Study of grain growth and thermal stability of nanocrystalline RuAl thin films deposited by magnetron sputtering

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(1)Thin Solid Films 527 (2013) 1–8. Contents lists available at SciVerse ScienceDirect. Thin Solid Films journal homepage: www.elsevier.com/locate/tsf. Study of grain growth and thermal stability of nanocrystalline RuAl thin films deposited by magnetron sputtering M.A. Guitar a,⁎, K. Woll a, E. Ramos-Moore a, b, F. Mücklich a a b. Functional Materials, Materials Science Department, Saarland University, Saarbrücken D-66123, Germany Facultad de Fisica, Pontificia Universidad Catolica de Chile, Santiago 7820436, Chile. a r t i c l e. i n f o. Article history: Received 12 March 2012 Received in revised form 9 December 2012 Accepted 10 December 2012 Available online 23 December 2012 Keywords: Aluminides RuAl Thermal stability Grain growth Thin films. a b s t r a c t Special properties of RuAl regarding temperature stability and resistance against corrosion and oxidation, make this intermetallic material suitable to be used under high-temperature and chemically aggressive working conditions. In the present work, the microstructure, thermal stability and grain growth are studied by means of scanning transmission electron microscopy, peak broadening analysis of X-ray diffraction and electron backscatter diffraction. Isothermal kinetics analysis revealed that annealing at temperatures as low as 650 °C does not activate the grain growth, whilst at 700 °C and times lower than 6 h, the growth is induced by a 3D curvature-driven evolution characterised by a monomodal grain size distribution. Higher temperatures and/or annealing times show also abnormal grain growth characterised by a bimodal distribution of the grain sizes, which is correlated with the presence of impurities and grain-orientation-specific driving forces. The corresponding grain boundary activation energy is 145± 14 kJ/mol, indicating that the growth is highly dominated by grain-boundary diffusion due to the high volume fraction of grain boundaries compared with coarse-grained polycrystalline materials. © 2012 Elsevier B.V. All rights reserved.. 1. Introduction The special interest to study RuAl started when a strong resistance of this intermetallic was observed against aqueous corrosion in very aggressive mediums, like HNO3, aqua regia, FeCl3 and HF [1], and also its resistance to oxidation up to at least 900 °C [2]. These properties are due to the formation of an Al2O3 protective scale, which has been found to be dense and compact even after oxidising during 100 h at 1000 °C [3]. Further, the coefficient of thermal expansion of RuAl is nearly equal to that of Al2O3 in a large temperature range, avoiding the detachment of the protective scale, which makes favourable the use of RuAl-based intermetallics as protective coatings in applications that demand oxidation resistance [4,5]. For these reasons and also due to its high melting point (about 2050 °C) and high temperature strength, the intermetallic compound RuAl is regarded as a potential material for high-temperature applications in aggressive environments [1]. On the other hand, RuAl coatings can also act as working layers for Glass Moulding Dies (GMD) in order to increase their service life. The nature of the glasses used to produce precision lenses is limited by the properties of the materials used in the pressing method. Normal optical glasses need to be heated up to 550 °C–750 °C in order to be ⁎ Corresponding author at: Dept. Materials Science & Engineering, Saarland University, Campus D3.3, room 3.13, D-66123 Saarbruecken, Germany. Tel.: +49 681 302 70521; fax: +49 681 302 70502. E-mail address: a.guitar@mx.uni-saarland.de (M.A. Guitar). 0040-6090/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2012.12.046. pressed. At that temperature, the chemical activity of the glass is greatly increased, thereby increasing the adhesion between the moulding tool and the glass surfaces [6]. For these reasons, the working layers must have satisfactory thermal stability and oxidation resistance to obtain good dimensional accuracy of the glass components. NiAl has been studied as a working layer for GMD by Zhong et al. [5,7] due to its good oxidation resistance and non-wettability in the presence of the molten glass. RuAl, like NiAl, has a B2 structure, but possesses a higher melting point and better oxidation resistance. These comparative advantages make RuAl a promising material for applications that need protective layers such as glass moulding dies and thermal actuators [8,9]. Nanocrystalline (NC) materials present grains smaller than 100 nm and a large volume fraction of interfacial area formed by all atoms situated at the grain boundaries [10]. Many structural and functional characteristics of NC materials depend strongly on thermal, chemical or phase stability. In particular, thermal stability of NC materials is of considerable importance for all high-temperature applications [11,12]. The phase formation in multilayers and electrical conductivity studies have been reported for RuAl thin films [9,13]. The thermal stability of nano-grained RuAl produced by mechanical alloying was previously studied by Liu et al. [14,15]. For the thin-film application of the intermetallic compound RuAl, it is of high interest to know the phase, the grain size and also the thermal stability at elevated temperatures. Any studies on thermal stability RuAl thin films have been found in the literature. In the context of the previously mentioned applications, the present work targets on the phase and thermal stability of NC RuAl thin films..

(2) 2. M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. For this purpose, RuAl thin films were produced by vapour deposition on stainless steel substrates. The samples were subsequently isothermally annealed at three different temperatures and the kinetics of grain growth was studied in the isothermal regime. Complementary techniques were used in order to study the grain evolution as a function of temperature and time. Image analysis (I-A) of scanning transmission electron microscopy (STEM) pictures and peak broadening analysis of X-ray diffraction (XRD) results were used to analyse the grain evolution. Electron backscatter diffraction (EBSD) was also performed in order to analyse the evolution of the texture and grain size distribution. 2. Experimental procedures The RuAl thin films were deposited from high-purity (99.99%) Al and Ru Targets, onto 15 mm× 15 mm AISI 316L stainless steel substrates. The deposition was carried out by AC/DC magnetron sputtering using a Von Ardenne PPS-A 200 source. The plasma was ignited on both targets at the same time whilst the sample plate rotated. In order to achieve better adhesion, a Ti layer (50 nm) was sputtered onto the substrate previous to the Ru–Al deposition. The deposition sequence started with Al, each layer was around 1 nm–2 nm thick and, at the end, a 600 nm thick multilayer was obtained. After the deposition, the samples were annealed under vacuum (~10 −3 Pa) in a tube furnace at 600 °C during 1 h in order to obtain RuAl single-phase films. The heating and cooling rate was 1 °C/min. The very short period of the multilayer, as well as the annealing parameters, leads to the direct formation of single-phase RuAl in the absence of other intermediate phases and stratified geometry [13]. The microstructure of the fabricated samples was analysed in a dual-beam focused ion beam/scanning electron microscopy workstation (FEI Helios NanoLab 600), which also allows STEM imaging. The phase and thermal stability of the fabricated RuAl thin films were studied by XRD using Cu Kα radiation (λ =0.1542 nm) at 40 kV and 40 mA in a PHILIPS X'Pert Pro MPD X-ray diffractometer. The diffraction geometry was θ–θ, and the incident and reflected optics consisted of parallel beams. The thermal treatments were carried out under vacuum (1× 10−3 Pa), and the samples were heated and cooled at 20 °C/min from room temperature up to 650 °C, 700 °C and 750 °C. For each of these temperatures, the samples were annealed during different intervals of time. The I-A consisted in the reconstruction of the grain boundaries, and the calculation of the grain sizes was performed using the software Aquinto-A4i. The average size and distribution of the grain size were calculated by measuring the mean feret, which is defined as the distance between two tangents to the contour of the particle, averaged over 36 different directions. EBSD was carried out on the surface of the samples (20 μm × 20 μm) using a voltage of 15 kV, a current of 11 nA, and a step size of 20 nm. No filter was applied to the raw data and the grains were defined as at least two adjacent points with similar orientation within a range of 5° of misorientation. Grain size distribution and inverse pole figures were generated for different grain populations (size partitions) in order to complement the information obtained by I-A and XRD. 3. Results and analysis 3.1. Phase stability and grain growth Fig. 1 shows the XRD patterns of the isothermally annealed samples at 650 °C, 700 °C and 750 °C. The diffraction curves were normalised using the most intense peak {100}. RuAl single-phase films were observed to be stable, even after annealing the samples for 6 h at 650 °C, 700 °C and 750 °C. As shown in Fig. 1, the diffractograms present the {100} and {200} peaks for all the temperatures and during the whole annealing process. As shown in the STEM micrograph of Fig. 2 (a), after the single-phase transformation RuAl presents grains with an equiaxed structure and an. Fig. 1. RuAl powder diffraction pattern (PDF) and room-temperature XRD patterns for non-treated sample (RT), and those isothermally annealed for 6 h at 650 °C, 700 °C and 750 °C.. average grain size of 24± 2 nm. Fig. 2 (b) and (c) shows the STEM micrograph of the sample annealed for 3 h at 750 °C and the corresponding grain boundary reconstruction, respectively. Similar grain boundaries can be observed in both, thus relevant and reliable information such as grain size distribution and grain area average could be obtained. The grain structure of the RuAl films annealed at 650 °C, 700 °C and 750 °C for different annealing times was characterised by STEM. The I-A calculated grain sizes as functions of the annealing temperature are plotted in Fig. 3. An increase in the mean grain size with increasing temperature was observed, reaching the values of 57± 2 nm, 74 ± 4 nm, 102 ± 5 nm and 107 ±8 nm after annealing at 750 °C for 1 h, 2 h, 4 h and 6 h, respectively. Fig. 4 shows the STEM micrographs of the samples annealed at 700 °C for 6 h and at 750 °C for 3 h, 4 h and 6 h. Grains associated with normal and abnormal grain growth are observed in the crosssection. The grains that showed abnormal growth have reached sizes which exceed the half thickness of the film, as shown in Fig. 4 (c) and (d). The histograms of grain size distribution are shown in Fig. 5. The grain size distribution, f(d), was fitted using the log-normal function described by Eq. (1) [16]: −1. f ðdÞ ¼ ðd⋅σ Þ. −1=2. ð2πÞ. h i 2 2 ⋅ exp –ð lnd−μ Þ =2⋅ σ. ð1Þ. The parameter d is the grain size, and the fitting parameters μ and σ correspond to the mean value of ln(d) and the standard deviation, respectively. All fitted curves showed a correlation higher than 92%. The log-normal distributions of grain size are typical for many nano-materials in initial and annealed states [17,18]. Fig. 5 (b) shows the average grain size for the samples annealed at 650 °C for 6 h. This distribution indicates that a normal grain growth occurs at these annealing parameters. A similar behaviour was observed for the samples annealed at the same temperatures, but for lower times. The specimens treated at 700 °C do not present abnormal grains when they are annealed for few hours (up to 4 h). As presented in Fig. 5 (c), the grain size distribution of the sample annealed at 700 °C during 1 h shows normal growth. Abnormal growth was observed in the samples annealed at 700 °C after 4 h, as shown in Fig. 5 (d). The samples annealed at 750 °C.

(3) M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. 3. Fig. 2. (a) STEM image of a RuAl sample after the heat treatment to induce single-phase formation, (b) STEM image of an annealed sample at 750 °C for 3 h and (c) grain boundary reconstruction of image (b) using the software A4i.. showed a normal grain growth up to 3 h of annealing (Fig. 5 (e)) and after, the grain growth was abnormal, as shown in Fig. 5 (f) (6 h at 750 °C). 3.2. Grain growth kinetics The kinetics of normal grain growth in pure and single-phase materials is modelled by a power growth law [13,16–19], which assumes that the driving force for grain growth is proportional to the grain boundary curvature. The law is based on empirical parameters and is presented in Eq. (2). n. n. Dðt Þ –Dð0Þ ¼ k⋅t. ð2Þ. In Eq. (2), the term D0 is the initial average grain size, D(t) is the grain size after annealing for a time t, k is a constant describing the grain boundary mobility, and n is an empirical grain growth exponent. k ¼ k0 ⋅ expð−Q =RT Þ. ð3Þ. Eq. (3) describes the constant k from Eq. (2). The term k0 is a weakly temperature-dependent constant, Q is the activation energy for grain growth, R is the universal gas constant and T is the temperature in Kelvin. Consequently, the n value and the growth rate must be determined from the temporal evolution of grain size. Moreover,. Fig. 3. Grain size obtained by using the software A4i as a function of the annealing temperature, for 1 h, 2 h, 4 h and 6 h.. the grain growth activation energy is calculated from the temperature dependence of k. In previous works [16,19,20], the grain growth exponents (n) were found to be in the n =2–10 range, and in some systems [19,20], temperature dependent. In Fig. 6, the isothermal grain growth data, obtained by I-A, and the corresponding fit using Eq. (1), are shown. The term n was obtained from these data by fitting the grain sizes to Eq. (1), and it was found to be 6.1 ±0.1, 4.1 ±0.1 and 3.2 ± 0.1, for the annealing temperatures of 650 °C, 700 °C and 750 °C, respectively. The grains in NC metals grow until a certain grain size limit is reached, and then the grain growth ceases. A stagnation of the grain growth has not been observed after the diverse annealing times used in this study, which would be depicted by a flattening of the curves shown in Fig. 6. The samples annealed at 650 °C show only a weak increase in the grain size. In the case of the samples annealed at 700 °C and 750 °C, the increase in grain size with time becomes more pronounced when the temperature is increased. In both cases, a more significant change in grain size was observed, reaching values up to about 55 nm and 107 nm, respectively. The I-A of STEM images used in the previous analysis is a direct method, which allows only a local analysis in a cross-section of the film. Otherwise, a much larger area of the sample can be analysed using an indirect method based on XRD. Since the information is extracted from the interaction volume within the sample (~1 μm3), the XRD line broadening is related with the average of the micro-grain sizes and microstrains. For this reason, the grain size was also calculated from XRD data using the Warren–Averbach (W–A) method [21,22] and compared to those obtained from I-A. The W–A analysis was carried out using the Bragg diffraction planes {100} and {200} shown in Fig. 1. Fig. 7 compares the data obtained from the I-A and the W–A methods for the samples annealed at 750 °C. The grain size values obtained by the W–A method are 10% to 12% lower than those calculated by the I-A method. Usually, for those samples showing a grain size smaller than 35 nm, the W–A method shows bigger values than the I-A method, probably due to the overlapping of two or more grains observed in the STEM micrographs. The grain size values and their distribution were also studied using EBSD on the surface of the samples. This technique allows obtaining information of the grain sizes and distribution in a large area of the sample (~1 mm2). As shown in Fig. 7, the mean grain size obtained by this technique is comparable to those values obtained previously by I-A and by XRD measurements using the W–A method. Fig. 8 shows the distribution of the grain sizes for samples annealed at 700 °C for 6 h and at 750 °C for 3, 4 and 6 h. As shown in Fig. 9, the grain size distributions obtained by EBSD present a bimodal distribution due to the normal and abnormal grain growth. For samples annealed at 750 °C for 3 h, 4 h and 6 h, grains that exceed diameters of 1.3 μm, 0.4 μm and 0.6 μm, respectively, were considered abnormal. Whereas grains that exceed 0.3 μm were also considered abnormal for the sample annealed.

(4) 4. M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. Fig. 4. STEM images showing the abnormal grain growth: (a) bright field image of a sample annealed 6 h at 700 °C; (b) bright field of a sample annealed for 3 h, (c) bright field of a sample annealed for 4 h, and (d) dark field of a sample annealed for 6 h at 750 °C.. at 700 °C for 6 h. Inverse pole figures were estimated considering the grain partition indicated with dash lines in Fig. 9. It is worth noting that, independently from the annealing conditions, the abnormal grains show a different growth direction with respect to the normal grains. The normal grains grow along the 001 crystallographic orientation, whilst the abnormal grains tend to grow towards the 111 crystallographic orientation.. The activation energy for grain growth Q can be estimated measuring the grain growth rate k of the average grain size at different temperatures. The growth rate is ruled by the Arrhenius relationship presented in Eq. (4) [23]:. k ¼ k0 ⋅ expð−Q =RT Þ or lnðkÞ ¼ lnðk0 Þ þ ð−Q=RT Þ:. ð4Þ. Fig. 5. Grain size distribution for the samples: (a) after RuAl single-phase formation; (b) after annealing at 650 °C for 6 h; (c) and (d) after annealing at 700 °C for 1 h and 6 h, respectively; (e) and (f) after annealing at 750 °C for 1 h and 6 h, respectively..

(5) M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. 5. of ln(k) as a function of 1 / T. The activation energy of the grain growth (Q = 145± 14 kJ/mol) was estimated from the slope of the linear minimum-squares fit shown in Fig. 10 in the temperature range between 650 °C and 750 °C. 4. Discussion. Fig. 6. Grain size of RuAl thin films as a function of annealing time, for annealing temperatures of 650 °C, 700 °C and 750 °C.. The term k0 is a constant, R is the universal gas constant and T is the temperature in Kelvin. In the present work, the rate of grain growth, k, was calculated considering the grain growth rate for the samples annealed at 650 °C, 700 °C and 750 °C for 6 h. Fig. 10 shows the plot. Fig. 7. Grain size evolution estimated by the I-A, W–A and EBSD methods, for samples annealed at 750 °C. The errors on the W–A grain sizes in panel (a) cannot be seen because of their small values, whereas the errors from the EBSD measurements are only shown in panel (b) to permit a better comparison view.. For an average grain size much smaller than the film thickness, the isothermal growth is induced by a 3D curvature-driven evolution, characterised by a monomodal grain size distribution. This growth process is also observed in bulk materials and obeys Eq. (2), as explained by Thompson [16]. For the RuAl films studied in this work, the ratio of the average grain size (24 nm) and the film thickness (600 nm) is 0.04; thus the grain growth kinetics has been found to match this behaviour for most of the temperature and time ranges. In fact, the grain size distribution showed a fit correlation of ≥92% with the log-normal distribution (Fig. 5), and agrees with similar experiments performed in NC materials in the initial and annealed states [17,18]. For annealing temperatures of 700 °C and times equal to or longer than 6 h, an abnormal growth was observed, showing a grain distribution with a bimodal behaviour. This behaviour was also observed after annealing the samples 4 h at 750 °C (Fig. 5). Abnormal grain growth can be the result of: (1) a change of the grain boundary energies at the surface and/or the film/substrate interface [16,24]; (2) the presence of impurities and/or precipitates [14,25]; and (3) grain-orientation-specific driving forces. The intersection of a grain boundary with the surface of the films can lead to the formation of grooves, which can result in grain growth stagnation [26]. At this point, those grains whose crystallographic orientation favoured the minimisation of the surface energy will preferentially growth. The presence of impurities or second phase particles can also restrict the mobility of the grain boundary, allowing the growth of a subpopulation of grains. In this case, abnormal grain growth occurs when the grain boundaries between abnormal and normal grains have much higher mobility than between normal grains. The grains in a strongly textured material show low angle boundaries and consequently low mobility. Grains presenting sufficiently different crystallographic orientation relative to the normal matrix grains have high angle boundaries, with high mobility, and therefore the grain growth in different directions is more probable [26,27]. Preliminary energy dispersive X-ray analyses performed on the centre of the RuAl thin film (not presented in this paper) showed the existence of ~ 3.0 at.%–3.5 at.% of Fe and ~ 1.5 at.%–2.0 at.% of Cr for the samples annealed at 700 °C, and ~ 6.0 at.% of Fe and ~ 3.0 at.% of Cr for those annealed at 750 °C. Fe and Cr impurities in the RuAl thin films are the consequence of a diffusion process of substrate elements from the Ti adhesion interlayer and/or the substrate. Nevertheless no precipitates containing Fe and Cr were observed throughout all the performed analyses. Thus, the impurities might remain segregated at grain boundaries or in the form of RuAl(Fe) of RuAl(Cr) [14]. When the grain boundaries move, the segregated atoms will attempt to remain in the boundary, forcing the grain border to drag its impurity load, so that it can only migrate as fast as the slowly moving impurities [28]. The presence of impurities retards or stops the normal behaviour, and with sufficient time, the abnormal grains will change the complete crystallographic orientation of the sample, whose tendency was observed in Fig. 9. The grain growth exponents n of the RuAl films were found to be dependent on the temperature with values equal to 6.1 ± 0.1, 4.1 ± 0.1 and 3.2 ± 0.1 for 650 °C, 700 °C and 750 °C, respectively. For normal growth in coarse-grained materials, the parameter n is expected to be 2, but for NC materials, variations have been observed [19,20,25]. The general trend is to increase towards 2 when increasing the annealing temperature [29], until the growth reaches a maximum size, and then the grain growth ceases. Many factors influence grain boundary mobility, leading to the deviation of n. These include.

(6) 6. M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. Fig. 8. EBSD scans of the surface of the samples annealed for (a) 6 h at 700 °C, for (b) 3 h, (c) 4 h and (d) 6 h at 750 °C. The different colours on the picture indicate determinate crystallographic orientations.. impurities (solute drag), precipitate or pore drag (temperature sensitive), surface grooving, grain boundary energy anisotropy, and initial texture [16,19,20,30]. Due to the fact that grain growth in NC materials occurs in a different manner than that in coarse-grained (CG) polycrystalline materials [31], it may be difficult to identify the grain growth mechanism on the basis of the exponent n alone [17,23,31,32]. The grain size of RuAl thin films was calculated by I-A using STEM images and compared to the results obtained by XRD methods and EBSD measurements. All the methods showed results for grain size within the same order of magnitude. However, the grain sizes calculated through XRD methods were found to have smaller values than those obtained via the I-A of STEM pictures. It has been indicated by other authors [21,33] that the grain size obtained by the W–A method looks, in almost cases, smaller than the size obtained by TEM or STEM observations. When a bimodal grain size distribution is analysed, the grain sizes obtained by the W–A method are always lower than the ones obtained from other measurements. A possible reason for this disagreement is that the origins of diffraction line broadening in a real sample are not only due to effects of small crystallite size and microstrain, but are also due to lattice imperfections [22]. The reason for which for small grains (≤35 nm) the grain size values calculated by W–A method are larger than those obtained by I-A, is probably the overlapping of grains when an image is made from a STEM foil. This is a usual problem in the case of NC materials when the grain size is smaller than the foil thickness (≤100 nm). The mean grain size obtained by EBSD measurements is comparable to the values obtained previously by I-A and by XRD measurements using the W–A method, as shown in Fig. 7. However, EBSD showed a. bimodal distribution and the average size obtained by this method presents a significant standard deviation due to the presence of abnormal grains. Indeed, 60% to 70% of the grains are concentrated in a range which does not exceed 80 nm and the rest is divided, in different percentages, between grains smaller than 45 nm and bigger than 90 nm. Grain sizes larger than 3 μm were observed for samples annealed at 750 °C for more than 4 h, whereas the samples annealed at 700 °C for 6 h and at 750 °C for 4 h have reached values of about 1.7 μm and 1.1 μm, respectively. These particle sizes are in some cases up to three or four times larger than those observed by STEM. This is probably due to the cutting direction and location of the STEM foil, which can result in making the image projection smaller than the real size of the grain. The activation energy for grain growth in RuAl films for a temperature range between 650 °C and 750 °C has been found to be Q= 145 ±14 kJ/mol, which is comparable to the value observed in NC NiAl (166.7 kJ/mol) [23]. Values lower than 145 ±14 kJ/mol (39 kJ/mol for the temperature range of 600 °C–750 °C) have been found previously by Liu et al. [14] for the activation energy in nano-RuAl produced by mechanical alloying. They observed that the grain stagnation occurs after annealing the samples during 3 h due to the presence of impurities (Mn, Cr and up to 15 at.% Fe), which retards the grain growth by “solute drag” during the movement of the grain boundaries. Impurity segregation on the grain boundary reduces the specific grain boundary energy and therefore the driving force for grain growth. Activation energy data for grain growth in NC materials are often compared to the activation energy associated with the lattice (Ql) and the grain boundary (Qgb) diffusion of the polycrystalline compound, in order to determine the mechanisms involved in the grain growth process. For the most pure.

(7) M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. 7. Fig. 10. Arrhenius plot for the estimation of the activation energy of grain growth, for a temperature range between 650 °C and 750 °C.. metals, the ratio of the activation energy for lattice and grain-boundary diffusion (Qgb/Ql) ranges between 0.4 and 0.6 [29]. The activation energy value for the grain growth determined in this work is much lower than the activation energy for lattice self diffusion of conventional CG RuAl (236± 18 kJ/mol) [34]. The ratio of both is 0.61±0.01, suggesting that grain growth in NC RuAl is dominated by grain boundary diffusion. This is expected to occur, since in NC films the fraction of grain boundaries is higher than in CG materials.. 5. Concluding remarks. Fig. 9. Grain size distributions and inverse pole figures obtained from EBSD analysis. The figures were estimated by considering a bimodal distribution that presents normal (left) and abnormal (right) grain growth. The [001] direction corresponds to the normal of the sample surface and the normalised intensity scale is linear between blue (minimum) and red (maximum).. Single-phase RuAl intermetallic thin films have been produced from Ru–Al multilayers with individual thickness of 1–2 nm, followed by annealing under vacuum at 600 °C during 1 h. The B2-RuAl intermetallic presented a preferred crystallographic orientation along the {100} planes parallel to the sample surface. Using complementary techniques (I-A, XRD and EBSD), it was found that the grain size values estimated by XRD were smaller than the ones obtained by I-A. The values calculated by EBSD were also comparable to those obtained by the previously mentioned methods. A bimodal distribution of the grain size was noticed with this method, which includes normal and abnormal grains. Thermal studies showed that temperatures as low as 650 °C are not sufficient to activate the grain growth. However, annealing at 700 °C and 750 °C increased the grain size from about 24 nm at room temperature up to 54 nm and 107 nm, respectively. The isothermal grain growth of RuAl thin films obeys the generalised power grain-growth model and the exponent n was found to be dependent on the temperature. For most of the temperatures and times used, normal grain growth of RuAl films has occurred. However, abnormal growth behaviour was observed for 6 h at 700 °C and 4 h and 6 h at 750 °C. From the three possible mechanisms of abnormal growth, we conclude that the presence of impurities and grain-orientation-specific driving forces are both responsible for the observed abnormal behaviour. The corresponding activation energy for NC RuAl grain growth has been found to be 145 ± 14 kJ/mol, which is in agreement with reported values for other B2 NC intermetallics. The ratio of the activation energy for lattice (literature) and grain boundary diffusion (this work) was 0.61 ± 0.01, indicating that the grain growth in NC RuAl is highly dominated by grain-boundary diffusion due to the high volume fraction of grain boundaries compared with CG bulk materials..

(8) 8. M.A. Guitar et al. / Thin Solid Films 527 (2013) 1–8. Acknowledgements This study was funded within a research project MU 959/24-1 of the Deutsche Forschungsgemeinschaft (DFG). The authors would like to thank the EFRE Funds of the European Commission for support of activities within the AME-Lab project. The authors are also grateful to Prof. Seidel, from the Department of Mechatronics, Saarland University, for the use of the magnetron sputtering devices. A. Guitar is grateful to the German Academic Exchange Service (DAAD) for the financial support. References [1] I. Wolff, JOM 49 (1997) 34. [2] F. Mücklich, N. Ilić, K. Woll, Intermetallics 16 (2008) 593. [3] F. Soldera, N. Ilić, S. Brännström, I. Barrientos, H. Gobran, F. Mücklich, Oxid. Met. 59 (2003) 529. [4] B. Tryon, T.M. Pollock, M.F.X. Gigliotti, K. Hemker, Scripta Mater. 50 (2003) 9. [5] M. Jackson, US4980244—Protective Alloy Coatings Comprising Cr–Al–Ru, 1990. [6] F. Klocke, T. Bergs, K. Georgiadis, H. Sarikaya, F. Wang, in: Proc 7th Inter Conf “The Coatings in Manufacturing Engineering”, Greece, 2008, p. 209. [7] D. Zhong, G.G. Mustoe, J. Moore, J. Disam, Surf. Coat. Technol. 146–147 (2001) 312. [8] D. Zhong, E. Mateeva, I. Dahan, J.J. Moore, G.G.W. Mustoe, T. Ohno, Surf. Coat. Technol. 133–134 (2000) 8. [9] J.A. Howell, C.L. Muhlstein, B.Z. Liu, Q. Zhang, S.E. Mohney, J. Microelectromech. S 20 (2011) 933.. [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]. X. Song, J. Zhang, L. Li, K. Yang, G. Liu, Acta Mater. 54 (2006) 5541. K.W. Liu, F. Mücklich, R. Birringer, Intermetallics 9 (2001) 81. R. Andrievski, J. Mater. Sci. 8 (2003) 1367. N. Zotov, K. Woll, F. Mücklich, Intermetallics 18 (2010) 1507. K.W. Liu, F. Mücklich, Acta Mater. 49 (2001) 395. K.W. Liu, F. Mücklich, Intermetallics 13 (2005) 373. C.V. Thompson, Ann. Rev. Mater. Sci. 20 (1990) 245. J. Palmer, C.V. Thompson, H. Smith, J. Appl. Phys. 62 (1987) 2492. R. Dannenberg, E. Stach, J.R. Groza, B.J. Dresser, Thin Solid Films 379 (2000) 133. M. Iordache, S. Whang, Z. Jiao, Z. Wang, Nanostruct. Mater. 11 (1999) 1343. L. Wang, X.Y. Qin, W. Xiong, L. Chen, M.G. Kong, Mater. Sci. Eng., A 434 (2006) 166. Z. Zhang, F. Zhou, E. Lavernia, Metall. Mater. Trans. A 34 (2003) 1349. T. Ungar, J. Gubicza, TMS Ann. Meet. (2002) 595. L.Z. Zhou, J.T. Guo, Scripta Mater. 40 (1998) 139. D.A. Porter, K.E. Easterling, Phase Transformations in Metals and Alloys, CRC Press, 1992. , (ISBN 0-7487-5741-4). C.V. Thompson, Ann. Rev. Mater. Sci. 30 (2000) 159. C.V. Thompson, R. Carel, J. Mech. Phys. Solids 44 (1996) 657. A. Rollett, Acta Metall. 37 (1989) 1227. G. Gottstein, L.S. Shvindlerman, Grain Boundary Migration in Metals: Thermodynamics, Kinetics, Applications, CRC Press, 2009. , (ISBN 978-1-4200-5435-4). J.C.M. Li, Microstructure and Properties of Materials, World Scientific Publishing Co Pte Ltd, 2000. , (ISBN 981-02-4180-1). M. Huang, Y. Wang, Y.A. Chang, Thin Solid Films 449 (2004) 113. M. Datta, S. Pabi, B. Murty, Mater. Sci. Eng., A 284 (2000) 219. C.E. Kril, R. Birringer, Philos. Magn. A 77 (1998) 621. T. Ungar, J. Gubicza, G. Ribárik, A. Borbély, J. Appl. Crystallogr. 34 (2001) 298. K. Woll, C. Holzapfel, H.A. Gobran, F. Mücklich, Scripta Mater. 57 (2007) 1..

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