Modifications and Growth Mechanisms of Ultrathin Aluminum Oxide

Mar 29, 2010 - Ivey , M. M.; Layman , K. A.; Avoyan , A.; Allen , H. C.; Hemminger , J. C. J. Phys. Chem. B 2003, 107, 6391. [ACS Full Text ACS Full T...
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Modifications and Growth Mechanisms of Ultrathin Aluminum Oxide Films on NiAl in Water Vincent Maurice,* Ine`s Bennour, Sandrine Zanna, Lorena H. Klein, and Philippe Marcus* Laboratoire de Physico-Chimie des Surfaces, CNRS (UMR 7045), Ecole Nationale Supe´rieure de Chimie de Paris (Chimie ParisTech), UniVersite´ Pierre et Marie Curie, 11 rue Pierre et Marie Curie, 75005 Paris, France ReceiVed: February 9, 2010; ReVised Manuscript ReceiVed: March 18, 2010

X-ray photoelectron spectroscopy (XPS) and atomic force microscopy (AFM) have been combined to study the modifications and growth of alumina thin films on NiAl(001) induced by exposure to water vapor at low pressure (10-6 mbar) and by immersion in ultrapure liquid water at room temperature. The film thickness and stoichiometry, as well as the alloy composition changes, were analyzed by angle-resolved XPS measurements. Ultrathin (∼0.7 nm) hydroxylated native oxide films were compared to thicker (∼4 nm) anhydrous and oxygendeficient thermal oxide films grown at 900 °C under low air pressure (10-4 mbar). The results show that immersion in liquid water causes the formation of one to two equivalent additional monolayers of aluminum oxide at the interface between the native oxide film and the alloy, suggesting anion transport along the oxide/ alloy interface after entry at intergranular oxide sites. The hydroxylated oxide surface remains unchanged. In contrast, immersion of the thicker thermal oxide films in liquid water causes the formation of ∼3 equivalent additional aluminum oxide monolayers at the oxide/water interface, indicating cation transport through the film. AFM suggests a preferential growth at the boundaries between the oxide grains due to faster ion transport. The ∼4 nm thick thermal oxide film is inert when exposed to water at low vapor pressure (10-6 mbar) confirming the pressure gap for the water-induced modifications of transitional alumina films. Exposure to ambient air partially annihilates the oxygen deficiency observed after formation as shown by the variation of the O-Al stoichiometry and core level shift effects. Introduction Water-induced chemical and structural modifications of aluminum oxide surfaces, interfaces, and thin films are key issues in applications such as corrosion protection, nano- and microelectronics, MEMs processing, sensors, and catalysis.1,2 Model studies have been performed on ultrathin ( 1 Torr at room temperature.4-7,11 At higher pressure, a cooperative mechanism between adjacent adsorbed H2O molecules involving nearest neighbor hydrogen-bonding interactions would lower the kinetic barrier for water dissociation11,12 and thus account for the pressure gap of the reactivity.3 At intermediate pressure (10-7 to 10-1 Torr), water-induced modifications have also been observed on ∼1 nm thick transitional films grown on Ni3Al(111),13 Ni3Al(110),7,13-15 NiAl(001),16 and NiAl(110)15 at room temperature. They have been assigned to the cooperative (i.e., pressure dependent) mechanism between adjacent adsorbed molecules and are initiated at the surface defect sites.3 The reaction with water induces a marked rearrangement of the structure resulting in * To whom correspondence should be addressed. E-mail: (V.M.) [email protected]; (P.M.) philippe-marcus@ chimie-paristech.fr. Tel: +33 (0)1 44276738. Fax: +33 (0)1 46340753.

the loss of long-range order and in surface roughness increase observed by low energy electron diffraction (LEED) and scanning tunneling microscopy (STM).7,13-16 Hydroxylation has been concluded from X-ray photoelectron spectroscopy (XPS) measurement of the O to Al atomic ratio16 although no evidence of a high binding energy shoulder in O 1s spectra was measured with polychromatic sources.7,13,14,16 The amorphization reaction initiates at the oxide/vapor interface at lower H2O exposures and propagates toward the oxide/alloy interface with increasing pressure and time.13,16 The loss of long-range order is also accompanied by an increase in average oxide film thickness measured by XPS16 and Auger electron spectroscopy (AES),15 which is consistent with the high field mechanism of oxide growth (as described by Mott et al.17,18) in which the electric field established by electron transfer between metal atoms at the oxide/alloy interface and dissociatively adsorbed molecules (water or oxygen) at the gas/oxide interface assists ion migration through the film. Comparison of the thickness increase measured on Ni3Al(110) and NiAl(110) and density functional theory (DFT) calculations of the Al vacancy formation energy led to the conclusion that the rate determining step in the oxide thickening process (for ultrathin films) is Al diffusion from the substrate bulk to the oxide/alloy interface,15 in agreement with previous conclusions for aluminum oxide growth on other alloys (e.g., Ti-Al19,20). Here we report on the modifications and growth mechanism of alumina thin films formed on NiAl(001) upon exposure to water at low vapor pressure (10-6 mbar) or to liquid water (at ambient temperature). Ultrathin native oxide films of thickness (18 MΩ · cm, pH 6.0) at room temperature from 1 min to 24 h in cumulated steps. This required the transfer of the sample from the XPS spectrometer to the immersion cell and vice versa through air. The modifications of the thermal oxide films induced by this ex situ transfer were also analyzed. A VG ESCALAB 250 spectrometer (Thermo Scientific) was used for XPS analysis. An AlKR monochromatized radiation (hν ) 1486.6 eV) was employed as X-ray source. The binding

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energies (EB) were calibrated by setting the C 1s signal corresponding to hydrocarbon contamination (-CH2-CH2-) at 285.0 eV. Survey spectra were recorded with a pass energy of 100 eV and high resolution spectra of the Al 2p, Ni 3p, Ni 2p, O 1s, and C 1s core level regions were recorded with a pass energy of 20 eV. Angle-resolved measurements were made by varying the takeoff angles of the analyzed photoelectrons with respect to the surface plane. Peak fitting was performed using the AVANTAGE software and a Shirley background. The procedure allowing us to calculate the equivalent thickness of the alumina layer from the 3+ 0 to Al2p peak intensity ratio and the composition of the Al2p 0 to modified alloy region underneath the oxide from the Ni3p 0 Al2p peak intensity ratio (assuming an homogeneous depth distribution in the metallic phase) has been described previously.19 The O-Al stoichiometric ratio of the alumina layer was calculated from the intensity ratio of the O1s to Al3+ 2p peaks. This required the calibration of the transmission factor ratio (TO1s/ TAl2p) of the spectrometer that was made using a bulk boehmite (AlOOH) sample. Scofield photoionisation cross sections were used and the values of the photoelectron attenuation lengths were calculated using the method of Tanuma et al.29 An Agilent 5100 AFM/SPM microscope was used for AFM analysis. Intermittent contact (i.e., tapping) imaging was performed ex situ in the laboratory air using Si tips with a nominal radius of curvature of 10 nm and mounted on Si cantilevers vibrating at their resonance frequency. The topographic images were acquired at constant damped oscillation amplitude. Results and Discussion Pristine Oxide Films. Figure 1 shows the X-ray photoelectron (XP) Al 2p-Ni 3p and O 1s core level spectra for the NiAl(001) surface covered by the native oxide film. Binding energies and full width at half maximum (FWHM) values of the component peaks are compiled in Table 1. The Al 2p core level was fitted by two peaks at 72.3 and 72.7 eV corresponding 0 , 3/2-1/2 spin-orbit to metallic aluminum in the alloy (Al2p doublet, intensity ratio fixed at 2:1) and one peak at 74.2 eV 3+ , unresolved corresponding to the oxidized aluminum (Al2p 30-32 The Ni 3p core level was fitted 3/2-1/2 spin-orbit doublet). 0 ) in the with four peaks corresponding to metallic nickel (Ni3p alloy (3/2-1/2 spin-orbit doublet, 2:1 intensity ratio) and their satellites.16,33-36 The absence of oxidized nickel was confirmed 0 by the single Ni2p 3/2 peak observed at 853.6 eV (Table 1), showing the selective oxidation of aluminum on β-NiAl in these conditions, in agreement with previous oxidation experiments performed at room temperature.34,36 The observed increase of 3+ 0 the Al2p to Al2p intensity ratio with decreasing takeoff angle is expected for an oxide film covering the substrate surface (Figure 1). The equivalent thickness of the oxide film for various samples was calculated to be 0.7 ( 0.1 nm from measurements at 90 or 45° takeoff angles. The thickness is in agreement with previously reported values for alumina thin films formed on NiAl at room temperature34,36 and amounts to a 3-oxygen layer thick oxide film. The Al atomic concentration of the alloy below the oxide was calculated to be 43 ( 1 and 40 ( 2 at % at 90 and 45° takeoff angles, respectively, the bulk alloy being 50 at %. The Al depletion measured below the oxide is consistent with the selective transfer of the Al atoms to form the oxide film and the absence of rehomogenization of the modified alloy surface owing to extremely slow self-diffusion in the bulk at room temperature. A gradient with stronger Al depletion at the oxide/ alloy interface is indicated by the angle-resolved measurements.

Figure 2. AFM topographic images of the NiAl(001) surface covered by (a) the native oxide layer and (b) the thermal oxide layer formed at 900 °C. (a) ∆Z ) 9.3 nm; (b) ∆Z ) 24 nm. Some grains are pointed in (a).

The O 1s core level was fitted with three peaks at 530.5(O1sA), 532.0(O1sB), and 533.4 eV (O1sC) assigned to O2oxide ions, OH- hydroxide ions, and adsorbed water molecules, respectively.30-32 At lower takeoff angles, the O1sB to O1sA intensity ratio increases showing that the OH- and O2- ions are concentrated in the outer and inner layers of the oxide film, respectively. The O1sC to O1sB intensity ratio does not show a marked angular dependence indicating a similar depth distribution of the H2O and OH- species, both concentrated in the outer layers of the oxide film. The O-Al atomic ratio in the oxide film was calculated to be 2.35 (2.15 excluding the O1sC component). This value is higher than that for Al2O3 (1.5) and intermediate between those for AlOOH and Al(OH)3 compounds (2 and 3, respectively). This confirms that the ultrathin native oxide film cannot be described by a single compound homogeneous in depth but predominantly contains hydroxide species in the outer layers and at the surface and oxide species in the inner layers. Figure 2a shows the AFM topographic image of the alloy surface covered by the native oxide film. The morphology observed between the polishing grooves appears granular with a lateral grain size of 50 ( 8 nm and a depression depth between the grains of 0.4 ( 0.2 nm consistent with the film thickness. However, this morphology is not assigned to the native oxide film but rather to the substrate surface. Indeed, STM data

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Figure 3. XP Al 2p-Ni 3p and O 1s core level spectra recorded at 90° takeoff angle for NiAl(001) covered by the thermal oxide layer formed at 900 °C before and after exposure to p(H2O) ) 1 × 10-6 mbar for 15, 30, and 90 min at room temperature.

obtained after room temperature oxidation in UHV conditions of atomically flat NiAl(001) surfaces have shown that the lateral dimensions of the oxide grains were 2-3 nm and intergranular boundaries ∼0.5 nm deep.36 Such oxide grains are clearly beyond the resolution limit of AFM on the present as prepared surface. Figure 3 shows the XP Al 2p-Ni 3p and O 1s core level spectra for the NiAl(001) surface oxidized at 900 °C for 5 min in p(air) ) 1 × 10-4 mbar (see EB and FWHM values of component peaks in Table 1). A striking difference of the Al2p-Ni3p spectrum with that obtained for the native oxide film 0 is the large increase of the Al3+ 2p to Al2p intensity ratio, evidencing a much thicker film. No oxidation of nickel was observed (single Ni02p peak at 853.9 eV, Table 1) as expected for β-NiAl in these oxidizing conditions.3,8,10,13,16,33,35-41 Some dispersion in equivalent thickness, ranging from 3.3 to 5.2 nm, was obtained with different preparations of the pristine oxide film. However, the Al atomic concentration of the alloy below the oxide films was repeatedly measured to be 55-58 at %, showing that below the thicker oxide film formed at higher temperature the surface of the alloy is now enriched in Al. Al interfacial enrichment was also observed below thinner ∼1 nm alumina layers grown on NiAl(001) at 750 °C and low oxygen pressure (5 × 10-7 mbar).16,33,35,36 Thus it appears that increasing the temperature not only induces a faster growth of the oxide film as expected for a diffusion-limited growth mechanism of the oxide but also rehomogenizes the alloy surface modified by selective oxidation. The rehomogenization mechanism implies the segregation of

J. Phys. Chem. C, Vol. 114, No. 15, 2010 7135 metallic Al from the bulk alloy to the oxide/alloy interface and, in parallel, the diffusion of nickel from the oxide/alloy interface to the bulk alloy. The diffusion coefficients of these two elements are identical in β-Ni50Al50 at 900 °C.42,43 Thus, the alloy surface depleted in Al by selective oxidation becomes depleted in Ni after rehomogenization. Of course, rehomogenization of the alloy surface can only be obtained if the oxide growth rate is sufficiently slow, as observed under low oxygen pressure at 900 °C in the present study and at 750 °C previously.16,34-36 The thermal oxide film is also characterized by a marked modification of the O 1s core level spectrum that can be fitted with a single component peak at 532.6-532.7 eV having the same width as those of the native oxide film (Table 1). This peak, assigned to the oxide anions of the oxide layer, is observed at higher binding energy (+2.1-2.2 eV) than on the native oxide film, also confirming previously reported data for thermal oxide 3+ peak at 75.8 eV films on NiAl and Ni3Al.13,16,33-36,44 The Al2p is also shifted to higher binding energy (+1.6 eV) in comparison with the native oxide. The formation of a Schottky barrier resulting from band bending at the interface between a defective oxide and the alloy is a likely explanation of the core level 3+ and O1s peaks, not shifts.44 The different shift of the Al2p 45 predicted by the band-bending model, also suggests a retarding field (i.e., charging) less effective on higher kinetic energy photoelectrons. We also note shifts of +0.8, +1.1, and 0 0 , Ni3p , and +0.3-0.4 eV to higher binding energies of the Al2p 0 Ni2p core levels, respectively, for the alloy surface below the thermal oxide compared to that below the native oxide. The similar direction of the shifts excludes an effect of the modified stoichiometry on the charge transfer between Al and Ni atoms at the intermetallic surface below the thermal oxide, but rather points to some electron depletion (i.e., oxidation) of both the Al and Ni interfacial atoms. The calculated values of O to Al atomic ratio were 1.0-1.2 and 1.2-1.5 at 90 and 45° takeoff angles, respectively, showing a different depth distribution of the two elements in the thermal oxide film. The values are significantly lower than 1.5 showing an O substoichiometric (or Al overstoichiometric) deviation in the defective oxide film. The value at higher takeoff angle indicates a larger deviation in the inner part of the oxide film. This stoichiometry deviation is consistent with the presence of oxygen vacancies and/or Al interstitials, and thus with an anionic and/or cationic growth mechanism of the oxide film, respectively. Such mechanisms are observed on NiAl in atmospheric pressure conditions at T > 1050 °C (growth of the R phase) and at 750 °C < T < 950 °C (growth of the transitional θ phase), respectively.21-27 The temperature selected in the present study is consistent with a regime of cationic oxide growth and thus with the presence of Al3+ interstitials. However, the Al enrichment observed in the alloy below the oxide film is indicative of the slower growth rate at low oxygen pressure than at atmospheric pressure, which implies the presence of O2vacancies and thus a mixed ionic growth mechanism of the oxide. The AFM topographic image of the alloy surface covered by the thermal oxide film also shows a marked difference with comparison to the surface covered by the native oxide layer (Figure 2). It reveals a granular nanostructure that can unambiguously be assigned to the thermal oxide film. The lateral grain size is 22 ( 5 nm and the depression depth between the grains is 2 ( 1 nm. The measured grain boundary depth is consistent with the absence of pinholes in the oxide of ∼4 nm equivalent thickness, suggesting complete surface coverage by

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Figure 4. XP Al 2p-Ni 3p and O 1s core level spectra recorded at 90 (a) and 40° (b) takeoff angles on NiAl(001) covered by the native oxide layer and immersed 24 h in ultrapure liquid water.

the oxide layer. Deeper depressions (reaching ∼7 nm in average) are locally observed and are preferentially formed along the polishing grooves as revealed by lower magnification images (see below). There are assigned to a local effect of the residual surface stress of the polished surface on the rehomogenization of the alloy surface modified by Al depletion. The densification of dislocations could promote locally the alloy rehomogenization mechanism and form these local depressions. Inertness at Low Water Vapor Pressure. Figure 3 also shows the XP Al 2p-Ni 3p and O 1s core level spectra for the thermal oxide film exposed to low pressure water vapor (1 × 10-6 mbar) at room temperature. The spectra have been 0 peak. Their perfect normalized to the intensity of the Ni3p superimposition indicates that the surface is unchanged which is confirmed by peak fitting and unchanged intensity ratios (Table 1). The ∼4 nm thick alumina film is thus inert when exposed to low pressure water vapor. This contrasts with the thinner (∼ 1 nm) thermal oxide films grown on NiAl(001) previously studied and for which core level shifts and increases of the thickness and O-Al atomic ratio were observed in the same low pressure water exposure conditions.16,36 This also contrasts with the thickness increase reported for the 0.7 nm thick transitional films on NiAl(110) and Ni3Al(110) reported in the 10-6 Torr water vapor pressure range.15 The atomic surface structure of the oxide film was not determined in the present study. However, it seems quite unlikely that the thermal oxide film grown has a well-ordered defect-free termination. The AFM images reveal a granular nanostructure suggesting the presence of a high density of defects at the surface of the oxide that should be reactive in this pressure range as observed previously.3,13-16 The absence of chemical modifications (core level shifts and/or stoichiometry variations) observed in the ∼4 nm thick films produced on

NiAl(001) in the present study is thus assigned to the thickness dependence of the reactivity to water at low pressure. Modifications in Atmospheric Air. The modifications of a pristine thermal oxide film exposed to air were also investigated. After a 10 min exposure to laboratory air (at room temperature), the XP Al 2p-Ni 3p and O 1s core level spectra could be fitted with the same components as prior to air exposure (Table 1). The oxide film equivalent thickness and the Al concentration of the alloy below the oxide were calculated to be 4.5 nm and 56% versus 4.3 nm and 57% prior to exposure, respectively, showing no significant changes. A slight difference was the absence of marked angular-dependence of the O to Al atomic ratio in the oxide film measured to be 1.40 and 1.34 at 90 and 45° takeoff angles, respectively. This shows a similar depthdistribution of these elements after exposure to air, suggesting that the stoichiometry deviation of the pristine thermal oxide film has been preferentially corrected in the inner part after exposure to air at 1 atm. The slight decrease (-0.2-0.3 eV) of 3+ and O1s EB values suggests the reduction of the band the Al2p bending at the oxide/alloy interface contributing to the core level shift,44,45 in agreement with the stoichiometry correction. These modifications imply that, after oxygen uptake at the surface resulting from the dissociative adsorption of O2 and/or H2O molecules at the surface, O2- ions diffuse inward through the oxide layer. The transfer of Al3+ ions from the alloy to the inner part of the film should result in a decrease instead of an increase of the O:Al stoichiometric ratio, and can thus be rejected. The diffusion of O2- ions seems quite unlikely through the oxide network at room temperature. However, it could be promoted by the relatively high concentration of O2- vacancies suggested by the initial stoichiometry deviation and by the poorly crystallized oxide network.

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Figure 5. Effect of exposure to liquid water on the native oxide film. (a) Equivalent thickness of the oxide film, (b) Al atomic concentration of the alloy below the oxide, and (c) atomic O to Al ratio. The values were calculated from spectra recorded at 90 and 40° takeoff angle.

Growth in Liquid Water. Figure 4 shows the Al 2p-Ni 3p and O 1s core level spectra of the surface covered by the native oxide film and exposed for 24 h to liquid water at room temperature. Table 1 shows that the components used for peak fitting are the same as prior to exposure. Identical EB and FWHM values (within (0.1 eV) were obtained at intermediate exposure. This indicates that the nature of the species on the surface is unchanged after exposure to liquid water. Note 3+ however by comparison with Figure 1 the increase of the Al2p 0 to Al2p intensity ratio after exposure indicating thickening of the oxide layer. This is accompanied by the decrease of the 0 0 to Ni3p intensity ratio indicating further depletion in Al of Al2p the alloy surface below the oxide. The angular-dependence of 0 0 to Ni3p intensity ratio confirms the gradient of Al the Al2p depletion below the oxide. Figure 5 shows the variation of the calculated equivalent thickness of the oxide film and Al concentration at the oxide/alloy interface up to 24 h of

immersion in liquid water. The increase of the oxide layer thickness is paralleled by Al depletion increase in the alloy surface below the oxide, showing the slow but continuous further growth of the aluminum oxide layer by selective oxidation of aluminum in the alloy (nickel remaining metallic as confirmed by the XP Ni 2p core level, Table 1). The XP O 1s core level spectrum shows the same O1sA (O2-), O1sB (OH-) and O1sC (H2O) components as prior to exposure to liquid water (Figure 4). However the intensity ratio of the O1sB (OH-) to O1sA (O2-) peaks decreases after exposure. In Figure 5c, the O-Al atomic ratio in the oxide film (excluding the O1sC component) is observed to decrease upon increasing exposure to water. The angular-dependence still observed after 24 h excludes that a defined Al2O3 or AlOOH ultrathin layer homogeneous in depth is formed. The combination of these data shows that the growth of the oxide layer mostly results from the thickening of the inner part enriched in oxide species. This

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Figure 6. AFM topographic images of the NiAl(001) surface covered by (a) the native oxide layer and (b) the thermal oxide layer formed at 900 °C and immersed 24 h in ultra pure liquid water. (a) ∆Z ) 28 nm; (b) ∆Z ) 14 nm.

suggests the penetration of oxide ions below the extreme surface of the oxide film after dissociative adsorption of the water molecules and their transport to the oxide/alloy interface where they combine to Al ions to form new oxide layers. The equivalent of one to two additional monolayers of oxide is thus formed by anionic growth in these exposure conditions. Figure 6a shows the AFM topographic image of the alloy surface covered by the native oxide film after exposure to liquid water for 24 h. The lateral grain size is 49 ( 14 nm, that is, not markedly different from the value prior immersion. The depression depth between the grains increases to 1.5 ( 0.5 nm, which better reveals the granular morphology. These depressions exceeding the oxide film equivalent thickness suggest however that the observed morphology is not that of the modified oxide film and mostly corresponds to the alloy surface. Thus, it can be concluded that the NiAl surface covered by the ∼0.7 nm thick native aluminum oxide layer is reactive when exposed to liquid water. A slow growth of the film, reaching ∼1.0 nm after 24 h of exposure, is observed and results from the formation of one to two additional oxide monolayers in the inner part of the film by predominant anion transport to the oxide/alloy interface. The growth by selective oxidation of the alloy is confirmed by the increasing depletion in metallic Al observed underneath the oxide film. The surface and outer layers of the film remains hydroxylated and hydrated.

Maurice et al. Figure 7 shows the XP Al 2p-Ni 3p and O 1s core level spectra recorded on the NiAl(001) surface covered by a 4.3 nm thermal oxide modified in air and exposed for 24 h to liquid water at room temperature. The components used for peak fitting are the same as prior to exposure (Table 1). Whereas identical EB and FWHM values (within (0.1 eV) are obtained for the 0 components, we note a marked shift to lower binding Al2p 3+ and O1sA (O2-) peaks, energies (-0.4-0.5 eV) of the Al2p confirming the reduction of the band bending effect observed after exposure to air and suggesting further correction of the stoichiometry deviation in the film. Figure 8a shows the variation of the calculated equivalent thickness of the oxide film exposed to water for up to 8 h of exposure. A small thickness increase of the oxide layer reaching 4.9-5.0 nm is observed rapidly before stabilization. It is accompanied by the decrease of the Al concentration of the alloy, more pronounced at the surface of the alloy than in the deeper layers as indicated by the angle-resolved measurements (Figure 8b). Figure 8c shows the variation of the calculated O to Al atomic ratio in the oxide film. A rapid increase to 1.5-1.6 is calculated from the spectra recorded at 40° takeoff angle. A similar variation, although less pronounced, is also observed at 90° takeoff angle, the value reaching 1.4-1.5. These data evidence a selective transfer of Al atoms from the alloy to the oxide film induced by exposure to liquid water, and a resulting thickness increase of 3 equivalent oxide monolayers. The stoichiometry variation reaching the Al2O3 value in the outer part of the film is consistent with the uptake of oxygen at the extreme surface and the growth of the new oxide at the oxide/ liquid water interface. The lower O:Al stoichiometry value reached in the inner part of the film suggests that the injected Al3+ ions are not yet all combined to O2- anions in agreement with a mechanism of oxide growth by cation transport. Figure 6b shows the AFM topographic image of the alloy surface covered by the thermal oxide film and immersed in liquid water for 24 h. The lateral grain size is 21 ( 10 nm, not markedly different than prior immersion. The depression depth between the grains is 2.0 ( 0.5 nm which remains consistent with the absence of pinholes in the film. The average value is not different than prior immersion. However, the smaller local variation ((0.5 instead of (1 nm) suggests preferential growth of the oxide at the intergranular boundaries where the initial thickness was lower. These data show that the ∼4 nm thick thermal oxide film is also reactive and thickens by the equivalent of 3 oxide monolayers when exposed to liquid water. The comparison with its inertness at low pressure water vapor confirms for a ∼4 nm thick film the pressure gap of reactivity previously reported for thinner films of transitional alumina,3 and suggests that water dissociation is indeed promoted by interaction between coadsorbed molecules. The injection of Al3+ ions in the inner part of the film required for oxide growth is indicative of the establishment of an electric field between the oxide/alloy and the oxide/water interfaces, in agreement with the high field mechanism of oxide growth.17,18 It also confirms that substrate effects mediated by electron transfer, as recently observed with gold particles,46 must not be neglected when studying the reactivity of oxide surfaces modeled by ultrathin films. In contrast with the native oxide film, the growth of the thermal oxide in liquid water predominantly occurs at the oxide/ water interface as deduced from the angular dependence of the stoichiometry, indicating a predominant cation transport in the film. This difference suggests a combined effect of thickness and nanostructure. The intergranular sites of the thicker thermal

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Figure 7. XP Al 2p-Ni 3p and O 1s core level spectra recorded at 90 (a) and 45° (b) takeoff angles for NiAl(001) covered by the thermal oxide layer formed at 900 °C and exposed to p(H2O) ) 1 atm for 24 h at room temperature.

oxide film apparently do not reach the oxide/alloy interface thus preventing direct access of the water molecules to the buried interface. Oxide growth must then proceed by ion transport through the oxide network and we observe that this predominantly occurs by cation transport. The AFM data suggest preferential growth at the interfaces between the oxide grains which can be explained by accelerated ion transport at grain boundaries.47,48 According to the STM data of Fre´my,36 the intergranular sites of the thinner oxide film grown at room temperature do reach the oxide/alloy interface. Direct anion entry at the oxide/alloy interface can then take place after water dissociation and the present results suggest that they migrate along the oxide/alloy interface to grow the equivalent of one to two oxide monolayers. The absence of further growth of the native oxide film in liquid water evidence that this additional oxide layers form a much more effective barrier against ion transport than the thicker defective thermal oxide film grown in this study. Conclusion The modifications and growth induced by exposure to water vapor at low pressure (1 × 10-6 mbar) and/or immersion in liquid water at room temperature of ultrathin native and thermal aluminum oxide films on NiAl(001) have been studied. Angleresolved XPS was applied to analyze the thickness and stoichiometry of the oxide film and the surface concentration of the modified alloy below the oxide. The oxide film morphology was characterized by AFM. The NiAl surface covered by the ∼0.7 nm thick native aluminum oxide layer formed in laboratory air at room temperature is reactive when exposed to liquid water. A slow growth of the film, reaching ∼1.0 nm after 24 h of immersion, is observed and results from the formation of one to two additional aluminum oxide monolayers in the inner part of the film by anion transport to the oxide/alloy interface. The growth by selective oxidation of the alloy is confirmed by the increase

of the initial depletion in metallic Al observed underneath the oxide film. The surface and outer layers of the film remain hydroxylated and hydrated as indicated by the three component peaks resolved in the O 1s core level. The nanoscale morphology of the oxide film could not be resolved by AFM on the as prepared surface. However, previous STM data suggest that the anions could access the oxide/alloy interface at intergranular sites between the oxide grains and migrate along the oxide/ alloy interface to grow the new oxide monolayers. A thicker (∼4 nm) thermal aluminum oxide film was grown by heating the NiAl surface at 900 ( 50 °C for 5 min under low air pressure (1 × 10-4 mbar). Because of the slow growth of the oxide, the alloy surface is rehomogenized and enriched in Al segregating from the bulk (and Ni diffusing toward the bulk). The oxide film is anhydrous and markedly oxygendeficient with a stoichiometry deviation more pronounced in 3+ the inner layers. Marked positive core level shifts of the Al2p 2and O1s peaks are assigned to band bending at the interface between the oxygen-deficient film and the alloy and to charging. 0 0 0 , Ni3p , and Ni2p peaks also Positive core level shifts of the Al2p indicate electron depletion of the Al and Ni atoms at the alloy surface. The oxide morphology is nanogranular with grains ∼20 nm wide in average. This oxide-covered surface is inert when exposed to water at low pressure (1 × 10-6 mbar) in contrast with thinner (∼1 nm) transitional alumina thin films also grown on NiAl,3,15,16,36 showing a thickness-dependent mechanism of water dissociation at low pressure. Exposure to ambient air does not modify the oxide layer thickness nor the alloy composition below. However, the oxygen deficiency of the film is reduced as shown by the 3+ 2and O1s change of stoichiometry and the decrease of the Al2p core level shifts, indicating oxygen penetration and transport through the initially highly defective oxide network. Immersion in liquid water at room temperature causes oxide growth accompanied by Al consumption in the alloy. Three new oxide monolayers (+0.7 nm) are formed predominantly at the oxide/

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Figure 8. Effect of exposure to liquid water on the thermal oxide film. (a) Equivalent thickness of the oxide film, (b) Al atomic concentration of the alloy below the oxide, and (c) atomic O to Al ratio. The values were calculated from spectra recorded at 90 and 40° takeoff angles.

water interface indicating cation transport through the film. The decrease of the local variation of the depth measured by AFM at intergranular boundaries of the oxide film suggests preferential growth at the interfaces between the oxide grains in agreement with accelerated ion transport at grain boundaries.47,48 References and Notes (1) Padture, N. P.; Gell, M.; Jordan, E. H. Science 2002, 296, 280. (2) Chambers, S. A.; Droubay, T.; Jennison, D. R.; Mattsson, T. R. Science 2002, 297, 827. (3) Kelber, J. Surf. Sci. Rep. 2007, 62, 271. (4) Elam, J. W.; Nelson, C. E.; Cameron, M. A.; Tolbert, M. A.; George, S. M. J. Phys. Chem. B 1998, 102, 7008. (5) Liu, P.; Kendelewicz, T.; Brown, J. G., Jr.; Nelson, E. J.; Chambers, S. A. Surf. Sci. 1998, 417, 53.

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