely used in various applications due to their inherent properties. For example, Au is the noblest and most chemically inert of the known elements [1], so alloys with a high Au content can be used in harsh environments (oxidizing and corrosive). Au is also known to have excellent catalytic properties [2] making it useful for energy generation or conversion. It is also highly sought after for its aesthetic value [3], making it popular in jewelry and ornamental alloys. Pure gold is biocompatible and, when combined with Cu and Ag as an alloy, has been reported to promote the corrosion of Cu and Ag, thus inducing antibacterial function [4]. Combining the excellent inherent properties of Au-based alloys with shape memory properties opens up new opportunities for personal, healthcare, home, energy and transportation applications. “Spangold” is an 18-carat gold jewelry alloy (Au7Cu5Al4 or Au43.75Cu31.25Al25 in at.%) that undergoes a reversible martensitic transformation (martensite transformation start temperature (Ms) = 22 to 30 °C, reverse transformation
⇑ Corresponding
authors; e-mail addresses: [email protected]; [email protected]; [email protected] à Present address: Facility for Analysis Characterisation Testing and Simulation (FACTS), School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore.
finish temperature (Af) = 77 to 90 °C) and exhibits shape memory properties [5–11]. The reversible martensitic transformation is attributed to the b-AuCuAl parent phase, which has a body-centered cubic (bcc) structure [8] which undergoes structural transformation to a martensite phase. Depending on the composition and temperature, the latter may be indexed with orthorhombic or monoclinic unit cells [9]; however, to a first approximation, it may also be considered as a body-centered tetragonal (bct) structure [10]. Here, for simplicity, the bct martensite structure will be used for the purpose of identifying the characteristic X-ray diffraction peaks. Previous studies revealed that the parent phase shows atomic ordering (B2, DO3, L21) that is dependent on the thermal history of the alloys, and it was revealed that L21-type ordering was critical for the alloy to exhibit shape memory behavior [8,11]. Functional properties resulting from its reversible martensitic transformation [12] makes these alloys attractive for use as actuators, superelastic materials and elasto-caloric materials. Examining the published phase diagram of the system Au–Cu–Al reveals that the stability region of the b-AuCuAl phase extends beyond the “Spangold” composition [13–15]. Accordingly, some compositions with reversible phase transformation besides “Spangold” have already been identified [16,17]. However, most of the studies made on this system were on bulk alloys of a limited compositional range, and focused mainly on the “Spangold” composition
http://dx.doi.org/10.1016/j.actamat.2014.11.035 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
P.J.S. Buenconsejo, A. Ludwig / Acta Materialia 85 (2015) 378–386
[5–11,13–28]. The compositional variations reported in the literature are limited to ternary variations (Au30Cu49Al21, with Ms about 350 °C, and Au29Cu47Al24, with Ms about 300 °C) [16], and quaternary additions of Ir and Co over a limited range (Ms < 100 °C) [17]. Moreover, very few studies have been done on melt-spun ribbons [29] and thin films [30]. Although b-AuCuAl exists over a wide range in the phase diagram, no systematic investigation of compositional variation has been reported. The relationship between composition and phase transformation behavior is still unknown. The main purpose of this paper is to explore the Au–Cu–Al ternary system in a large compositional region for potential new thin film shape memory alloys (SMAs) and to establish functional phase diagrams showing the correlations between composition, processing, structure and functional properties by using the thin film combinatorial materials science approach. This approach is an effective way to systematically explore a large compositional region, and it has been successfully applied to explore multi-component systems for the development of new thin film SMAs [31–38]. Although the results of the present study are based on thin films, thin film composition “hits” such as compositions with special and optimized functional properties, discovered by thin film combinatorial study have been successfully up-scaled to bulk scale [36]. Thus the results of this study could also be used as a guide for further development of bulk Au–Cu–Al SMAs. 2. Experimental details The Au–Cu–Al thin film materials library was prepared by using a wedge-type multilayer sputter deposition technique from elemental targets (>99.9% purity). A detailed description of thin film materials library synthesis using this method can be found in Ref. [31]. The substrate was a 4-inch-diameter oxidized Si wafer, with 1.5 lm SiO2 acting as a diffusion and reaction barrier. The surface of the substrate was patterned with photo-resist crosses, which were removed by lifting them off after deposition. The steps created were used for thickness measurements. Each element was deposited sequentially as a thickness wedge. The thickness gradient directions were oriented 120° away from each other, creating a compositional variation over the surface of the substrate. The precursor (Au/Cu/Al)30 multilayer film was annealed in UHV (>1 10 8 torr) at 500 °C for 1 h in order to induce interdiffusion, alloying and phase formation. After annealing, the materials library was left in the chamber to cool to room temperature. The thicknesses measured across the materials library using a profilometer varied between 438 and 870 nm, corresponding to an Au-poor region and an Au-rich region, respectively. The measurement regions of 4.5 mm 4.5 mm were laid over the materials library. Composition analysis was performed by automated energy-dispersive X-ray analysis (EDX) using INCA-EDX software and a scanning electron microscope (JEOL JSM 5800LV) equipped with an Oxford INCA X-act EDX detector (INCA 250 system). X-ray diffraction (XRD) patterns were collected from each measurement region using a fully automated XRD system (PANalytical X’pert PRO system with Cu Ka radiation and an X-ray probe size of 5 mm 3 mm) equipped with an x–y–z stage. Data visualization of the combinatorial data set was carried out using XRDsuite software [39]. Phase
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transformations were characterized by temperature-dependent electrical resistance (R(T)) measurements using a fully automated four-point probe set-up [40]. The temperature cycle range was limited to 20 to 250 °C (heating/cooling rate at 5 K min 1), so phase transformations occurring beyond these limits could not be measured. After the highthroughput screening process, the materials library was diced to separate each measurement region. This was necessary to carry out further experimentation where it is not possible to accommodate the complete materials library. The phase-transforming compositions were selected and were characterized further by temperature-dependent X-ray diffraction (XRD(T)), using a heating–cooling attachment (Anton Paar 450 + LNC) for the XRD system. Selected samples having different transformation behavior were examined further in order to identify their thermal cyclic stability by R(T) measurements.
3. Results and discussion 3.1. Phase transformation properties The Au–Cu–Al materials library was screened for its phase-transforming properties by cyclic R(T) measurements in the range of 20 to 250 °C. Fig. 1 shows typical non-linear R(T) curves confirmed in the materials library, which are indicating reversible phase transformations. Compositions exhibiting a reversible phase transformation have sharp inflection points on heating and cooling, corresponding to the transformation temperatures. The martensite transformation start (Ms) and reverse transformation finish (Af) temperatures were determined by the tangential method. For non-transforming compositions only a linear trend was confirmed. A strong variation of transformation behavior for different compositions is obvious in the R(T) curves, such as high or low transformation temperatures and wide or narrow hysteresis. The compositions showing a reversible transformation as indicated by the R(T) measurements were investigated by XRD(T) to identify the type of structural transformation occurring. Two representative XRD(T) results of
Fig. 1. Typical R(T) curves of selected compositions showing phase transformations. The transformation points, such as Ms and Af are indicated by arrows. DT is defined here as the difference between Ms and Af. The curves are offset vertically for clarity.
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phase-transforming compositions are shown in Fig. 2. The parent phase is indexed as b-AuCuAl [8] and the martensite phase is indexed as bct-AuCuAl [10]. The composition Au57Cu32Al11 (Fig. 2a) was thermally cycled between 50 and 250 °C, and exhibited high transformation temperatures (Ms = 130.6 °C and Af = 163.6 °C). XRD patterns were collected at 25 °C intervals on heating and cooling. It is confirmed that the (2 2 0) peak of b-AuCuAl (parent phase) appeared for temperatures >150 °C, while the (2 0 2) peak of bct-AuCuAl (martensite) completely disappeared at 175 °C. On cooling, the (2 0 2) peak of bctAuCuAl started to appear from 125 °C, while the (2 2 0) of b-AuCuAl disappeared completely at 100 °C. For the composition Au49Cu30Al21 (Fig. 2b), the same trend was confirmed but the structural transformation occurred at lower temperatures (Ms = 1.9 °C and Af = 19.9 °C). The structural transformation temperatures in XRD(T) corresponds very well with the transformation temperatures as determined by R(T) measurement. Fig. 3 summarizes the phase transformation properties confirmed in the Au–Cu–Al materials library: values of Ms, Af and DT are plotted in a ternary composition diagram. Ms varies from 15 to 149 °C, Af varies from 5 to 183 °C and DT varies from 8 to 66 K. The variation of Ms is strongly influenced by the Al content: alloys with <15 at.% Al exhibit high transformation temperatures. A similar trend could be observed for Af temperatures with respect to the Al content. Compositions with Al >15 at.% and Au <50 at.% have higher Af temperatures compared to compositions with Au >50 at.%. Thus there is a slight dependence of DT (=Af Ms) with respect to the Au content. Narrow hysteresis values were confirmed for alloys with Al content close to 15 at.%, whereas wide hysteresis values were measured for alloys with Au <50 at.% and Al
>15 at.%. The complete phase-transforming composition region as determined within the limits of this study ( 20 °C to 250 °C) is confined within the range of 35–70 at.% Au and 9–23 at.% Al, with the balance Cu. It is interesting to note that the Ms temperature changes strongly with Al content. The AuCuAl alloy is considered to be a b electron compound [8]. Therefore its stability is influenced by the valence electron per atom ratio (e/a). The valence electron numbers for Au and Cu are both equal to 1, but that of Al is 3. Thus the Al content changes the e/a ratio of the alloy. The Ms and Af temperatures for all phase-transforming compositions are plotted as a function of their e/a ratio in Fig. 4. The figure clearly shows an interesting trend whereby at e/a < 1.35 the transformation temperatures increases almost linearly with decreasing e/a, while at e/a > 1.35 the transformation temperatures reach a plateau. The scatter in the data plots may be explained by the existence of mixed phases causing a shift in the nominal composition of the b-AuCuAl phase. For alloys with a higher concentration of valence electrons, the b-AuCuAl phase is stabilized, thus Ms decreases with increasing e/a. The range where it reaches a plateau may indicate the critical concentration of valence electrons in the alloy. The composition range of potential SMAs confirmed in this study is wider than previously reported [5–11,16,17]. Interestingly, the transformation properties can be tuned by adjusting the composition, making these alloys suitable for a variety of potential applications, like devices operating at body temperature or device components operating at elevated temperatures (>100 °C). The transformation hysteresis (DT) value is also tunable between 8 and 66 K, where compositions with a narrow DT are suitable for high-speed microactuators while compositions with a wide DT are excellent for superelastic applications.
Fig. 2. XRD(T) profiles of the compositions (a) Au57Cu32Al11 and (b) Au49Cu30Al21 undergoing reversible structural transformations on heating and cooling between a b-AuCuAl parent phase and a bct-AuCuAl martensite phase.
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Fig. 3. Phase transformation properties measured on the Au–Cu–Al materials library, which was annealed at 500 °C for 1 h. Color-coded plots of (a) Ms, (b) Af and (c) DT over composition visualize the compositional dependence of the phase transformation properties. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 4. Transformation temperatures (Ms, Af) as a function of the e/a ratio (valence electron per atom ratio).
3.2. Phase constitution A high-throughput XRD analysis was performed on the Au–Cu–Al materials library in order to investigate the underlying relationship between functional properties and phase constitution. The results are visualized as colorcoded peak intensity profiles of the predominant phases according to their characteristic Bragg diffraction peaks. In all of the peak intensity profiles it should be noted that it is difficult to draw out exact phase boundaries due to:
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phase mixtures, similarity of diffraction peaks of different phases, peak shifts due to alloying and “shoulder effects” as influenced by neighboring peaks. For this reason, the peak intensity profiles in Fig. 5 are sometimes attributed to more than one phase, but they are differentiated according to their composition. For example in Fig. 5a the high peak intensity for compositions with <40 at.% Au is due to (1 1 1) Cu3Au. This peak intensity profile is also due to (2 1 1) AuAl and (2 2 0) b-AuCuAl for compositions having Au contents between 35 and 50 at.%. Nevertheless, the results clearly show the existence of predominant phases in certain composition ranges, and it is sufficient in this high-throughput study to make qualitative composition– structure–property relationships. For reference, the phasetransforming composition boundary as determined by R(T) screening is represented by a black dashed boundary line plotted on each peak intensity profile. The phase transforming phase is the b-AuCuAl (Fig. 5a) parent phase, which undergoes martensitic transformation to a bct-AuCuAl (Fig. 5b) martensite product. In bulk alloys the b-AuCuAl phase was reported to undergo order–disorder transformation from a disordered bcc (A2) at high temperatures to B2, L21 or DO3, depending on the thermal treatment history, and it has been suggested that L21 ordering was necessary for the appearance of martensitic transformation [8,11]. However, it was difficult to determine the type of ordering from the XRD patterns due to preferred orientation occurring in thin films and overlapping with neighbor diffraction peaks. To identify this phase regardless of the ordering state, the peak intensity profile for the (2 2 0) b-AuCuAl phase at 2h = 41.3° is selected since it is a diffraction peak that is common to them all and is the most prominent. This phase is predominant in the upper half of the phase-transforming composition range. This is because the compositions within this region have Af below or near room temperature (Fig. 3b). The martensite phase is indexed as a bct-AuCuAl phase, so the transformation from parent to martensite will split the (2 2 0) b-AuCuAl peak into (2 0 2) and (2 2 0) bctAuCuAl peaks. The (2 0 2) bct-AuCuAl peak was selected since it was most prominent (Fig. 5b), and it is predominant in the lower part of the phase-transforming composition range. This composition range shows the martensite phase since the Ms > room temperature (Fig. 3a). The composition range of the b/bct-AuCuAl phase found in this study did not coincide exactly with the phase diagram [13–15]. Possible reasons for this are as follows: (i) compositions measured by the EDX method have an error of ±1–2 at.%, causing a slight shift in the actual composition of the alloys; or (ii) the formation of phases by annealing a multilayer thin film precursor undergoes a different phase formation mechanism and the resulting phases do not necessarily correspond to the same phases formed by bulk metallurgical methods. Within the phase-transforming composition range, the b/bct-AuCuAl phase coexists to a large extent with secondary phases. For example, in Fig. 2, the XRD pattern shows the coexistence of several phases for the compositions Au57Cu32Al11 and Au49Cu30Al21. Outside the phasetransforming composition boundary, non-transforming phases are predominant, such as the intermetallic compounds of the Au–Al and Cu–Al systems, and Au–Cu solid solution. These were identified using an XRD database [41]. In the Au–Al part of the system, the intermetallic phases Au2Al (Fig. 5b, e and g), AuAl (Fig. 5a and c), AuAl2
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(Fig. 5d), Au4Al (Fig. 5e) and Au8Al3 (Fig. 5f and g) were confirmed. These phases formed on the upper and right side of the phase-transforming composition boundary. In the Cu–Al part of the system, the intermetallic phase Cu4Al (Fig. 5i) was confirmed, and this formed on the left side of the phase-transforming composition boundary. In the Au–Cu part of the system, face-centered cubic (fcc)-AuCu solid solution phases (Fig. 5a, b, e, f, g, h and i) were confirmed. They form on the lower side of the phasetransforming composition boundary. The fcc-AuCu phase is represented with different peak intensity plots due to a significant peak shift. The large difference in the atomic radii of Au and Cu causes a significant peak shift due to change in lattice constant with composition [42]. For example, the (1 1 1) fcc peak varies from about 39° at the Au-rich side to about 42° at the Cu-rich side. Au–Cu compositions at Au = 70, 50 and 33 at.% also undergo ordering, but in this study the ordering of the binary phases was not investigated further. In conjunction with R(T) screening and the XRD mapping results, it is concluded that the phase-transforming
composition range is limited by the formation of a large amount of non-transforming binary intermetallics and solid solutions. 3.3. Alloying behavior of b-AuCuAl The alloying behavior of b-AuCuAl is investigated in more detail by determining the influence of composition on the lattice constants. The XRD patterns collected at room temperature show either the b-AuCuAl phase or the bct-AuCuAl phase, depending on the Ms temperature. The phase-transforming compositions were selected after dicing the materials library. XRD patterns were recorded at 250 °C, at which temperature all the samples are in the b-AuCuAl phase (>Af). A preferred orientation was observed for all samples, therefore the lattice constant of the b-AuCuAl phase was estimated using only a single peak. The preferred orientation of most samples was (2 2 0), so the lattice constant was evaluated using this peak. However, on two samples the (2 2 0) peak did not appear due to a very strong (2 2 2) peak, so this peak was used
Fig. 5. Phase constitution map of phases confirmed in the Au–Cu–Al materials library. Peak intensities at characteristic 2h positions were plotted with respect to composition. The peak intensity color codes are as follows: red = high, green = medium, blue = low. The color-coded scheme identifies the most prominent phase with respect to composition. The black dashed line encircles the phase transforming region as identified with R(T). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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instead. Fig. 6 shows the plot of b-AuCuAl lattice constants (ab) measured at 250 °C with respect to composition. The lattice constants at 250 °C were all mostly higher than the reported values of 0.616–0.619 nm, which were measured at lower temperatures and for different compositions [8]. The increased values are due to the thermal expansion of the unit cell and a higher Au content. The plotted results show the compositional dependence of ab only, since the temperature was fixed at 250 °C. As a function of Au content (Fig. 6a), ab follows a decreasing trend for Au <50 at.% and plateaus for Au >50 at.%. As a function of Cu content (Fig. 6b), ab follows a decreasing trend for Cu >30 at.% while it plateaus for Cu <30 at.%. However, there are no clear trends confirmed with respect to Al content (Fig. 6c). The inflection point (Fig. 6a and b), where the trend changes from plateau to decreasing when Au <50 at.% and Cu >30 at.%, indicates a nominal composition of Au50Cu30Al20 for the b-AuCuAl phase. To observe a more meaningful relationship, ab was plotted as a function of the Au/Cu ratio (Fig. 6d). A clear trend is confirmed with respect to the Au/Cu ratio: ab decreases for Au/Cu <2 and it saturates to about 0.625 nm for Au/Cu >2. In a solid solution phase, like b-AuCuAl, changes in lattice constants are strongly dependent on the solubility of constituent elements. Thus the observation that the lattice constants reach a plateau for >50 at.% Au could be interpreted as the maximum amount of Au soluble in the b-AuCuAl phase. Moreover, the decreasing trend of ab for <50 at.% Au means that the Au site in the unit cell is substituted with a smaller atom. Comparing the metallic radius [43] of the constituent elements, Au (144 pm) and
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Al (143 pm) are similar, but significantly larger than Cu (128 pm). It is therefore reasonable that ab decreases with increasing substitution of Cu for Au. This is also consistent with the reported site occupancy in the crystal structure of b-AuCuAl phase [8], where the Au-rich site (C site) has the following site occupancy: Au = 0.653, Cu = 0.328, Al = 0.019. This means that Cu substitution for Au at the Au-rich site is more likely than Al substitution for Au. Furthermore, the ratio of Au/Cu site occupancy is 1.99, which coincides perfectly with the inflection point in Fig. 6d. 3.4. Visual appearance of the Au–Cu–Al materials library The aesthetic value of Au-based alloys is characterized by their color, so a photograph of the Au–Cu–Al thin film materials library annealed at 500 °C after complete R(T) screening was taken (Fig. 7a). The materials library was placed on a white paper with markers. The photograph was taken under neutral color temperature in order to be able to do a qualitatively correct color interpretation. The measurement region where phase transformations were confirmed is within the dashed black boundary. A clear color variation is observable over the materials library indicating a composition gradient. This is a well-known phenomenon in Au-based jewelry and in this alloy system, where color changes as a result of changes in the composition and phase constitution [3]. The dark blue area in the Cu-rich region could be attributed to Cu-oxide that formed on the surface as a result of repeated thermal cycling (up to 250 °C) during R(T) screening. The yellow-gold color is dominant for regions close to the Au-rich side. The color
Fig. 6. Lattice constants (ab) of b-AuCuAl measured at 250 °C plotted as a function of (a) Au content, (b) Cu content, (c) Al content and (d) Au/Cu ratio.
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Fig. 7. (a) Photograph of the Au–Cu–Al materials library after the R(T) screening step. Crosses mark the corners of the measurement regions (samples). Within the dashed boundary is the region where phase transformations were confirmed. (b) The corresponding plot of Ms temperature over the physical coordinates of the materials library.
becomes less yellow and more white when going towards the Al-rich side. In this direction a region of reddish yellow color is apparent, which falls partly within the phasetransforming composition range. When compared with the Ms temperature (Fig. 7b), the regions with high transformation temperatures exhibit a yellow-gold color, which is due to the higher Au–Cu content. On the other hand, the regions with lower transformation temperatures have a wider variety of colors, such as yellow-gold, light yellow-gold, whitish gold and reddish yellow-gold. These alloys could be used for various ornamental and jewelry applications. Functions derived from reversible martensitic transformation, such as shape memory effect, superelasticity and elasto-caloric effects, could be incorporated to provide new applications, such as smart ornaments and jewelry. 3.5. Thermal cyclic stability One of the most important considerations for SMA application is the cyclic stability of its functional property. During the R(T) screening process the whole materials library experiences about 52 thermal cycles. For screening purposes, the in-house R(T) test stand is equipped with a probe head having five sets of four-point electrical resistance probes, where each set can measure areas with 4.5 mm separation. It can thus measure five compositions simultaneously. The materials library is fixed by vacuum on the
hot/cold plate, so all compositions are simultaneously subjected to thermal cycling. This is one limitation of the R(T) screening test stand. To investigate further the stability of the observed reversible phase transformations, the materials library was diced to separate each measurement region corresponding to one composition (4.5 mm 4.5 mm size). Selected compositions were individually subjected to further thermal cycles using a probe head with a single set of fourpoint electrical resistance probes. Fig. 8 summarizes the R(T) thermal cycle test made after the R(T) screening. Raw data are given in Supp. Fig. 1. The selected compositions are Au55Cu34Al11, which has a high transformation temperature, and Au44Cu43Al13, which has a low transformation temperature. The difference between the first and fortieth cycles for Au55Cu34Al11 is minimal, but the hysteresis has narrowed slightly. Compared to the R(T) curve measured during R(T) screening, the difference in the transformation curve is likewise minimal. For Au44Cu43Al13, the first to third cycles exhibited a noisy signal (Supp. Fig. 1) around the transition point, so the fourth cycle was used for comparison instead. The change in the R(T) curves was minimal for the fourth and fortieth cycles, and likewise when compared to the R(T) curve measured during R(T) screening. The transformation temperatures with respect to thermal cycles for both compositions are compared in Fig. 8c. It can be clearly seen that for both compositions the transformation temperatures did not significantly change with thermal
Fig. 8. Thermal cycling test results of the selected compositions (a) Au55Cu34Al11 and (b) Au44Cu43Al13. The R(T) curve measured during the R(T) screening test is also shown as a reference. (c) The transformation temperatures (Ms and Af) are plotted with respect to thermal cycles. The measured values obtained during the R(T) screening are plotted as the zeroth cycle.
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cycles. Furthermore, it should be noted that during the R(T) screening all samples experienced several thermal cycles. For comparison, the transformation temperatures measured during R(T) screening are plotted as the zeroth cycle. The very small difference (Ms decreased by 4.8 °C for Au55Cu34Al11 and 4.1 °C for Au44Cu43Al13) between the zeroth cycle and the succeeding 40 cycles suggests that the phase transformation behavior of these alloys shows excellent thermal stability. The origin of this excellent thermal stability could not be explained within the scope of this study, but it will be a topic for future investigations on this new SMA system.
4. Conclusions The Au–Cu–Al system annealed at 500 °C for 1 h was systematically investigated by the thin film combinatorial materials science approach. As a result, very promising new SMAs with tunable functional properties and very stable functional properties were identified and characterized. This includes SMAs that are functional at room temperature and in a hot environment (>100 °C). The following conclusions are drawn from this study: 1. Reversible phase transformation over a wide range of compositions (35–70 at.% Au, 9–23 at.% Al and Cu balance) were identified by R(T) measurements. The phase transformation temperatures and hysteresis values span the following ranges: Ms = 15 to 149 °C, Af = 5 to 183 °C and DT = 8 to 66 K. The variation of these properties is strongly composition dependent. This means that it is possible to tune the desired property for a wide variety of potential applications. 2. The phase constitution across the materials library was investigated by high-throughput XRD analysis. The results revealed the correlation of phase constitution to composition and functional properties. Reversible phase transformation is attributed to the presence of b-AuCuAl and its martensite product. The formation of a large amount of non-transforming phases – Au– Al, Cu–Al and Au–Cu binary phases – was found to limit the range of compositions that show reversible phase transformations. 3. The lattice constants (ab) of the b-AuCuAl phase at 250 °C were found to vary with Au and Cu content. For compositions with >50 at.% Au and <30 at.% Cu, the lattice constant does not significantly change. However, for compositions with <50 at.% Au and >30 at.% Cu, the lattice constant decreases. This behavior is due to preferential substitution of Cu for Au on the Au-rich site of the b-AuCuAl unit cell with composition change. 4. The visual appearance of thin films within the transforming region varies with composition and shows different functional properties. The compositions that show high transformation temperatures are yellow-gold in color due to high Au and low Al contents. The compositions that show a low transformation temperature have various colors, such as light yellow-gold, reddish yellow-gold and whitish gold. 5. The thermal cyclic stability of two selected compositions with high (Au55Cu34Al11) and low (Au44Cu43Al13) transformation temperatures were investigated by R(T) thermal cycling test up to 40 thermal cycles. Both compositions show minimal change in their transformation
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temperatures during thermal cycles, suggesting that they have excellent thermal cyclic stability. Acknowledgements The authors would like to express their gratitude for support from the German Research Foundation (DFG) within the funding for the research unit FOR 1766 (High temperature shape memory alloys).
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