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Thin Solid Films 580 (2015) 36–44

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Aluminium–copper–nickel thin film compositional spread: Nickel influence on fundamental alloy properties and chemical stability of copper Martina Hafner, Wolfgang Burgstaller, Andrei Ionut Mardare ⁎, Achim Walter Hassel Christian Doppler Laboratory for Combinatorial Oxide Chemistry, Institute for Chemical Technology of Inorganic Materials, Johannes Kepler University Linz, Altenberger Str. 69, 4040 Linz, Austria

a r t i c l e

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Article history: Received 13 May 2014 Received in revised form 5 March 2015 Accepted 5 March 2015 Available online 12 March 2015 Keywords: Aluminium–copper alloys Thin film combinatorial libraries Co-evaporation Scanning droplet cell microscopy Scanning Kelvin probe Cyclic voltammetry Downstream analytics Corrosion

a b s t r a c t An Al–Cu–Ni thin film compositional spread was deposited by thermal evaporation and investigated in order to study the Ni influence on the overall properties. The chemical composition was detected by energy dispersive X-ray spectroscopy and showed a compositional spread of approximately 20 at.% Ni. Decreasing the Ni content in the Al–Cu–Ni thin films resulted in an increased grain size and characteristic surface microstructure evolution. Scanning Kelvin probe measurements were performed to investigate the surface potential variation along the compositional gradient, and a distinct surface potential drop was observed between Al–Cu–7 at.% Ni and Al–Cu–13 at.% Ni. The results of the X-ray photoelectron spectroscopy surface analysis and Auger electron spectroscopy as well as the electrochemical investigations by cyclic voltammetry evidenced mainly the presence of Al2O3 but also CuO and Cu2O together with metallic Cu were clearly identified along the compositional gradient. Chemical dissolution experiments have shown that Ni is enhancing the chemical stability of Cu, excepting inside the compositional region between 7 and 13 at.% Ni. © 2015 Elsevier B.V. All rights reserved.

1. Introduction Cu and Cu alloys are presenting a widely varied field of application. Besides the predominant use of pure Cu for electrical conductors a number of alloys possess high industrial importance. Due to their high electrical conductivity and attractive physical properties, Al–Cu based thin film alloys are greatly relevant for electronic applications being used as interconnects with good resistance toward electromigration [1]. Cu based alloys such as brass or Isotan® are used for a large number of applications in different industries due to their high mechanical strength, good thermal conductivity and excellent corrosion resistance. The combination of these elements makes the Al–Cu–Ni alloys (in either bulk or thin film form) to become very interesting materials, particularly for applications requiring high electrical stability and good corrosion resistance. The corrosion resistance of bulk Al–Cu–Ni is generally attributed to the self-passivation of the Al which is forming a protective Al2O3 layer on the surface. This is improving the alloy corrosion behaviour in neutral solutions. In the particular case of chloride containing solutions, where Al is prone to pitting corrosion, the presence of Ni improves the corrosion resistance of Al–Cu–Ni as compared with Al–Cu alloys. In alkaline solutions the protection against corrosion is attributed to the Cu which forms a duplex layer of CuO and Cu2O mixed with Al2O3 [2]. Additionally, the incorporation of Ni in Cu2O decreases the number of ⁎ Corresponding author.

http://dx.doi.org/10.1016/j.tsf.2015.03.018 0040-6090/© 2015 Elsevier B.V. All rights reserved.

cation vacancies which leads to an inhibition of the corrosion process [3]. These two processes are influencing the electrochemical behaviour of the alloys surface resulting in a significant improvement of the corrosion resistance in different solutions. However, stabilization of the Al–Cu–Ni alloys in sulphate–chloride solutions is limited for Ni contents up to 30 wt.%, while for higher Ni concentrations an increase of the corrosion rate is observed. In solutions containing chloride (e.g. seawater) the stability is increasing with the Ni content [4,5]. Badawi et al. showed that addition of Ni, Zn and/or Al to Cu leads to synergistic effects triggering specific changes in the final alloys properties. Therefore, ternary Cu alloys such as Cu–Ni–Zn, Al–Cu–Zn or Al–Cu–Ni can be used in chloride containing solutions due to their ability to form a protective passive layer [6]. However, in all bulk investigations previously mentioned, only discrete alloys are analysed. A combinatorial approach resulting in obtaining thin film alloys with a graded composition simultaneously deposited on the same substrate represents a more suitable route for basic properties mapping as a function of composition. Vapour phase deposition by thermal evaporation is a suitable method for thin film deposition which allows formation of material gradients by using two or more evaporation sources simultaneously. Due to the fact that the evaporation rates can be independently controlled via the applied electrical power on each evaporation source, the method allows the formation of precisely tuned thin film compositional spreads [7]. In this work an Al–Cu–Ni thin film compositional spread was produced by thermal

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co-evaporation using a two-evaporation-source geometry. Investigations were done in order to characterize the thin film compositional spread as a function of the Ni content. Apart from microstructural and crystallographic characterization, the influence of the naturally formed oxide on various Al–Cu–Ni alloys was electrochemically investigated and the chemical dissolution behaviour of Al–Cu–Ni alloys was assessed. 2. Experimental details Al–Cu–Ni thin films were obtained from vapour phase deposition using a two-source evaporation geometry in order to create a thin film compositional spread. A state of the art self-developed evaporation system was used for this purpose. The system includes a steel chamber with a base pressure of 10−5 Pa and a pneumatically removable base plate containing pure Cu electrical feedtroughs supporting up to three individual evaporation sources, while typically only one source is employed. Each evaporation source is electrically heated by independent DC power sources (N2 kW each). For heating of the materials to be deposited, both W-boats and W-baskets holding BN crucibles (5 ml) can be used. The evaporation process is monitored by high resolution crystal quartz micro-balances (QCM) (with a noise level below 3 pm s−1) which are placed directly above the evaporation sources. A specially designed aluminium shield ensured that each QCM will detect only the material coming from its corresponding evaporation source. The power sources are controlled by LabView software with integrated proportional-integral-differential controller using the feedback input from the crystal quartz. The raw materials used for the co-deposition of Al–Cu–Ni thin films were high purity Al and Cu pellets (used for formation of Al–Cu bulk) and Ni powder (N99.95%). Self-cast bulk Al–Cu alloy (Al–72.4 at.% Cu) was used as one evaporation source from a W-boat, 250 W being sufficient for reaching the required evaporation temperature (N820 °C). The other deposition source used in the co-evaporation process was a W-basket holding a BN crucible containing the Ni powder (Alfa Aesar). The Ni source needed 600 W for reaching the evaporation temperature of 1072 °C. The bulk Al–Cu was fabricated in a pivotable induction furnace (Linn—High Therm) with a maximum power of 10 kW. The ratio of Al–Cu was chosen so that the final alloy composition is Al–72.4 at.% Cu. The furnace chamber was evacuated down to 1 Pa and then flushed with Varigon® H6 gas (Linde) to remove the oxygen. This procedure was repeated several times to avoid oxide formation during the melting process. Once the materials were completely molten, the furnace chamber was turned into a vertical position by its pivotal function in order to cast the material into a stainless steel mould (20 × 20 × 200 mm3). The produced ingots were machined into small pellets (approximately 10 mm3) to serve as evaporation source for the thin film deposition. The distance between the centres of the evaporation sources was approximately 86 mm, which means that a compositional spread will be formed along a substrate due to the cosine law governing the thermal evaporation processes. The thin film deposition was performed simultaneously on two borosilicate glass substrates (microscope slides, 26 × 76 mm2, VWR International GmbH) in order to provide a sufficiently large area for further investigations. This approach presents the advantage that two identical samples with the same compositional gradient and same history are produced during the same deposition. The evaporation distance was kept constant for both sources at 120 mm. Before deposition the substrates were cleaned using an ultrasonic bath containing sequentially acetone, ethanol and deionized water followed by N2-drying. A length scale was defined for each sample along its 76 mm edge for an easier properties mapping. At a position of X = 0 mm the Ni source was located while the Al–Cu source was located at X = 76 mm. The evaporation rate corresponding to each source was in situ adjusted in order to obtain a compositional spread centred on Al–Cu– 10 at.% Ni. The deposition was stopped when a thickness of 450 nm

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was indicated by both QCMs in an additive manner. The Al–Cu–Ni thin film deposition was carried out at room temperature in a pressure of 5 × 10−4 Pa. Due to the radiation during the evaporation process the substrates have reached a temperature as high as 200 °C. The distance between each source and substrate was kept constant at 120 mm for the thin film formation. The evaporation rates used for obtaining the desired compositional spread were 0.6 nm s−1 for Al–Cu and 0.055 nm s−1 for Ni. Additionally, Al–Cu thin films evaporated from a single source containing the bulk Al–72.4 at.% Cu were separately deposited as reference samples in order to comparatively study the effect of Ni during the Al–Cu–Ni thin film formation. Identical to the Al–Cu–Ni thin films, the location of the Al–Cu source was also X = 76 mm in order to reproduce the same geometrical conditions. For precisely determining the film thickness distribution, atomic force microscope (AFM) measurements (Nanosurf easy scan 2) were performed at various locations along the compositional gradient of the Al–Cu–Ni thin film compositional spread. The measurements were performed in contact mode and indicated an overall thickness distribution between 320 and 630 nm along the entire 76 mm length of the substrates. Microstructural investigations of the Al–Cu–Ni thin film were done using a field emission Zeiss Gemini 1540 XB scanning electron microscope (SEM) with 5 kV acceleration voltage and in-lens detector. The chemical composition along the Al–Cu–Ni thin film compositional spread was determined using an energy dispersive X-ray spectroscopy (EDX) built in the SEM system (Oxford INCA) with incident energy of the electron beam of 20 kV. Crystallographic properties of the deposited thin films were studied using an X-ray diffraction (XRD) system (PANalytical X'pert Pro) with a Cu-Kα source without monochromator. The XRD measurements were done in Bragg–Brentano geometry. All diffractograms were recorded in a 2θ ranging between 4 and 140°. The XpertHigh software was used for data analysis and peak identification. Scanning Kelvin probe (SKP) measurements were performed to investigate the surface potential of the thin film. The SKP was operating with a CrNi alloy tip (tip diameter of 300 μm) in automatically controlled constant height mode with a tip–sample distance of 110 μm [8]. The investigations were done by measuring a centred 10 mm wide area along the entire compositional spread. To minimize environmental influences, all measurements were performed under controlled humidity and temperature conditions. A qualitative analysis of the surface oxides present along the Al–Cu–Ni compositional spread was done using XPS and Auger spectroscopy (ThermoFisher—Theta Probe). In the XPS measurements Al-Kα radiation (1486.6 eV) was used with an acceleration voltage of 15 kV and the emission current was 6.7 mA. The spot size was 400 μm in diameter. The electrochemical behaviour of the Al–Cu–Ni thin film was studied by cyclic voltammetry using a scanning droplet cell microscope (SDCM). Due to its capability of performing localized investigations, together with its built in surface scanning function, the SDCM is an electrochemical tool ideal for thin film libraries characterization and surface properties mapping [9]. A schematic drawing of the SDCM probe is presented in part (a) of Fig. 1. The main body of the cell is made from a 2.5 mm in diameter boron-silicate glass capillary. A tip with a diameter of 420 μm was obtained by thermally pulling the capillary with the help of a PC-10 puller (Narishige, Tokyo, Japan) and subsequent grinding. A glass capillary based Ag/AgCl micro-reference electrode was inserted in the main body of the cell in close proximity of the cell tip [10]. The potential of the reference electrode, as measured for calibration against a commercial Ag/AgCl electrode, was determined to be 0.208 V vs. the standard hydrogen electrode (SHE). A 100 μm Au-wire used as counter electrode was wrapped around the reference electrode. The electrolyte inlet was attached to the upper part of the SDCM body using two component epoxy glue. In order to precisely define the wetted area on the surface of the Al–Cu–Ni thin films (working electrode, WE) a soft silicone sealing was formed at the SDCM tip by dipping it in liquid silicone followed by N2-drying. This allowed operation of SDCM in contact mode [11]. All

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3. Results and discussion

Fig. 1. (a) Schematic drawing of the SDCM. (b) Schematic drawing of the V-type SDCM.

cyclic voltammograms (CVs) were recorded at room temperature using a potentiostat (EG&G Instruments, Model 283) at a scan rate of 15 mV s−1. Starting at a potential of 0 V vs. SHE, the potential was swept up to 0.7 V vs. SHE followed by a down sweep to −0.4 V vs. SHE before reaching the initial potential of 0 V vs. SHE. Up to 5 consecutive cycles were recorded for each investigated spot. The electrolyte used for the electrochemical investigations was 0.1 M acetate buffer prepared from p.a. grade acetic acid, p.a. grade sodium acetate and deionized water (pH = 6.0). The chemical stability of the Al–Cu–Ni thin film compositional spread was studied by investigating the chemical dissolution of Cu. The dissolution behaviour of the Al–Cu–Ni thin film was tested using a flow type scanning droplet cell microscope (FT-SDCM) channelling 0.1 M HNO3 toward a small addressed area on the investigated surface. A schematic representation of the FT-SDCM is given in part (b) of Fig. 1. Using a 3D printing technique, a V shape flow profile was obtained in a polymer block allowing the acid to contact the Al–Cu–Ni surface in a spot with a 2 mm diameter. A small Viton O-ring glued at the tip of the plastic block allowed an effective sealing of the acid within the addressed volume permitting the dissolved species to be transported away through the outlet channel. A peristaltic pump was used to flow the acid at a constant flow rate of 0.75 ml min−1 into the cell and further into an atomic absorption spectrometer (AAS) (HITACHI Z-8230 polarized Zeeman atomic spectrophotometer with background correction). In order to analyse the amount of dissolved Cu, a copper hollow cathode lamp (λ = 324.8 nm) was used as radiation source. The atomization unit was an acetylene-air flame (2.1 L min−1 acetylene, 15.0 L min−1 air). The detected signal was digitalized by an 18-bit analogue to digital interface (Labjack U6 Pro) with a resolution of 0.15 mV. Data evaluation of the converted signal was performed by a self-developed LabView data acquisition software including graphic data processing. Calibration of the AAS was done by measuring blank and standard solutions which are based on a stock solution of copper (1000 mg l−1).

In order to characterize the alloys stoichiometry along each substrate obtained during the co-evaporation of Al–Cu and Ni, EDX investigations were done for mapping the compositional gradient. In Fig. 2 these results are summarized. Each elemental composition is plotted as a function of the position at which the EDX measurement was done along the substrate length (76 mm). Additionally, an EDX analysis of Al–Cu reference films deposited from the single alloy source under identical conditions as the Al–Cu–Ni compositional spread was used for investigating the effect of the Ni on the Al/Cu ratio in the thin films. In part (a) of Fig. 2 such EDX mapping is presented as measured along the entire X axis. Both Al and Cu showed a compositional gradient of approximately 2 at.% with a Cu enrichment directly above the evaporation source (X = 76 mm). The composition of the Al–Cu thin film varied from Al–59.1 at.% Cu to Al–61.3 at.% Cu. This corresponds to an Al/Cu ratio of 0.63 at the Cu rich side (X = 76 mm) increasing to 0.69 at the Al rich side (X = 0 mm). Along the entire Al–Cu film only a 3% variation of this elemental ratio was therefore measured. Since the deposition distance (120 mm) was identical with the one used for the Al–Cu–Ni thin film compositional spread, this Al/Cu ratio variation can be considered as a reference for further discussion. Even though both Al and Cu were evaporated from the same source (Al–72.4 at.% Cu alloy) their compositional gradients vary. This can be attributed to the different atomic masses of Al and Cu. Since both species will receive the same energy as soon as the evaporation temperature is reached, they will have different momentums. As a result of gas phase collisions between Al and Cu atoms, the lighter Al will be scattered leading to a Cu enrichment at the sample edge positioned directly above the Al–Cu source. The Al–Cu thin film composition is therefore strongly affected by the deposition distance, as previously observed

Fig. 2. EDX results of (a) Al–Cu thin film evaporated from a single source and (b) Al–Cu–Ni thin film deposited from two sources, while d represents the film thickness.

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when studying Al–Cu thin films. Additionally, a compositional deviation of approximately 10 at.% Cu is observed in the Al–Cu thin film as compared to the bulk evaporation source. This may be due to the presence of both Al2Cu and pure Cu in the source having different evaporation temperatures [7]. The EDX of the Al–Cu–Ni compositional spread is presented in part (b) of Fig. 2. Both Cu and Ni showed compositional gradients of approximately 20 at.% while the Al gradient was much weaker. The effect of the cosine law governing the compositional distribution during the codeposition process can be directly observed in the shape of the curves presented in Fig. 2(b). In the present case no linear compositional gradient was observed due to the special evaporation geometry used and to the rather high exponents of the cosine law (N 4) previously measured for both Al–Cu and Ni sources. Along the Al–Cu–Ni thin film compositional spread the Ni content varied between 24.8 and 5.7 at.%. The presence of Ni atoms in the gas phase during the Al–Cu–Ni thin film formation is expected to change the energetic balance and the atomic collision process. This is comparatively evidenced by the different Al/Cu ratio, which in the case of Al–Cu–Ni varies from 0.46 at the Cu rich side and 0.56 at the Ni rich side as observable from Fig. 2(b). As compared to the composition of the films deposited from the Al–Cu source alone, the Al–Cu–Ni films show an additional 10% decrease of the individual Al and Cu compositions along the entire spread. The current 10% variation is therefore attributed to the presence of the Ni

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atoms. This means that the Ni gradient within the Al–Cu–Ni compositional spread is the main influencing factor upon the final alloys composition. In order to investigate the possible film modifications triggered by the various thicknesses, in part (b) of Fig. 2 the overall thickness profile of the Al–Cu–Ni compositional spread as measured by AFM is also presented. The film thickness of 630 nm obtained directly above the Al–Cu source at X = 76 mm gradually decreases to 320 nm directly above the Ni source at X = 0. This strong thickness variation is unavoidable due to the desired low amounts of Ni in the compositional spread. The very different evaporation rates used for both Al–Cu and Ni sources are playing a role in this thickness evolution, together with the individual cosine laws governing the thickness distributions obtained from the individual sources. In most cases, such situation may be avoided by splitting such compositional spread across several different samples. The drawback of this approach is the necessity of analysing several different samples possibly affecting the direct comparison of results coming from various alloys due to possible different thin film formation conditions and history. The microstructure of the Al–Cu–Ni compositional spread was investigated by SEM at various locations along the samples. In Fig. 3 a collection of surface images corresponding to various Ni contents is presented. At low Ni concentration (5.7 at.%) a compact surface structure can be identified with grain sizes with equivalent diameter ranging

Fig. 3. SEM images of selected compositions on the Al–Cu–Ni thin film compositional spread.

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from 30 to 100 nm. This microstructure is similar with the one observed on the surface of the reference Al–Cu thin films (not shown here), the small amount of Ni not having a visible influence on the grain shapes or sizes. Additionally, no visible change was observed on the Al–Cu reference thin film surface along the elemental Al and/or Cu gradient. This indicates that the Al–Cu microstructure is invariable for the given thin film formation conditions in the absence of Ni. Gradually increasing the amount of Ni in the Al–Cu–Ni thin film compositional spread resulted in a smooth decrease of the grain sizes. In the same time the equivalent diameter of the grains starts to be more uniform, approximately 13 at.% Ni resulting in a smooth surface with uniform and slightly elongated 80 nm grains. Higher amounts of Ni identified in the Al–Cu–Ni compositional spread resulted in the formation of almost spherical grains, their average diameters decreasing as low as 50 nm. However, the overall Al–Cu–Ni film thickness gradient (presented in part (b) of Fig. 2) may also play a role in the surface grain evolution. The lower Ni content surface (showing the largest grains) corresponds to the thicker film (N600 nm) while the smaller grain sizes observed for the high Ni amounts correspond to the thinner Al–Cu–Ni films (320 nm). The fact that the grain size decreases approximately four times with increasing the Ni amount from 5 to 25 at.% may not be triggered only by the thickness which decreases by a factor of two. In this idea, the presence of Ni could influence the surface microstructure evolution. Otherwise, a preferential in-plane grain growth has to be concluded which is enhanced at higher Al–Cu–Ni film thicknesses solely due to the particularities of the film growth. The crystallographic structure for the Al–Cu–Ni compositional spread was mapped along the compositional gradient using XRD. Various locations on the compositional spread surface were sequentially irradiated by the X-ray beam and the diffraction was analysed. In Fig. 4 the X-ray diffractograms measured for a few selected Al–Cu–Ni compositions are presented. In all cases the XRD patterns were measured in a 2θ range between 4 and 140°. However, in Fig. 4 only the range of 30–110° is presented due to the lack of peaks outside this range. As in previous discussions, the Ni amount is followed in the XRD mapping of the compositional spread. A clear main XRD peak can be observed in all diffractograms together with a smaller peak observable at higher angles. At low Ni compositions (5.7 at.%) the main XRD peak is centred at a 2θ angle of 42.8° and a small shoulder can be observed at 43.8° as a secondary peak. A similar trend was observed when investigating Al–Cu compositional spreads, when the shoulder was due to the presence of pure Cu [7]. Increasing the amount of Ni in the Al–Cu–Ni thin film alloys, the main peak shifts toward higher diffraction angles while the shoulder disappears above approximately 10 at.% Ni. This trend is commonly encountered in thin film compositional spreads when the alloys have all the same crystallographic symmetry or when a change between

Fig. 4. XRD results of selected compositions on the Al–Cu–Ni thin film compositional spread.

two similar symmetries occurs [12,13]. However, there is no clear rule defining this behaviour. Such smooth diffraction peak shift could also be related to a crystallographic change between similar symmetries. The best XRD database match was identified in Al7Cu23Ni where the main peak characterizes a (0 0 18) orientation. Similar results were previously found during the investigation of Al–Cu–Ni bulk alloys when more peaks attributed to Al7Cu23Ni were identified [14]. The fact that in the present study only one major peak was evidenced is often encountered when comparing crystallographic properties of bulk and thin film alloys. This is generally due to higher texturing obtained in thin films triggered by the film growth particularities, e.g. deposition temperature, atom energy, substrate surface, sticking coefficients, etc. The smaller XRD peak observable at low Ni contents around 94° has also a composition dependent evolution. At low Ni amounts, its intensity is maximized while with the increase of Ni content, its intensity decreases and finally no peak can be identified above 10 at.% Ni. This peak may be identified as belonging to cubic AlCu4 (which is approximately the composition of the bulk source used for the film formation) or Al4Cu9. This is also related to its disappearance at high Ni contents where only the main XRD peak remains. In conclusion, the most likely symmetry of the Al–Cu–Ni thin film compositional spread is based on a combination of monoclinic and cubic structures (as suggested by the Al7Cu23Ni and AlCu4 or Al4Cu9, respectively). The secondary peak described as a shoulder of the main XRD peak can be attributed either to the cubic AlCu4 or to the γ1 phase of the bulk Al–Cu–Ni alloy [15]. This conclusion is based on previous bulk studies of the ternary Al–Cu–Ni system. The high temperature bulk Al–Cu–Ni phase diagram shows a sequence of five phases as a function of the Ni content (between 5 and 25 at.%). The first transition occurs at 6.5 at.% Ni when a mixed γ1 + bcc phase changes to bcc + B2. With the increase of the Ni amount, pure B2 phase is observed above 11 at.% Ni and finally above 14 at.% Ni a mixture of B2 and fcc was revealed. Moreover, quenching resulted in a clear γ1 phase at low Ni contents [16]. In the present study the same trend is observed, the γ1 phase is identified as Al4Cu9 and occurs only at low Ni concentrations. Amounts of Ni above approximately 10 at.% resulted in a complete suppression of both the secondary peak/shoulder and the high angle small peak concomitant with the shift of the main XRD peak towards higher angles. To characterize the electrical surface properties of the Al–Cu–Ni thin film compositional spread, the contact potential difference (CPD) was mapped along the compositional gradient using a SKP. The CPD depends on various surface parameters such as adsorption layers, oxide formation, structural changes, surface defects, etc. SKP mapping is a sensitive, non-destructive and non-contact method for investigating conducting and semi-conducting surfaces in an attempt to identify relations between the sample surface and the electrochemical (Volta) potential [17]. The CPD probed by the SKP tip can be used for describing the analysed materials work function (Φm) if the tip work function (ΦSKP) is known. The proportionality relationship between the work function difference and the CDP is described as CPD = (Φm − ΦSKP) / e, where e is the elementary charge. Any change in CPD measured along the sample surface directly reflects the surface work function variation [18]. In Fig. 5 the results of the CPD mapping are shown in a cross sectional view as well as in a 3D plot. In the Al–Cu–Ni thin films, a strong potential drop can be observed along the compositional gradient. The region corresponding to this drop is located between 7 and 13 at.% Ni and a decrease of the CPD by approximately 200 mV was measured. This tremendous decrease could be related to the formation of Cu2O and CuO phases on the sample surfaces. By keeping constant the environmental conditions, any local change of the Al–Cu–Ni thin films and its natural oxides can be directly indicated by a change in the work function and hence in the CPD. The dramatic CPD change along the compositional spread can be related to the semiconducting properties of Cu2O and has been described in previous studies [19,20]. In addition, the presence of Ni influences the carrier generation and transport in thin films [21]. The fact that the minimum CPD was found at a composition matching

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Fig. 5. SKP scan along the Al–Cu–Ni thin film compositional spread.

the stabilization of the monoclinic phase above 10 at.% Ni may not be a coincidence. However, since the CPD is a surface property, the connection to the crystallographic dynamics along the Al–Cu–Ni compositional spread can be done only through the surface natural oxides. A qualitative analysis of the Al–Cu–Ni thin film compositional spread naturally oxidized surface was done by XPS and Auger spectroscopy. In both cases, the binding energy scale is presented as measured without the typical C1s calibration due to the small deviations obtained. Along the entire Al–Cu–Ni compositional spread, the naturally formed surface oxide mainly contained Al2O3 but also metallic and oxidized Cu together with trace amounts of NiO. This was evidenced by the binding energy corresponding to the XPS peaks. The integration of the XPS peaks allowed a quantification of the components of the surface oxide. As expected due to the valve metal behaviour of Al, the amount of Al2O3 on the surface was high, slightly oscillating around the value of 84 at.% along the entire compositional spread. The integration of the Cu2p peak revealed amounts averaging 10 at.% with a maximum of 15 at.% observed in the middle of the compositional zone corresponding to the maximum CDP drop previously measured by SKP. Only at the highest amounts of Ni in the compositional spread, the surface oxides contained up to 1 at.% of NiO while below 10 at.% Ni in the film no NiO could be detected on the surface anymore. Since the amount of Al2O3 was almost independent on the Ni content of the compositional spread, a careful study of Cu and Ni species could reveal more details about the chemical stability of the Al–Cu–Ni alloys. A few representative compositions were analysed and the results obtained are summarized in Fig. 6. In part (a), the CuLMM Auger scans are presented while in part (b) the Cu2p XPS results are shown. For low amounts of Ni (b 7 at.%) present in the compositional spread, the typical Cu2+ shake-up satellite is very weak, as can be observed in the XPS data. This indicates that very little amounts of CuO could be found on the thin film surface. The shakeup satellite becomes stronger with increasing the amount of Ni in the Al–Cu–Ni compositional spread, but completely disappears at the maximum amount of Ni (18.26 at.%) in the investigated alloys. In principle, this is to be expected since the Cu2O is less stable as compared to CuO. However, this behaviour also suggests that within the special compositional zone previously identified in the SKP investigations (Fig. 5), Ni may promote the surface formation of Cu2O instead of CuO in the metallic natural oxide. All Auger spectra show two maxima visible indicating that two different oxidation states are always present along the surface of the Al–Cu–Ni compositional spread. The use of the modified Auger parameter (MAP) allows the interpretation of the results independent of possible surface charging or X-ray energy [22,23]. The MAP analysis indicates the presence of metallic Cu on the surface, independently on the amount of Ni found in the investigated thin film alloys. Adding all this information it can be concluded that the surface of the Al–Cu–Ni

Fig. 6. (a) The LMM Auger spectra of Cu for selected compositions along the Al–Cu–Ni thin film. (b) XPS spectra of Cu for selected compositions along the Al–Cu–Ni thin film compositional spread.

thin film compositional spread is formed from a mixture of metallic Cu and Cu2O throughout the entire compositional spread, except for most of the special compositional zone (7–11 at.% Ni). Here, metallic Cu is found mainly in mixture with CuO but small amounts of Cu2O are still possible to coexist. Due to the fact that several species are simultaneously available on the Al–Cu–Ni compositional spread surface, an oxidation/reduction electrochemical study was performed for identifying the most active species. The electrochemical behaviour of Al–Cu–Ni thin film compositional spread was investigated using cyclic voltammetry with a neutral pH electrolyte (acetate buffer). In order to asses both oxidation and reduction of the metallic alloys, potentials ranging from −0.4 to 0.7 V vs. SHE were used with a potential scanning rate of 15 mV s−1. For this purpose a SDCM with a tip diameter of 400 μm was used in the contact mode (see Fig. 1(a)). The surface of the Al–Cu–Ni compositional spread was mapped using the automatic scanning mode of the SDCM [24] and the oxidation/reduction characteristics of various metallic alloys were investigated. A few selected CVs are plotted in Fig. 7 and the Al–Cu–Ni compositions are indicated for each curve. In part (a) of Fig. 7 only the first CV cycle is presented for various compositions. For all addressed alloys, typical anodic current density peaks describing the oxidation process are found while no reduction peak could be identified. This effect could partially be attributed to the valve metal behaviour of Al present in the alloys which will be anodized during the anodic sweep of the CV measurements. Since the high field model of oxide growth on Al associates a constant current density plateau with the oxide formation process during a potentiodynamic sweep [24], the oxidation of Al will contribute to an increase in the background current density level during the measurements on Al–Cu–Ni thin films. Therefore the oxidation/reduction peaks observed are describing the behaviour of

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direct access to Cu, otherwise covered by passivating Al, cannot be completely excluded. A complete CV containing all five cycles is presented in Fig. 8(a) for 18.3 at.% Ni. This particular composition is suitable for illustrating the oxidation/reduction dynamics previously discussed as a function of the CV cycling since the highest amount of Ni in the Al–Cu–Ni thin films was found to be responsible for a suppression of Cu2+ from the surface (see Fig. 6). Apart from its dependence on Ni content, the main oxidation peak (Cu+ to Cu2+) decreases in height with the CV cycle number. In the same time both, the Cu0 to Cu2+ oxidation and the reduction peaks are increasing their height with the CV cycle number. This confirms the idea of irreversibility of the Cu+ oxidation developed during the analysis of the first CV cycle throughout the Al–Cu–Ni compositional spread (Fig. 7(a)). The amount of Cu+ present on the surface is consumed by oxidation leading to lowering of the main oxidation peak with time. The concomitant increase of the secondary oxidation and reduction peaks with time can be explained by an increase of the available amount of Cu0 on the surface due to the reduction process. Following this idea, the reduction process responsible for the observed peak would be Cu2+ to Cu0. In order to quantify the oxidation and reduction processes, a coulometric analysis was done by integrating the current dependence on time during the CV measurements. Both oxidation peaks and the reduction peak were numerically integrated using a standard baseline algorithm. These results are presented in part (b) of Fig. 8. The irreversible character of Cu+ to Cu2+ oxidation is quantified here and an overall charge decrease of approximately 75% can be observed after the fifth cycle as reported to the first. The nonlinear appearance of the measured charge during the Cu2O transformation into CuO may suggest that small amounts of Cu+ may also result from the reduction of Cu2+. This hypothesis is supported by the amount of charge transformed during the reduction process, which is slightly Fig. 7. (a) First CV cycle; (b) fourth CV cycle for selected compositions of the Al–Cu–Ni thin film.

the Cu and/or Ni species. The oxidation peaks centred approximately at 0.12 V vs. SHE are increasing in height with the Ni content and two additional smaller oxidation peaks are visible for the alloys corresponding to the transition compositional zone previously identified in the SKP investigations (Fig. 5). The main oxidation peak can be attributed to the oxidation of Cu+ to Cu2+ while the smaller oxidation peak present only for the 9.7 and 11.6 at.% Ni is most likely responsible for the presence of a mixed Cu+/Cu2+ oxide. This conclusion is supported by the previously discussed XPS surface analysis (Fig. 6) which revealed a mixture of Cu+/Cu2+ present on the surface. The absence of reduction peaks can indicate that either the oxidation of Cu2O to CuO is not reversible during the first CV cycle or the peak is smeared out and hidden. In order to further investigate the surface oxidation dynamics, several consecutive CV cycles were recorded. In part (b) of Fig. 7 the fourth CV cycle is presented for the same Al–Cu–Ni compositions previously discussed in Fig. 7(a). Apart from the main Cu+ to Cu2+ oxidation peak, a clear secondary oxidation peak appears together with a reduction peak in the backward scan. The broad secondary oxidation peak centred approximately at 0.25 V vs. SHE is most likely responsible for the oxidation of Cu0 to Cu2+ while the reduction peak centred approximately at 0 V vs. SHE indicates the reduction of Cu2+ to Cu0. Similar to the first CV cycle, all oxidation and reduction peaks are increasing in height with the amount of Ni present in the Al–Cu–Ni thin film compositional spread. The conclusion that the presence of Ni is stabilizing the Cu2O on the surface of the Al–Cu–Ni compositional spread (previously derived from the XPS investigation in Fig. 6) is confirmed by the height dynamic of the main oxidation peak. This is in accordance with previous investigations on bulk Al–Cu–Ni alloys where a void filling mechanism in Cu2O by Ni atoms was described [5,6]. However, the possibility that the presence of Ni would modify the surface at the nanoscale allowing

Fig. 8. (a) CV curves for the Al–Cu–18.3 at.%Ni alloy (b) charge density for the Cu+ and Cu0 oxidation and the Cu2+ reduction of the Al–Cu–18.3 at.%Ni alloy.

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higher than the charge measured during the Cu0 to Cu2+ oxidation corresponding to each CV cycle. The reduction charge shows an almost linear behaviour with time while the Cu0 oxidation shows a nonlinear behaviour. This nonlinearity may be related with an increased integration error due to the presence of the strong Cu+ oxidation peak affecting the integration baseline. Overall, the charge difference between Cu0 oxidation and reduction for each CV cycle is in average about 25 μC cm−2. This difference could only partially be attributed to the formation of small amounts of Cu+ since the oxidation of Al cannot be neglected. Due to its valve metal behaviour, Al will form a passivating oxide which will not be reduced, therefore also contributing to the observed charge density difference. The anodization of Al is evidenced in part (a) of Fig. 8 by the presence of a current density plateau above 0.4 V vs. SHE, in agreement with high field model for oxide growth [25]. Since no sharp current increase is observable, the possibility of an electrode transfer reaction occurring in electrolyte by charge tunnelling through the natural oxide can be excluded. The anodization process during the potential sweep results in an increase of the natural surface oxide with more than 50% reducing the tunnelling probability [25]. In order to evaluate the chemical stability of Al–Cu–Ni thin film compositional spread in acidic environments, local Cu dissolutions were performed using the FT-SDCM. The entire compositional spread was scanned with a resolution of approximately 1 at.% and downstream analytics via AAS was used for quantifying the amount of dissolved Cu. In Fig. 9 the AAS results are presented as depending on the Ni content which is expected to chemically stabilize the Al–Cu–Ni alloys [26]. The integration of dissolution curves, typical for AAS measurements, led to measuring the amount of Cu locally dissolved along the thin film compositional spread and its thickness normalized values are plotted in black (left scale) in Fig. 9. Additionally, the mass of dissolved Cu was normalized to the Cu composition for studying the dissolution behaviour independent on the Cu compositional gradient. This curve is also presented in Fig. 9 (blue curve, square symbol, right scale). For both presented curves an overall decrease of the Cu mass dissolved by 0.1 M HNO3 is observed with increasing the Ni content along the Al–Cu–Ni thin films. Within the previously defined special compositional zone (between 7 and 14 at.% Ni) a peak is present for both curves. This indicates a higher Cu dissolution rate in this zone leading to a decreased chemical stability of the investigated alloys. This behaviour can be related to the surface composition previously analysed by XPS (Fig. 6). The presence of both metallic Cu and CuO on the surface most likely has an influence over the observed decreased chemical stability. This result is surprising since CuO is more stable than Cu2O identified outside the special zone. This suggests that in this special zone the metallic Cu is more readily accessible as compared to the remaining film. Overall, the presence of Ni is decreasing the Cu dissolution rate. However, a

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slightly increased Cu dissolution rate was evidenced in the range of 7– 14 at.% Ni for an almost constant Al amount of 30 at.% along the investigated Al–Cu–Ni thin film ternary library. At a first glance, this result may appear contradictory to previous findings on bulk Al–Cu–Ni, where an increase in Ni amount resulted in a corrosion increase [3–6]. However, the fact that Cu dissolution decreases with increasing amount of Ni is supported by the mentioned studies where the increased overall corrosion is explained by mobile cation vacancies consumption in the Cu2O layer by Ni. The observed Cu2O stabilization by Ni results in a lower Cu release rate, in spite of the reported overall corrosion resistance decrease, with increasing amount of Ni [5]. 4. Conclusion In this work, thermally evaporated Al–Cu–Ni thin films were fabricated and investigated in a combinatorial approach with a compositional gradient from 5.7 at.% to 24.8 at.% Ni. The thin film formation was done by physical vapour deposition using a two-source evaporation geometry for creating the compositional spread, where Al–Cu (used as a prealloy) and Ni were evaporated from individual sources. The surface and microstructural as well as the electrochemical properties of the deposited thin films were carefully analysed. The surface morphology and microstructure of the Al–Cu–Ni show a smooth decrease of the grain size with increasing Ni content. A monoclinic symmetry was evidenced by XRD measurements along the entire Al–Cu–Ni compositional spread. A distinct compositional zone (Al–Cu–7 at.% Ni–Al–Cu–13 at.% Ni) was identified during SKP measurements where a significant surface potential drop was observed. This compositional zone was correlated to the presence of CuO on the surface, as evidenced by XPS and Auger spectroscopy, while the rest of the compositional spread surface contained only Al2O3 together with both metallic Cu and Cu2O. Electrochemical characterization by cyclic voltammetry on various spots along the Al–Cu–Ni gradient using SDCM showed a clear Cu+ to Cu2+ oxidation directly related to the Ni content. The chemical stability of the Al–Cu–Ni thin films in terms of Cu superficial release was tested by dissolution experiments in 0.1 M HNO3 using a FT-SDCM with downstream analytics. The Cu dissolution was enhanced in the previously identified compositional zone due to the presence of mixed Cu species. Acknowledgements The financial support by the Austrian Federal Ministry of Economy, Family and Youth and the National Foundation for Research, Technology and Development (COMBOX) is gratefully acknowledged. Also, the experimental assistance provided by Mr. Jiri Duchoslav for XPS and Auger measurements is gratefully acknowledged. References

Fig. 9. Dissolution curve for normalized Cu mass of dissolution experiments in 0.1 M HNO3. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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