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Growth and structure of ultrathin alumina films on the (110) surface of -Al4Cu9 complex metallic alloy
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 J. Phys.: Condens. Matter 26 485009 (http://iopscience.iop.org/0953-8984/26/48/485009) View the table of contents for this issue, or go to the journal homepage for more Download details: IP Address: 132.239.1.231 This content was downloaded on 11/05/2017 at 09:59 Please note that terms and conditions apply.
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Journal of Physics: Condensed Matter J. Phys.: Condens. Matter 26 (2014) 485009 (11pp)
doi:10.1088/0953-8984/26/48/485009
Growth and structure of ultrathin alumina films on the (1 1 0) surface of γ -Al4Cu9 complex metallic alloy M Warde´ 1,2 , J Ledieu3 , L N Serkovic Loli3 , M Herinx1,2 , M-C de Weerd3 , 1,2 ´ ´ 3 , S Le Moal1,2 and M-G Barthes-Labrousse V Fournee 1
Institut de Chimie Mol´eculaire et Mat´eriaux d’Orsay, UMR 8182 CNRS, Universit´e Paris-Sud, Bˆatiment 410, 91405 Orsay Cedex, France 2 CNRS, Orsay, F-91405, France 3 Institut Jean Lamour, UMR 7198 CNRS, Universit´e de Lorraine, Parc de Saurupt, 54011 Nancy Cedex, France E-mail: [email protected] Received 21 July 2014, revised 29 September 2014 Accepted for publication 10 October 2014 Published 6 November 2014 Abstract
The first stages of oxidation of the (1 1 0) surface of a γ -Al4 Cu9 complex metallic alloy were investigated by combining x-ray photoemission spectroscopy, low energy electron diffraction and scanning tunnel microscopy studies. Oxidation at room temperature in the 2 × 10−8 to 2 × 10−7 mbar oxygen pressure range occurs in two steps: a fast regime is followed by a much slower one, leading to the formation of a thin aluminium oxide film showing no long range order. Cu–O bonds are never observed, due to fast oxygen induced aluminium segregation. The low value of the estimated activation energy for aluminium diffusion (0.65 ± 0.12 eV at−1 ) was ascribed to the presence of two vacancies in the γ -Al4 Cu9 structure. Annealing at 925 K the oxide film formed at room temperature leads to the formation of small crystallized domains with a sixton structure similar to structures reported in the literature following the oxidation of Cu-9%Al(1 1 1), NiAl (1 1 0) and FeAl(1 1 0) surfaces as well as ultrathin Al films deposited onto Cu(1 1 1) or Ni(1 1 1) surfaces. Two contributions were observed in the O1s peaks, which have been ascribed to loosely bound oxygen species and oxygen belonging to the sixton structure respectively. Keywords: ultrathin alumina film, (1 1 0) γ -Al4 Cu9 , complex metallic alloy (Some figures may appear in colour only in the online journal)
In particular, a number of studies related to interactions with oxygen were driven by the excellent oxidation and corrosion resistance initially reported for Al–Cu–Fe and Al–Cu–Cr–Fe quasicrystals [4, 5]. Following this pioneer work, most investigations to date have been devoted to the first stages of oxidation of Alrich complex metallic alloys. Similar to more conventional intermetallics, oxidation of these materials in a vacuum environment leads to Al segregation and the formation of a passivating aluminium oxide layer, a few nm thick (for a review, see [6, 7] and the references therein). In most cases, the oxide layer has been found to be amorphous. However, following high temperature oxidation of the pentagonal
1. Introduction
Complex metallic alloys (CMAs) are characterized by unit cells containing from a few tens to several thousand atoms arranged into highly symmetric clusters. Quasicrystals represent the ultimate complexity, having an infinite number of atoms within the cell. CMAs have attracted a lot of interest during the last two decades, due to their potentially useful chemical and mechanical properties [1–3]. The interest in studying CMAs’ surface reactivity comes from their specific electronic structure, which is related to the existence of highly symmetric clusters and can affect interactions of the surface atoms with surrounding atoms and molecules. 0953-8984/14/485009+11$33.00
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© 2014 IOP Publishing Ltd Printed in the UK
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surface of i-Al–Pd–Mn and the 10-fold-symmetry surface of the decagonal Al–Co–Ni, the formation of nano-domains of hexagonal aluminium oxide thin layers aligned along a two-fold or five-fold symmetry direction of the substrate respectively has been reported [8–10]. The limit in the lateral size of the domains was ascribed to the lack of translational order in the structure of the quasicrystalline substrates used in these studies. To our knowledge, such ordered films have never been observed on the surface of periodic complex metallic alloys, although well-ordered oxide films can be easily obtained on a number of conventional binary alloy surfaces (see for example [11]). Due to the importance of Cu-rich Al–Cu alloys in catalysis and microelectronics, their oxidation behaviour has been extensively studied. In particular, oxygen adsorption in ultrahigh vacuum (UHV) conditions on the (1 0 0) surface of α-Cu–Al solid solutions containing increasing quantities of aluminium (0.15, 1.5, 5, 12 and 17 at% aluminium) has been followed by Hoffman and co-authors using Auger electron spectroscopy (AES), photoemission spectroscopies (XPS and UPS) and low energy ion spectroscopy (LEIS) [12–17]. Two stages are observed in the oxidation kinetics when oxygen exposure is performed at room temperature (RT). Initially, chemisorbed oxygen quickly reacts with surface aluminium sites, leading to the formation of a thin aluminium oxide with AlO stoichiometry. This fast process is observed up to 7–20 L of oxygen exposure, depending on the aluminium concentration in the alloy. It is followed by a much slower stage during which both Al–O and Cu–O bonds are formed, the formation of new Al–O bonds being ascribed to the diffusion of oxygen in the bulk and Al segregation to the surface. When the exposure to oxygen is for long enough, the Al2 O3 stoichiometry is reached. Annealing the oxidized alloy surfaces in the 520–725 K temperature range results in the disappearance of the Cu–O bonds, an increase in the aluminium segregation kinetics and the formation of a stable pure alumina layer. A number of studies have also been devoted to the structure of the aluminium oxide films formed on (1 1 1) surfaces of Cu–Al alloys containing more than 9 at% Al. Saturation with oxygen at 670 K leads to an amorphous oxide layer which crystallizes following annealing at 920 K, √ √ as indicated by the appearance of a ( 3 × 3) low energy electron diffraction (LEED) pattern. When oxidation is performed at higher temperatures (up to 995 K), a √ √ (7 3 × 7 3)R30◦ pattern appears which has been ascribed to the formation of a γ -Al2 O3 aluminium oxide film having the orientation relationship Al2 O3 (1 1 1)//Cu9%atAl(1 1 1) [18–23]. However, Napetschnig et al have shown that moderate oxidation at a high temperature (100 L at 953 K) leads to the formation of a complex LEED diagram which can be ascribed to the formation of a thinner aluminium oxide film (0.5 nm thick) similar to that observed on NiAl(1 1 0) and consisting of two layers having Al16 O24 and Al24 O28 stoichiometries respectively (overall stoichiometry Al2 O2.6 ) [24, 25]. To the best of our knowledge, the first stages of the oxidation of Al–Cu intermetallics have not been reported so
far. However, the behaviour of the quasicrystalline icosahedral i-Al62 Cu25.5 Fe12.5 and the crystalline tetragonal ω-Al7 Cu2 Fe phases has been studied by Rouxel et al who have shown that a thin aluminium oxide layer is formed under UHV conditions during the first stages of oxidation, whereas at higher oxygen exposure and temperature, the structural complexity of the icosahedral phase may delay or even prevent the nucleation and epitaxial growth of γ -Al2 O3 crystallites compared to the tetragonal phase [26–28]. In this paper, we will present results obtained for the first stages of oxidation of the (1 1 0) surface of γ -Al4 Cu9 at RT, and the effect of subsequent annealing at 925 K on the oxide film structure. The γ -Al4 Cu9 phase presents an intermediate complexity, having a simple cubic structure with a lattice parameter equal to 0.871 nm and 52 atoms per unit mesh (16 Al + 36 Cu) [29, 30]. With 0.127 valence electrons per unit volume, its electronic structure shows some similarity with the icosahedral Al–Cu–Fe phase (0.124 e- Å−3 ) [31]. It can be described by a bcc packing of 26-atom clusters or by the stacking of flat and puckered {1 1 0} planes along the [1 1 0] direction. The structure of the clean (1 1 0) surface of γ -Al4 Cu9 has been extensively investigated by Gaudry et al by combining LEED, XPS, UPS and scanning tunnel microscopy (STM) experiments with density functional theory (DFT) calculations [32]. These authors have shown that cleaning cycles of sputtering and annealing at 933 K lead to the selection of puckered planes as preferential surface termination. These planes correspond to a simple bulk truncation and contain 6 Al and 12 Cu atoms arranged in a rectangular unit mesh 8.76 Å × 12.4 Å. It will be shown in this paper that, due to Al segregation, an aluminium oxide layer showing no long range order is obtained when RT oxidation is performed on the (1 1 0) surface of γ -Al4 Cu9 . Subsequent annealing of this oxide layer at 925 K leads to the appearance of small crystallized oxide islands having a sixton structure similar to the structure obtained for aluminium oxide films formed on NiAl(1 1 0). 2. Experimental details
The preparation of the γ -Al4 Cu9 crystal has previously been described in detail [32]. Induction melting of the pure elements (Al-99.9% and Cu-99.99%) under an argon atmosphere is followed by homogenization by annealing at 1323 K for 4 h and slow cooling down to RT. A single grain is then oriented using Laue backscattering along the [1 1 0] axis and cut perpendicular to this direction into slices ∼1 mm thick. Two samples were prepared having surfaces of typically a few mm2 . They were polished using diamond paste down to 1/4 micron before being introduced in the UHV chambers. Three UHV systems were used in this study, all of them with a base pressure better than 2 × 10−10 mbar. The system located in Orsay is equipped with a rear-view LEED (VG ThermoFisher) and an XPS spectrometer (CLAM2, VG ThermoFisher). Both systems located in Nancy are equipped with rear-view LEED (Specta-LEED, Omicron), XPS spectrometers (EA 125, Omicron) and STM facilities (VT-STM, Omicron). 2
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Figure 1. Evolution of the (a) Cu3s and Al2s, (b) O1s, (c) Cu2p and (d) CuLMM peaks for the clean γ -Al4 Cu9 (1 1 0) surface and after subsequent O2 exposure at RT. Take-off angle θ = 90◦ .
curve for the aluminium oxide Alox 2s and O1s components. The precise fitting parameters (peak position, peak width and peak asymmetry) depended slightly on the analyser but were kept constant when comparing data monitored on the same analyser. 5% mean errors were estimated from the reproducibility of the peak fitting.
In all systems, the γ -Al4 Cu9 surfaces were prepared by cycles of Ar+ sputtering at RT and annealing in the temperature range 920–950 K until the disappearance of all contaminant features in the XPS data (mainly C1s and O1s). The temperature of the samples was measured using infrared optical pyrometers which were previously calibrated by coupling to a thermocouple. In all chambers, such a calibration led to setting the emissivity of the pyrometers to 0.35. In situ oxidation was performed by dosing the clean surface with pure (5 N) oxygen at pressures ranging from 2 × 10−8 to 2 × 10−7 mbar. All oxygen doses in this work will be expressed in Langmuir units (1 L = 10−6 Torr s ∼ 1.33 × 10−6 mbar s). As the geometries of the three vacuum chambers used in this study were different (in particular the location of the gauges with respect to the specimen and the pumps) and the gauges were not calibrated, gauge readings were dependent on the equipment used. Therefore, for the sake of clarity, all oxygen doses have been normalized by setting the change in the slope observed in the oxidation kinetics plots (which have been monitored in all equipments) at 500 L, corresponding to the value obtained in the experiments performed in Nancy. Modifications in the surface structure were probed by LEED (long-range order) and STM (local atomic structure) whereas changes in the surface chemical composition were followed by XPS using a non monochromatized Mg Kα (1253.6 eV) or Al Kα (1486.6 eV) radiation source. As the energy separation between Al2p and Cu3p3/2 is less than 3 eV, there is a strong overlap between the Cu3p3/2 photopeak and the feature corresponding to oxidized aluminium. Therefore the oxidation state was followed by studying the changes in the shape of the Al2s (at 117.8 ± 0.2 eV) and Cu3s (at 122.9 ± 0.2 eV) peaks throughout this study. Al 2s, Cu 3s, O1s, Cu2p and Auger CuLMM peaks were monitored following each oxygen dosing. All XPS spectra were recorded in the constant pass energy mode (50 eV for survey data; 20 eV for narrow windows) and analysed using the CasaXPS software. Peak fitting was performed on data collected in narrow windows using Shirley background subtraction and a Doniach-Sunjic fit curve for the Al2s and Cu3s peaks and a 70% Gaussian–30% Lorenzian fit
3. Results 3.1. RT oxidation
Figure 1 shows the evolution of the Cu3s, Al2s, O1s, Cu2p photopeaks and the CuLMM Auger peak as a function of oxygen exposure up to 1000 L. The binding energies for the Al2s, Cu3s and Cu2p3/2 peaks on the clean sample are 117.8 eV, 122.9 eV and 933.0 ± 0.2 eV respectively, in fair agreement with previous measurements [32]. The apparent binding energy of 335.0 ± 0.2 eV observed for the Auger Cu L3 M4,5 M4,5 (1G) transition corresponds to a kinetic energy of 918.6 eV, similar to the reference value usually ascribed to pure copper. The only differences which can be observed in the data following oxidation are the appearance of the O1s peak and a filling of the valley between the Al2s and Cu3s peaks. This latter effect can be ascribed to the appearance of a new feature corresponding to oxidized aluminium on the high binding energy side of the Al2s peak. Peak fitting of the Al2s–Cu3s region for clean and oxidized surfaces is shown in figure 2. The fitting parameters (position, shape, full width at half maximum (FWHM)) corresponding to Al2s and Cu3s peaks were first determined using data recorded on the clean surface (figure 2(a)). The spectra monitored following oxidation were then fitted by keeping the Al2s and Cu3s peak parameters constant and introducing a third peak (Alox 2s) for oxidized aluminium (figure 2(b)). The best fits were achieved with the Alox 2s feature at 119.7 ± 0.2 eV and 1.7 eV and 2.8 eV FWHM values for Al2s and Alox 2s peaks respectively. The ratio between the intensities of the Alox 2s and Al2s peaks was then used to follow the oxidation kinetics at RT. 3
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Figure 2. Decomposition of the Cu3s and Al2s photopeaks obtained (a) for the clean γ -Al4 Cu9 (1 1 0) surface and (b) after 3200 L O2
exposure at RT. Take-off angle θ = 90◦ .
Moreover, the thickness of the oxide film corresponding to the onset of the slow stage is very similar for all experiments: 8.5 ± 0.6 Å. This value has been deduced from the intensity of the Al2s and Alox 2s peaks using the Strohmeier relationship [33]: Nm λm I0 d = λ0 sin θ ln +1 N 0 λ0 I m where λm (λ0 ) the attenuation length of the photoelectrons in the γ -Al4 Cu9 phase (oxide layer), θ the take-off angle (the angle between the surface and the analyser), Nm (N0 ) the volume density of Al atoms in the γ -Al4 Cu9 phase (aluminium oxide) and Im (I0 ) the intensity (peak area) of the Al2s (Alox 2s) peak. In the present work, the values of the attenuation lengths λ0 = 28.6 Å (33.1 Å), λm = 19.2 Å (22.1 Å) for the Mg Kα (Al Kα) anode have been deduced using the TPP-2 M formula proposed by Tanuma et al to calculate the inelastic mean free paths [34] without any attempt to correct for the contribution of elastic scattering. Although the model unit mesh for the γ -Al4 Cu9 phase contains 31% Al atoms, it has been shown in previous studies and confirmed in the present experiments that the surface region is enriched with aluminium and contains 44% Al atoms as deduced from Al2s and Cu3s peak intensities, thus leading to Nm = 0.035 at Å−3 . N0 = 0.043 at Å−3 was used in our calculations, as the volume density in amorphous aluminium oxide films is usually assumed to be close to that of the oxide γ phase [33]. Turning to the O1s region, peak fitting could be achieved using a single component (referred to hereafter as O1sLBE ) at 531.4 ±0.2 eV (FWHM = 1.6 eV) for exposures to oxygen up to 120 L (see figure 4). However for longer exposures, it was necessary to introduce a second component at a higher binding energy (O1sHBE at 532.4 ± 0.2 eV; FWHM = 2.6 eV).
Figure 3. Evolution of the Alox 2s/Al2s intensity as a function of the oxygen exposure at 2 × 10−8 mbar (black square) and at 2 × 10−7 mbar (red dot). Take-off angle θ = 90◦ .
Figure 3 shows the variation of this ratio as a function of oxygen exposure when dosing is performed either at 2 × 10−8 or at 2 × 10−7 mbar oxygen pressure. Both plots are similar within experimental errors. A fast initial oxidation regime is observed up to oxygen exposure ∼500 L. It is followed by a slower stage during which the oxide film grows at a much lower rate. The similarity of the two plots obtained for different oxygen pressures provides some evidence that oxidation kinetics is driven by aluminium diffusion to the surface rather than by the quantity of adsorbed oxygen. Experiments performed in the three experimental chambers used in this study lead to similar results, a fast increase followed by a slower stage, hence allowing normalization of the gauge readings by using the break corresponding to the change in slope, as mentioned in section 2. 4
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Figure 4. Decomposition of the O1s photopeak obtained after (a) 120 L O2 exposure at RT and (b) 1000 L O2 exposure at RT. Take-off angle θ = 90◦ .
Figure 5. Intensity of Alox 2s, O1sHBE , O1sLBE and O1sHBE core levels as a function of the oxygen exposure at RT. Take-off angle θ = 90◦ .
Figure 6. LEED patterns (E = 60 eV) obtained after RT oxidation (P (O2 ) = 2 × 10−7 mbar) of the γ -Al4 Cu9 (1 1 0) surface.
Figure 5 shows the variation in the intensities of both components as a function of oxygen dosing, together with the variation in the intensity of the Alox 2s peak. It can be seen that all these plots follow the same trend, exhibiting a rapid initial increase up to 500 L, followed by a slower variation for longer exposures. Although the O1sHBE is not observed at low coverages, it is difficult to totally set aside its presence due to a strong overlap of the O1s peak with the Cu LMM Auger feature, which introduces large uncertainties in the peak decomposition for small O1s intensities. Figure 6 shows the LEED patterns observed at 60 eV following sequential dosing of the γ -Al4 Cu9 (1 1 0) surface with oxygen at RT and 2 × 10−7 mbar. The LEED pattern observed for the clean surface is similar to the one previously
reported [32] and corresponds to a rectangular surface unit√cell having parameters a = 8.7 Å and b = 12.4 Å (b/a = 2). Initial dosing with oxygen does not induce any change in the LEED pattern which remains well contrasted with a low background. For exposure longer than 60 L, the background increases and the spots fade away until they disappear for exposures longer than 500 L, corresponding to the change in the oxidation rate in figure 3. The very long persistence of bright spots and low background suggests either a very low sticking coefficient for oxygen or preferential adsorption in well-defined surface sites. STM images were recorded following various exposures to oxygen at RT. Figure 7 shows that exposure to oxygen 5
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J. Phys.: Condens. Matter 26 (2014) 485009
Figure 7. STM images of the γ -Al4 Cu9 (1 1 0) surface following exposure to oxygen at RT (a) 6 L O2 ((23 × 23 nm2 ), Sample bias voltage Vb = +1 V, Tunnelling current It = 0.04 nA); (b) 100 L O2 ((40 × 40 nm2 ), Vb = +1 V, It = 0.08 nA).
leads to the appearance of white spots, the concentration of which is increasing with oxygen dosing. These spots can be ascribed to oxygen adsorption or oxide nucleation. However, the very low density of these spots together with their even distribution indicate that oxygen randomly adsorbs on the γ -Al4 Cu9 (1 1 0) surface with a very low initial sticking coefficient. Further evidence of the low value of the sticking coefficient is provided when one considers that the oxygen dose required to reach the slow oxidation regime in our case (500 L, as deduced from figure 3) is much larger than the 7–20 L reported by Hoffman and co-workers for α-Cu–Al solid solutions [12–17]. However, extreme caution must be taken when comparing oxygen doses reported by different authors, due to the different experimental geometries and the lack of proper gauge calibration. 3.2. Annealing at 925 K
Annealing was performed at 925 K for 30 and 60 min following the initial dosing of the γ -Al4 Cu9 (1 1 0) surface with 500 L oxygen at RT. The most obvious change which appears in the XPS data is a modification in the relative intensities of the Cu3s and Al2s peaks (see figure 8), which is mainly related to the increase of the Al2s peak and can thus be ascribed to segregation and enrichment of the surface by aluminium atoms. Annealing of the oxide film formed at RT for 60 min at 925 K leads to the appearance of a faint and poorly defined LEED pattern, shown in figure 9(a). It has not been possible to improve the quality of the LEED pattern either by longer annealing or by annealing at different temperatures. Nevertheless, three sub-sets of spots with rectangular unit meshes can be identified, superimposed to a high background and weak streaks forming polar circles. The three sets of spots are schematically drawn in figure 9(b). The largest rectangular unit mesh (white rectangle) corresponds to the clean substrate spots (white dots). The two
Figure 8. Evolution of the Cu3s and Al2s photopeaks after 500 L O2
exposure at RT (red/bottom curve) and subsequent annealing at 923 K during 30 min (green/middle curve) and 1 h (pink/top curve). Take-off angle θ = 90◦ .
remaining sets of spots (blue and red spots and rectangles) can be ascribed to two domains of the same structure, having a rectangular unit mesh (a = 10.5 Å and b = 18.2 Å) and approximately rotated by ±23.5◦ with respect to the substrate. The re-appearance of the spots corresponding to the clean surface suggests that, after annealing, the surface is not fully covered by the oxide layer. In addition, the poor contrast observed in the LEED pattern probably reflects the small size of the ordered domains. These conclusions are supported by STM experiments performed following annealing of the oxide film. It proved very difficult to obtain images of the oxide film 6
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Figure 9. (a) LEED pattern (E = 60 eV) of the crystallized aluminium oxide obtained on γ -Al4 Cu9 (1 1 0) after 500 L O2 exposure at RT and subsequent annealing at 925 K. (b) Simulated LEED pattern. The white unit cell corresponds to the substrate (a = 8.7 Å, b = 12.4 Å); √ the blue and red unit cells correspond to the two domains of the sixton oxide (a = 10.5 Å, b = 18.2 Å, b /a ≈ 3).
Figure 10. (a) 5 nm × 4.5 nm high resolution STM image (Fourier filtered) of the crystallized Al oxide obtained on γ -Al4 Cu9 (1 1 0). Vb = −1.65 V and It = 0.37 nA. The rectangle corresponds to the oxide unit cell. Al atoms (empty circles) of the interfacial Al16 O24 plane are arranged in heptagon–pentagon pairs (light blue solid line); (b) FFT analysis.
be identified, having unit cell dimensions (10.5 Å × 18.2 Å) identical to those observed by Napetschnig et al [24] after dosing a Cu-9%Al(1 1 1) sample with 100 L at 953 K with an oxygen partial pressure of 1.3 × 10−7 mbar. In addition, the bright spots appearing in our STM images are organized in regular heptagonal–pentagonal arrangements similar to those corresponding to the interface Al layer obtained by Napetschnig et al [24] on Cu-9%Al(1 1 1) and Kresse et al [25] following oxidation of NiAl(1 1 0). In view of these similarities, we believe that the crystallized oxide islands which are obtained in the present work have the same structure as the aluminium oxide film on NiAl(1 1 0), whose unit cell has been described in terms of a stacking sequence of four layers Ali −Oi −Als −Os composed of 16 interface Al, 24 interface O, 24 surface Al and 28 surface O atoms respectively or by an interfacial Al16 O24 and surface Al24 O28 layers, leading to an Al2 O2.6 global oxygen deficient stoichiometry [25, 35]. It should be noted that the same structure has recently been observed by Pr´evot et al for the oxide film formed by oxidation at RT under 10−6 mbar oxygen followed by
with atomic resolution. The best image of the oxide islands which could be obtained is shown on figure 10(a), together with the corresponding FFT analysis (figure 10(b)). Bright dots are observed, which are organized in regular heptagonalpentagonal arrangements similar to those observed by Kresse et al following oxidation of NiAl(1 1 0) [25]. It should be noted that the unit mesh deduced from this image is slightly distorted and larger (a = 10.4 ± 0.4 Å and b = 19.6 ± 0.8 Å) than that obtained from the LEED pattern, due to the thermal drift during the experiment. 4. Discussion 4.1. Oxide structure
The results presented in this paper show some similarities with the oxidation behaviour reported in the literature for α-Cu–Al solid solutions. In particular, in spite of the poor quality of the complex LEED patterns observed in our studies after annealing the oxide film formed at RT at 925 K, two domains of oxide rotated 23.5◦ with respect to the substrate can 7
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Figure 12. Evolution of the Al2s + Alox 2s/(Al2s + Alox 2s + Cu3s) intensity as a function of the take-off angle θ for the clean γ -Al4 Cu9 (1 1 0) surface (black square) and for the γ -Al4 Cu9 (1 1 0) surface after 3200 L O2 exposure at RT (red dot).
Figure 11. Atomic arrangement of the sixton interface layer on the Al4 Cu9 (1 1 0) surface. Yellow/dark grey dots: Cu; light blue/light grey dots: Al in the substrate surface layer; dark blue/black dots: Al in the sixton interfacial plane; large white circles: O in the sixton interfacial plane. Cu rows of the underlying substrate are indicated by vertical lines. Drawing produced using VESTA [42].
work, only crystallized domains having a small size are formed. Moreover, the weak streaky polar circles observed in figure 9(a) suggest that the small oxide domains can present all orientations with respect to the substrate, in addition to the favoured 23.5◦ misorientation which gives rise to the main new diffraction spots.
annealing at 1000 K of 2.5 aluminium monolayers deposited on top of Ni(1 1 1) [36, 37]. It was argued by these authors that the specific ratio between the length of the two sides of √ the alumina unit cell ( 3), which corresponds to a sixton rectangle (i.e. a rectangle inscribed in a regular hexagon), can be attributed to the hexagonal arrangement of the oxygen atoms in the interfacial layer Oi and the Al atoms in the surface layer Als . Moreover, Pr´evot et al pointed out that the ordered alumina films obtained in previous work performed by various authors on NiAl(1 1 0) [38], FeAl(1 1 0) [39] and aluminium ultrathin films deposited onto Cu(1 1 1) [40] or Ni(1 1 1) [41] could also be interpreted in terms of the same sixton structure within small distorsions of the unit cell, thus suggesting that this structure could be the equilibrium state of a two-layer-thick alumina film supported on a metal substrate. The misorientation of ∼23.5◦ of the sixton structure with respect to the substrate observed in our studies looks very similar to the ±24◦ rotation reported for alumina ultrathin films formed on NiAl(1 1 0) and FeAl(1 1 0) [25, 38, 39] and can be ascribed, in the same way, to a ‘row matching’ effect. The distance between two pure Cu rows in the [1 −1 0] direction for the model γ -Al4 Cu9 (1 1 0) surface can be estimated to 4.13 Å. As shown in figure 11, a 23.5◦ misorientation of the sixton structure allows most of the Al atoms of the heptagonpentagon pairs in the interface layer (Ali ) to be located on top of these pure copper rows thus avoiding unfavourable Al-Al neighbouring. It should be noted that, although the oxide structure is incommensurate with the substrate, the strong intensity observed for some spots in the LEED patterns suggests the vectorial relationship 3b∗ ∼ 4b∗ −a ∗ , i.e. commensurability is nearly achieved in one direction. Finally, it must be kept in mind that, in view of the poor quality of the LEED patterns observed in this
4.2. Al segregation
The results presented in section 3 show that oxidation at RT of the γ -Al4 Cu9 (1 1 0) surface leads to the exclusive formation of Al–O bonds. Oxygen bonding to copper is never observed, even at low oxygen coverages, although the clean surface model mainly contains copper (12 Cu versus 6 Al atoms), thus suggesting easy segregation of aluminium towards the surface. For all Cu–Al alloy compositions and surface orientations studied so far, a strong oxygen induced segregation at elevated temperature has been reported, due to the energetically favourable formation of Al–O bonds [12– 20, 43, 44]. The large thickness of the oxide film obtained in this work for oxidation at RT (8.5 Å ± 0.6 Å) suggests that oxygen-induced segregation occurs even at RT. Further evidence of this segregation effect is given in figure 12. The variation in aluminium percentage, defined as the ratio of the sum of the Al2s + Alox 2s peak intensities (IAl ) upon the sum of the Al2s + Alox 2s + Cu3s intensities (Itot ), is plotted as a function of the take off angle for the clean surface and the surface saturated with 3200 L oxygen at RT. Figure 12 clearly shows a rather weak Al surface enrichment for the clean surface (in good agreement with the 44 at% obtained by Gaudry et al [32]) whereas oxidation leads to a strong enhancement in the aluminium surface concentration, due to aluminium diffusion towards the surface. In order to estimate the Al segregation energy, the clean surface has been saturated with oxygen at RT then annealed at a fixed temperature for increasing periods of time. Three annealing temperatures have been used (350, 365 and 385 K). Between 8
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Figure 14. Intensity of Alox 2s, O1sLBE and O1sHBE core levels as a function of the surface preparation (O2 exposure at RT or annealing at 923 K). Take-off angle θ = 90◦ . After each annealing at 925 K the LEED pattern exhibits both the sixton oxide structure (S) and the substrate structure (C). After each O2 exposure the LEED pattern only exhibits the sixton structure. The solid lines are only guides for the eye.
Figure 13. Evolution of Al as a function of the square root of
annealing time for several temperatures following 3200 L O2 exposure at RT. Take-off angle θ = 30◦ .
each temperature, the crystal was cleaned by the sputterannealing process described in section 2 then dosed again with 3200 L. Figure 13 shows the variations in the relative enrichment in Al (Al, defined as the difference between the Al percentage at a given annealing time and the initial Al percentage) as a function of the square root of the annealing time t 1/2 . Due to the high inertia of our specimen holder, annealing periods must be long enough to minimize errors and only a few experimental points can be obtained. In spite of these restrictions, an Arrhenius analysis of the variation of the initial slope of these plots as a function of 1/T leads to a rough estimation of the activation energy for Al diffusion of 0.65 ± 0.12 eV at−1 (64 ± 12 kJ mol−1 ). This value is much smaller than those usually ascribed to the vacancy diffusion mechanism of Al in crystalline Al-based intermetallics (∼1.6 eV) [45] or in Cu (1.81 eV) [43] although it is close to the value observed by Gil-Gavatz et al on i-Al–Cu–Fe for temperatures below 770 K (0.6 eV) [46], which has been ascribed to a phason flip mechanism. However the phason flip mechanism, being intrinsic to a quasiperiodic structure, cannot be invoked in our case. A tentative explanation for the low value observed in this study can be given when considering that the activation energy of diffusion is composed of a vacancy formation and a vacancy migration contribution to enthalpy, which have close values (0.68 and 0.62 respectively in pure Al) [46]. As the γ -Al4 Cu9 phase has a structure based on a 3 × 3 × 3 CsCl superstructure containing two vacancies [31, 32] only the vacancy migration enthalpy must be considered in our case, thus leading to a very low value of the activation energy for Al segregation.
the O1s line shape during oxidation of Al or Al-containing alloys have already been reported in the literature. In particular, two components in the O1s have been observed at 530.1 and 532.0 eV during oxidation of α-Cu–Al (1 0 0) surfaces and ascribed to Cu–O and Al–O bonds respectively [13, 16, 17]. On pure aluminium, components in the range of 530.6–531.3 eV and 532.2–532.7 eV were associated with chemisorbed and oxidic oxygen respectively [47–51] whereas the main 532.3–533.4 eV and the weak 534.3–535.1 eV features observed by Jeurgens et al are associated to oxidic species within the oxide film and surface oxide species respectively [52]. The presence of two O1s components at 531.2 and 532.8 eV has also been reported in studies performed on the ultrathin aluminium oxide film formed on Ni3 Al(1 1 1) and ascribed to oxygen atoms at the metal-oxide interface and in the surface layer respectively [53]. However, a high resolution core-level spectroscopy study performed by Martin et al on the ultrathin aluminium oxide film formed on NiAl(1 1 0) showed that the high binding energy component should be assigned to only part of the surface oxygen atoms having an aluminium atom directly located underneath and not residing close to another oxygen atom [54]. In the present study, the formation of Cu–O bonds can be disregarded in view of the lack of any modification occurring in either the Cu2p photopeak or the CuLMM Auger peak. Moreover, the similar general trends shown in figure 5 in the Alox 2s, O1sLBE and O1sHBE intensities suggest that both O1s features are related to Al–O bonds. To gain further insight into the origin of these two components, cycles of exposure to oxygen at RT followed by annealing for 1 h at 925 K were performed. Between each step, XPS data and LEED patterns were recorded and the results are shown in figure 14. Ox1 refers to initial 1000 L oxygen dosing at RT. At this stage, only a strong background with no diffraction spots is observed in the LEED pattern. As described in section 3.2,
4.3. Oxygen peak
It has been shown in section 3.1 that the O1s peak can be fitted with a single component (O1sLBE at 531.4 eV) for oxygen exposure up to 120 L but a second component (O1sHBE at 532.4 eV) is required for higher dosing. Changes in 9
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subsequent annealing for 1 h at 925 K leads to the appearance of the spots characteristic of the sixton structure (labelled S in figure 14) and the reappearance of the clean substrate diffraction spots (C). Changes in the photopeaks are also observed, with a decrease in Alox 2s and O1sLBE intensities and an increase of O1sHBE . Additional exposure to 6000 L at RT (Ox2) leads to the disappearance of the clean substrate spots C whereas the sixton spots S are still visible. The intensities of all photopeaks are increasing. Further annealing leads to the reappearance of the clean substrate diffraction spots (C) in addition to the S spots and the decrease in the intensities of all photopeaks. The same tendency is observed for the third cycle: the disappearance of only the clean substrate spots and an increase in the intensities of all photopeaks following oxygen dosing at RT (Ox3), the reappearance of the clean substrate diffraction spots (C) in addition to the S spots and a decrease in the intensities of all photopeaks following subsequent annealing. However, it should be noted that the variations observed in the intensity of the O1sLBE peak are much larger than for the O1sHBE feature. The observed oscillating behaviour of the photopeaks, together with the reappearance/disappearance of the C spots suggests that the O1sLBE peak corresponds to loosely bond O species which are adsorbed on the clean surface during exposure to oxygen at RT and partly desorb at high temperatures, thus liberating patches of clean surface and providing new adsorption sites for the next oxygen dosing. In view of the weak variations observed in the O1sHBE peak during the second and third cycles, together with the persistence of the S spots, we believe that this feature can be ascribed to oxygen atoms in the crystallized oxide islands. The presence of the O1sHBE feature even following the first dosing at RT means that very small sixton domains are already present at this stage, whose size is increasing following the first annealing, leading to a large increase in the corresponding signal. The small decrease observed in the intensity of the O1sHBE peak following annealing in cycles 2 and 3 is more difficult to explain. It can be ascribed either to desorption of some O species within the ordered oxide islands or desorption of part of these islands. It should be noted that the binding energy of the O1sHBE peak observed in our studies (532.4 eV) is close to the value ascribed by Martin et al to oxygen atoms in the ultrathin oxide film formed on NiAl(1 1 0) [54]. The additional component shifted 1.23 eV towards higher binding energies observed by these authors and assigned to specific oxygen atoms in the surface layer was not observed in our case. However, the possible presence of a peak at higher energy is difficult to precisely assess in our case, due to the strong overlap with the CuLMM peak (see figure 4). In any case, neither the ratio between the O1sHBE and the O1sLBE observed in our studies (>1, to be compared with 0.19 in [54]) nor the up and down behaviour during the dosing/annealing cycles can be explained by ascribing both our O1sHBE and O1sLBE signals to oxygen species belonging to the sixton structure. However, a high resolution core-level spectroscopy study would be necessary for a more precise assessment of the origin of both the O1s signals observed in our study.
5. Conclusion
RT oxidation of the (1 1 0) surface of the γ -Al4 Cu9 complex metallic alloy was performed in the range of 2 × 10−8 to 2 × 10−7 mbar oxygen pressure. A fast oxidation step is observed up to 500 L, followed by a slow regime during which the oxide film grows at a much lower rate. Cu–O bonds are never observed, in spite of the high Cu content of the clean (1 1 0) surface, due to a strong oxygen induced aluminium segregation. The activation energy for aluminium diffusion was estimated at 0.65 ± 0.12 eV at−1 (64 ± 12 kJ mol−1 ), a rather low value which was explained by the presence of two vacancies in the γ -Al4 Cu9 structure, thus reducing the activation energy to the sole contribution of vacancy migration enthalpy. When the oxide film formed at RT is annealed at 925 K, small crystallized domains with a sixton structure appear. The similarity of this sixton structure with structures reported in the literature following the oxidation of Cu-9%Al(1 1 1), NiAl (1 1 0) and FeAl(1 1 0) surfaces as well as ultrathin Al films deposited onto Cu(1 1 1) or Ni(1 1 1) surfaces, gives further evidence of the free standing nature of this ultrathin aluminium oxide layer. However, it is the first time to our knowledge that this peculiar structure has been observed on a surface having a complex structure with a large unit mesh. In addition, it seems that it is easier to form the ordered structure on such a surface, as it is directly obtained by annealing the oxide film formed at RT whereas, on all the other surfaces studied so far, its formation required exposure to oxygen at high temperature. Work is in progress to check if the sixton structure can be formed on the surface of other complex metallic alloys. Finally two contributions were observed in the O1s peaks at 531.4 and 532.4 eV, which have been ascribed to loosely bond oxygen species and oxygen belonging to the sixton structure respectively. Acknowledgments
The Agence Nationale pour la Recherche is acknowledged for its financial support (ANR-08-BLAN-0041). This work was performed within the frame of the European Centre for the Development of new Metallic Alloys and Compounds (CMAC, www.eucmac.eu). References [1] Dubois J-M 2005 Useful Quasicrystals (Singapore: World Scientific) [2] Thiel P A 2008 Annu. Rev. Phys. Chem. 59 129 [3] Barth´es-Labrousse M-G and Dubois J-M 2008 Phil. Mag. 88 2217–25 [4] Dubois J-M, Kang S S and Massiani Y 1993 J. Non-Cryst. Solids 153–154 443–5 [5] Dubois J-M, Proner A, Bucaille B, Cathonnet P, Dong C, Richard V, Pianelli A, Massiani Y, Ait-Yaazza S and Belin-Ferr´e E 1994 Ann. Chim.-Sci. Mater. 19 3–25 [6] Fourn´ee V, Barth´es-Labrousse M-G and Dubois J-M 2008 Solid State Phenomena (Switzerland: Trans Tec Publications) pp 407–50 10
M Warde´ et al
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[31] Dong C 1996 Phil. Mag. A 73 1519–28 [32] Gaudry E, Shukla A K, Duguet T, Ledieu J, deWeerd M C, Dubois J-M and Fourn´ee V 2010 Phys. Rev. B 82 085411 [33] Strohmeier B R 1990 Surf. Interface Anal. 15 51–6 [34] Tanuma S, Powell C J and Penn D R 1994 Surf. Interface Anal. 21 165–76 [35] Nishimura T, Hoshino Y, Okazawa T and Kido Y 2008 Phys. Rev. B 77 073405 [36] Pr´evot G, Naitabdi A, Bernard R and Borensztein Y 2010 Phys. Rev. B 81 085405 [37] Pr´evot G, Le Moal S, Bernard R, Croset B, Lazzari R and Schmaus D 2012 Phys. Rev. B 85 205450 [38] Stierle A, Renner F, Streitel R, Dosch H, Drube W and Cowie B C 2004 Science 303 1652 [39] Graupner H, Hammer L, Heinz K and Zehner D M 1997 Surf. Sci. 380 335–51 [40] Jeliazova Y and Franchy R 2002 Appl. Surf. Sci. 187 51–9 [41] Le Pevedic S, Schmaus D and Cohen C 2008 Surf. Sci. 602 67–76 [42] Momma K and Izumi F 2011 VESTA 3 for three-dimensional visualization of crystal, volumetric and morphology data J. Appl. Crystallogr. 44 1272–6 [43] Yu Y, Sagisaka K and Fujita D 2009 Surf. Sci. 603 723–6 [44] Lanzani G, Kangas T and Laasonen K 2009 J. Alloys Compounds 482 33–42 [45] Joulaud J L, Berbardini J, Bergman C, Dubois J-M, Calvayrac Y and Gratias D 1997 Phil. Mag. A 75 1287 [46] Gil-Gavatz M, Rouxel D, Pigeat P, Weber B and Dubois J-M 1999 Mater. Res. Soc. Symp. Proc. 553 75–80 [47] Bradshaw A M, Domcke W and Cederbaum L S 1977 Phys. Rev. B 16 1480–8 [48] Bradshaw A M, Hofmann P and Wyrobisch W 1977 Surf. Sci. 68 269–76 [49] Zhukov V, Popova I, Fomenko V and Yates J T J 1999 Surf. Sci. 441 240–50 [50] Popova I, Zhukov V and Yates J T J 1999 Appl. Phys. Lett. 75 3108–10 [51] Popova I, Zhukov V and Yates J T J 2000 J. Appl. Phys. 87 8143–7 [52] Jeurgens L P H, Sloof W G, Tichelaar F D and Mittemeijer E J 2002 Surf. Sci. 506 313–32 [53] Vesselli E, Baraldi A, Lizzit S and Comelli G 2010 Phys. Rev. Lett. 105 046102 [54] Martin N M et al 2011 Phys. Rev. B 83 125417
[7] Barth´es-Labrousse M-G, Beni A and Schmutz P 2011 Complex Metallic Alloys ed J-M Dubois and E Belin-Ferr´e (Weinheim: Wiley-VCH) pp 243–72 [8] Longchamp J-N, Burkardt S, Erbudak M and Weisskopf Y 2007 Phys. Rev. B 76 094203 [9] Burkardt S, Erbudak M and Maeder R 2009 Surf. Sci. 603 867–72 [10] Erbudak M, Mungan M and Burkardt S 2009 Appl. Surf. Sci. 256 1284–8 [11] Franchy R 2000 Surf. Sci. Rep. 38 195–294 [12] Hoffman A, Maniv T and Folman M 1988 Surf. Sci. 193 57–80 [13] Kravchuk T and Hoffman A 2007 Surf. Sci. 601 87–94 [14] Hoffman A and Folman M 1991 Vacuum 42 1133–6 [15] Shen W, Nyberg G L and Hoffman A 1995 Surf. Sci. 334 209–14 [16] Kravchuk T, Akhvlediani R and Hoffman A 2005 Nucl. Instrum. Methods Phys. Res. B 230 413–8 [17] Kravchuk T and Hoffman A 2006 Surf. Sci. 600 1252–9 [18] Yoshitake M, Bera S, Yamauchi Y and Song W 2003 J. Vac. Sci. Technol. A 21 1290–3 [19] Yoshitake M, Bera S and Yamauchi Y 2003 Surf. Interface Anal. 35 824–8 [20] Song W, Yoshitake M, Bera S and Yamauchi Y 2003 Japan. J. Appl. Phys. 42 4716–20 [21] Yoshitake M, Song W, Libra J, Masek K, Sutura F, Matolin V and Prince K C 2008 J. Appl. Phys. 103 033707 [22] Yoshitake M, Nemsak S, Skala T, Tsud N, Kim T, Matolin V and Prince K C 2010 Surf. Sci. 604 2150–6 [23] Yoshitake M, Song W, Libra J, Masek K, Sutura F, Matolin V and Prince K C 2010 Appl. Surf. Sci. 256 3051–7 [24] Napetschnig E, Schmid M and Varga P 2008 Surf. Sci. 602 1750–6 [25] Kresse G, Schmid M, Napetschnig E, Shishkin M, Koehler L and Varga P 2005 Science 308 1440–2 [26] Gil-Gavatz M, Rouxel D, Pigeat P, Weber B and Dubois J-M 2000 Phil. Mag. A 80 2083–97 [27] Rouxel D, Gil-Gavatz M, Pigeat P and Weber B 2005 J. Non-Cryst. Solids 351 802–9 [28] Rouxel D and Pigeat P 2006 Prog. Surf. Sci. 81 488–514 [29] Arnberg L and Westman S 1978 Acta Crystallogr. A 34 399–404 [30] Villars P and Calvert L D 1998 Pearson’s Handbook of Crystallographic Data for Intermetallic Phases (Materials Park, OH: ASM International)
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Growth and structure of ultrathin alumina films on the (110) surface of -Al4Cu9 complex metallic alloy
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 J. Phys.: Condens. Matter 26 485009 (http://iopscience.iop.org/0953-8984/26/48/485009) View the table of contents for this issue, or go to the journal homepage for more Download details: IP Address: 132.239.1.231 This content was downloaded on 11/05/2017 at 09:59 Please note that terms and conditions apply.
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Journal of Physics: Condensed Matter J. Phys.: Condens. Matter 26 (2014) 485009 (11pp)
doi:10.1088/0953-8984/26/48/485009
Growth and structure of ultrathin alumina films on the (1 1 0) surface of γ -Al4Cu9 complex metallic alloy M Warde´ 1,2 , J Ledieu3 , L N Serkovic Loli3 , M Herinx1,2 , M-C de Weerd3 , 1,2 ´ ´ 3 , S Le Moal1,2 and M-G Barthes-Labrousse V Fournee 1
Institut de Chimie Mol´eculaire et Mat´eriaux d’Orsay, UMR 8182 CNRS, Universit´e Paris-Sud, Bˆatiment 410, 91405 Orsay Cedex, France 2 CNRS, Orsay, F-91405, France 3 Institut Jean Lamour, UMR 7198 CNRS, Universit´e de Lorraine, Parc de Saurupt, 54011 Nancy Cedex, France E-mail: [email protected] Received 21 July 2014, revised 29 September 2014 Accepted for publication 10 October 2014 Published 6 November 2014 Abstract
The first stages of oxidation of the (1 1 0) surface of a γ -Al4 Cu9 complex metallic alloy were investigated by combining x-ray photoemission spectroscopy, low energy electron diffraction and scanning tunnel microscopy studies. Oxidation at room temperature in the 2 × 10−8 to 2 × 10−7 mbar oxygen pressure range occurs in two steps: a fast regime is followed by a much slower one, leading to the formation of a thin aluminium oxide film showing no long range order. Cu–O bonds are never observed, due to fast oxygen induced aluminium segregation. The low value of the estimated activation energy for aluminium diffusion (0.65 ± 0.12 eV at−1 ) was ascribed to the presence of two vacancies in the γ -Al4 Cu9 structure. Annealing at 925 K the oxide film formed at room temperature leads to the formation of small crystallized domains with a sixton structure similar to structures reported in the literature following the oxidation of Cu-9%Al(1 1 1), NiAl (1 1 0) and FeAl(1 1 0) surfaces as well as ultrathin Al films deposited onto Cu(1 1 1) or Ni(1 1 1) surfaces. Two contributions were observed in the O1s peaks, which have been ascribed to loosely bound oxygen species and oxygen belonging to the sixton structure respectively. Keywords: ultrathin alumina film, (1 1 0) γ -Al4 Cu9 , complex metallic alloy (Some figures may appear in colour only in the online journal)
In particular, a number of studies related to interactions with oxygen were driven by the excellent oxidation and corrosion resistance initially reported for Al–Cu–Fe and Al–Cu–Cr–Fe quasicrystals [4, 5]. Following this pioneer work, most investigations to date have been devoted to the first stages of oxidation of Alrich complex metallic alloys. Similar to more conventional intermetallics, oxidation of these materials in a vacuum environment leads to Al segregation and the formation of a passivating aluminium oxide layer, a few nm thick (for a review, see [6, 7] and the references therein). In most cases, the oxide layer has been found to be amorphous. However, following high temperature oxidation of the pentagonal
1. Introduction
Complex metallic alloys (CMAs) are characterized by unit cells containing from a few tens to several thousand atoms arranged into highly symmetric clusters. Quasicrystals represent the ultimate complexity, having an infinite number of atoms within the cell. CMAs have attracted a lot of interest during the last two decades, due to their potentially useful chemical and mechanical properties [1–3]. The interest in studying CMAs’ surface reactivity comes from their specific electronic structure, which is related to the existence of highly symmetric clusters and can affect interactions of the surface atoms with surrounding atoms and molecules. 0953-8984/14/485009+11$33.00
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J. Phys.: Condens. Matter 26 (2014) 485009
surface of i-Al–Pd–Mn and the 10-fold-symmetry surface of the decagonal Al–Co–Ni, the formation of nano-domains of hexagonal aluminium oxide thin layers aligned along a two-fold or five-fold symmetry direction of the substrate respectively has been reported [8–10]. The limit in the lateral size of the domains was ascribed to the lack of translational order in the structure of the quasicrystalline substrates used in these studies. To our knowledge, such ordered films have never been observed on the surface of periodic complex metallic alloys, although well-ordered oxide films can be easily obtained on a number of conventional binary alloy surfaces (see for example [11]). Due to the importance of Cu-rich Al–Cu alloys in catalysis and microelectronics, their oxidation behaviour has been extensively studied. In particular, oxygen adsorption in ultrahigh vacuum (UHV) conditions on the (1 0 0) surface of α-Cu–Al solid solutions containing increasing quantities of aluminium (0.15, 1.5, 5, 12 and 17 at% aluminium) has been followed by Hoffman and co-authors using Auger electron spectroscopy (AES), photoemission spectroscopies (XPS and UPS) and low energy ion spectroscopy (LEIS) [12–17]. Two stages are observed in the oxidation kinetics when oxygen exposure is performed at room temperature (RT). Initially, chemisorbed oxygen quickly reacts with surface aluminium sites, leading to the formation of a thin aluminium oxide with AlO stoichiometry. This fast process is observed up to 7–20 L of oxygen exposure, depending on the aluminium concentration in the alloy. It is followed by a much slower stage during which both Al–O and Cu–O bonds are formed, the formation of new Al–O bonds being ascribed to the diffusion of oxygen in the bulk and Al segregation to the surface. When the exposure to oxygen is for long enough, the Al2 O3 stoichiometry is reached. Annealing the oxidized alloy surfaces in the 520–725 K temperature range results in the disappearance of the Cu–O bonds, an increase in the aluminium segregation kinetics and the formation of a stable pure alumina layer. A number of studies have also been devoted to the structure of the aluminium oxide films formed on (1 1 1) surfaces of Cu–Al alloys containing more than 9 at% Al. Saturation with oxygen at 670 K leads to an amorphous oxide layer which crystallizes following annealing at 920 K, √ √ as indicated by the appearance of a ( 3 × 3) low energy electron diffraction (LEED) pattern. When oxidation is performed at higher temperatures (up to 995 K), a √ √ (7 3 × 7 3)R30◦ pattern appears which has been ascribed to the formation of a γ -Al2 O3 aluminium oxide film having the orientation relationship Al2 O3 (1 1 1)//Cu9%atAl(1 1 1) [18–23]. However, Napetschnig et al have shown that moderate oxidation at a high temperature (100 L at 953 K) leads to the formation of a complex LEED diagram which can be ascribed to the formation of a thinner aluminium oxide film (0.5 nm thick) similar to that observed on NiAl(1 1 0) and consisting of two layers having Al16 O24 and Al24 O28 stoichiometries respectively (overall stoichiometry Al2 O2.6 ) [24, 25]. To the best of our knowledge, the first stages of the oxidation of Al–Cu intermetallics have not been reported so
far. However, the behaviour of the quasicrystalline icosahedral i-Al62 Cu25.5 Fe12.5 and the crystalline tetragonal ω-Al7 Cu2 Fe phases has been studied by Rouxel et al who have shown that a thin aluminium oxide layer is formed under UHV conditions during the first stages of oxidation, whereas at higher oxygen exposure and temperature, the structural complexity of the icosahedral phase may delay or even prevent the nucleation and epitaxial growth of γ -Al2 O3 crystallites compared to the tetragonal phase [26–28]. In this paper, we will present results obtained for the first stages of oxidation of the (1 1 0) surface of γ -Al4 Cu9 at RT, and the effect of subsequent annealing at 925 K on the oxide film structure. The γ -Al4 Cu9 phase presents an intermediate complexity, having a simple cubic structure with a lattice parameter equal to 0.871 nm and 52 atoms per unit mesh (16 Al + 36 Cu) [29, 30]. With 0.127 valence electrons per unit volume, its electronic structure shows some similarity with the icosahedral Al–Cu–Fe phase (0.124 e- Å−3 ) [31]. It can be described by a bcc packing of 26-atom clusters or by the stacking of flat and puckered {1 1 0} planes along the [1 1 0] direction. The structure of the clean (1 1 0) surface of γ -Al4 Cu9 has been extensively investigated by Gaudry et al by combining LEED, XPS, UPS and scanning tunnel microscopy (STM) experiments with density functional theory (DFT) calculations [32]. These authors have shown that cleaning cycles of sputtering and annealing at 933 K lead to the selection of puckered planes as preferential surface termination. These planes correspond to a simple bulk truncation and contain 6 Al and 12 Cu atoms arranged in a rectangular unit mesh 8.76 Å × 12.4 Å. It will be shown in this paper that, due to Al segregation, an aluminium oxide layer showing no long range order is obtained when RT oxidation is performed on the (1 1 0) surface of γ -Al4 Cu9 . Subsequent annealing of this oxide layer at 925 K leads to the appearance of small crystallized oxide islands having a sixton structure similar to the structure obtained for aluminium oxide films formed on NiAl(1 1 0). 2. Experimental details
The preparation of the γ -Al4 Cu9 crystal has previously been described in detail [32]. Induction melting of the pure elements (Al-99.9% and Cu-99.99%) under an argon atmosphere is followed by homogenization by annealing at 1323 K for 4 h and slow cooling down to RT. A single grain is then oriented using Laue backscattering along the [1 1 0] axis and cut perpendicular to this direction into slices ∼1 mm thick. Two samples were prepared having surfaces of typically a few mm2 . They were polished using diamond paste down to 1/4 micron before being introduced in the UHV chambers. Three UHV systems were used in this study, all of them with a base pressure better than 2 × 10−10 mbar. The system located in Orsay is equipped with a rear-view LEED (VG ThermoFisher) and an XPS spectrometer (CLAM2, VG ThermoFisher). Both systems located in Nancy are equipped with rear-view LEED (Specta-LEED, Omicron), XPS spectrometers (EA 125, Omicron) and STM facilities (VT-STM, Omicron). 2
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Figure 1. Evolution of the (a) Cu3s and Al2s, (b) O1s, (c) Cu2p and (d) CuLMM peaks for the clean γ -Al4 Cu9 (1 1 0) surface and after subsequent O2 exposure at RT. Take-off angle θ = 90◦ .
curve for the aluminium oxide Alox 2s and O1s components. The precise fitting parameters (peak position, peak width and peak asymmetry) depended slightly on the analyser but were kept constant when comparing data monitored on the same analyser. 5% mean errors were estimated from the reproducibility of the peak fitting.
In all systems, the γ -Al4 Cu9 surfaces were prepared by cycles of Ar+ sputtering at RT and annealing in the temperature range 920–950 K until the disappearance of all contaminant features in the XPS data (mainly C1s and O1s). The temperature of the samples was measured using infrared optical pyrometers which were previously calibrated by coupling to a thermocouple. In all chambers, such a calibration led to setting the emissivity of the pyrometers to 0.35. In situ oxidation was performed by dosing the clean surface with pure (5 N) oxygen at pressures ranging from 2 × 10−8 to 2 × 10−7 mbar. All oxygen doses in this work will be expressed in Langmuir units (1 L = 10−6 Torr s ∼ 1.33 × 10−6 mbar s). As the geometries of the three vacuum chambers used in this study were different (in particular the location of the gauges with respect to the specimen and the pumps) and the gauges were not calibrated, gauge readings were dependent on the equipment used. Therefore, for the sake of clarity, all oxygen doses have been normalized by setting the change in the slope observed in the oxidation kinetics plots (which have been monitored in all equipments) at 500 L, corresponding to the value obtained in the experiments performed in Nancy. Modifications in the surface structure were probed by LEED (long-range order) and STM (local atomic structure) whereas changes in the surface chemical composition were followed by XPS using a non monochromatized Mg Kα (1253.6 eV) or Al Kα (1486.6 eV) radiation source. As the energy separation between Al2p and Cu3p3/2 is less than 3 eV, there is a strong overlap between the Cu3p3/2 photopeak and the feature corresponding to oxidized aluminium. Therefore the oxidation state was followed by studying the changes in the shape of the Al2s (at 117.8 ± 0.2 eV) and Cu3s (at 122.9 ± 0.2 eV) peaks throughout this study. Al 2s, Cu 3s, O1s, Cu2p and Auger CuLMM peaks were monitored following each oxygen dosing. All XPS spectra were recorded in the constant pass energy mode (50 eV for survey data; 20 eV for narrow windows) and analysed using the CasaXPS software. Peak fitting was performed on data collected in narrow windows using Shirley background subtraction and a Doniach-Sunjic fit curve for the Al2s and Cu3s peaks and a 70% Gaussian–30% Lorenzian fit
3. Results 3.1. RT oxidation
Figure 1 shows the evolution of the Cu3s, Al2s, O1s, Cu2p photopeaks and the CuLMM Auger peak as a function of oxygen exposure up to 1000 L. The binding energies for the Al2s, Cu3s and Cu2p3/2 peaks on the clean sample are 117.8 eV, 122.9 eV and 933.0 ± 0.2 eV respectively, in fair agreement with previous measurements [32]. The apparent binding energy of 335.0 ± 0.2 eV observed for the Auger Cu L3 M4,5 M4,5 (1G) transition corresponds to a kinetic energy of 918.6 eV, similar to the reference value usually ascribed to pure copper. The only differences which can be observed in the data following oxidation are the appearance of the O1s peak and a filling of the valley between the Al2s and Cu3s peaks. This latter effect can be ascribed to the appearance of a new feature corresponding to oxidized aluminium on the high binding energy side of the Al2s peak. Peak fitting of the Al2s–Cu3s region for clean and oxidized surfaces is shown in figure 2. The fitting parameters (position, shape, full width at half maximum (FWHM)) corresponding to Al2s and Cu3s peaks were first determined using data recorded on the clean surface (figure 2(a)). The spectra monitored following oxidation were then fitted by keeping the Al2s and Cu3s peak parameters constant and introducing a third peak (Alox 2s) for oxidized aluminium (figure 2(b)). The best fits were achieved with the Alox 2s feature at 119.7 ± 0.2 eV and 1.7 eV and 2.8 eV FWHM values for Al2s and Alox 2s peaks respectively. The ratio between the intensities of the Alox 2s and Al2s peaks was then used to follow the oxidation kinetics at RT. 3
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Figure 2. Decomposition of the Cu3s and Al2s photopeaks obtained (a) for the clean γ -Al4 Cu9 (1 1 0) surface and (b) after 3200 L O2
exposure at RT. Take-off angle θ = 90◦ .
Moreover, the thickness of the oxide film corresponding to the onset of the slow stage is very similar for all experiments: 8.5 ± 0.6 Å. This value has been deduced from the intensity of the Al2s and Alox 2s peaks using the Strohmeier relationship [33]: Nm λm I0 d = λ0 sin θ ln +1 N 0 λ0 I m where λm (λ0 ) the attenuation length of the photoelectrons in the γ -Al4 Cu9 phase (oxide layer), θ the take-off angle (the angle between the surface and the analyser), Nm (N0 ) the volume density of Al atoms in the γ -Al4 Cu9 phase (aluminium oxide) and Im (I0 ) the intensity (peak area) of the Al2s (Alox 2s) peak. In the present work, the values of the attenuation lengths λ0 = 28.6 Å (33.1 Å), λm = 19.2 Å (22.1 Å) for the Mg Kα (Al Kα) anode have been deduced using the TPP-2 M formula proposed by Tanuma et al to calculate the inelastic mean free paths [34] without any attempt to correct for the contribution of elastic scattering. Although the model unit mesh for the γ -Al4 Cu9 phase contains 31% Al atoms, it has been shown in previous studies and confirmed in the present experiments that the surface region is enriched with aluminium and contains 44% Al atoms as deduced from Al2s and Cu3s peak intensities, thus leading to Nm = 0.035 at Å−3 . N0 = 0.043 at Å−3 was used in our calculations, as the volume density in amorphous aluminium oxide films is usually assumed to be close to that of the oxide γ phase [33]. Turning to the O1s region, peak fitting could be achieved using a single component (referred to hereafter as O1sLBE ) at 531.4 ±0.2 eV (FWHM = 1.6 eV) for exposures to oxygen up to 120 L (see figure 4). However for longer exposures, it was necessary to introduce a second component at a higher binding energy (O1sHBE at 532.4 ± 0.2 eV; FWHM = 2.6 eV).
Figure 3. Evolution of the Alox 2s/Al2s intensity as a function of the oxygen exposure at 2 × 10−8 mbar (black square) and at 2 × 10−7 mbar (red dot). Take-off angle θ = 90◦ .
Figure 3 shows the variation of this ratio as a function of oxygen exposure when dosing is performed either at 2 × 10−8 or at 2 × 10−7 mbar oxygen pressure. Both plots are similar within experimental errors. A fast initial oxidation regime is observed up to oxygen exposure ∼500 L. It is followed by a slower stage during which the oxide film grows at a much lower rate. The similarity of the two plots obtained for different oxygen pressures provides some evidence that oxidation kinetics is driven by aluminium diffusion to the surface rather than by the quantity of adsorbed oxygen. Experiments performed in the three experimental chambers used in this study lead to similar results, a fast increase followed by a slower stage, hence allowing normalization of the gauge readings by using the break corresponding to the change in slope, as mentioned in section 2. 4
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Figure 4. Decomposition of the O1s photopeak obtained after (a) 120 L O2 exposure at RT and (b) 1000 L O2 exposure at RT. Take-off angle θ = 90◦ .
Figure 5. Intensity of Alox 2s, O1sHBE , O1sLBE and O1sHBE core levels as a function of the oxygen exposure at RT. Take-off angle θ = 90◦ .
Figure 6. LEED patterns (E = 60 eV) obtained after RT oxidation (P (O2 ) = 2 × 10−7 mbar) of the γ -Al4 Cu9 (1 1 0) surface.
Figure 5 shows the variation in the intensities of both components as a function of oxygen dosing, together with the variation in the intensity of the Alox 2s peak. It can be seen that all these plots follow the same trend, exhibiting a rapid initial increase up to 500 L, followed by a slower variation for longer exposures. Although the O1sHBE is not observed at low coverages, it is difficult to totally set aside its presence due to a strong overlap of the O1s peak with the Cu LMM Auger feature, which introduces large uncertainties in the peak decomposition for small O1s intensities. Figure 6 shows the LEED patterns observed at 60 eV following sequential dosing of the γ -Al4 Cu9 (1 1 0) surface with oxygen at RT and 2 × 10−7 mbar. The LEED pattern observed for the clean surface is similar to the one previously
reported [32] and corresponds to a rectangular surface unit√cell having parameters a = 8.7 Å and b = 12.4 Å (b/a = 2). Initial dosing with oxygen does not induce any change in the LEED pattern which remains well contrasted with a low background. For exposure longer than 60 L, the background increases and the spots fade away until they disappear for exposures longer than 500 L, corresponding to the change in the oxidation rate in figure 3. The very long persistence of bright spots and low background suggests either a very low sticking coefficient for oxygen or preferential adsorption in well-defined surface sites. STM images were recorded following various exposures to oxygen at RT. Figure 7 shows that exposure to oxygen 5
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Figure 7. STM images of the γ -Al4 Cu9 (1 1 0) surface following exposure to oxygen at RT (a) 6 L O2 ((23 × 23 nm2 ), Sample bias voltage Vb = +1 V, Tunnelling current It = 0.04 nA); (b) 100 L O2 ((40 × 40 nm2 ), Vb = +1 V, It = 0.08 nA).
leads to the appearance of white spots, the concentration of which is increasing with oxygen dosing. These spots can be ascribed to oxygen adsorption or oxide nucleation. However, the very low density of these spots together with their even distribution indicate that oxygen randomly adsorbs on the γ -Al4 Cu9 (1 1 0) surface with a very low initial sticking coefficient. Further evidence of the low value of the sticking coefficient is provided when one considers that the oxygen dose required to reach the slow oxidation regime in our case (500 L, as deduced from figure 3) is much larger than the 7–20 L reported by Hoffman and co-workers for α-Cu–Al solid solutions [12–17]. However, extreme caution must be taken when comparing oxygen doses reported by different authors, due to the different experimental geometries and the lack of proper gauge calibration. 3.2. Annealing at 925 K
Annealing was performed at 925 K for 30 and 60 min following the initial dosing of the γ -Al4 Cu9 (1 1 0) surface with 500 L oxygen at RT. The most obvious change which appears in the XPS data is a modification in the relative intensities of the Cu3s and Al2s peaks (see figure 8), which is mainly related to the increase of the Al2s peak and can thus be ascribed to segregation and enrichment of the surface by aluminium atoms. Annealing of the oxide film formed at RT for 60 min at 925 K leads to the appearance of a faint and poorly defined LEED pattern, shown in figure 9(a). It has not been possible to improve the quality of the LEED pattern either by longer annealing or by annealing at different temperatures. Nevertheless, three sub-sets of spots with rectangular unit meshes can be identified, superimposed to a high background and weak streaks forming polar circles. The three sets of spots are schematically drawn in figure 9(b). The largest rectangular unit mesh (white rectangle) corresponds to the clean substrate spots (white dots). The two
Figure 8. Evolution of the Cu3s and Al2s photopeaks after 500 L O2
exposure at RT (red/bottom curve) and subsequent annealing at 923 K during 30 min (green/middle curve) and 1 h (pink/top curve). Take-off angle θ = 90◦ .
remaining sets of spots (blue and red spots and rectangles) can be ascribed to two domains of the same structure, having a rectangular unit mesh (a = 10.5 Å and b = 18.2 Å) and approximately rotated by ±23.5◦ with respect to the substrate. The re-appearance of the spots corresponding to the clean surface suggests that, after annealing, the surface is not fully covered by the oxide layer. In addition, the poor contrast observed in the LEED pattern probably reflects the small size of the ordered domains. These conclusions are supported by STM experiments performed following annealing of the oxide film. It proved very difficult to obtain images of the oxide film 6
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Figure 9. (a) LEED pattern (E = 60 eV) of the crystallized aluminium oxide obtained on γ -Al4 Cu9 (1 1 0) after 500 L O2 exposure at RT and subsequent annealing at 925 K. (b) Simulated LEED pattern. The white unit cell corresponds to the substrate (a = 8.7 Å, b = 12.4 Å); √ the blue and red unit cells correspond to the two domains of the sixton oxide (a = 10.5 Å, b = 18.2 Å, b /a ≈ 3).
Figure 10. (a) 5 nm × 4.5 nm high resolution STM image (Fourier filtered) of the crystallized Al oxide obtained on γ -Al4 Cu9 (1 1 0). Vb = −1.65 V and It = 0.37 nA. The rectangle corresponds to the oxide unit cell. Al atoms (empty circles) of the interfacial Al16 O24 plane are arranged in heptagon–pentagon pairs (light blue solid line); (b) FFT analysis.
be identified, having unit cell dimensions (10.5 Å × 18.2 Å) identical to those observed by Napetschnig et al [24] after dosing a Cu-9%Al(1 1 1) sample with 100 L at 953 K with an oxygen partial pressure of 1.3 × 10−7 mbar. In addition, the bright spots appearing in our STM images are organized in regular heptagonal–pentagonal arrangements similar to those corresponding to the interface Al layer obtained by Napetschnig et al [24] on Cu-9%Al(1 1 1) and Kresse et al [25] following oxidation of NiAl(1 1 0). In view of these similarities, we believe that the crystallized oxide islands which are obtained in the present work have the same structure as the aluminium oxide film on NiAl(1 1 0), whose unit cell has been described in terms of a stacking sequence of four layers Ali −Oi −Als −Os composed of 16 interface Al, 24 interface O, 24 surface Al and 28 surface O atoms respectively or by an interfacial Al16 O24 and surface Al24 O28 layers, leading to an Al2 O2.6 global oxygen deficient stoichiometry [25, 35]. It should be noted that the same structure has recently been observed by Pr´evot et al for the oxide film formed by oxidation at RT under 10−6 mbar oxygen followed by
with atomic resolution. The best image of the oxide islands which could be obtained is shown on figure 10(a), together with the corresponding FFT analysis (figure 10(b)). Bright dots are observed, which are organized in regular heptagonalpentagonal arrangements similar to those observed by Kresse et al following oxidation of NiAl(1 1 0) [25]. It should be noted that the unit mesh deduced from this image is slightly distorted and larger (a = 10.4 ± 0.4 Å and b = 19.6 ± 0.8 Å) than that obtained from the LEED pattern, due to the thermal drift during the experiment. 4. Discussion 4.1. Oxide structure
The results presented in this paper show some similarities with the oxidation behaviour reported in the literature for α-Cu–Al solid solutions. In particular, in spite of the poor quality of the complex LEED patterns observed in our studies after annealing the oxide film formed at RT at 925 K, two domains of oxide rotated 23.5◦ with respect to the substrate can 7
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Figure 12. Evolution of the Al2s + Alox 2s/(Al2s + Alox 2s + Cu3s) intensity as a function of the take-off angle θ for the clean γ -Al4 Cu9 (1 1 0) surface (black square) and for the γ -Al4 Cu9 (1 1 0) surface after 3200 L O2 exposure at RT (red dot).
Figure 11. Atomic arrangement of the sixton interface layer on the Al4 Cu9 (1 1 0) surface. Yellow/dark grey dots: Cu; light blue/light grey dots: Al in the substrate surface layer; dark blue/black dots: Al in the sixton interfacial plane; large white circles: O in the sixton interfacial plane. Cu rows of the underlying substrate are indicated by vertical lines. Drawing produced using VESTA [42].
work, only crystallized domains having a small size are formed. Moreover, the weak streaky polar circles observed in figure 9(a) suggest that the small oxide domains can present all orientations with respect to the substrate, in addition to the favoured 23.5◦ misorientation which gives rise to the main new diffraction spots.
annealing at 1000 K of 2.5 aluminium monolayers deposited on top of Ni(1 1 1) [36, 37]. It was argued by these authors that the specific ratio between the length of the two sides of √ the alumina unit cell ( 3), which corresponds to a sixton rectangle (i.e. a rectangle inscribed in a regular hexagon), can be attributed to the hexagonal arrangement of the oxygen atoms in the interfacial layer Oi and the Al atoms in the surface layer Als . Moreover, Pr´evot et al pointed out that the ordered alumina films obtained in previous work performed by various authors on NiAl(1 1 0) [38], FeAl(1 1 0) [39] and aluminium ultrathin films deposited onto Cu(1 1 1) [40] or Ni(1 1 1) [41] could also be interpreted in terms of the same sixton structure within small distorsions of the unit cell, thus suggesting that this structure could be the equilibrium state of a two-layer-thick alumina film supported on a metal substrate. The misorientation of ∼23.5◦ of the sixton structure with respect to the substrate observed in our studies looks very similar to the ±24◦ rotation reported for alumina ultrathin films formed on NiAl(1 1 0) and FeAl(1 1 0) [25, 38, 39] and can be ascribed, in the same way, to a ‘row matching’ effect. The distance between two pure Cu rows in the [1 −1 0] direction for the model γ -Al4 Cu9 (1 1 0) surface can be estimated to 4.13 Å. As shown in figure 11, a 23.5◦ misorientation of the sixton structure allows most of the Al atoms of the heptagonpentagon pairs in the interface layer (Ali ) to be located on top of these pure copper rows thus avoiding unfavourable Al-Al neighbouring. It should be noted that, although the oxide structure is incommensurate with the substrate, the strong intensity observed for some spots in the LEED patterns suggests the vectorial relationship 3b∗ ∼ 4b∗ −a ∗ , i.e. commensurability is nearly achieved in one direction. Finally, it must be kept in mind that, in view of the poor quality of the LEED patterns observed in this
4.2. Al segregation
The results presented in section 3 show that oxidation at RT of the γ -Al4 Cu9 (1 1 0) surface leads to the exclusive formation of Al–O bonds. Oxygen bonding to copper is never observed, even at low oxygen coverages, although the clean surface model mainly contains copper (12 Cu versus 6 Al atoms), thus suggesting easy segregation of aluminium towards the surface. For all Cu–Al alloy compositions and surface orientations studied so far, a strong oxygen induced segregation at elevated temperature has been reported, due to the energetically favourable formation of Al–O bonds [12– 20, 43, 44]. The large thickness of the oxide film obtained in this work for oxidation at RT (8.5 Å ± 0.6 Å) suggests that oxygen-induced segregation occurs even at RT. Further evidence of this segregation effect is given in figure 12. The variation in aluminium percentage, defined as the ratio of the sum of the Al2s + Alox 2s peak intensities (IAl ) upon the sum of the Al2s + Alox 2s + Cu3s intensities (Itot ), is plotted as a function of the take off angle for the clean surface and the surface saturated with 3200 L oxygen at RT. Figure 12 clearly shows a rather weak Al surface enrichment for the clean surface (in good agreement with the 44 at% obtained by Gaudry et al [32]) whereas oxidation leads to a strong enhancement in the aluminium surface concentration, due to aluminium diffusion towards the surface. In order to estimate the Al segregation energy, the clean surface has been saturated with oxygen at RT then annealed at a fixed temperature for increasing periods of time. Three annealing temperatures have been used (350, 365 and 385 K). Between 8
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Figure 14. Intensity of Alox 2s, O1sLBE and O1sHBE core levels as a function of the surface preparation (O2 exposure at RT or annealing at 923 K). Take-off angle θ = 90◦ . After each annealing at 925 K the LEED pattern exhibits both the sixton oxide structure (S) and the substrate structure (C). After each O2 exposure the LEED pattern only exhibits the sixton structure. The solid lines are only guides for the eye.
Figure 13. Evolution of Al as a function of the square root of
annealing time for several temperatures following 3200 L O2 exposure at RT. Take-off angle θ = 30◦ .
each temperature, the crystal was cleaned by the sputterannealing process described in section 2 then dosed again with 3200 L. Figure 13 shows the variations in the relative enrichment in Al (Al, defined as the difference between the Al percentage at a given annealing time and the initial Al percentage) as a function of the square root of the annealing time t 1/2 . Due to the high inertia of our specimen holder, annealing periods must be long enough to minimize errors and only a few experimental points can be obtained. In spite of these restrictions, an Arrhenius analysis of the variation of the initial slope of these plots as a function of 1/T leads to a rough estimation of the activation energy for Al diffusion of 0.65 ± 0.12 eV at−1 (64 ± 12 kJ mol−1 ). This value is much smaller than those usually ascribed to the vacancy diffusion mechanism of Al in crystalline Al-based intermetallics (∼1.6 eV) [45] or in Cu (1.81 eV) [43] although it is close to the value observed by Gil-Gavatz et al on i-Al–Cu–Fe for temperatures below 770 K (0.6 eV) [46], which has been ascribed to a phason flip mechanism. However the phason flip mechanism, being intrinsic to a quasiperiodic structure, cannot be invoked in our case. A tentative explanation for the low value observed in this study can be given when considering that the activation energy of diffusion is composed of a vacancy formation and a vacancy migration contribution to enthalpy, which have close values (0.68 and 0.62 respectively in pure Al) [46]. As the γ -Al4 Cu9 phase has a structure based on a 3 × 3 × 3 CsCl superstructure containing two vacancies [31, 32] only the vacancy migration enthalpy must be considered in our case, thus leading to a very low value of the activation energy for Al segregation.
the O1s line shape during oxidation of Al or Al-containing alloys have already been reported in the literature. In particular, two components in the O1s have been observed at 530.1 and 532.0 eV during oxidation of α-Cu–Al (1 0 0) surfaces and ascribed to Cu–O and Al–O bonds respectively [13, 16, 17]. On pure aluminium, components in the range of 530.6–531.3 eV and 532.2–532.7 eV were associated with chemisorbed and oxidic oxygen respectively [47–51] whereas the main 532.3–533.4 eV and the weak 534.3–535.1 eV features observed by Jeurgens et al are associated to oxidic species within the oxide film and surface oxide species respectively [52]. The presence of two O1s components at 531.2 and 532.8 eV has also been reported in studies performed on the ultrathin aluminium oxide film formed on Ni3 Al(1 1 1) and ascribed to oxygen atoms at the metal-oxide interface and in the surface layer respectively [53]. However, a high resolution core-level spectroscopy study performed by Martin et al on the ultrathin aluminium oxide film formed on NiAl(1 1 0) showed that the high binding energy component should be assigned to only part of the surface oxygen atoms having an aluminium atom directly located underneath and not residing close to another oxygen atom [54]. In the present study, the formation of Cu–O bonds can be disregarded in view of the lack of any modification occurring in either the Cu2p photopeak or the CuLMM Auger peak. Moreover, the similar general trends shown in figure 5 in the Alox 2s, O1sLBE and O1sHBE intensities suggest that both O1s features are related to Al–O bonds. To gain further insight into the origin of these two components, cycles of exposure to oxygen at RT followed by annealing for 1 h at 925 K were performed. Between each step, XPS data and LEED patterns were recorded and the results are shown in figure 14. Ox1 refers to initial 1000 L oxygen dosing at RT. At this stage, only a strong background with no diffraction spots is observed in the LEED pattern. As described in section 3.2,
4.3. Oxygen peak
It has been shown in section 3.1 that the O1s peak can be fitted with a single component (O1sLBE at 531.4 eV) for oxygen exposure up to 120 L but a second component (O1sHBE at 532.4 eV) is required for higher dosing. Changes in 9
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subsequent annealing for 1 h at 925 K leads to the appearance of the spots characteristic of the sixton structure (labelled S in figure 14) and the reappearance of the clean substrate diffraction spots (C). Changes in the photopeaks are also observed, with a decrease in Alox 2s and O1sLBE intensities and an increase of O1sHBE . Additional exposure to 6000 L at RT (Ox2) leads to the disappearance of the clean substrate spots C whereas the sixton spots S are still visible. The intensities of all photopeaks are increasing. Further annealing leads to the reappearance of the clean substrate diffraction spots (C) in addition to the S spots and the decrease in the intensities of all photopeaks. The same tendency is observed for the third cycle: the disappearance of only the clean substrate spots and an increase in the intensities of all photopeaks following oxygen dosing at RT (Ox3), the reappearance of the clean substrate diffraction spots (C) in addition to the S spots and a decrease in the intensities of all photopeaks following subsequent annealing. However, it should be noted that the variations observed in the intensity of the O1sLBE peak are much larger than for the O1sHBE feature. The observed oscillating behaviour of the photopeaks, together with the reappearance/disappearance of the C spots suggests that the O1sLBE peak corresponds to loosely bond O species which are adsorbed on the clean surface during exposure to oxygen at RT and partly desorb at high temperatures, thus liberating patches of clean surface and providing new adsorption sites for the next oxygen dosing. In view of the weak variations observed in the O1sHBE peak during the second and third cycles, together with the persistence of the S spots, we believe that this feature can be ascribed to oxygen atoms in the crystallized oxide islands. The presence of the O1sHBE feature even following the first dosing at RT means that very small sixton domains are already present at this stage, whose size is increasing following the first annealing, leading to a large increase in the corresponding signal. The small decrease observed in the intensity of the O1sHBE peak following annealing in cycles 2 and 3 is more difficult to explain. It can be ascribed either to desorption of some O species within the ordered oxide islands or desorption of part of these islands. It should be noted that the binding energy of the O1sHBE peak observed in our studies (532.4 eV) is close to the value ascribed by Martin et al to oxygen atoms in the ultrathin oxide film formed on NiAl(1 1 0) [54]. The additional component shifted 1.23 eV towards higher binding energies observed by these authors and assigned to specific oxygen atoms in the surface layer was not observed in our case. However, the possible presence of a peak at higher energy is difficult to precisely assess in our case, due to the strong overlap with the CuLMM peak (see figure 4). In any case, neither the ratio between the O1sHBE and the O1sLBE observed in our studies (>1, to be compared with 0.19 in [54]) nor the up and down behaviour during the dosing/annealing cycles can be explained by ascribing both our O1sHBE and O1sLBE signals to oxygen species belonging to the sixton structure. However, a high resolution core-level spectroscopy study would be necessary for a more precise assessment of the origin of both the O1s signals observed in our study.
5. Conclusion
RT oxidation of the (1 1 0) surface of the γ -Al4 Cu9 complex metallic alloy was performed in the range of 2 × 10−8 to 2 × 10−7 mbar oxygen pressure. A fast oxidation step is observed up to 500 L, followed by a slow regime during which the oxide film grows at a much lower rate. Cu–O bonds are never observed, in spite of the high Cu content of the clean (1 1 0) surface, due to a strong oxygen induced aluminium segregation. The activation energy for aluminium diffusion was estimated at 0.65 ± 0.12 eV at−1 (64 ± 12 kJ mol−1 ), a rather low value which was explained by the presence of two vacancies in the γ -Al4 Cu9 structure, thus reducing the activation energy to the sole contribution of vacancy migration enthalpy. When the oxide film formed at RT is annealed at 925 K, small crystallized domains with a sixton structure appear. The similarity of this sixton structure with structures reported in the literature following the oxidation of Cu-9%Al(1 1 1), NiAl (1 1 0) and FeAl(1 1 0) surfaces as well as ultrathin Al films deposited onto Cu(1 1 1) or Ni(1 1 1) surfaces, gives further evidence of the free standing nature of this ultrathin aluminium oxide layer. However, it is the first time to our knowledge that this peculiar structure has been observed on a surface having a complex structure with a large unit mesh. In addition, it seems that it is easier to form the ordered structure on such a surface, as it is directly obtained by annealing the oxide film formed at RT whereas, on all the other surfaces studied so far, its formation required exposure to oxygen at high temperature. Work is in progress to check if the sixton structure can be formed on the surface of other complex metallic alloys. Finally two contributions were observed in the O1s peaks at 531.4 and 532.4 eV, which have been ascribed to loosely bond oxygen species and oxygen belonging to the sixton structure respectively. Acknowledgments
The Agence Nationale pour la Recherche is acknowledged for its financial support (ANR-08-BLAN-0041). This work was performed within the frame of the European Centre for the Development of new Metallic Alloys and Compounds (CMAC, www.eucmac.eu). References [1] Dubois J-M 2005 Useful Quasicrystals (Singapore: World Scientific) [2] Thiel P A 2008 Annu. Rev. Phys. Chem. 59 129 [3] Barth´es-Labrousse M-G and Dubois J-M 2008 Phil. Mag. 88 2217–25 [4] Dubois J-M, Kang S S and Massiani Y 1993 J. Non-Cryst. Solids 153–154 443–5 [5] Dubois J-M, Proner A, Bucaille B, Cathonnet P, Dong C, Richard V, Pianelli A, Massiani Y, Ait-Yaazza S and Belin-Ferr´e E 1994 Ann. Chim.-Sci. Mater. 19 3–25 [6] Fourn´ee V, Barth´es-Labrousse M-G and Dubois J-M 2008 Solid State Phenomena (Switzerland: Trans Tec Publications) pp 407–50 10
M Warde´ et al
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[31] Dong C 1996 Phil. Mag. A 73 1519–28 [32] Gaudry E, Shukla A K, Duguet T, Ledieu J, deWeerd M C, Dubois J-M and Fourn´ee V 2010 Phys. Rev. B 82 085411 [33] Strohmeier B R 1990 Surf. Interface Anal. 15 51–6 [34] Tanuma S, Powell C J and Penn D R 1994 Surf. Interface Anal. 21 165–76 [35] Nishimura T, Hoshino Y, Okazawa T and Kido Y 2008 Phys. Rev. B 77 073405 [36] Pr´evot G, Naitabdi A, Bernard R and Borensztein Y 2010 Phys. Rev. B 81 085405 [37] Pr´evot G, Le Moal S, Bernard R, Croset B, Lazzari R and Schmaus D 2012 Phys. Rev. B 85 205450 [38] Stierle A, Renner F, Streitel R, Dosch H, Drube W and Cowie B C 2004 Science 303 1652 [39] Graupner H, Hammer L, Heinz K and Zehner D M 1997 Surf. Sci. 380 335–51 [40] Jeliazova Y and Franchy R 2002 Appl. Surf. Sci. 187 51–9 [41] Le Pevedic S, Schmaus D and Cohen C 2008 Surf. Sci. 602 67–76 [42] Momma K and Izumi F 2011 VESTA 3 for three-dimensional visualization of crystal, volumetric and morphology data J. Appl. Crystallogr. 44 1272–6 [43] Yu Y, Sagisaka K and Fujita D 2009 Surf. Sci. 603 723–6 [44] Lanzani G, Kangas T and Laasonen K 2009 J. Alloys Compounds 482 33–42 [45] Joulaud J L, Berbardini J, Bergman C, Dubois J-M, Calvayrac Y and Gratias D 1997 Phil. Mag. A 75 1287 [46] Gil-Gavatz M, Rouxel D, Pigeat P, Weber B and Dubois J-M 1999 Mater. Res. Soc. Symp. Proc. 553 75–80 [47] Bradshaw A M, Domcke W and Cederbaum L S 1977 Phys. Rev. B 16 1480–8 [48] Bradshaw A M, Hofmann P and Wyrobisch W 1977 Surf. Sci. 68 269–76 [49] Zhukov V, Popova I, Fomenko V and Yates J T J 1999 Surf. Sci. 441 240–50 [50] Popova I, Zhukov V and Yates J T J 1999 Appl. Phys. Lett. 75 3108–10 [51] Popova I, Zhukov V and Yates J T J 2000 J. Appl. Phys. 87 8143–7 [52] Jeurgens L P H, Sloof W G, Tichelaar F D and Mittemeijer E J 2002 Surf. Sci. 506 313–32 [53] Vesselli E, Baraldi A, Lizzit S and Comelli G 2010 Phys. Rev. Lett. 105 046102 [54] Martin N M et al 2011 Phys. Rev. B 83 125417
[7] Barth´es-Labrousse M-G, Beni A and Schmutz P 2011 Complex Metallic Alloys ed J-M Dubois and E Belin-Ferr´e (Weinheim: Wiley-VCH) pp 243–72 [8] Longchamp J-N, Burkardt S, Erbudak M and Weisskopf Y 2007 Phys. Rev. B 76 094203 [9] Burkardt S, Erbudak M and Maeder R 2009 Surf. Sci. 603 867–72 [10] Erbudak M, Mungan M and Burkardt S 2009 Appl. Surf. Sci. 256 1284–8 [11] Franchy R 2000 Surf. Sci. Rep. 38 195–294 [12] Hoffman A, Maniv T and Folman M 1988 Surf. Sci. 193 57–80 [13] Kravchuk T and Hoffman A 2007 Surf. Sci. 601 87–94 [14] Hoffman A and Folman M 1991 Vacuum 42 1133–6 [15] Shen W, Nyberg G L and Hoffman A 1995 Surf. Sci. 334 209–14 [16] Kravchuk T, Akhvlediani R and Hoffman A 2005 Nucl. Instrum. Methods Phys. Res. B 230 413–8 [17] Kravchuk T and Hoffman A 2006 Surf. Sci. 600 1252–9 [18] Yoshitake M, Bera S, Yamauchi Y and Song W 2003 J. Vac. Sci. Technol. A 21 1290–3 [19] Yoshitake M, Bera S and Yamauchi Y 2003 Surf. Interface Anal. 35 824–8 [20] Song W, Yoshitake M, Bera S and Yamauchi Y 2003 Japan. J. Appl. Phys. 42 4716–20 [21] Yoshitake M, Song W, Libra J, Masek K, Sutura F, Matolin V and Prince K C 2008 J. Appl. Phys. 103 033707 [22] Yoshitake M, Nemsak S, Skala T, Tsud N, Kim T, Matolin V and Prince K C 2010 Surf. Sci. 604 2150–6 [23] Yoshitake M, Song W, Libra J, Masek K, Sutura F, Matolin V and Prince K C 2010 Appl. Surf. Sci. 256 3051–7 [24] Napetschnig E, Schmid M and Varga P 2008 Surf. Sci. 602 1750–6 [25] Kresse G, Schmid M, Napetschnig E, Shishkin M, Koehler L and Varga P 2005 Science 308 1440–2 [26] Gil-Gavatz M, Rouxel D, Pigeat P, Weber B and Dubois J-M 2000 Phil. Mag. A 80 2083–97 [27] Rouxel D, Gil-Gavatz M, Pigeat P and Weber B 2005 J. Non-Cryst. Solids 351 802–9 [28] Rouxel D and Pigeat P 2006 Prog. Surf. Sci. 81 488–514 [29] Arnberg L and Westman S 1978 Acta Crystallogr. A 34 399–404 [30] Villars P and Calvert L D 1998 Pearson’s Handbook of Crystallographic Data for Intermetallic Phases (Materials Park, OH: ASM International)
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