Growth, Composition, and Structure of Ultrathin Vanadium Films

The XPD results rule out the formation of a V−Sn mixed oxide. The simulation of the experimental XPD curves by means of single-scattering cluster (S...
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J. Phys. Chem. B 2000, 104, 3121-3129

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Growth, Composition, and Structure of Ultrathin Vanadium Films Deposited on the SnO2(110) Surface† A. Atrei,*,‡ U. Bardi,§ C. Tarducci,§ and G. Rovida§ Dipartimento di Scienze e Tecnologie Chimiche e dei Biosistemi, UniVersita` di Siena, 53100 Siena, Italy, and Dipartimento di Chimica, UniVersita` di Firenze, 50129 Firenze, Italy ReceiVed: August 31, 1999; In Final Form: NoVember 30, 1999

The growth mechanism of vanadium films deposited on the SnO2(110) surface at room temperature and the structure of the phases formed after annealing at 800 K were studied using X-ray photoelectron spectroscopy (XPS), low-energy ion-scattering (LEIS), low-energy electron diffraction (LEED), and X-ray photoelectron diffraction (XPD). The vanadium films were deposited by thermal evaporation in ultrahigh vacuum (UHV) conditions on an oxygen-deficient SnO2 surface prepared by cycles of sputtering and annealing. In the initial stages of vanadium deposition a redox reaction occurs at the metal-oxide interface leading to the formation of vanadium oxide and metallic tin. Upon increasing the amount of deposited vanadium, we observed the growth of islands of metallic vanadium. XPS and LEIS data show that the surface of the vanadium film is partially covered by metallic tin diffusing from the interface. When the deposited vanadium films are heated to 800 K, the reoxidation of metallic tin and the oxidation of all metallic vanadium take place, with no detectable diffusion of vanadium into the SnO2 bulk. The XPD results rule out the formation of a V-Sn mixed oxide. The simulation of the experimental XPD curves by means of single-scattering cluster (SSC) calculations performed for various structural models indicate that the product of the oxidation is an oxide of formula VOx≈2 with a structure close to that of the rutile VO2(110) surface. The vanadium oxide phase is covered by layers of disordered tin oxide.

I. Introduction The interest in vanadium films deposited on SnO2 is due to the use of mixed Sn-V oxides as gas sensors. Semiconducting (n-type) SnO2 is widely employed as the active material for sensing hydrocarbon gases and other toxic or explosive vapors.1,2 The gas detection mechanism in these systems is based on the electrical conductivity increase when the gas reacts with oxygen species adsorbed on the SnO2 surface. To increase the sensitivity and selectivity of the sensors, metals (palladium, for instance1,3) or other oxides (e.g., vanadium oxides4) are added to the tin dioxide base material. The conductivity variations upon the reaction between the gas and the oxide are determined, to a large extent, by the properties of the metal/SnO2 and the oxide/ SnO2 interfaces. It is evident that the composition and structure of these interfaces are fundamental parameters for a better understanding of the gas sensing mechanisms. For a detailed characterization at the atomic scale of these interfaces there is the need for model systems consisting of metal and metal oxide films deposited on well-characterized surfaces of a SnO2 single crystal. In the present work, we investigated by several surface science techniques the mechanism of growth and structure of ultrathin vanadium films deposited on a SnO2 single crystal oriented along the (110) surface. The vanadium oxide layers were obtained by solid-state reaction between the deposited film and the SnO2 substrate. Preliminary results about the structure and composition of vanadium oxide layers on the SnO2(110) surface are published elsewhere.5 In the present paper, the †

Part of the special issue “Gabor Somorjai Festschrift”. * Corresponding author (E-mail: [email protected]). ‡ Universita ` di Siena. § Universita ` di Firenze.

growth mechanism of vanadium and the chemical states at the metal/SnO2 interface were determined by means of X-ray photoelectron spectroscopy (XPS) and low-energy ion scattering (LEIS). The structures of the oxide phases formed upon annealing of the vanadium films were investigated by means of low-energy electron diffraction (LEED) and X-ray photoelectron diffraction (XPD). The last technique is particularly suited for investigating the growth of ultrathin films, as shown by previous applications for the investigation of metal-on-metal films6 and metal-on-oxide surfaces.7 Despite the technological relevance of their applications, metals deposited on tin dioxide surface have not been as systematically investigated as have metals on TiO2 surfaces.8,9 One of the reasons is the difficulty in preparing stoichiometric SnO2 surfaces in normal vacuum chamber for surface science experiments, whereas stoichiometric and well-ordered TiO2 surfaces can be prepared by sputtering and annealing.8,9 To our knowledge, only Pd and Sn ultrathin films on SnO2 have been investigated by means of surface science techniques.3,10,11 In those studies surface analysis techniques were used together with in situ conductance measurements of the Pd/SnO2(110) model sensors exposed to gases in order to correlate surface composition and resistivity.3,10 II. Experimental Methods The experiments were performed in an ultrahigh vacuum (UHV) system with a base pressure in the low 10-8 Pa range. The chamber was equipped with a hemispherical electron/ ion energy analyzer used for XPS, XPD, and LEIS. The chamber was also equipped with two-grid LEED optics. Nonmonochromatized Mg KR (1253.6 eV) was used as excitation source for XPS and XPD measurements. The angle between the X-ray

10.1021/jp9930853 CCC: $19.00 © 2000 American Chemical Society Published on Web 01/29/2000

3122 J. Phys. Chem. B, Vol. 104, No. 14, 2000 source and the analyzer axis was 55° and the semicone angle of acceptance of the analyzer was 5°. The XPS spectra were collected in the fixed analyzer transmission mode with a pass energy of 44 eV. The binding scale was calibrated by setting the Ag 3d5/2 peak at 368.3 eV.12 LEIS spectra were measured using a beam of 1 keV He+, impinging on the surface at an angle of 45° along the [11h0] azimuth and with a scattering angle of 135°. The density current was of the order of 1 × 10-8 A/cm2. The SnO2 sample was mounted on a manipulator capable of polar and azimuthal rotations. The sample holder allowed the sample to be heated to 1000 K. The temperature was measured by means of a chromel-alumel thermocouple spot-welded on one of the clamps holding the sample. XPD curves were collected by rotating the sample around the manipulator axis (polar curves) and around the normal to the surface (azimuthal curves). The intensity of the Sn 3d5/2 (768 eV kinetic energy), O 1s (723 eV kinetic energy), and V 2p3/2 (737 eV kinetic energy) photoemission peaks was monitored as a function of the emission angles. The XPD intensities were measured as the peak height at the maximum with respect to the background. The background was estimated by measuring the intensity of the XPS signal at a kinetic energy 5-10 eV higher than the value for the photoemission peak. The tin dioxide single-crystal sample was cut within 1° along the (110) plane from a naturally grown crystal (cassiterite) and polished down to 0.1 µm.13 The resulting sample was a plate 10 × 10 × 2 mm. The SnO2 was prepared by cycles of Ar ion sputtering (2 keV, 5 µA) and annealing up to 900 K in a vacuum or in oxygen atmosphere (5 × 10-4 Pa). After several cycles of sputtering and annealing of the SnO2(110) surface at 800 K, no impurities could be detected within the limit of the sensitivity of XPS and LEIS. The vanadium layers were deposited from a high-purity V wire (99.999% purity) by means of a micro electron-beam evaporator. After prolonged degassing, the pressure in the chamber remained in the 10-8 Pa range during the evaporation. The evaporation rate was estimated from the attenuation of the XPS signal from the substrate upon vanadium deposition. The experimental XPD curves were reproduced by means of calculations performed on the basis of the single-scattering cluster spherical wave (SSC-SW) model.14 The scattering by the oxygen and vanadium atoms is described in terms of phase shifts derived from a muffin-tin potential for rutile VO2. Because the scattering of the photoelectron is determined mainly by core electrons, there is no effect of the vanadium chemical state on the calculated phase shifts. The phase shifts were calculated with the Van Hove and Barbieri program package.15 In the calculations the crystal structure is represented by a cluster, the dimensions of which are determined by a maximum length of the path between the emitter and the scatterer. In the present calculations this length was set to 12 Å. Larger values of this parameter did not produce any significant variation of the calculated XPD curves. III. Results Clean SnO2(110) Surface. XPS and LEIS spectra were measured for the ion-bombarded surface and after annealing at various temperatures. We did not find any significant variation of the O/Sn ratio between the sputtered and annealed surface. Previous studies indicated that the preparation procedure consisting of ion sputtering and annealing in a vacuum produced an oxygen-deficient surface.16-20 Indeed, the composition of the surface after several cycles of sputtering and annealing in a vacuum, as determined by XPS using the sensitivity factors reported in ref 21 for the Sn 3d5/2 and O 1s peaks (SnO1.7), is

Atrei et al. consistent with an oxygen-depleted SnO2 surface. The chemical state of tin in this oxide cannot be derived by XPS since the chemical shift of the Sn 3d peaks and MNN Auger lines is the same in SnO2 and SnO.16-18,12 Indeed, no changes in the peak shape and energy locations of the Sn 3d5/2 (486.7 eV) line and the MNN Auger transitions were observed after annealing the sputtered surface in different conditions. In the initial stages of the present investigation we observed only a (1 × 1) LEED pattern in the whole annealing temperature range, with sharper spots and lower background upon increasing the annealing temperature. This LEED pattern corresponds to the (110) bulktruncated structure of SnO2 (rutile-type structure). Only after several cycles of sample preparation and vanadium depositions could we detect a faint c(2 × 2) superstructure after annealing to about 700 K, and a (4 × 1) superstructure upon increasing the temperature to 900 K. These LEED patterns were observed also in previous works in similar temperature ranges.17,19 No significant changes in the surface composition of the sample, as determined by XPS and LEIS, were detected for the surface showing the (1 × 1) LEED pattern in comparison to the surfaces showing the c(2 × 2) and (4 × 1) superstructures. These different LEED patterns may be attributed to ordered oxygen vacancies, as indicated by a recent scanning tunneling microscopy (STM) investigation of this surface.22 Although the results shown here are for vanadium layers deposited on surfaces showing a (1 × 1) periodicity, tests performed on (4 × 1) surfaces indicate a very similar behavior. XPS and LEIS Results for the As-Deposited Vanadium Films. The deposition of vanadium films was monitored by means of XPS and LEIS. The Sn 3d, O 1s, and V 2p spectra measured for increasing amounts of vanadium deposited are shown in Figure 1a and b. The Sn 3d spectra clearly show that upon deposition of vanadium a component appears on the low BE side of the oxide main lines. The BE energy of this component, determined by a curve fitting procedure, corresponds well to that of metallic tin (484.5 eV).12,21 The Sn 3d5/2 component of the oxide shifts slightly over a range of about 0.2 eV depending on vanadium coverage. This shift is probably the effect of the interface formation. The O 1s peak intensity drops significantly when the vanadium film thickness is increased. The O 1s BE does not change (within ( 0.2 eV) with respect to the value of the clean surface (530.6 eV). The amount of vanadium deposited was estimated from the decrease of the oxide component intensity in the Sn 3d5/2 peak. For this calculation, we used an attenuation length in metallic vanadium of 14 Å, determined using the Tanuma, Powell, Penn method.23 The evaporation rate was found to be 0.4 Å/minute. The evaporation rate is transformed in layer equivalent (LE)/ min. assuming a layer-by-layer growth and that the vanadium film grows with the (110) plane parallel to the substrate surface (interlayer distance 2.0 Å). Because of the reaction and diffusion at the V/SnO2(110) interface (see below), the film thickness determined in this way should be considered only as a rough estimation of the actual value. In the early stages of growth, the main component of the V 2p3/2 peak is located at 515.5 eV. This value does not correspond to metallic vanadium (512.4 eV) but to vanadium in oxides.12,21,24 The V 2p3/2 BE would be compatible with oxides containing V(III) or V(IV) according to the data reported in the literature.12,24 The nature of this oxide phase will be discussed later on the basis of the XPD results. The metallic vanadium component (at 512.5 eV) becomes dominant for thicker films. Because the V 2p1/2 peak overlaps with the O 1s KR3,4 satellite, before further curve fitting analysis the spectrum

Ultrathin Vanadium Films

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Figure 2. Peak areas of the metallic tin, tin oxide, metallic vanadium, and vanadium oxide components in the Sn 3d5/2 and V 2p peaks as a function of the V deposition time.

Figure 1. XPS spectra measured for increasing amount of vanadium deposited. (a) Region of Sn 3d peak. (b) Region of O 1s and V 2p peak. An intensity offset was added to the curves for a better reading. The thickness of the vanadium films is expressed in layer equivalents (LE).

of the clean surface was subtracted, after the appropriate intensity reduction, from those measured for the vanadium-

covered surfaces. From the curve fitting analysis it turns out that the vanadium oxide component is much broader than that for metallic vanadium (the full widths at half-maximum of the V 2p3/2 are 3.2 and 2.2 eV, respectively). This could indicate the existence of more than one oxidation state of vanadium in the oxide phase. However, the BE of the vanadium oxide component (515.5 ( 0.2 eV), as well as its width, remains the same for the whole range of film thickness. The results of the curve fitting analyses are summarized in Figure 2 where the areas of the XPS peaks corresponding to oxidized vanadium, metallic vanadium, tin in SnO2, and metallic tin are plotted as a function of the vanadium deposition time. We found that the fraction of oxide in the vanadium film is larger when the same amount of vanadium was evaporated in a larger number of shorter doses instead of in a single, longer dose. This observation can be explained considering that with smaller doses and during the time needed to collect the XPS and LEIS spectra, vanadium can react with the substrate producing metallic tin and vanadium oxide. The LEIS spectra measured after vanadium deposition are shown in Figure 3. The Sn signal decreases upon increasing the amount of vanadium in the early steps of deposition and reaches a nearly constant value for thicker films. On the other hand, the oxygen signal increases slightly after the deposition of 2-3 LE of vanadium and remains constant for thicker vanadium deposits. The plot of the tin LEIS signal (normalized to 1 for the clean surface) versus the vanadium coverage is shown in Figure 4. The behavior of this curve indicates that the growth mechanism of the vanadium film, as oxide in the early steps and as metal later, is rather complex. If the vanadium oxide formed by the reaction at the V/SnO2(110) interface grows as 2D islands covering the substrate surface, the tin LEIS signal

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Figure 3. LEIS spectra measured for increasing amounts of vanadium deposited. The ion beam was impinging on the sample with an angle of 45° along the [11h0] azimuth and the scattering angle was 135°. The signal intensity is reported versus the E/Eo ratio, where E and Eo are the final energy and initial energy of the ions, respectively.

should decrease linearly as a function of the vanadium coverage and eventually go to zero when the whole surface is covered by the deposited film. For the formation of 3D islands (considering a reasonable shape of the islands) the LEIS signal from the substrate should also be zero when several LE of vanadium are deposited. In the present case, the tin LEIS signal is only halved and it remains constant when more than 10 LE of vanadium are deposited. Therefore, on the basis of the LEIS data we can conclude that the growth mechanism of the vanadium films involves the diffusion of material from the interface onto the film surface. To determine if the vanadium islands are covered by metallic tin or by tin oxide, we measured XPS spectra at normal and grazing (Θ ) 70°) emission angle. The Sn 3d spectra clearly show that the metallic tin component increases on going from normal to grazing emission (Figure 5). This is evidence that the metallic tin, formed by the reaction of vanadium with the SnO2, migrates onto the vanadium islands and partially covers them. The determination of the fraction of vanadium surface covered by metallic tin is prevented by the difficulties in quantifying LEIS data when different matrixes are involved.25 The oxygen LEIS signal visible also for the thicker vanadium films may be attributed to oxygen chemisorbed on the vanadium surface. This oxygen appears to be the product of both the dissociative adsorption of carbon monoxide from the gas phase (a carbon signal was observed in the XPS spectra measured for the deposited films) and of the diffusion of oxygen atoms from the substrate. XPS and LEIS Results for the Annealed Vanadium Films. The as-deposited vanadium films 3 and 6 LE thick were

Atrei et al.

Figure 4. Sn LEIS signal (normalized to the intensity of the Sn signal for the clean surface) versus the amount of vanadium deposited (expressed in layer equivalents).

annealed at 570 and 800 K for 10 min. The thermal treatments of the vanadium films of different thickness led to the same results. The XPS spectra recorded after annealing show that the metallic component in the Sn 3d peaks is strongly reduced (Figure 6, top). For instance, the metal/oxide area ratio for the Sn 3d5/2 is 0.35 and 0.08 for the 6 LE film before and after the annealing, respectively. After annealing, the area of the Sn 3d oxide component increases with respect to that corresponding to the as-deposited film. However, the Sn 3d intensity measured for the annealed surface is significantly lower compared to that of the clean surface. Furthermore, the O 1s peak intensity is lower in the spectrum for the annealed surface with respect to that for the clean surface. The thermal treatment produces the complete oxidation of the vanadium film (Figure 6, bottom). The V 2p3/2 and O 1s BEs in the oxide are 515.5 and 530.6 eV, respectively. The V 2p3/2 BE and width are the same as those determined for the oxide phase formed upon vanadium deposition at room temperature. Note that the total area of the V 2p peaks does not decrease after the annealing, indicating that no significant diffusion of vanadium in the SnO2 bulk takes place. On the basis of the XPS data, it is not possible to distinguish between the formation of V2O3 and VO2 because the V 2p3/2 BE measured for this vanadium oxide phase falls in the range of values determined for V2O3 (515.1-515.7 eV) and VO2 (515.4-516.3 eV).12,24 Moreover, in this range of composition, a series of oxides VnO2n-1 with 4 e n e 8 exists.26 Unfortunately, no XPS data for these oxides are available in the literature. The V 2p3/2 BE of the oxide is close to the value (515.2 eV) measured for vanadium oxide layers obtained by depositing vanadium on the TiO2(110) surface27 in the presence of an oxygen atmosphere. A similar value of the V 2p3/2 BE was determined for submonolayers of vanadium deposited

Ultrathin Vanadium Films

Figure 5. Sn 3d XPS spectra measured at normal (θ ) 0°) and grazing (θ ) 70°) emission angle after deposition of 6 LE of vanadium.

on TiO2(110) and annealed at 473 K. These values were interpreted with the incorporation of vanadium in the titanium oxide matrix.28 A higher V 2p3/2 BE (516.5 eV) was measured for VO2 layers deposited on TiO2(110).7 A weak vanadium signal is visible in the LEIS spectrum measured after annealing (Figure 7). Even taking into account the larger ion scattering cross section of tin compared to vanadium, the LEIS results indicate that only a residual amount of vanadium is present in the outmost layer after the thermal treatment. The O/Sn LEIS signal ratio for this surface is larger (0.16) than the value for the clean SnO2 surface (0.10). This change in the LEIS signal is to be attributed to a structural modification of the tin oxide (i.e., a topmost surface exhibiting a higher amount of oxygen) rather than to the formation of a phase richer in oxygen than the starting surface. The total amount of oxygen, as determined by XPS, is not compatible with the existence of SnO2, assuming either the formation of VO2 or V2O3. On the basis of the XPS results we can conclude that the annealing in a vacuum at temperatures ranging from 500 to 800 K produces the reoxidation of the metallic tin formed during the deposition of vanadium and the oxidation of metallic vanadium. Moreover, the LEIS results clearly indicate that the vanadium oxide phase formed upon the thermal treatment is in the subsurface region of the oxide, covered by tin oxide. No change in the surface composition was observed by XPS and LEIS after an oxygen exposure of 10 min at a pressure of 1 × 10-4 Pa with the sample at 800 K. XPD Results for the Annealed Vanadium Films. The surfaces obtained by room temperature deposition of vanadium do not show long-range order, because the LEED pattern of the clean substrate surface becomes progressively weaker and disappears upon increasing the metal coverage. The annealing

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Figure 6. Comparison of the O 1s, V 2p (bottom), and Sn 3d (top) XPS peaks measured for the as-deposited 6 LE vanadium film and after annealing to 800 K.

process of the vanadium films does not restore the long-range order, because after the thermal treatment only a diffuse background was observed in LEED. To determine whether the oxide phases may exhibit short-range order and/or be epitaxially oriented, XPD measurements were performed. The XPD azimuthal curves of the O 1s, V 2p3/2, and Sn 3d5/2 measured at θ ) 55° for the annealed phase are shown in Figure 8 together with those of O and Sn measured for the clean surface. To improve the signal-to-noise ratio and to compensate for intensity asymmetries, all curves were symmetry averaged exploiting the mirror plane in the [001] direction. The comparison of these two sets of data shows that the O 1s curves of the annealed and of the clean surface are very similar as far as peak positions and anisotropy are concerned. The relatively high modulation of V 2p3/2 shows that the vanadium oxide formed in the annealing process is ordered and epitaxially oriented. The Sn 3d5/2 XPD curve of the annealed phase shows extremely weak intensity modulations. These results rule out the incorporation of vanadium atoms in the SnO2 lattice, producing a V-Sn mixed oxide. If such mixed oxide forms, both the V and Sn XPD curves should show intensity modulations. This result is in agreement with the conclusions of previous works on bulk vanadium and tin oxide mixtures, where it was shown that V-Sn mixed oxide exists only for V atomic fractions less than 0.02.29 Indeed, the weak vanadium signal observed in LEIS may well correspond to the saturation concentration of vanadium in the tin oxide layers at the surface. The decrease of the anisotropy in the Sn 3d XPD curve after annealing indicates that the largest contribution to the Sn 3d peak comes from tin atoms in a matrix which is either amorphous or randomly oriented, the residual

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Atrei et al.

Figure 7. Comparison of the LEIS spectra measured for the 6 LE vanadium film deposited at room temperature and after annealing to 800 K.

intensity modulations being due to the SnO2 substrate. The oxide film at the surface, consisting of layers of disordered tin oxide on top of the epitaxial vanadium oxide phase, appears to be thick enough to effectively attenuate the signal from the SnO2 substrate. This means also that the O 1s XPD modulations are due to oxygen atoms in the vanadium oxide lattice with minor contributions from the SnO2 substrate. We estimated the thickness of the vanadium oxide and of the tin oxide layers utilizing the ratio of the V 2p3/2 and Sn 3d5/2 polar XPD curves. The polar curves measured in the [11h0] and [001] azimuths were averaged to reduce the intensity modulations due to diffraction effects. On the basis of the XPS and LEIS results, it seems reasonable to assume that a continuous layer of SnO covers a layer of VO2, having a thickness dV. The V 2p/Sn 3d intensity ratio as a function of the polar angle θ is given by R(ϑ) )

IV(ϑ)

)

ISn(ϑ)

{

[ ]}

n2λ2 1 - exp k

{

[

n1λ1 1 - exp -

dSn

λ1 cos(ϑ)

dV

]} [ ] [ ] [ dSn

‚exp -

λ2 cos(ϑ)

+ n3λ3 exp -

λ1 cos(ϑ)

dV

λ2 cos(ϑ)

‚exp -

dSn

]

λ1 cos(ϑ)

where k is a constant including the cross-sections for photoelectron emission and instrumental factors; n1, n2, and n3 are the atomic concentrations of vanadium and tin in SnO, VO2, and SnO2, respectively; λ1(17), λ2(19), and λ3(20) are the electron attenuation lengths (in Å) for the photoelectrons in VO2, SnO, and SnO2, respectively. The attenuation lengths were calculated using the Tanuma, Powell, Penn method.23 Because

Figure 8. O 1s, V 2p3/2, and Sn 3d5/2 azimuthal XPD curves measured for the annealed V/SnO2(110) surface. The O and Sn XPD curves obtained for the clean SnO2(110) surface are shown (dashed line) at the top of each panel. The polar angles are measured from the normal to the surface. The azimuthal XPD curves were collected at a polar emission angle of θ ) 55° azimuth. The azimuthal angle φ ) 0° corresponds to the [001] direction of the substrate. The intensity scale of the curves was normalized to improve the clarity. For each of the azimuthal XPD curves the maximum value of the anisotropy defined as ∆I ) (Imax - Imin)/Imax%, where Imin and Imax are the minimum and maximum intensity values in the XPD curve, is reported.

the kinetic energies of the V 2p and Sn 3d photoelectrons are close, we calculated the attenuation lengths for the V 2p and the Sn 3d at the same energy (750 eV, the average of the kinetic energies for the two photoemission peaks). The numerator of this expression represents the intensity from the VO2 layers attenuated by the SnO layers. The first term of the denominator is the intensity of the SnO layers at the surface and the second term is the intensity of bulk SnO2, attenuated by the SnO and VO2 layers. The fitting of the experimental R(ϑ) curve with the calculated curve was accomplished by a least-squares procedure. The best fit was obtained for a thickness of 6 and 14 Å for the SnO and VO2 layers. Considering the uncertainties about the stoichiometry of the oxides, the film morphology, and the calculated attenuation lengths, these figures must be considered simply as an estimation of the thickness of the oxide layers at the surface. Although this slab (consisting of tin oxide and vanadium oxide layers) is thick enough to attenuate effectively the substrate signal, an additional contribution to the reduced anisotropy of the Sn 3d XPD curves may come from a disordered region at the vanadium oxide/SnO2 interface. This disorder can be produced by the diffusion of lattice oxygen at

Ultrathin Vanadium Films

Figure 9. Schematic structural models for the VO2(110) (left) and the V2O3(101h2) surface.

the surface, leaving an oxygen-depleted region below the vanadium oxide layers. For a better characterization of the vanadium oxide phase, the experimental XPD curves were compared with those calculated for various surfaces of vanadium oxides. As basic structural models for the XPD calculations, we considered the structure of bulk-truncated surfaces of VO2 and V2O3 (Figure 9). For temperatures higher than 340 K, VO2 has a tetragonal structure (rutile-type) whereas for lower temperatures a monoclinic phase is stable.30 Considering the similarity of the XPD curves of Sn in SnO2(110) and V in the oxide phase, VO2(110) appeared to be a good candidate for the XPD calculations. The differences of the lattice parameters between SnO2(110) and VO2(110) are about 4% and 11% in the [11h0] and [001] directions, respectively.30 The other structural model tested in the XPD calculations was the (10-12) surface of V2O3 (having the corundum type structure). Although the lattice parameters of the V2O3(101h2) (a ) 5.105 Å, b ) 5.646 Å) surface do not match those of SnO2(110), this orientation of the crystal was chosen because it has a rectangular unit cell and is the cleavage plane of V2O3.8 In the calculations we considered clean surfaces of the oxides, although the measurements were performed for vanadium oxide covered by disordered tin oxide layers. Thus, we assumed that the effect of the disordered tin oxide layer is only the attenuation of vanadium and oxygen (from the vanadium oxide layers) signals.

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Figure 10. Comparison of experimental azimuthal XPD curves with the SSC calculations performed for the V2O3(101h2) surface (top) and the VO2(110) surface (bottom). In the calculations the V2O3 cell is oriented in such a way that the longer side of the unit cell is parallel to the [001] direction of the substrate, whereas the VO2(110) cell is oriented in the same manner as that of the SnO2.

The results of the calculations for V2O3(101h2) and VO2(110) are shown in Figure 10. The XPD curves calculated for the VO2(110) model provide a sufficient level of agreement with the experimental intensities, whereas the model corresponding to the V2O3(10-12) surface can be ruled out (Figure 10). Furthermore, for the polar XPD curves measured in the [1-10] and [001] azimuth the best agreement is obtained for the VO2(110) model.5 Although V2O3 could be oriented along different crystallographic planes, this possibility appears unlikely because the (0001) and the (101h2) surfaces are the most stable orientations of this oxide. On the other hand, the V2O3(0001) surface is not compatible with the rotational symmetry of the experimental XPD curves. As refinement of the VO2(110) model, we set the sides of the VO2 unit cell equal to those of the substrate (to simulate a pseudormorphic film) and varied the interlayer distances. The variations of these structural parameters did not lead to a better agreement between experimental and calculated XPD curves, as shown by the visual comparison and by the reliability factor (R-factor) analysis performed using different kinds of R-factors.31 The discrepancies between the experimental and calculated XPD curves can be explained with the formation of a defective vanadium oxide VOx≈2 with a structure close to that of the rutile VO2 phase. The vanadium oxides VnO2n-1 with 4 e n e 8 have large unit cells but the local structure around the vanadium atoms is a distortion of that found in rutile VO2.32 Hence, the XPD results

3128 J. Phys. Chem. B, Vol. 104, No. 14, 2000 suggest that one member of this series of oxides may be the product of the thermal treatment of vanadium films on SnO2(110). The presence of such an oxygen-deficient phase would explain also why the V 2p3/2 BE is lower with respect to the values for stoichiometric VO2 surfaces.24 IV. Discussion An important aspect in determining the growth mechanism of a metal overlayer on an oxide surface is the effect of interfacial reactions. Metals with a higher affinity toward oxygen than the metal in the oxide substrate should grow as oxide layers and at the same time cause the partial reduction of the substrate. Because of the strong interaction at the interface, during the initial stages of growth the oxide film is expected to form 2D islands covering the substrate surface. On the other hand, less reactive metals are expected to grow as 3D clusters. This picture turns out from systematic studies of metals (including vanadium) having different affinities for oxygen deposited on TiO2(110).9,27,33 For V/SnO2(110), in addition to the interfacial redox reaction, it is necessary to consider the possibility of migration of metallic tin from the interface onto the film surface. Our results do not rule out the possibility that, in the early stages of metal deposition, the vanadium oxide film grows as 2D islands. However, the presence of tin in the topmost layer alters the curve of the Sn LEIS signal versus vanadium coverage (which would be proportional to the fraction of the free substrate surface if no substrate material diffusion occurs). The growth of metallic vanadium islands follows the formation of vanadium oxide upon an increase in the metal coverage. The presence of metallic tin in the topmost layer of thicker vanadium film indicates that tin acts as a surfactant diffusing through the defect of the film and partly covering its surface. Note, however, that the present results concern the deposition of vanadium layers on an oxygendeficient surface prepared by sputtering and annealing and a different behavior could be expected on a stoichiometric SnO2 surface. For instance, the amount of metallic tin produced by the interface reaction between SnO2 and Pd increases with an increase in the oxygen deficiency of the substrate surface.3,10,11 However, the “in situ” preparation of a stoichiometric SnO2 surface requires oxidation in 100 Pa of oxygen at 700 K,16,17 something that was not possible in the experimental setup of the present study. The thermal treatment of the vanadium films produces the reoxidation of metallic tin and the complete oxidation of vanadium. The oxidation is carried out by lattice oxygen atoms diffusing from the bulk of SnO2. Thus, a region of oxygen vacancies must be present in the bulk of the substrate oxide, beyond the depth probed by the photoelectrons. The present results indicate that SnO2 and vanadium oxide do not form a mixed oxide but remain as separated phases at the surface, with the vanadium oxide phase covered by a layer of composition SnOx≈1. The weak intensity modulations of the Sn XPD curves after annealing would be compatible with a tin-oxide monolayer on top of the vanadium oxide phase. The tin atoms in the 2D oxide should not give any detectable intensity modulations because the forward scattering intensity enhancement along the emitter-scatterer direction can be observed only at grazing emission angles. In this model, a disordered tin oxide layer should be present at the vanadium oxide/SnO2 interface to attenuate the contribution from the SnO2 substrate to the XPD modulations. However, the XPS results indicate a thickness of the tin oxide layer significantly larger than one monolayer. Thus, the most plausible model consists of layers of amorphous or

Atrei et al. randomly oriented SnO. The in-depth distribution of the oxides may result from the oxidation of metallic tin which is located at the surface of the vanadium film. However, the LEIS spectra measured for vanadium oxide-SnO2 films used in gas sensing devices also indicate that a very small amount of vanadium is present on the surface.34 These results provide some fundamental data on the composition of mixed tin/vanadium oxide surface which may be used as a basis for the understanding the effect of selectivity increase in vanadium oxide/SnO2 sensors. Because we found only traces of vanadium in the outermost surface layer on both real sensors34 and model systems (present work), we can speculate that the tin oxide layer covering the vanadium oxide layers is the active species in determining the higher selectivity of the gas sensing surfaces. V. Conclusions The growth mechanism of vanadium films on SnO2(110) is characterized by an interfacial reaction which occurs in the initial stages of deposition, leading to the formation of vanadium oxide and metallic tin. Upon further deposition we observed the formation of a metallic vanadium overlayer. The mechanism of growth of the vanadium films involves also the diffusion of metallic tin onto the film surface. The main results for the surface obtained after annealing of the vanadium films at 800 K show that the thermal treatment leads to the reoxidation of metallic tin and to the complete oxidation of metallic vanadium, with no diffusion of vanadium in the SnO2 bulk. We can conclude that the thermal treatment does not produce a vanadium-tin mixed oxide but a vanadium oxide phase covered by layers of tin suboxide. The vanadium oxide phase, having a structure close to that of rutile VO2, is epitaxially oriented with the (110) plane parallel to the substrate surface. Acknowledgment. This work was supported by the Ministero della Ricerca Scientifica e Tecnologica (MURST) through the fund “Programmi di Ricerca di rilevante interesse nazionale” and by the Consiglio Nazionale delle Ricerche (CNR). References and Notes (1) Semancik, S.; Cavicchi, R. E. Appl. Surf. Sci. 1993, 70/71, 337. (2) Azad, A. M.; Akbar, S. A.; Mhaisalkar, S. G.; Birkefeld, L. D.; Goto, K. S. J. Electrochem. Soc. 1992, 139, 3690. (3) Fryberger, T. B.; Semancik, S. Sens. Actuators, B 1990, 2, 305. (4) Angelucci, R.; Boarino, L.; Cardinali, G. C.; Cavani, F.; Critelli, C.; Dori, L.; Parisini, A.; Pizzochero, G.; Poggi, A.; Trifiro`, F. Thin Solid Films 1997, 29, 743. (5) Atrei, A.; Bardi, U.; Cortigiani, B.; Rovida, G. Surf. ReV. Lett., in press. (6) Egelhof, W. F. In Ultrathin Magnetic Structure I; Bland, J., A. C., Heinrich, B., Eds.; Springer-Verlag: Berlin, 1994. (7) Sambi, M.; Sangiovanni, G.; Granozzi, G. Phys. ReV. B 1997, 55, 7850. (8) Henrich, V. E.; Cox, P. A. The surface science of metal oxides; Cambridge University Press: Cambridge, 1994. (9) Diebold, U.; Pan, J.; Madey, T. E. Surf. Sci. 1995, 331, 845. (10) Semancik, S.; Fryberger, T. B. Sens. Actuators, B 1990, 1, 97. (11) Cavicchi, R.; Semancik, S. Surf. Sci. 1991, 257, 70. (12) Handbook of X-ray Photoelectron Spectroscopy; Chastain, J., Ed.; Perkin-Elmer: Eden Prairie, MN, 1992. (13) Sample prepared by the Surface Preparation Laboratory, FOM, Amsterdam, Netherlands. (14) Fadley, C. S. In Synchrotron Radiation Research: AdVances in Surface Science; Bachrach, R. Z., Ed.; Plenum: New York, 1991. (15) Van Hove, M. A.; Barbieri, A. Phase shift package, private communication and http://electron.lbl.gov/software. (16) Cox, D. F.; Fryberger, T. B.; Semancik, S. Phys. ReV. B 1988, 38, 2072. (17) Cox, D. F.; Fryberger, T. B.; Semancik S. Surf. Sci. 1989, 225, 121. (18) Egdell, R. G.; Eriksen, S.; Flavell, W. R. Surf. Sci. 1987, 192, 265.

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