Copper and nickel ultrathin films on metal-oxide crystal surfaces

Science of Ceramic Interfaces II J. Nowomy (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

473

Copper and nickel ultrathin films on metal-oxide crystal surfaces Preben J. Moller* Department of Chemistry, University of Copenhagen, DK-2100 Copenhagen 13, Denmark

Universitetsparken 5,

Abstract The synthesis and in situ characterization of ultrathin films of copper and nickel on single-crystalline surfaces of metal oxides, and in particular the development in the epitaxial and electronic structures of the metal particles and films during atomic layer epitaxial growth, are reviewed. We consider recent results for lowindex surfaces ofa-AleO s, CaO, a-Fe2Os, LaA1Os, MgO, NiO, SrTiOs, TiO2, ZnO and yttria-stabilized ZrO2, using low-energy electron diffraction and electron spectroscopies, and discuss the problem of electron-impact induced surface charging. Analysis by combined electron spectroscopy and photodesorption methods of the reactivity of the metal deposits is briefly discussed for the cases of CO exposure to Ni-deposited TiOe and to Cu-deposited ZnO crystal surfaces.

1. I N T R O D U C ~ O N Studies of single-crystal surfaces of the metal oxides are of fundamental importance in the understanding of a wide range of both basic and applied research, and the interfaces of these surfaces with metal particles and thin films are of decisive importance in many areas of chemistry and materials science: such as in heterogeneous catalysis, e.g. the synthesis of methanol or the oxidation of a range of hydrocarbons; in ceramics, e.g. joining, and the synthesis and reactivity of high-Tc superconductors and of electrical contacts on these, or the synthesis special layered sandwich constructions and composits; in micro-electronics, e.g. electronic-packaging systems and the formation of metal contacts to the oxide layers whose properties may be of semiconductor, insulator or even metal nature, or in synthesis and application of selective gas sensors; in metallurgy, e.g. metal corrosion resistance. In order to u n d e r s t a n d well and, in particular, to control and in some cases to improve the conditions of many of these processes it is fruitful to carry out studies on well-defined single crystal surfaces and their interfaces with the metal deposits. With well-defined surfaces we mean surfaces that reproducible in atomic scale is determined with regard to both geometric, i.e. two-dimension-ally crystallographic, and electronic structure. Not least due to its importance in many large-scale

474 industrial applications this type of research has gained a strong m o m e n t u m over the last decade. The structure and properties of the metal oxide surfaces have been discussed previously in a review [1], metals on oxides were reviewed in 1987 [2], and recently the bulk properties of the transition metal oxides have been reviewed [3]. In the present review we will discuss recent experimental results on geometric and electronic structures of clean MO, M02, M203 and MaMbO 3 metal-oxide crystal surfaces and of the adsorption of copper and nickel on a range of metal-oxide single crystal surfaces, and also give some examples from studies on the reactivity of these surfaces toward carbon monoxide. After a brief section on experimental methods (Sect.2) and a discussion of e-beam induced surfaces charging (Sect. 3) we consider the following six crystal classes: (i) rocksalt (Sect.4), (ii) rutile (Sect. 5), (iii) corundum (Sect. 6), (iv) perovskite (Sect. 7), (v) wurtzite (Sect. 8) and (vi) fluorite (Sect. 9).

2. EXPERIMENTAL METHODS The experimental methods that were used in experiments that we discuss here were all ultra-high-vacuum (UHV) based, i.e. at base pressures below 10 .7 Pa. The base pressure was usually 6x10 9 Pa. Low-(and very-low)-energy electron diffraction (LEED and VLEED, respectively) were used for the characterization of the surface geometrical structure. Very recently the first results were obtained by scanning tunnel microscopy, for the rutile surface, though. That method will give much wanted and new light on the surface geometry. There are severe difficulties, though, due to the fact that most metal oxides are hardly good conductors. The cleanness, stoichiometry [4] and the electronic structures in detail were elucidated by Auger (AES), electron and highresolution electron energy-loss (EELS and HREELS, respectively), laboratory (ultraviolet (UPS) and X-ray (XPS)) and synchrotron-radiation-based photoemission (PES) [5] and target current spectroscopy (TCS) [6], and finally, in some gas-reactivity measurements, temperature-programmed desorption (TPD) and (photo- or laser-induced) desorption mass spectrometry. The metal was deposited in submonolayer-steps at a rate of 1/~Jmin either from an electron beam evaporator or from a Knudsen cell. In all cases the thickness d of the deposited metal layers was monitored by an oscillating quartz crystal microbalance (QCM) and calibrated by AES. For all the above methods there is often much difficulty, as compared to metal crystals, in carrying out experiments on the metal oxides and on submonolayermetal-covered metal oxide surfaces, due to several reasons, the most serious being surface charging problems, particularly for the large-band-gab semiconductors and insulators. Quite often the threshold incident electron energies are low. The range of usable electron energies is quite restricted, therefore. Secondly, the concentration of defects at most of these surfaces is considerably larger t h a n for metals. These experimental difficulties have probably been major reasons for the later development of the surface chemistry for the oxide crystal surfaces. We have succeeded in obtaining AE and EEL spectra and LEED patterns for many large-

475 band-gap metal oxides by creating, by flash, a small (not AES-surfacestoichiometry-detectable) amount of oxygen vacancies in the surface, giving effective positively charged holes t h a t can receive the electrons and thus give sufficient conductivity to allow for the AES, EELS and LEED experiments. For most metal oxide surfaces it is necessary to heat the samples to high temperatures in order to obtain clean and ordered surfaces. The heat t r e a t m e n t may be carried out in air if one makes precautions not to expose the surface to moisture while mounting in the UHV (using flow of nitrogen). When under UHV the surfaces are flashed in s i t u to ~ ~ ~ T--rcleanness and surface crystallographic order using a 1 sharply focussed beam of light from an external Xe lamp. 3. E I , E ~ N - B E A M

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INDUCED CHARGING

Let us start the discussion of the experimental inve stigatio ns of me tal-o n- me tal-oxide syste ms with an investigation of the troublesome charging-up behavior of many of the oxide surfaces upon incidence of an electron or light beam, the former yielding a negatively charged surface upon scattering, the latter a positively charged surface. Positive charges are fairly easily removed by electron flooding with a mild cloud of electrons from a 'flood gun' while the negative charges may not similarly be removed as easily (by positive ions) without damage. In some cases [7] negative charged may be removed by continuously using a second electron gun (operating at 1-2 keV) at glancing incidence and thereby changing the "secondary emission crossover", the point where the ratio of the secondary-electron-emission current to the incidentelectron-beam current passes one. We have found it possible to carry out electron diffraction and spectroscopy if the surfaces initially is flashed to high temperatures to create a few (oxygen) vacancies whereby it becomes possible for the electrons to tunnel away to a ground connector. In Figure 1 we demonstrate a familiar behavior of an AE spectrum upon irradiation with an electron beam [8]. Here the electron beam is incident perpendicularly onto the MgO(100) target surface at room temperature with an energy E, of 3.1 keV. The sample had been cleaved in air and then immediately, under

1000 1200 1400

Electron energy (eV)

Figure 1. AES from Cu/MgO(100). See text.

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Figure 2. Au~ger surface potential shift with Ep. A: clean MgO(100); [3:30 A Cu; o" 60 A Cu; o: isolated holder. nitrogen, mounted under UHV where it was flashed with the beam of light. The spectra were recorded when sufficient time (a few minutes) necessary to reach an equilibrium (saturated) potential shift had elapsed. First we note t h a t there is no charging problems for the clean surface (curve 1), but upon deposition of small amounts of Cu onto the surface we note t h a t the spectra are shifting towards higher energies (curves 2-5). We can thus follow a shift of the Cu(849eV) peak. At an average Cu thickness dcu of about 15 A (1/k - 0.1 nm), corresponding to about 4 monolayers (ML) of Cu, the Mg and 0 peaks disappeared. This observation agrees reasonably w e l l with the known value of the electron escape depth in Cu which is about 10 A. We have also found t h a t a threshold value of Ep exists above which the charges build up, measured as the potential shift of the spectra, and that no shift was found for Ep < .1 keV (Figure 2). On the other hand, it is necessary t h a t Ep is sufficiently large to produce enough emitted electrons from the surface. We have observed similar behavior on other metal-on-metal- oxide systems as well, such as Cu and Y on LaAI03 (100) [9]. An exponential relationship was found for the shift with de,. When the primary electron beam is aimed at a fixed particular point in the film, the shift is not very stable with beam exposure time; a dynamical behavior of the shift in potential is revealed although that change in the shift is small in comparison to the uncertainties of the data in Figure 2. As to the threshold in E, (Figure 2) we know t h a t electrons will penetrate deeper into the solid the higher Ep . We suggest t h a t the surface will be loaded to the critical threshold value, about 1.6 keV for MgO (and 1.4 keV for LaA10s [9]), consistent with results for many insulating materials [10]. Since the insulator MgO has an energy gap E~ of 7.8 eV, a work function of 3.14.4 eV and the vacuum level located in the gap, there are no states available to hold the arriving incident electrons which then are scattered from the surface,

477 thus explaining the lack of charge buildup on the clean MgO surface even for an Ep of 5 keV. We have explained charging-up results for metal deposition on metal oxides by a three step model [8]. (Step 1): Strong surface charging during ultrathin (a few tenths of ML coverage) deposition of Cu on MgO(100) and MgO(111) surfaces due to electron trapping centers in the interface (as explained by Cu 3d impurity levels located in the MgO band gap). (Step 2): Building up continuous Cu films which can accommodate local conduction bands. (Step 3): When reaching thicker layers (in the order of hundreds of/k and above), charging-up as for an isolated metallic target exposed to an incident electron beam.

4. COPPER AND NICKEL ON ROCKSALT METAL-OXIDE STRUCTURES

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Figure 3. Changes in Cu(LMM)/ O(KLL) and Cu(LMM)/Mg(KLL) AES intensity ratios with dcu.

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The high-symmetry cubic rocksalt crystals have the advantage that they cleave very easily and often give cleavages of high quality, i.e. having only relatively few defects. The ideal surface is non-polar (charge neutral) and atomically fiat. Particularly the basal (100) plane of these cubic crystals are often close to ideal for 8surfaces studies. It is terminated at surface in a simple bulk truncation. It is preferable if cleavages can be carried out in the UHV but for the rocksalt crystals, particularly MgO and NiO, there is only little difference compared to surfaces cleaved in air. One should take care, however, to minimize exposure to air (interaction with water vapor ) and therefore make the cleavages in a N2 box. We have carried out synthesis and combined AES-EELS-XPSLEED characterization of ultrathin Cu or Ni films on four rocksalt-type crystals: The alkaline-earth d o oxide surfaces MgO(100) [8,11-13], M g O ( l l l ) [8] and CaO(100) [14,15] were exposed to Cu, and the transition-metal oxide surface NiO(100) to Ni [16]. The system

478 Cu/MgO(100) has also been studied previously by electron microscopy and diffraction [17-18], and recently by AES and HREELS [19]. 4.1. Cu r MgO(100) a n d MgO(111) Magnesium oxide is an insulator with a wide bandgap of 7.8 eV (bulk) [1] and 7.0 eV (surface) [20], and it has a lattice constant of 4.21 A. In section 3 above the preparation of the surfaces of this crystal is described. During deposition of submonolayer amounts of Cu at room t e m p e r a t u r e (RT) onto well-defined (1• (i.e. simple bulk-truncated) MgO(100) and MgO(111) surfaces [11], the originally sharp substrate LEED-pattern becomes gradually blurred until the substrate p a t t e r n is eliminated at a dcu of about 8/k. When dcu has reached about 15/k, a week Cu(100) superstructure p a t t e r n appears for the case of freshly air-cleaved MgO(100) and a, also weak, Cu(111) superstructure for the case of the fresh MgO(111) surface. The p a t t e r n s sharpened with increasing amounts of Cu. In Figure 3 is shown the gradual change in chemical composition, i. e. of the AES signal intensity ratios divided by their respective AES sensitivity factors, of the surface layer with the Cu layer thickness dcu. The growth behavior indicates a Stranski-Krastanov mechanism: two stages, a monolayer followed by 3dimensional islands. This growth mode was later confirmed by Conrad et al. [ 19]. Large spectral changes are observed by EELS during deposition of Cu particles. Strong resonance peaks at 2.6 eV for MgO(100) and at 2.2 eV for M g O ( l l l ) were both eliminated during deposition of submonolayer deposits of Cu, and two new peaks appeared 2.2 eV and 4.5 eV. We have shown [21] t h a t this most probably is due to binding of the first small Cu deposited particles to the oxygen ligands as ions, sticking to Mg 2§ vacancy sites causing the surface defect peak to disappear and two new resonance peaks which have character of cuprous oxide to appear due to Cu impurity-level related electronic transitions. The MgO(111) spectrum only shows a small difference in the 6.1 eV surface-state related peak in comparison to the spectrum of the clean MgO(100) surface. The intensity of the surface-defect peak is a little weaker and the surface-state related peak a little stronger. We also note t h a t the initial deposition has the same influence on the EEL structure as does the change from the (100) to the polar and Mg-richer (111) surface, supporting our claim t h a t the initial Cu deposits have a high possibility of interaction with Mg § sites, i. e. the Vs color centers (holes trapped at a metal ion vacancy, with 5 oxygen ligands in the case of the MgO surfaces). We should point out, though, t h a t the finding of the Vs being the physical origin of the electronic loss structures by no means implies t h a t Fs§ centers are non-existing on oxide surfaces. Already at average coverage near 5 / k the spectra show some characteristics of bulk copper energy-loss features, and at higher coverage, above 15 /k, corresponding to a few monolayers, the spectrum is a bulk copper one. Furthermore we found t h a t the epitaxial copper film after oxidation by CO agrees very well with t h a t from oxidized copper crystals. We have also investigated the effect of heat t r e a t m e n t and reoxydation. Figure 4 [12] thus demonstrate for the Cu/MgO(100) system the effect of oxidation of a previously flashed surface (curve 2), deposition of a very small a m o u n t (0.5/k ) of Cu (curve 3) and the result of heating the latter surface, with the Cu on (curve 4). We see t h a t the loss structure at 2.6 eV still remains after heat t r e a t m e n t in 10 .5 Torr oxygen for 2 h at 430 ~ discarding the loss as an origin of a Fs§ center.

479 Deposition of 0 . 5 / k on MgO(100) produces loss spectra at 2.2 and 4.2 eV, and removes the peak at 2.6 eV. To give further evidence for our above claim of oxygen-bonded copper deposits the latter surface was heated to 300 ~ (curve 4). The 2.2 eV peak was enhanced ( further heat treatment to 410~ for 15 rain did not change the spectrum further) and consequently leaves out the possibility that the peak has originated from Cu clusters which have electronic characteristics different from bulk Cu. We believe that the peak is the well-known loss structure of the semiconductor Cu20, corresponding to a threshold transition from the Cu 3d valence band to the Cu 3s conduction band, in agreement with the above

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Figure 4. EELS (at Ep= 97 eV) from Cu/MgO(100). (1) clean, 400~ in UHV, (2) clean, 02 at 430~ (3) 0.5 ,s Cu on (2)-surface, (4) the (3)surface heated to 300~ for 15 rain.

480 observation that the spectrum after oxidation corresponds to oxidized Cu. The 4.2 eV peak may be due to the O 2p band to the Cu 3s levels. Furthermore, the above discussed electron trapping by copper deposits also supports the ionized-copperstate mechanism. These experimental results, supported for the (100) surface by theoretical calculation [22], clearly indicate that magnesium vacancies on the (1• 1)MgO(100) surface are trapping centers for the initial copper deposits, and it is hence reasonable to expect that these trapped copper ions are active sites for a subsequent nucleation, so that the copper atoms bind to the oxygen ions on the surface in a bridge configuration subsequently leading to a Cu(100) epitaxial growth through formation of clusters around the active center. Comparing to previous spectroscopic work on Cu(I) and Cu(II) oxides [23] we attribute the appearance of the 2.2 and 4.5 eV peaks in the 0.7/k Cu/MgO(100) EEL spectrum to Cu(II) states fitting into the Vs centers, and propose that the copper atoms forming the bridge configuration to the oxygen ions are in the Cu(I) state. This conclusion was confirmed by XPS measurements in which we used examination of the Modified Auger parameter d of copper (the sum of the 2ps~ binding energy and the Auger LMM kinetic energy), during the submonolayer growth [11], and also by Conrad et al. [19] using HREELS. For the (1• 1)MgO(111) surface we similarly find that the cation vacancies will be occupied by the copper bounded to three anions and thereby forming nucleation centers for Cu(111) epitaxial growth. The above proposed epitaxial growth was confirmed [11] by (somewhat diffuse) LEED patterns, indicating partial epitaxy. 4.2. N i cm MgO(100)

Ultrathin-fi]m deposition of nickel onto MgO crystal surfaces has been investigated by reflection high-energy diffraction (RHEED) [17] and by electron microscopy imaging (moir~ fringes and dark field images) and selected-area diffraction techniques [18]. A preferred epitaxial relationship was concluded [17] for both Cu and Ni deposits as Cu, Ni(110) [111] II MgO(100)[011]. 4.3.

Cu ~m

CaO(100)

The other alkaline-earth oxide crystal we will look at is calcium oxide. It is also an wide-bandgap insulator. It has a bandgap of 7.0 eV (bulk) [1] and 6.2 eV (surface) [20] (the bandgaps get smaller down through the group MgO, CaO, SrO, BaO) and a lattice constant of 4.81 A (the lattice constant increases in the group with the size of the metal ion). The sample was heated in air to 900~ before mounting in UHV, and after flashing with the light-beam a sharp (1• 1) LEED pattern from an AES clean CaO(100) is obtained whereupon EELS were carried out at Ep > 116 eV [14,15]. LEED was previously obtained for the clean CaO(100) surface, and 1% contraction between the outermost two layers and a vertical displacement of the Ca atoms with regard to the O atoms in the surface of less than 2 % were found [24]. Also EELS was obtained for the clean surface [25] but the spectra were only partly interpreted. After the appearance [26] of a theoretical determination of the density-of-states diagram for CaO, Figure 5, we are now able to assign interpretation [14] to the energy losses obtained. We thus assign losses at 5.2, 8.0, 10.0, 18.5, 26, 28.5, 31.0 and 36.4 eV to interband transitions, and losses at 14.2

481 eV to a surface plasmon (cos) and 28.5 eV to 2 cos The losses at 5.2, 8.0 and 14.2 eV were previously assigned (1,24]. We do not agree, however, to a previous [24] assignment of the loss at 36.4 eV to a volume plasmon (cop) but assign to a Ca 3pto- 3d interband transition[15]. Upon deposition of copper at RT onto CaO(100) the 5.2 eV loss is eliminated already after 2/~ of deposited Cu, and a strong loss at 4.3 eV appears in stead. v-bands

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Figure 5. Densities-of-state diagram from CaO. From ref. [26].

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Figure 6. Growth of Cu on CaO(100). See text.

Cu

The other losses of the substrate become attenuated. The surface plasmons at 14.2 and 28.5 eV are strongly attenuated, supporting the interpretation as surface plasmons. The interpretation much resembles t h a t for the Cu/MgO(100) surface. The mode of growth was studied boy AES for depositions 0 < de, < 10 A. For clean CaO surfaces there are no charging-up when using the s t a n d a r d Ep of 3 keV, but for Cudeposited surfaces it was necessary because of the charging to reduce Ep to less t h a n 2 keV. As we see from Figure 6 [14,15], which shows the change in Cu and Ca AES intensity with clc,, a clear break occurs at 2.1 /k, corresponding to completion of one ML. At higher coverage, the substrate intensity is higher and the

482 Cu intensity lower t h a n what would correspond to a F r a n k - van der Merwe (layer-by-layer) growth, hence we have a Stranski-Krastanov (1 or 2 ML followed by 3D-islands) growth mode for this system. In an alternative plot [14,15] the change in the calcium and oxygen intensities (both normalized with respect to the signal from the clean surface) with the Cu intensity is given. Here the break is also clearly shown. A calculation [15] of the thickness of the Cu in the first layer from these a t t e n u a t i o n curves, based on the known inelastic m e a n free p a t h for electrons in Cu, shows agreement with the values of dcu determined by the quartz crystal microbalance. When the surfaces with the deposited Cu layers is heated to 250 ~ coalescence of the Cu particles occurs. They move together on surface, forming individual islands in the equilibrium stage, hence they now after annealing show a VolmerWeber (pure 'islanding') growth mode, demonstrating the difference in diffusion of the Cu atoms with substrate temperature. 4.4. N i cm NiO(100)

Nickel oxide is also a good insulator, a 3d s transition-metal oxide with a band gap of 3.8 eV, and as with the previous rocksalt alkaline-earth oxides the crystal surfaces charge up upon electron irradiation. And also, like these, it cannot be reduced by chemical reduction. But again, high-temperature annealing, in this case to 600 ~ of the front surface as mounted in UHV, produces a sharp (1• LEED p a t t e r n as shown in Figure 7a [16]. In AES it is necessary, though, to reduce ED to 2 keV like for the CaO case. We have studied the growth of Ni on (1• in the 0 < c~i < 145 deposit range at substrate t e m p e r a t u r e s between 125 and 185 ~ At RT, no LEED could be obtained due to charging (the Ep threshold energy required to produce enough emittance of neutralizing electrons from the surface rose from 120 eV to more t h a n 300 eV with dNi), and at t e m p e r a t u r e s higher t h a n 215 ~ no Ni(100) p a t t e r n could be obtained with increasing the deposit dNi up to 80 A, probably due to decreasing sticking probability of the Ni deposits with increasing temperature. A Ni superstructure appeared at a deposit of about 20 ~,, and it became more clear with increasing c~i. Figure 7(b-d) [16] shows the overlayer structure at 145 A of Ni which we have attributed to a Ni(100) layer whose growth direction is parallel to the Ni[010] and [001] directions along NiO[010] and [001], respectively. A comparable result was earlier obtained by reducing a NiO(100) in H 2 and obtaining Ni(100) islands [27]. The epitaxy of the Ni overlayer can be described by using a site-coincidence preference or axial commensurate growth model [16, and refs. therein] in which every 6 unit cells of the Ni lattice are in registry with 5 substrate unit cells along a preferred direction (Figure 8). The strong modulation of the surface potential of the NiO(100) substrate forces Ni atoms to line up along its symmetrical directions, resulting in misfit dislocation of the Ni layer at the interface. The mode of growth was studied by over the 0 < d~ < 16 A range at substrate t e m p e r a t u r e s in the 20 to 135 ~ range. Figure 9 shows the AES results for 20 and 70 ~ respectively. As seen from the figure, the growth follows the Stranski-

483

Figure 7. Ni epitaxy at 460~ and Ni[001] II NiO[001.

Ni(100) IINiO(100), with Ni[010] IINiO[010]

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Figure 8. Misfit dislocation at a Ni/NiO(100) interface, viewed along [110]. Atoms in the (110) plane: filled symbols. Atoms in next-lower layer: open symbols.

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Figure 9. Change in Auger O(KLL) intensity from Ni/NiO(100)with dNi at 293 and 343 K. (curve b was lifted 20 units for clarity). See text

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Krastanov mechanism with a monolayer point at d~i ~ 2/k [16]. In the range 2
5. COPPER AND N I C K E L ON RU'I'H~ STRUCTURF~ Rutile, one of the (three) naturally occurrmog TiO 2 minerals, is tetragonal with lattice dimensions ao=4.594 A and Co=2.958 A. Stoichiometric TiO2 is a wide-(3.1 eV bulk, 1.9 eV surface [20])-bandgap 3d ~ transition-metal oxide, and is in its stoichiometric form a good insulator. By surface h e a t - t r e a t m e n t or by Ar + sputtering [ 1] oxygen vacancies may be created, causing it to be conducting n-type [51], and the surface t i t a n i u m 3d orbitals be partially populated. Due to applications in photocatalysis and as gas sensor it has been much studied. The (110) surface is - due to the high degree of coordination of the metal ions by oxygen - the most stable of its surfaces, and is ideal for surfaces studies of the crystal metal oxides. Its EELS structure is well known [1,29-31], and so is its low energy part of its H R E E L spectrum [30], and its valence-band structure along the F-Z-M line has recently been studied by angle-resolved photoemission with a comparative bulk-band-structure calculation [32], and most recently by angleresolved inverse-photoemission [33] which suggests t h a t the nearly perfect (110) surface is t e r m i n a t e d by a bridging oxygen layer. Despite the stability, Ti02(110) surfaces may nonetheless be t u r n e d into a (1• reconstruction upon heating in UHV, and we have found [34] t h a t there is a

485

Figure 10. LEED from TIO2(110). (a) 550~ for 3 hr --> lx2 reconstrution; (b) 700~ for 1 hr -~ l x l surface; 550~ for +30 rain --> coexistence of ( l x l ) and (lx2) surface structures; (d) + 80 rain; (e) +30 rain (total 140 rain) (ref. 34).

reversible relation between the (lx l) and (lx2) surfaces (the (lx2) TiO2(110)was first seen Kao et al. [35]). As shown in Figures 10a and 10b, respectively, a low

annealing temperature (550 ~ produces a (lx2) structure at the surface, while further annealing at 700 ~ gives a perfect unreconstructed ( l x l ) surface with sharp and bright integral LEED spots. Charging-up is avoided on this otherwise insulating sample since the slight reduction by heating in UHV increases the

486 conductivity, a procedure t h a t is equivalent to n-type doping [34]. The brightness of the weaker half-order spots are dependent on the annealing time (Figure 10c-e). In the initial period of the annealing, a weak line linking the bright spots is

Figure 11. Ball model for TiO2(ll0)-lx2.Large balls: oxygen, small balls: t i t a n i u m (ref. 34).

observed. After extended annealing over hours the lines have changed via coexistence of spots and lines into the haft-order spots. This behavior can be interpreted as coexistence with (lx2) domains at the surface, leading us to the proposal t h a t the (lx2) can be considered as a missing-row model as illustrated in Figure 11. We explain the process as follows. During the annealing there is competition between diffusion of bulk-lattice oxygen and desorption of surface oxygen, i.e. a t h e r m a l equilibrium. That the desorption of the oxygen anions is favored may be due to the fact that each protruded surface oxygen anion is coordinated with two sublayer titanium cations instead of three t i t a n i u m cations as in bulk TiO2, hence some surface oxygen may escape, forming (lx2) domains. At higher t e m p e r a t u r e s the diffusion is faster and reach equilibrium with desorption, resulting in a (1• 1) surface structure. This surface crystallographic arrangement, postulated on the basis of the electron spectroscopy and the LEED results, may be further tested by atomic-scale scanning tunnel microscopy measurements. This missing-row model is supported by corresponding changes in the Auger intensity ratios, and also by the appearance of a Ti s§ state in the bandgap as shown by UPS at 1 eV binding energy, Figure 12. The more (lx2) domains, the more Ti s§ intensity, which makes sense, since when a row of atoms are missing then the Ti 4§ states with sixfold coordination numbers is reduced to a Ti s§ state

487 with fourfold coordination. Very recent synchrotron-radiation based angle-resolved resonance photo emission (RESPE) results [36] indicate, however, that the reconstructed surface does not exhibit Ti 3d states at the Fermi level which would be expected in the case of a 'simple' vacancy model [37], so even though 'off-resonance' valence-band spectra support the above model the missing-row reconstruction model appears to be of more complex n a t u r e and will need refinement. r ..E =

5.1. Cu on TiO2(110) Due to the open structure of the (110) surface, with u n s a t u r a t e d oxygen and t i t a n i u m atoms, we may expect a less inert surface. We consider the cases of adsorption of III I 1 ! !i I1~'"1!111 copper and nickel, respectively, onto C' C D E A' TiO2(110) surfaces. i i I I l I I i I I I l The growth of copper on the EF 2 4 6 8 10 unreconstructed ( l x 1)TiO2(110) Binding energy (eV) surface at RT was studied over the deposition range 0 < dcu < 80/k [27]. After depositing dcu = 7 A a Figure 12. UPS He(I) from TiO2(l10). (a) hexagonal overlayer gradually (lx2) surface t h a t was heated at 550~ appeared, and by further for a few hrs.; (b) (lx2) surface initially deposition the hexagonal overlayer heated to 700~ for 440 min followed by spots became bright and sharp annealing at 550~ for a few hrs; (c) a while the substrate rectangular (Ix 1) unreconstructed surface (ref 34). p a t t e r n further a t t e n u a t e d until it disappeared completely at dcu = 50 /k (Figures 13a-c). Upon annealing the 50 A Cu-covered surface at 160 ~ for 5 rain the substrate p a t t e r n reappeared [Figure 13d] while UPS showed coexistence of a sharp Cu 3d band characteristic of a bulk Cu metal and O 2p of the substrate. This suggests an agglomeration of the Cu atoms r a t h e r t h a n diffusion of the Cu atoms into the substrate lattice during annealing at 100 ~ The copper atoms w_ere found to be positioned in registry with the substrate lattice along the [110] direction, and with the lattice constant of the Cu superlattice in k space being 3 times larger t h a n t h a t of the substrate lattice, as confirmed by the fact t h a t no splitting or distortion of the diffraction spots at coincident points in the LEED p a t t e r n s are observed. The superstructure lattice constant is 2.498/k, thus only 2% smaller t h a n the 2.556/~ of the Cu(111) lattice.

488 The TiO2(ll0) surface is terminated by oxygen atoms, and the surface oxygen atoms are stretched outward from the layer of highest density, forming a onedimensional row along the [001] direction [1]. There is a row of Ti atoms, with 5fold coordinated oxygen neighbors in the subsurface layer, in the middle between each one-dimensional oxygen row. Based upon our LEED observations and symmetry considerations we have concluded that each Cu hexagon in the first monolayer is located between protruded oxygen rows (Figure 14), and since the second Cu layer is expected [38] to adsorb on each b u l k hollow site of the first adsorbed Cu layer, the Cu superlattice will grow in the ABCABC mode leading to the fcc structure. The copper superlattice is stable to at least 80 A thickness.

Figure 13. LEED from Cu/TiO2(l10) for different dcu with Cu deposited at RT. (a) dcu = 7 A; (b) dcu = 15 A; (c) dcu = 50/k; (d) is the (c)-surface heated to 160~ for 5 rain (ref 34).

489 In the same m a n n e r as for the previous metal-on-metal oxide systems the growth mode was followed by AES. Figure 15 shows a clear break at dcu = 3.5 /k, corresponding to one ML coverage. The substrate LEED p a t t e r n was still visible at dcu = 40 A. The growth thus seems to be island formation on top of the initial monolayer, but since the substrate p a t t e r n is still visible at 40 A [34] the growth mode apparently is a Stranski-Krastanov type, islands on top of fairly large monolayer patches ( due to the sharp monolayer LEED pattern). Very recently [39] the growth for Cu on Ti02(110) has been investigated by lowenergy ion scattering which concludes a Volmer-Weber (three-dimensional islands) type of growth. It is at present difficult to judge between the results since the two methods of surface impact are very different, the former [34] using 3 keV electrons (with carefully optimized beam irradiation time versus the accuracy for the measurements]) and the latter [39] using a 1.5 keV He ion beam (with care to limit damage during the measurements). Let us finally suggest an explanation as to why the superlattice LEED spots are

Figure 14. Hexagonal close-packed Cu superstructure on TiO2(110). Shaded balls: Cu atoms; large white balls: oxygen anions; small black balls: Ti cations.

not visible until dcu ~ 7 ,~ has been reached. This thickness corresponds to two ML, and since the copper grows incommensurately on the surface the protruded one-dimensional oxygen rows may interfere with the diffraction from the first monolayer of Cu atoms located between the surface oxygen rows. We end this more detailed discussion of the growth mode by concluding t h a t the copper grows in a slightly contracted hexagonal superlattice in a combination of a one-dimensional c o m m e n s u r a t e (in registry) growth along one substrate direction and an i n c o m m e n s u r a t e growth in the two-dimensional (1• 1) surface, and t h a t the

490 protruded oxygen rows are i m p o r t a n t in the formation of the superlattice.

[]

140

120

D

_A

[]

..~ C=

[]

A

,4

n Cu (M23VV)

A 0 (KLL)

L_

~0 1 0 0

n

a

I.l.I

'" z

A

80

O

o Ti (L3M1M23)

DO

AD OA

A~ .m..

=

O

60

~X~

A

...m

O O

-

40

o ~

A o

D 20

Z3

0

O0 0

0

-

000

0

r'!

9

O

0

I

a

I

&

I

I

0

I

I

2

4 6 8 Cu deposition thickness

I

0

I

I

10 (A)

0

I

I

12

Figure 15. Changes in Auger Cu, O and Ti intensities w i t h increasing dcu 293-K deposition on TiO2(110).

When we follow the deposition of Cu on the ( l x l ) surface we note a clear evolution of two p e a k s and a shoulder in the valence b a n d He I s p e c t r u m (Figure 17) which are m a i n l y contributed by O 2p bands. These b a n d s are a t t e n u a t e d gradually w i t h do, while a new shoulder, which subsequently becomes a peak, appears in the oxide b a n d g a p due to Cu 3d emission, and we note t h a t the Cu 3d binding energy shifts continuously until reaching a pure copper s p e c t r u m at the thicker layer (where we have the hexagonal superstructure). We see (Fig~ore 17) t h a t the copper surface state S.S. is not clearly established u n t i l about 10 A of Cu (the figure also shows a He satellite).

491

dcd~,)

=E

Simultaneously we note a gradual attenuation of the substrate electronenergy losses while there appear two new interband-transition losses (A and B) and a plasmon loss (C) which shift with dcu. Besides, the doublet Auger Cu 3p spinorbit interaction-peak merges [40] as it does in Cu20 because of the weakened and broadened Cu 3d band overlapping the Oinduced 2p states. These latter bonds are rather weak due to the lower affinity between the Cu and the protruding O t h a n between Ti and O , but still sizeable enough to give a measurable charge transfer in the interface.

era

5

.

EF

.

.

10

~

3

~

~

2

~

,

-

~

.

12

A

Binding energy (eV)

Figure 16. Change in UP He(I) spectra with do,, from Cu/TiO2(ll0) at 293 K (ref. 40).

""I L l

0

A

02 ,

In Figure 18 is shown the change in charge transfer with the binding energy calculation, using the assumption cop- ~]n of the plasma frequency relation with the electron density and associating the change with a decrease of the free-electron density in the adsorbed layer for the O-Cu interaction in the surface. Comparing to the somewhat similar Cs-on-Cu(111) system [41] we may expect this model to extend to submonolayer coverage where only two-dimensional islands exist, since the density of the Cu on Ti02 is calculated to 1.23 • cm 2 from the StranskiKrastanov growth-mode results [34], the density being considerably higher t h a n for

. 1

,

;i/21 ,

1

tie

{

40

{

3'o

{

2'o

{

,'o

{

I

o

Energy loss (eV)

Figure 17. Changes in EEL spectra with dcu from Cu/Ti02(ll0) at RT (ref. 40).

492 the Cs/Cu system. By comparing Figure 18 to results from the Cu-on-Cu~O [42], which indicated a charge transfer of 0.31 e/Cu adatom during formation of Cu20, and to XPS results for the deposition of Cu on MgO(100) [12] in which formation of Cu(I) states were o b s e r v e d , then we conclude t h a t the initial deposited Cu atoms probably exist as Cu(I) states on the TiO2 surface.

,_ 0.3 r-

-

0.2

=,O0

0.1 O I

2.6

2.7

I

I

I

I

2.8 2.9 3.0 3.1 Binding energy (eV)

31

.2

Figure 18. Indication of charge transfer between Cu islands and TiO2(110) substrate v s . EB(CU 3d).

A,~c ~,

AREELS

[II I\ ll! i \

Cu/Ti02 (110)

II ~ i \ III I\

I ,~

\1i ....-.... oo e-

,4 ......

e,,era

\

fVa

= ~1

I

=

I

,tT',/; I

I

\~

\

Eo:'OOeV

o; ooo

\

u

/~/

\l/~

o

/ ~

~

95

~ I

clean

Energy loss, EL (eV)

Figure 19. Changes in AREEL spectra with dcu from T i Q ( l l 0 ) at RT. Absorption edge and loss values are marked (ref. 46).

However, we should be aware t h a t this state is quite different from t h a t of bulk Cu20 where we have Cu 3d levels involved in bonds to O, hence not behaving as isolated core states. In a comparative experiment using the catalyst CuC1 as adsorbent on the (l• found [43,44] a rehybridization between the Cu 3d (downwards oriented in the CuC1induced (4• 1) reconstructed surface) and the O 2p levels. The assignment of the oxidation states of Cu and the bonding relations among the various population levels are of considerable interest in these systems, e. g. in the elucidation of the initial stages in some catalytic reactions. We thus conclude t h a t interactions between the Cu 3d levels are very limited, and t h a t the bonding across the interface mainly is contributed by 4sp electrons, resulting in a charge transfer. The role of defect sites is also of great interest. The (1• surface of

493

/• I/

TiO~(110) contains intrinsic defects located mainly on the top few layers (in contrast to Ar+-sputtered surfaces where Ti § states, assigned to 3 d 1 polaronic states localized at crystallographic shear planes [45], are found deeply in the substrate). We have found from UPS results [40] on the (1• surfaces t h a t more charge transfer occurs from the initial deposited Cu atoms (which may be considered as extrinsic defects, playing the role of Cu (§ surface donors) to the defect sites, and t h a t copper multistates exist on the (lx2) surface.

He(I) UPS Cu/Ti 02(110)

J

....,..

I:=

I

a

f

/

~,,~

4.0

oo

( I I I

J Binding energy (eV)

Figure 20. UPS from CufriO2(ll0) at RT. The surface state of Cu hexagonal superlattice film is m a r k e d (ref 46).

The copper atoms t h a t are placed around defect sites probably correspond to higher binding-energy states, and may - due to the poor screening of the titanium-ion pair at the defect site -yield a higher capability of the defect site to take more electrons from copper adatoms. Hence the defect sites seem to have capability to trap electrons donated from copper atom and thus modifying the interface properties. For the Cu/TiO2(110) system, which is the one system of these t h a t we have most intensively studied, we finally discuss some fundamentally interesting, we think, phononelectron angle-resolved-EELS (AREELS) results in conjunction with UPS m e a s u r e m e n t s [46] from deposition of Cu on the (1• surface, using a HREELS

D 7.25 eV

E0= 40 0 eV

A

e ~ dcu(~,) 51 0

=2

560

39 5

660

. m

51 0 56.0

40

660

1'o 1'2

1'6

Energy loss, EL (eV)

Figure 21. Changes in AREEL spectra with scattering angle from Cu/TiO2(110) monolayer and thick film (ref. 46).

494 i n s t r u m e n t in the specular-reflection geometry (where the m o m e n t u m - t r a n s p o r t parallel to surface (kl) is nearly zero, whence the loss functions are analogous to those derived from optical data [46]) at 22 eV < Ep < 40 eV as a function of scattering angle in the range 490 < 0 < 66 ~ and the analyzer angle at 0 ~ in the UPS measurements. Phonon-assisted interband transitions are known to occur in several semiconductor materials, such as Si and Ge, in which the m i n i m u m in the conduction band and the m a x i m u m in the valence band are not located at the same point in the Brillouin zone (BZ), and in optical absorption experiments these indirect transitions require assistance from phonons to m a i n t a i n conservation of momentum. Upon deposition of Cu over the 0 < dcu < 40 A range at E p - 40.0 eV and scattering angle 0 = 56.00 [46] we observe (Figure 19) the appearance of four copper-related losses (A-D) while the substrate losses are attenuated. The intensity of the clean-surface elastic peak decreased almost one order of magnitude during the initial depositions and started increasing again at about the monolayer point d c ~ - 4/k. Also, it was observed t h a t the copper absorption-edge shifted, first toward lower energies and then, passing through a minimum, toward higher energies. The absorption edge could be determined as the cross-point of the extrapolation of a low-energy-power-law-dependent curve and the linear extrapolation of the absorption-edge curve [46]. The 3 losses (A,B and C) of lowest energies are allocated to interband transitions with Cu 3d levels, as described above in the results of the angle-integrated experiments, although there only 2 losses were resolved because an appreciable contribution from (k I ~ 0) scattering in the angleintegrated data smears out the structure of the (k I - 0) loss function [48]. Below 4 A the shift of the Cu 3d band dominates and is interpreted as previously discussed in terms of charge transfer. The fourth loss (D) is, as previously, assigned to a plasmon excitation.For thick (dcu- 39.5/k) overlayers an absorption of 2.08 eV was measured, in agreement with optical data for Cu. This edge was seen to move toward lower energies at coverage near the monolayer point which was surprising since the Cu 3d levels remain unshifted above the monolayer point as observed by UPS (Figure 20). To exclude the possibility of diffraction effects AREEL spectra were obtained (Figure 21) for different 0 both for the monolayer film and for the thick film. Figure 21 indicates both absorption edges and the other losses, A-D. As discussed in the above paragraph, indirect transitions require assistance from the phonons, hence the absorption edge for those semiconductor materials shifts Eg- Eph r a t h e r t h a n to Eg, where Eph is the phonon energy. The similarity between the present case and the case of indirect transition caused us to the proposal [46] of the existence of a substrate-phonon-assisted Cu-overlayer-interband transition, i.e. across the interface. This type of transition has to our knowledge not been reported earlier in the literature. The proposal is supported by a detailed analysis [ref. 44 and references therein] of the experimentally observed electronic transitions from copper, involving both occupied and unoccupied states (we do not need to consider the substrate electronic structure because the interaction to the substrate is weak as discussed above [34]); in brief the analysis is as the following. Due to the many possibilities of vertical transitions inside the BZ the discussion is complicated. The loss features may be dominated more by transitions near the

495

,ix', .-.

w3 Jl

2

X4

0

r~ i

FERMI LEVEL

,,z,

Wl

LU

-

F25,

F

h

X

Z

W

Q

L

A

F

2:

K

X

k (2n/a)

Figure 22. Bandstructure for Cu. We have marked examples of the three types of transitions; see text (ref 46 and refs. therein).

zone boundary t h a n by those near the center (due to the much larger area in the former region), in the most simple case, but to discuss the present case we have suggested a classification of three different types of transitions upon which a direct comparison with simulated energy bands can be made along the BZ symmetry axes The three types, illustrated in Figure 22, are (i) critical-point transitions from the high-symmetry points at which the m o m e n t u m matrix elements are so high t h a t they can contribute to the structure in the loss function; (ii) the transitions t h a t occur between occupied states and the Fermi surface and (iii) the transitions, which we refer to as a volume effect, t h a t occur in the region near the BZ Fermi surface and t h a t are allowed by the dipole-selection rule. The analysis reveals, with the help of Figure 22, t h a t the 2.08 eV absorption edge corresponds to transitions from the osculating points, while the loss intensities around 2.72 eV are contributed both from the osculating points and by the volume effect from the large 5-+6 and 4-+6 band-to-band transitions near X inside BZ and near the Fermi surface around the L neck, and the broad loss B mainly is caused by the volume effect from the large 3-+6, 4-+6 and 5-+6 transition regions. The transitions 6-+7 between sp states and 1-+6 from the bottom of the 3d band dominate loss C. In the monolayer case the three-dimensional (3-D) BZ becomes a two-dimensional (2-D) BZ, and the possibility for vertical transitions is lowered very much, in agreement with the one-order AREELS intensity decrease from the case of the thick layer to the case of the monolayer. A similar analysis as above is made for the 2-D case. In the UPS results of Figure 20 we note t h a t the submonolayer deposition of Cu

496 causes a shift toward a higher binding energy, as we discussed earlier, in contrast to the usual findings of the Cu 3d levels shifting toward EF in comparison to the bulk d band in isolated monolayer films, hence the blunt absorption edge for the monolayer deposit cannot be allocated to pure electronic transitions from the osculating points, but r a t h e r to a substrate-phonon-assisted Cu-interband transition. The energy difference between the two absorption edges is 0.3 eV, within the region of an oxide-substrate phonon, where the f u n d a m e n t a l modes and their combinations yield significant intensity up till at least 0.4 eV in our observations (a Cu phonon cannot cause such a large absorption-edge shift). The intense substrate optical phonons causes a coupling with electronic fluctuations in the copper overlayer, producing long-range disturbance into the vacuum and hence causing inelastic scattering of the incoming electrons by the dipole field into a small angle around the specular direction. Simultaneously the thermally excited phonons are annihilated, and the edge must shift down (phonon creation would need energy and thus have shifted the edge up). We therefore conclude t h a t we have revealed a substrate-phonon-assisted electronic Cu-interband transition at the interface. 5.2. Ni (m TiO2(ll0) Deposition of the transition metal nickel onto Ti02(110) was already early the subject for investigation due to its importance in m e t h a n a t i o n by heterogeneous catalysis, particularly after it had been found [49] t h a t Ni on a rutile-anatase mixture has a considerably stronger activity t h a n on other (SiO2 and AleO s) oxide supports. Kao et al. [35] thus found for deposition at RT over the 0.23 < 0 < 15.8 coverage range on reduced (i.e. heat-treated to release oxygen) a charge transfer from Ti02 to Ni (the Ni atoms become negatively charged) varying between -0.13 and -0.07 e per adatom at 0 - 0.5, by observing the shift in the Auger Ni(LMM) peak and the 2pa ~ core level of Ni, and t h a t the amount of charge transfer was dependent of the surface p r e t r e a t m e n t (annealing / sputtering). This behavior is often ascribed to a characteristic of the important strong metal-support interaction (SMSI) process in heterogeneous catalysis (the effect is characterized by a suppression or loss of the ability of the metal to chemisorb (hydrogen, carbon monoxide) when positioned on a reproducible oxide support, leading to selectivity properties). Later Onishi et al. [50] found t h a t for Ni deposits on a non-reduced TiO2(110) there was a small, 0.1 e per adatom, and as the overlayer grows more dense, the lateral interaction between the Ni adatoms predominates and inhibits the electron transfer through the interface. No epitaxial relationship was found for the Ni overlayer in these studies. We have found [51], however, t h a t epitaxy did in fact occur, but the epitaxial growth of Ni on the (1• 1)TiO2(110) surface cannot be predicted by any existing theory. Two types of nickel islands are formed, a hexagonal structure oriented parallel to the substrate, and a hexagonal structure whose direction of growth is inclined with reference to the substrate plane. The growth of u l t r a t h i n layers of Ni on the (l• surface over the 0 < dNi < 80 A range at RT shows attenuation of the substrate L_EED p a t t e r n up until 10/k where three weak parallel lines along the substrate [110] direction begin to appear, and upon further deposition the intensity of the lines increases, and for dNi > 30 /k the substrate p a t t e r n has completely disappeared. After slightly annealing, the three-line LEED p a t t e r n s changes into many extra spots lying in the lines. After careful analysis of a series of the p a t t e r n s three sets of diffraction

497 spots were analyzed, and it was observed t h a t among these two of the sets are symmetrical with respect to the substrate {001] line in reciprocal space. These two sets move in complex way with E, such t h a t corresponding symmetrical diffraction spots move away from each other, within the line, with increasing Ep. This usually indicates t h a t the overlayer is tilted away from the substrate plane. After a computer simulation based upon Ewald-sphere analysis [16] we noticed t h a t the observed complex movement of the diffraction spots only occurs along the [110] direction in the substrate. The relation between the spots in the [001] direction is somewhat similar to the case (above described) of Cu on the same substrate, hence t h a t Ni is packed in hcp, so we used a hexagonal lattice in the simulation. We found t h a t the best fit for all the p a t t e r n s in 10 eV steps over the 50 < Ep < 180 eV range was obtained for angle of 270 between the ( l l l ) N i plane and the substrate with the tilt axis in the [001] direction, indicating t h a t two types of Ni islands are formed simultaneously upon the initial monolay_er. The first type grows parallel to the substrate in registry with the [110] direction but incommensurately in the other 2-D directions of the substrate, as in the case for Cu on this substrate [34], leading to normal growth of fcc Ni. We found a lattice constant of 2.50/k, i.e. lightly larger t h a n the bulk 2.49/k. For the 270 tilted plane we found t h a t the (131)plane satisfies the requirement: the angle between the two planes is 29.5 ~ in good agreement with the observed 27 ~ The two fcc structured types of islands can be characterized as: N i ( l l l ) II TiO2(ll0), Ni[101] II TiO2[001] and Ni(131) II TiO2(ll0), Ni[101] II WiO2[001] 60

-\ O

9~

50

,4

_O

\

... 40 ..--...

z

o O~KtLI

_ 0 0

\

-

30

~ 20

13 Ti(L3M23M23 )

R

O,o

-%\

o,%

e-

~ 3 u ._~.._ . . E l ~ . ~ ~ . .h. . ~

_

10

~

0~

_

_

O~D

0

2 4 6 8 10 12 14 Ni deposition thickness (A)

16

Figure 23. Changes in Auger Ti and O intensities with dNi on TiO2(ll0) at RT. (a) quasi-isotropic growth model; (b) anisotropic growth model (ref. 16).

The growth mode for Ni on the substrate at RT was studied over the 0 < dNi < 16 range by AES (Figure 23). We find a sharp monolayer breakpoint at dui = 2.5 A. Since the substrate was still visible at 25 A we exclude a layer-by-layer growth mode and assign it to a S transkiKrastanov growth mode. With the same assumptions as for the growth of Ni on NiO(100), as described above, we have analyzed the data using the quasi-isotropic growth model [28] and found t h a t only the isotropic growth model fits correctly the growth behavior. We expect t h a t after the first hcp Ni monolayer with sixfold symmetry, the anisotropic features of the

498 substrate is weakened so much t h a t the observed isotropic growth can occur, as observed. As for the nickel oxide case we have also here determined the density of the Ni islands before coalescence and find 2.9 • 10 is cm 2. The 3-D growth of Ni (after monolayer coverage) hence is described well in both cases by the quasiisotropic growth model. At RT, CO molecules do not adsorp on a clean TiO2(110) surface, but at a nickelpromoted surface it occurs. In a combined valence-band UPS and vibrational HREELS experiment [52] a distinct peak in He(II) UPS at EB=I 1.0 eV below E F grows, and further CO exposure induces two new CO-derived peaks appear at 8.0 and 11.0 eV (as was, independently, found in an UPS study by Onishi et al. [50]) These two peaks can be attributed to the emissions from (1~ + 5(~) and 4(~ states of CO, respectively. The CO adsorbs with its carbon-end oriented towards the surface. The existence of molecular CO at the Ni/TiO2(110) surface (in agreement earlier studies [53] showing associative CO adsorption on a Ni crystal at RT) was confirmed by HREELS spectra [52] on 1-, 2-, 3- and 4-/k Ni-deposited TiO2(l10) surface at RT. The saturation coverage of CO increases with dNi, indicating that CO molecules bind to the Ni atoms rather t h a n to the substrate atoms. At saturation coverahe, CO molecules adsorb simultaneously on the 2-fold bridge-sites and the terminal sites on the (111)-oriented Ni islands on the TiO2(110) support. The occupation of the edge-sites of the Ni islands gives rise to an exeptionally low C-O stretching vibration of 152 meV. This frequency, indicative of a weakened C-O bond, suggests existence of a precursor to the dissociated state, which is important inthe understanding of the bevior of the catalyst supported on the metal oxide. 5.3. Ni (m TiO~(100) The (100) surface of TiO2 reconstructs upon 500~ annealing to a (1• at 800~ to a (1• and at 1200~ to a (1• LEED p a t t e r n [29], and also this subs~raLc surface shows Ti § interband transition upon Ar § sputtering by removing surface oxygen [55]. The interface with Ni is of particular interest in the process of CO hydrogenation. Kao et al. [54] first used a TiO2(100) as substrate for Ni deposition as a model catalyst and found by XPS t h a t Ni atoms at the interface are negatively charged. The growth of Ni upon TiO2(100) was found [56] by AES and confirmed (using the sequential-layer-sputtering model [57]) by secondary ion mass spectrometry (SIMS) to follow a layer-by-layer growth for the first three Ni layers. Then islands of Ni grow. For a non-stoichiometric support no island growth is found and the Ni is diffusing into the bulk even at RT, depending on the amount of oxygen vacancies present [58].

6. COPPER AND NICKEL ON CORUNDUM STRUCTURRS Of the (trigonal-lattice) corundum structures, we will here describe recent results for adsorption of Cu and Ni on a non-transition-metal oxide, alumina, a-A120 s (saphire), and of Cu on a transition-metal oxide, a-Fe2Os (haematite). A1 and Fe are here surrounded by six oxygen ligands to their (distorted) octahedral sites. Among these, the a-A1203 is by far the one t h a t have been studied most intensively due to its very wide applications in fields as diverse as ceramics, composites, microelectronics, refractories, catalysts, membranes and cements. It is unique also with respect to its electronic properties; it has a very wide band gap

499 of 8.7 eV [59]. After heating to 900~ in air followed by h e a t - t r e a t m e n t in UHV to 700~ a sharp lx 1 surface structure is observed [60] (heating the crystal in UHV to higher temperatures, up to 1400~ yields rotated ~]3x~/3, 3~]3x3~]3 or ~]31x~/31 reconstructed AleOs(0001) surfaces [61]). The electronic structure of a-alumina is now well understood, both from a semiempirical (extended Hiickel) calculation [62] and from an embedded-cluster calculation [63]. The loss structure has been studied experimentally by EELS [63-65]. Work on adsorption of Cu and Ni (and other metals) onto single-crystal surfaces of these oxides are scarce, however, perhaps due to experimental difficulties with charging effects as with the earlier discussed MgO. 6.1. Cu r AI208(0001) The crystallographic structure of the Cu/a-A1203(0001) interface was studied by tra_nsmission_electron microscopy [66], and epitaxial (111)Cu II a-A12Os(0001), [211]Cu II [2110]a-AleOs(0001) relationship was demonstrated by back-reflection Laue X-ray diffraction [67], ion-etch characteristics and scanning electron microscopy [68], and the strong adhesion of copper to this surface obtained as a result of an ion-sputtering t r e a t m e n t was investigated by XPS [69,70]. Theoretically the metal-sapphire shear strength was investigated by serfconsistent-field X-alpha scattered-wave cluster molecular-orbital models, and it was found t h a t a chemical bond is established between the metal d-electrons and the nonbonding 2p-electrons of the oxygen anions on the A12Os surface [71], and that the bond strength between the oxygen anions to the close-packed firsttransition-metals is placed between the empirical values for stronger metal-oxygen bonds and weaker bulk-oxide bonds [72]. We have studied [73] the growth of Cu on aAleO~(0001 ) in the 0 < dou < 130 A range, and have found a Cu(11 1)-R30 ~ s u p e r structure on the (lx 1)-aA120~(0001) surface after having heated a 53-/k film to 650~ for 30 rain (Figure 24). The same LEED p a t t e r n as the one shown in Figure 24 could for the same 53-/k surface be obtained also at lower temperatures, for instance 180~ at 15 rain, but then it was necessary to use rather high Ep because of charging effects. By AES, EELS and UPS (in the latter Figure 24. LEED showing rotationally we compared to an A120s aligned epitaxial Cu(111) on a-A1203(O001) film formed on an AI(111) after 650~ heating, surface and made sure that charging effects did not interfere since the O 2p peak

500 did not shift while the Cu peak at the same I I I I l I time shifts tohigher binding energies with dcu [73]) we have found [60] t h a t already from dCu (/~) the initial low-end submonolayer coverages a chemical 27 bond can be established between copper and the 1• smooth substrate, and our data (Figure 25), which show no Cu(0) signal in the initial submonolayer stage, support an interpretation in terms of a weak charge1.7 transfer process where charge is donated from the copper particles to substrate oxygen, so that initially Cu(I) 0.7 states and t h e n Cu(0) states are seen with increasing dcu (by presputtering an A1203(0001) surface it 0 (clean a-AI203) was found [70] t h a t the interface involve Cu(I) d 1~ configuration). Due to an observed I ! I I' = = more rapid decrease of 20 30 40 50 60 70 the aluminum-toKinetic energy (eV) oxygen Auger intensity ratio for the lower (dcu Figure 25. Changes in Cu Auger fine structure < 2/k) deposits it is with dcu. likely t h a t we see initial copper atoms positioned in a-top position of the aluminum" sites which are in 3-fold hollow positions created by the oxygen atoms (Figure 26). Ideally the outermost atoms of the 1• 1 a-AleOs(0001) surface are aluminum, but here the outermost atoms are expected to be oxygen due to the high-temperature oxidation p r e t r e a t m e n t of the sample. This was confirmed by AES. ,..-,,..

L'u .,,....

cIp

.0,.,, e-

,._ Q.I

501

Figur 26. Surface geometric analysis of LEED pattern (Figure 24) from Cu/a-AleOs(0001),

1.3 A Cu/~-AI203(0001) 0.6 ._o

0.4 .'~_ oo ,,Ik

,_

9 "-

Cu/AI AI/0

=-

Cu/0

02 9 I

0

100

i

I

I

~"''~--=-i

200 300 400 500 Annealing temperature T('C)

4

i

600

700

Figure 27. AES on the thermal stability of Cu small islands on a-A12Os(0001).

502 From the changes of the Cu-to-O and Cu-to-A1 Auger intensity ratios it is difficult clearly to evaluate the growth mechanism at RT. The ratios initially change linearly with dcu, followed by a slow exponential change, hence a StranskiKrastanov mode, a copper ML followed by subsequent island (cluster) nucleation [60], is indicated. However, later XPS results [75] conclude a Volmer-Weber type of growth, pure cluster nucleation with no monolayer formation at RT (and also for the high-temperature ~]31• reconstructed surface). As in our case, no superstructure LEED pattern was observed at RT.

/~

dcu(/~) 3.3

T/~ 670

330

3.3

220

3.3

25

A

v

vJ

e-. D

~__

0 (clean)

I

I

I

I

I

25

I

20 30 40 50 60 70 Kinetic energy (eV)

Figure 28. Fine-structure Auger analysis on the t h e r m a l stability of 3.3-A Cu thick deposits on a-AleOs(O001).

503 Very recently, though, an Angle Resolved XPS investigation [76] finds t h a t copper initially grows as uniform layers on l x l A12Os(0001 ) followed by formation of clusters after 2-3 atomic layers of Cu deposited at RT, thus agreeing with our conclusion on growth mode (and with the mechanism of charge transfer). A recent combined HREELS and AES investigation [77] also finds t h a t the first monolayeforms a smooth overlayer at 95 K on a thin A12Os film t h a t was formed on an AI(111) surface, i.e. Stranski-Krastanov-type growth, and t h a t very little change was observed by heating the layer to 400 K. We have investigated also the t h e r m a l stability of the copper film [60, 65]. As seen from Figure 27, the interface composition is quite constant until about 400~ This further indicates a chemical bond at the interface between copper and the clean (~-A12Os(0001) surface. We see some decrease for a 2-/k film in the Cu-to-O and Cuto-A1 ratios with increasing temperature; there might be three possible causes for this: (i) interdiffusion between copper and substrate, (ii) copper agglomeration at the surface, and ('di) desorption or evaporation of Cu from the surface. We consider evaporation to be insignificant at t e m p e r a t u r e s lower t h a n 400~ and the r a t h e r unchanged ratios over the RT to 400~ range for monolayer copper on the substrate indicate t h a t interdiffusion has not happened at the interface. Rather this indicate t h a t small copper islands have increased in size with temperature, resulting in the decrease in these Auger ratios. Figure 28 show the changes with t e m p e r a t u r e of the Cu(M2.3VV) transition for a 3.3-A deposition [65]. The heatt r e a t m e n t (for 15 rain) at 200~ caused a split into two peaks (at 55 and 58 eV) in this Auger line compared to t h a t of the u n h e a t e d sample where the kinetic energy for this transition is 57 eV (A1 and Cu(I). We believe t h a t the split is caused by a mixture of Cu(I) and Cu0) states. Furthermore, we note t h a t the peak position is shifted to a lower kinetic energy, relative to the transition in copper mc~al, indicating t h a t agglomeration has occurred as a result of the heat treatment. F u r t h e r heating to higher t e m p e r a t u r e s results in a "~ r--'--] i"-'-'-I decrease of the Cu(0) intensity, here at 58 eV, when compared to the 55-eV peak which is ascribed to Cu(I), suggesting that the Cu(I) on the surface at these F--] I--~ I---] t e m p e r a t u r e s is more stable t h a n Cu(0). b L . v/~ EELS results for a 0.7-A deposit [65] shows t h a t the new RT-copper-induced loss at 5.5 eV, which suggests t h a t m Cu(I) can be established in a charge-transfer process for 1 thicknesses dcu< 2/k, is stable until around 400~ and is decreased by heat t r e a t m e n t to 600~ in agreement with the AES results. 1-3 N Figure 29 illustrates schematically the different situations on the substrate at RT for 2-D and 3-D clusters and after annealing, respectively. o' - AI 2 0 3 Cu In the case of submonolayer-copper thicknesses a quite Cu (111) stable chemical bond between copper and the substrate is thus established, resulting initially in copper in a Figure 29. Growth Cu(I) state, while for thicker films the copper diffuses in model for Cu/athe direction of the surface lattice to form nucleation A12Os(0001). See centers. text.

504 In the above referred HREELS-AES investigation [77] it was found t h a t heating to 700 K caused an inward diffusion of the copper overlayer. It should be noted, however, as also the authors do, t h a t since the AleOJAI(111) substrate possesses metallic a l u m i n u m beneath, and likely intermixed with the AltOs films, the thermodynamic driving force for the diffusion of copper into the bulk is the formation of a Cu-A1 alloy [77,78] in t h a t case (their films were only -4 and -8 /~ thick). 30-/~ thin A1203 films, made by oxydation of thin A1 foils, were also investigated as to the effect of cluster size on the shifts in core-level binding energy and in the kinetic energy [79]. A charging-up behavior in the m e a s u r e m e n t s at certain energies, as previously discussed (Sect. 3), was naturally found also for this strong-insulator system. a-Al2Os(0001) The energy-loss structure for RT deposits of Ni on a-A12Os(0001) (model for the uses in a range of catalytic processes) was obtained for thicknesses of 0.5, 2 and 3/~ [64], and it was concluded t h a t the stoichiometric surface generate chemical bonding with a l u m i n u m dangling bonds as trapping centers (a similar conclusion was obtained for the (i012) surface [64]). A combined LEED and XPS investigation [80] did not observe detectable reaction at RT, but when Ni was deposited at 800~ in UHV the alumina was seen partially reduced upon Ni deposition, and when the deposition was carried out at that t e m p e r a t u r e in the presence of 02 a NiA1204 spinel phase was formed. Here the XPS showed the shake-up peak characteristic of the oxidation state of nickel, and LEED showed an ordered overlayer structure with sixfold s y m m e t r y with the spots from the NiAleO 4 at about half the distance from t h a t of sapphire (there is an approximate factor of 1.7 between the lattice constants of NiA1204 and a-A12Os, and a lattice mismatch of 3.5% between its (111)plane and the substrate surface). Postannealing (in separate RHEED chamber) to t e m p e r a t u r e s up to l l00~ showed dissociation of NiO until complete metallic Ni was recovered without observation of a NiA1204. Sixfold epitaxy was observed both in the UHV and in 5• 10 .7 Tort 02. LEED was very much hampered due to charging, though. Decrease in the intensity in Kikuchi bands (lines) indicated t h a t the overlayers were strained. An ab initio, cluster, unrestricted Hartree-Fock calculation on the electronic bonding properties of the Ni/Al~Os has found [80] an average bond strength of 5.3 eV and an average charge transfer of 1.5 e per adatom from Ni to antibonding states on the alumina surface and proposes t h a t the catalytic action of aluminasupported nickel is in part connected with the presence of nickel as a positive ion rather t h a n as Ni ~ (and thus seems to contradict earlier experimental findings [80] that the catalytic turnover number measured on a single-crystal Ni, with site density derived from the Ni atoms of a (100) plane, is in good agreement with values reported from high-area (alumina) supported catalysts, with site density derived from chemisorption data). As in the case of Cu/AleO3(0001) we will also for the Ni/AleO s system compare briefly to results from deposition of Ni onto thin films of A1203 formed upon clean AI(111) surfaces. In an AES-HREELS investigation on Ni deposited onto thin A120 s films grown by oxidation of an AI(111) surface it was found [83] t h a t the deposited Ni atoms form 3-D clusters on the A12OJAl(lll) substrate at 200 K, and t h a t thermally 6.2. N i r

505 induced smoothing and diffusion of the Ni overlayer occur at characteristic t e m p e r a t u r e s in the 200-700 K range. From a layer with an A1203 thickness of about 3 A it was tentatively postulated t h a t at t e m p e r a t u r e s below 400 K the results suggest a smoothing of the 3-D Ni cluster layer to a more uniform Ni overlayer occurs, and at t e m p e r a t u r e s above 400 K t h a t inward diffusion of the Ni overlayer p r e d o m i n a n t l y occurs. The Ni/A12Os system has been used as catalyst for m e t h a n a t i o n of CO, and no Ni diffusion had been reported for the 200-700 K t e m p e r a t u r e range. However, it was found now t h a t surface segregation of metallic A1 toward the Ni overlayer occurs during the Ni penetration process, explained by an interdifhasion of Ni and A1~ in macroscopic channels within the A12Os layer, p e r m i t t i n g the formation of a Ni-A1 ahoy. It was thus found t h a t the metallic a l u m i n u m b e n e a t h the AleO 3 film plays a major role in driving the inward diffusion of the Ni overlayer. Recently we have investigated this system by AES, XPS and EELS and have found [81] evidence by correlated AES and XPS m e a s u r e m e n t s for a StranskiKrastanov type of growth (a 2-D layer followed by 3-D clusters) for Ni deposited at RT on a 7-AleO s thin film grown on an Al(111) surface. We have calculated the equivalent thickness of the Ni deposits at the end of the 2-D p a r t of the growth by using the Al(111) plane as basis, and the calculated thickness corresponds to the formation of a layer of Ni on a l u m i n a (with a sticking coefficient near one), as also evidenced [84] by t r a n s m i s s i o n electron microscopy (TEM). XPS showed unaffected A1 core-level features, and the appearance of a new peak at 3.5 eV below the final E F, and the peak shifted with dNi towards a position located 2.0 eV below EF. The Ni 2p peaks shift toward lower binding energies. The curves EB(Ni 2p3~) vs. dNi and EK(L2.sMV) vs. dNi each contains 3 steps which we consider for the two regions, I: dsi < 0.4 A and II: dNi > 0.4 A. In region I, in which the 2-D monolayer is formed, we find large changes. EB decreases 2.8 eV, the modified Auger p a r a m e t e r a' decreases 2.2 eV while E K increases 0.6 eV. The variation in a' at the initial stage we interpret as due to a low-nucleation-rate formation of small clusters on the surface, i.e. a size effect, as observed by TEM [85]. The large shifts in EB we assign to a combined effect of the presence of clusters and a formation of a new compound. F o r m a t i o n of nickel oxide is in agreement with an observed large change in the initial effect, ae = -3.9 eV. The negative sign of ae indicates a positive charge for the nickel clusters, which we explain as a charge transfer from Ni 3d toward ligand states, as we described above for the Cu/TiO2 system. Using the method of Kao et al. [84] we found a positive m e a n charge of 0.39 e / a d a t o m with regard to the final deposit, indicating a selective interaction between Ni and O, and t h u s find positively charged Ni as suggested in the above mentioned ab initio calculation [81], although we find a lower positive charge t h a n the calculated value of 1.5 e per adatom. For region I we have thus found formation of 2-D clusters of oxidized nickel and have elucidated the ionic character of the Ni-A1 bond. In region II, E B and E K still approach the bulk values, and we see 3 oscillations in the evolution curves for 0.4 A < dNi < 8 A. The oscillations in the Auger signal we interpret as instability of the deposit, i.e. some of the 2-D clusters nucleate to 3-D structures. This might explain an observed decrease of the Ni m e a n charge toward a value of 0.21 e per adatom. With reference to the previous discussed work [82] of Ni/AleOs/Al(111) we can deduce t h a t metallic Ni are located inside the aggregates, but the importance of the initial effect indicates t h a t the electronic

506 structure is quite different from t h a t of the metallic state. The m e a n value of a' gradually approaches the Ni ~ value, and the amplitude of the oscillations decreases, indicating establishment of the Ni ~ state. For dNi > 3/k the 3-D clusters are growing in size, and the initial effect decreases. The electronic structure seems to be established before reaching 2 ML. When comparing the chemistry in the interfaces of the Cu and Ni on bulk aA1903(0001) crystal substrates with those of the few-/k-thick A120s films formed by oxidation of an A1(111) surface there is a m a r k e d difference in particular with regard to the diffusion behavior, where the A1 in the latter is a strong driving force. However, much useful information have been obtained which help in understanding the chemistry for both systems. As referred to (Sect. 6.1) for the similar case with Cu it was shown theoretically by Johnson and Pepper [71] from cluster model calculations t h a t a primarily covalent bond can be established between metal and the oxygen anions on the A120 s surface. It was also shown t h a t its strength decreases in the series Fe, Ni, Cu, and Ag. The reduction in strength in this series they attribute to primarily to the increasing occupation of antibonding orbitals established by the metaloxygen interaction. Their calculation was expanded to show [72] a uniform decrease from about 5 eV to about 1 eV for binding energies per interracial 0 when an O-covered (0001) basal plane of a-AleO~ binds oxidatively to close-packed metal surfaces of the first transition series, and that, except perhaps for Cu, adhesive failure of interfaces of the structure modeled will occur between the first and second metal-atom layers. The hypothesis and calculations hence have found experimental support. a-Fe2Os(0001) Iron oxide and iron-based materials are very important catalyst used in many reactions such as the Fischer-Tropsch and the ammonia synthesises, in photoassisted electrolysis of water and in gas-sensor applications. The 3d 5 (magnetic) hematite Fe2Os is an insulator also, and its crystal surfaces [87-89] and their reactivity have been of considerable interest recently, but adsorption of copper onto the oxide surfaces has only recently been in the focus_of interest [90,91]. We have studied the clean and copper-deposited a-Fe203(1012) [90] and (0001) [91] by AES, EELS, LEED and HREELS. Figure 30 shows the HREEL spectrum from a clean a-Fe2Os(0001) surface obtained at room t e m p e r a t u r e at 600 off normal incidence. We observe the surface phonon frequencies 61,62 and 6s together with multiple and combination phonon losses. The crystal structure of aFe2Os is anisotropic, and we may expect [92] a loss function 6.3. C u r

P(r

- r 1 Im {'[sl (co) ~1 (co)+ 1] ~}

where ~l and ~1 are the dielectric functions parallel and perpendicular to the crystal c axis, respectively. Here, where we have 11 meV FWHM resolution, we can still see the 62 feature, and even the fourth multiple loss 463 is obtained. These modes are typical Fuchs-Kliewer (surface optical) phonons whose properties are mainly determined by bulk p a r a m e t e r s such as the transverse optical phonon frequency C0TO[93]. The results from calculations using the above function will compare with the experimental HREELS results.

507

'~ / x

3,1os

...-,...

:3

2~''~S

I / //

v/ /

~//

~

\

~

-4

~'~

4~ 3

.,, A,,

,:o:

I

\

I

0

I

100 200 Energy loss (meV)

I

I

300

400

Figure 30. HREELS from clean a-Fe2Os(0001). See text.

Upon depositing Cu on the (1• 1) a-Fe203(0001) substrate surface [91] the surface phonons are gradually changed (Figure 33) while at the same time a hexagonal rotationally aligned Cu (111) superstructure [91] appears and become gradually brighter and sharper until finally, after depositing for 210 s (~ 35 A) only the C u ( l l l ) structure is present. We note t h a t upon the initial deposition of copper a new vibrational feature was found at 18-20 meV (Figure 33). This vibrational peak was maintained even until 120 s deposition (- 20 A) and indicates an Cu-O interaction as judged from a comparison to calculated and experimental results [94-96], suggesting formation of a Cu(I) state at the interface, similarly as was found at RT also for the copper-alumina system [60,74]. This indicates a chargetransfer process as we saw earlier (Sects. 4.1 and 5.1). m

6.4. Cu r a-Fe203(1012) Also growth of Cu on cz-Fe2Os(1012) surfaces have been carried out [90]. The samples are artificially grown small crystal flakes and are more difficult to investigate experimentally t h a n the quite large (0001) growth faces of n a t u r a l haematite crystals. The LEED structure of the (1012_) surface depends upon the surface t r e a t m e n t [89]. Upon flash-annealing the (1012)surface in situ, Cu grows in a Volmer-Weber-type growth mode and no clear superstructure LEED p a t t e r n is observed for Cu deposited at RT. The substrate 1• 1 LEED p a t t e r n gradually becomes a t t e n u a t e d while the substrate plasmon and inter-band-transition energylosses attenuate, and the familiar low-energy interband-loss due to Cu appears.

508 7. COPPER ON PEROVSKITE STRUCTURF_~

Figure 31. LEED patterns from SrTi03(100) and segregatedcalcium superstructures after heating to 900~

Among the perovskite structures we will discuss another Ti(IV) compound, the ternary d o transition-metal cubic SrTiOs. It is, like other MaMbOs perovskite structures where the cation Mb is Ti(1V) and M a is bivalence cation, closely related in its properties to rutile, TiO2. Electronicallly [97,98], as demonstrated by RESPE [99], but also geometrically, since the atomically flat and nonpolar SrTiOs(100) surface may be terminated either by a Ti02 or by a SrO plane [100] (patches of both terminations will coexist, butpreparations may be carried out so that one of them dominates [99]). SrTiOs may for 2D surfaces experiments be considered as a simple cubic structure at RT since its distortions from the stable hightemperature cubic phase are very small. Of the perovskite clean crystal surfaces, only the (100) and (111) surfaces of SrTiOs and the (100) surface of BaTiOs have been investigated The HREEL spectrum from SrTi03(100) is thus well determined [101,102], but with regard to Cu and Ni adsorption only Cu has been reported [102,103] as deposited on the SrTiOs(100), in both cases stimulated by its qualities as a good matching substrate for the YBa2Cu3OT.x high-Tc oxygen-deficient perovskite superconductor of which high-quality epitaxial films had just been grown on that surface. The processing of synthesis of epitaxial YBa2CusOT.x films involves deposition on a 400-600~ SrTiO3(100) substrate followed by annealing in oxygen at 900-950~ Before we proceed with a discussion on the copper adsorption we will briefly discuss an impurity reconstruction that may take place just around the 900~ annealing temperature. In the synthesis of single-crystal SrTi03 it is difficult

509 completely to remove traces of calcium. Even though the bulk analytical content of Ca is only a few ppm in commercial crystals a detailed AES-LEED investigation has showed [104] t h a t Ca will segregate to the surface and change the originally 700~ sharp 1• 1 surface (Figure 31, upper) to a superstructure (Figure 31, mid) (which we attribute to a multiple scattering across the topmost layers) with about 3% Ca in the surface layer, or with only traces of (( 1% Ca to a p(2• reconstruction (Figure 31, lower) when the air-annealing temperature is 890~ and the humidity below 40%. Since the radius of a Ca 2§ ion is not very different from that of Sr 2§ we suggest t h a t the calcium ions exchange with the ions in the surface region, causing a change in the surface potential which induces the p(2• surface reconstruction (segregation of Ca to the surface of oxides has been studied also in the case of MgO, in which AES and LEISS results agree with a theoretical prediction based upon the surface heat of segregation [105]). A model for the occurrence of the p(2• structure could perhaps consist of (1) a combination of a relaxation of the Sr atom in every second unit cell in the topmost layer and (2) a segregation of ca atoms towards the second- or third-layer region, a model that will agree with our observation of the Ca-to-Sr and Ca-to-Ti ratios during development of the p(2• structure. In every second surface unit cell, the Sr atoms would then be displaced upwards, and downwards in the other half of the unit cell. The validity of the model could be explored in direct experiments on displacements of surface atoms. 7.1. Cu (m SrTiO3(100) In the two (independent) investigations [102,103] it was found t h a t Cu at RT grows in a Volmer-Weber type mode, individual clusters forming at the surface, and at high coverage a coalescence was indicated. For a deposits corresponding to 10 ML, AES results indicate [102] t h a t the copper growth and formation of islands are related to the TiO2 layer termination of the surface, and t h a t the copper growth mainly occurs on the 80% of the surface found to be TiO2 layers. No Cu epitaxial superstructures were found at the RT substrate. Growth of Cu onto a SrTiO3(100) substrate at 600 o has, however, shown [106] by RHEED (in synthesis of DyBa2CusOT.x films by MBE) t h a t copper was depositing primarily with (100) and (110) planes parallel to the interface. Here growth of metallic copper as islands was also indicated, and confirmed by AES during depositing a total of 150 A of Cu. The EEL spectrum from the clean SrTiOs(100) surface at RT shows [102] losses at 6.5, 10.2 and 13.1 eV, which are attributed to interband valence-band transitions between occupied O 2p bands and unoccupied Ti 3d levels [97], and a loss at 22.8 eV which can be assigned either to an O 2p-Ti 3d transition or to a plasmon loss (a bulk plasmon loss is suggested at 26.4 eV [97]), and a loss at 28.8 eV which may also be a plasmon loss. Upon deposition of Cu two new peaks appear at 4.5 and 7.0 eV [102]. These are assigned to an interband transition with Cu 3d as an initial state and to a copper plasmon loss, respectively. The plasmon losses and interband-transition losses are a t t e n u a t e d with increasing de,. The results from UPS (He(I)) show (Figure 32), for the clean surface, energy emission from O 2p levels at 5.2 and 7.3 eV. In the band gap there are some occupied states which are assigned to the He(I) "ghost peak", to carbon states or to Ti 3d states. The O 2p is gradually a t t e n u a t e d with dcu while the copper related to pure Cu 3d states appear as an impurity in the bandgap, close to the valence band maximum.

510

dcu( 10

.,,,-..

5

.E= 2 ..-,,,.

Z

1

0 I

Ef-O

1

1

I

-2 -4 -6 Binding energy (eV)

,.

1

-8

Figure 32. UPS He(I) from Cu/SrTiOs(100). Cu is growing within the bandgap (ref. 102).

The copper emission peak appears at a distance of 2.0 eV and 4.2 eV, respectively, from the main O 2p emission peaks. Within experimental error, no band bending is observed. From the UPS results we also obtain the change in work function A#, found to be 0.4 eV (with a quite large uncertainty of • 0.1 eV) during deposition of Cu until saturation just before dcu = 5 A at the # for pure Cu, 4.48-5.10 eV [107], based upon a measured ~ of 4.2 eV [108] for the SrTiOs(100). The positive A~ and the absence of banding indicate a weak interaction and only little charge transfer, and thus only very weakly bound copper to SrTiOs(100). XPS results [103] found no changes in 8r, Ti or O core levels with doll at RT surfaces, and no evidence of Cu-O compound formation, but instead indication of Cu cluster formation. Analysis of the Cu 2ps ~ binding energy showed a gradual shift of 0.2 eV in binding energy, associated with changing cluster size. By a sputter-profile analysis of a 100-A Cu overlayer the XPS furthermore showed

511 quickly appearance of substrate species and it took longer for the Cu 2p signal to disappear, suggesting an irregular substrate coverage by the coalesced Cu overlayer. No oxygen out-diffusion was observed. Annealing of the 100-/k Cu overlayer at 500~ led to significant changes in atomic distributions with sufficient oxygen diffusion to convert the metal overlayer to an oxide [102]. Also Sr outdiffusion was significant. The annealing caused further clustering in the Cu overlayers. The two investigations [102,103] hence show basically the same conclusion as to the behavior of the Cu deposits during growth at RT and annealing.

8. COPPER AND NICKEL ON WURTZITE STRUCTURF~ The only metal oxide crystallizing in the wurtzite structure is ZnO, but then it has been investigated very intensively for several years due to its high importance particularly in heterogeneous catalysis as substrate for copper, but also in gas sensor, varistor and solar cell applications. In the wurtzite structure the zinc ions are coordinated tetrahedrally with the oxygen ions. It is an insulator, E ~ 3 . 4 eV, Its electronic band structure is quite well investigated [109-111]. The valence band is mainly contributed by O 2p levels and the conduction band by Zn 4s and 4p levels. The Zn 3d band lies below the O 2p band, which is u n u s u a l in comparison to the band structure of other transition metal oxides. Electronic levels within the bandgap of clean ZnO surfaces have not been revealed; in stead, the location of states due to oxygen and zinc dangling bonds have been confirmed within the valence band and conduction band, respectively [ 1 12], corresponding to the 7.4-eV peak in the EEL spectrum [113]. It is conducting as ntype semiconductor in its stoichiometric form. It has equilibrium surface concentrations of intrinsic surface defects, interstitial Zn or oxygen vacancies, which deviate from corresponding bulk values, and extremely small concentrations of surface defects, easily obtained under UHV conditions, Figure 33. LEED from a thermally induce degenerate e t c h e d ZnO(0001) (F_v= 63 eV). accumulation layers

512 leading to "metallic surface conductivity" [ 114]. Investigations on the ZnO crystal surfaces have concentrated on the two hexagonal lo_w-index polar surfaces, Znterminated ZnO(0001) and O-terminated ZnO(0001), that have coor_dinatively u n s a t u r a t e d zinc or oxide ions, respectively, and the non-polar ZnO(1010)surface that has eq_ual numbers of zinc and oxide ions in dimer sites. Most recently, also the ZnO(1120) has been studied. Their electronic and structural properties is well known. Very recently, though, an ordered array of sub-sextets around each of the main hexagonal diffraction spots from a ZnO(0001) surface was observed by LEED after prolonged t h e r m a l t r e a t m e n t at 800 K, Figure 33 [115]. The origin of the observed sub-sextets may be understood in terms of diffracted reflection from a ZnO(0001) surface containing hexagonal pits created by a t h e r m a l etching during the annealing [ 115]. The chemistry, particularly toward CO, of Cu particles or thin films on these catalytically important low-index ZnO surfaces has been object for many investigations over the last decade [116], due to the great importance of the lowpressure synthesis of the basic chemical compound methanol (used for further synthesis of a large group of compounds and for direct fuel as substitute for gasoline) by use of binary Cu/ZnO or, usually, ternary Cu/ZnO/A12Os or Cu/ZnO/Cr20 s catalysts, from carbon monoxide and hydrogen (2H2+CO --> CHsOH), but also in the water-gas shift process (H20 + CO --> H2 + COx) and in methanol steam reforming (CH3OH + H90 -~ CO2 + 3H20). The Cu/ZnO system functions as a gas sensor, as well [117]. We will here concentrate on the geometry and electronic structure during growth of ultrathin layers of copper on the above ZnO crystal surfaces. 8.1. Cu ~m ZnO(0001) Growth of copper on the oxygen-terminated ZnO(0001) surface at RT was studied by XPS, AES, ISS and LEED [118] and also by UPS, XPS, Valence Band PES and REPES [116]. It was found that Cu grows, at least for the first ML (2 A) in a 2-D p(l• 1) overlayer mode (Cu atoms in 3-fold hollow sites created by the outer layer of oxygen atoms), followed by formation of rotational]y-aligned epitaxial 3D Cu(111) microcrystals, randomly distributed over the surface but not observable by LEED before about 3 ML (6-8/k) coverage [116,118]. Increasing temperature favors more agglomeration o[118], producing both sharpened overlayer and substrate patterns for a 8-A surface at 523 K (but only substrate-pattern sharpening for a 0.3-ML surface [116]). The growth can thus be described as following a Stranski-Krastanov-type mode for this surface. This is further supported by the ISS results. Very recently [ll9]_it was found t h a t at 130 K the Cu is cationic at tiny coverages at the ZnO(0001) surface, but becomes nearly neutral at coverages beyond a few percent, forming monolayer-clusters until about 505 surface is covered whereafter the Cu islands grow thicker without filling the gaps between the islands. By annealing to 850 K further clustering takes place. XPS found a mild decrease (about 0.6 eV) of the Cu(2p) binding energy with increasing dcu within the first ML, indicating that the Cu species are not completely metallic [116]. The PES at h v = 120 eV shows an increase in intensity at the top edge of the substrate 0 2p valence band near binding energies of 3-4 eV with increasing dcu, due to Cu 3d states which have a very high PES cross section relative to the O 2p levels at 120 eV, and the Cu 3d levels shift 0.4-0.5 eV to lower binding energy (relative to the ZnO features), similar to a relaxation shift m

513 observed for the core level. Furthermore it is noted that a well-defined Cu 4s peak is not observed, indicating that the low-coverage supported Cu is neither extensively oxidized (4s still present) nor atomic in nature and suggest that the 4s level is somewhat delocalized [116]. The Cu 2p core level binding energies and lack of shakeup satellites indicate that the highly dispersed copper is either Cu ~ or Cu § a point which has been controversial for quite some time, and which is important for elucidating the complex mechanisms in the above synthetic processes. By REPES it was found [116] that the relative energy splitting of the dispersed Cu is not significantly different t h a n in Cu metal, showing that the dispersed copper is Cu ~ and not Cu§ However, detailed analysis of at the satellite peak at the Cu 3p --->4s edge shows that the dispersed copper displays unique REPES features showing t h a t the copper is not purely metallic, atomic or oxidized [116]. The deposition resulted initially in downward, followed by upward, band bending and similarly in the work function for this surface (leading finally to the Cu bulk value). The downward part is indicative of charge donation from the surface to the bulk. However, due to the low number of surface sites expected to be depleted by the Cu, this is expected to have only little effect on the chemical nature of the copper overlayer atoms [116]. The contact potential difference between the Cu and ZnO creates a Schottky barrier of 0.5-0.6 eV as measured by the total change in band bending. 8.2. Gu r ZuO(O001) In the growth of Cu on the_zinc-terminated ZnO(0001) there are many similarities with the Cu/ZnO(0001) case. The similarity of the spectroscopic results (Cu 2psa final-state relaxation shifts and the narrowing of the Cu 2psa peak with dcu, and in the Zn 2pa a normalized intensity ratios) thus point to the same growth mode for at least the first monolayer.However, no Cu overlayer structure was observed. Furthermore, the Cu 3d levels in a He(II) experiment appear at the top edge of the oxide 2p Band, not overlapping the oxide band as strongly as in the Cu/ZnO(0001) case. In REPES it is found again that the copper is not purely metallic, atomic or o_xidized and the data for 0.3-ML Cu is almost indistinguishable from the ZnO(0001) surface. The band bending is only upward and the work function only increasing with dcu, however [116]. Deposition of 0.3 ML of Cu onto the above thermally etched ZnO(0001) surface leads [120] to elimination of the sub-sextet structure and to a broadening and attenuation of the integral diffraction spots, and a new hexagonal diffraction appeared gradually appeared near a coverage of 2 ML and it became predominant at an average of 6 atomic layers almost completely covering the surface, showing an epitaxial Cu(111) ordered island formation, with an observed ratio of the unitcell dimensions in reciprocal space acu/az~o - 1.24 which is slightly smaller t h a n the similar (1.27) p a r a m e t e r for bulk Cu. In a low-energy spectroscopy experiment [120], the target current spectrum [6] (S(E) = dJ(E)/dE vs. incident energy E) at perpendicular incidence from the clean ZnO(0001) surface gives a fine structure, where the maxima may be analyzed in terms of the matching method for determining elastic reflection coefficients [6]. The location of the extrema coincides with the location of band-structure critical points in the Brillouin zone for empty electronic states [ 121,122]. The results demonstrates the expected agreement with calculated density-of-states, DOS between the experimental maxima and the DOS-

514 structure in corresponding critical points (L,M,F,A). A higher-energy m a x i m u m is connected with reflection variation due to diffraction from thermally etched pits [115]. From the TCS we also obtain (peak A, the vacuum-level position of the sample) the work function ~ of the ZnO(0001) surface to be about 3.5 eV, in agreement with the literature value [114]. With Cu deposition of 0.3 ML of Cu we obtain A~ = 0.5 eV and with 6 Cu layers ~= 4.5 ev, i. e. smaller t h a n ~ for bulk C u ( l l l ) , thus showing that the surface is not completely covered by C u ( l l l ) islands. The patch-like coverage is also indicated by the broadening of the primary peak. The dispersed copper on the (0001) surface is readily incorporated into the ZnO lattice as a Cu § site with strong CO chemisorption abilities and is, therefore, a likely possibility for high catalytic activity [116]. 8.3. Cu (m ZnO(1010)

19

1

2

...-,,..

c

t'~

3 v

4 e(D r" m

5 I 40

I 30

I 20 Energ/

1 10 loss

I 0 E L(eV)

Figure 34. EELS from Cu/ZnO (1010). eV. See text (ref. 13).

Ep=97

Deposi_tion of Cu on the ZnO(1010) surface was studied by EELS over the 0 < dc, <- 60 A r a n g e at RT [13]. The most striking results are those from the submonolayer region which is shown in Figure 34. Spectrum 1 is from a clean surface annealed in 10 .5 Tort 02 at 320 ~ and spectra 2 and 3 are obtained after deposition of 0.5 and 1/k of Cu, respectively, on the ZnO(1010) surface of (1). Spectrum 4 shows the clean surface yields loss features at 1.9, 4.0, 7.4, 8.8, 12.5, 18,2 and 24.5 eV. These agrees with previous work [109], except for the 1.9 eV peak which was unfortunately considered as induced by the effect of double derivatives of the loss signal. We note t h a t the 1.9-eV loss structure Vstill remains after the heat t r e a t m e n t in O2 (as did the corresponding 2.6-eV loss structure for the MgO(001) surface [13]). This experimental evidence strongly discards the 1.9-ev loss structure (and the 2.6

515 -eV loss for MgO(001)) as candidates for F, § centers [21], but r a t h e r suggests t h a t a surface V, center, a hole trapped at a metal ion vacancy with 3 oxygen ligands (and 5 oxygen ligands in the MgO(001) case) as the physical origin of the loss structure (although it by no means implies t h a t F, § centers do not exist on oxide surfaces). Upon deposition of 0.5 A of Cu a new structure at 2.0 eV arises, and further deposition of 0.5 A of Cu brings up a weak peak at 4.3 eV (similarly appear peaks at 2.2 and 4.3 eV, respectively). We may ask, w h a t is the nature of these peaks. It is not from the transition Cu 3d ~ 4s since in EELS studies on clean Cu, the loss structure was found. The loss feature disappears from the (1010) system upon deposition of 3.7/k, and it does not appear in the spectrum from a continuous Cu thin film [13]. To elucidate the cause of the peak, the 1-A Cu/ZnO(1010) surface was heated (spectrum 4 in Figure 34) [13], and the peak became enhanced, consequently negating the possibility t h a t the 2-eV peak has originated from small Cu clusters, which have electronic properties different from copper (the peak should have decreased in intensity in case of a very small cluster (it has been reported t h a t bulk metal properties may arise even from clusters consisting only 4 copper atoms [ 123]). In a test experiments of deposition of similar small amounts on a clean Si(111) surface, no 2-eV loss was seen. We believe t h a t the 2-eV loss is the well known loss structure of the semiconductor Cu20, corresponding to a threshold transition from the Cu 3d valence band to the Cu 3s conduction. To support this interpretation of the 2-eV peak, a 1-/k-Cu-deposited surface was exposed to 02 (spectrum 5, Figure 34). This exposure decreased the intensity of the 2-eV peak (and in the MgO case elimination of it completely), confirming t h a t the copper, which corresponds to the 2-eV loss peak, is in the oxidized state. In ZnO where the bandgap is 3.4 eV, the Cu 3d impurity levels overlap the 0 2p band. A Cu 4s level, which has donated one electron, therefore superimposes the conduction band. Hence the 4.3-eV peak of the 1-A Cu/ZnO(1010) probably originates from a copper cluster, as also confirmed by the oxydation experiment (and by the peak-growth at higher coverages [13] (in the MgO case, the 4.3-eV peak may be due to the O 2p --->Cu 3d). We have therefore concluded t h a t the initially deposited copper (at the low end of the sub-monolayer coverages) exists in an ionized state and is bonded to oxygen ions in the surface of the substrate, and t h a t this state probably is the most stable one, Cu(I), corresponding to the induced 2-eV loss peak, and also t h a t the 1.9-eV peak originates from electronic resonance of V, centers acting as copper-deposit trapping-interaction centers. As for the ZnO(0001) no Cu-overlayer superstructure has been observed, but the similar photoelectron spectroscopic evidence as for the (0001) surface was found upon Cu deposition [116]. At 15-20/k coverage a new set of diffraction spots were observed t h a t appeared to be rotationally aligned with the rectangular substrate pattern, but the spots were too diffuse to determine the precise symmetry. The Cu 3d levels develop much the same as on the (0001), with the levels appe_aring at the top edge of the oxide 2p band and not overlapping as strongly as (0001). The band bending and the work function changed with dc~ as for the (0001); thus also for this surface a small charge donation from the surface to the bulk is indicated [116]. Dispersed copper is also on this ZnO(1010) surface, as on the zinc-terminated ZnO(0001) surface, readily oxidized and annealed into the ZnO lattice as a Cu(I) site, a coordinatively u n s a t u r a t e d Cs~ site (that is the only copper center found to

516 adsorb CO w i t h high affinity, 88 kJ/mol, and the dispersed atomic Cu and Cu clusters chemisorb CO with approximately the same affinity as copper m e t a l [116]). UPS He(II) e x p e r i m e n t s carried out in the low 110-130 K t e m p e r a t u r e range [116] have indicated t h a t highly dispersed copper on the ZnO(0001) a n d (1010) surfaces chemisorb CO w i t h a b o u t the same affinity as copper metal, while chemisorption on the Cu/ZnO(0001) was m u c h weaker, a n d high-affinity Co chemisorption, often associated w i t h the catalytic active site, was shown to occur at a coordinatively u n s a t u r a t e d t e d r a h e d r a l Cu § site created on the Cu/ZnO(0001) surface upon a n n e a l i n g in oxygen, and furthermore, chemisorption to the Cu § site p e r t u r b s the CO electronic sstructure m u c h more t h a n chemisorption to either Cu ~ or Zn 2§ due to a stronger a and ~ interactions as indicated from valence-band PES. A Cu-CO active site m u s t provide a lower energy p a t h w a y in the synthesis of m e t h a n o l t h a n a similar Zn2+-CO complex to bring about the k n o w n lower activation b a r r i e r in this synthesis, but u n t i l a definite m e c h a n i s m has been elucidated for this reaction, the n a t u r e of copper activation will be dab_ated [116]. L a t e r [124], in a TPD investigation of CO adsorption on Cu/ZnO(0001) at 130 K, a strong CO TPD-peak at about 160K was found for thicker Cu films (or films t h a t have been a n n e a l e d to high t e m p e r a t u r e s to induce 3D clustering), characteristic for adsorption sites t h a t are Cu(111)-like, and loss in CO adsorption capacity was found to be not as great as rthe loss in Cu surface area. This was i n t e r p r e t e d [124] as caused by a CO-induced redispersion of 3D Cu clusters into 2D_islands. A TPD investigation was recently carried out on the Cu/ZnO(1010) surface as well [125]. At RT, submonolayer a m o u n t s were found to enhance the ability of ZnO to adsorb CO drastically. A small a m o u n t of C02 desorbed from the surface as a result of CO reaction w i t h surface oxygen. A desorption order of 2 a n d a desorption activation energy of a_bout 60 kJ/mole were obtained for CO desorption from the Cu-deposited ZnO(1010) surface, which is in good a g r e e m e n t w i t h the binding energy of CO adsorbed on Cu(100) surfaces. Also the effect of light was investigated on one of these lowest-indexed surfaces .......

e-

=. 1.0 r

= 0.5 m

~PJ

i

0

I

I

i

I

I

500 Time of flight, t

i

I

i

I

1000 (ps)

Figure 35. Time-of-flight distributions of CO2 from CO/Cu/ZnO(0001). (1) clean ZnO(0001); (b) dcu = 0.3 ML (ref. 126).

517 [126]. In the adsorption at RT of CO Cu/ZnO(0001) a double-component time-offlight distribution in laser-induced desorption of CO2 was observed (Figure 35) from CO adsorbed on ZnO(0001), and a strong increase of the a m o u n t and a lowering of the threshold for C09 desorption was caused by a 0.3-ML Cu deposition. m 8.4. Cu cm Z n O ( l l 2 0 ) The growth of Cu on ZNO(1120) and the related c h a n g e s in the electronic structure were investigated for deposits until dcu= 40 A [127]. At RT, the growth shows initially a linear decay with dc,, for dc, < 1.2 A, and then an exponential decay, for 3 < dc, < 40 A (Figure 36), agreeing with the monolayer-simultaneousmultflayer (MSM) model (as classified by Argile and Rhead [128]) for thermodynamic unstable growth with a pre-mono]ayer breakpoint. The model is close to a Volmer-Weber type growth mode, i.e. "islanding". The curve for ideal growth without a pre-monoloayer breakpoint, shown as a thin line in Figure 39, exhibits a break at dc,- 2.0 A (it could perhaps be argued t h a t two breaks may be seen prior to the exponentially decaying curve, but that conclusion would lead to drastic discrepancies in the attenuation lengths of the layers). We consider the 2 A to be the deposition thickness corresponding to monolayer coverage. This growth mode is similar to wh_at very recently has been found for Cu growth on the oxygenterminated ZnO(0001) surface [ 119], and to some other meta]/meta]oxide surfaces as well [128,11,102], and one notices that the pre-monolayer breakpoint obtained by Ernst et al. [119] for the (0001) surface practically coincides with the present value. We did not see any Cu superstructure, only gradual attenuation of the substrate pattern, in agreement with the MSM growth mode. The EELS results for the Cu growth on this surface [127] agrees with our previous results [13]_for Cu growth on the ZnO(1010) surface at RT. Coverage-dependentinterband 6O transitions were allocated at 2.2+0.1 and 4.2+0.1 eV, and a Cu-related plasmon increases in energy from 6.7 to 7.0 eV itq~~ 40 with dcu. The plasmon ""40 ~176 frequency COp did not for this surface show a significant shift with dcu at RT for dcu < 6/k. -20 . . . . . At higher coverages, COp approaches the 7.1-eV value of clean Cu, and this behavior at Oi II I I I I I I RT we believe is primarily 0 10 20 30 40 reflecting the surface-to-bulk dcu (A) development of the copper film. To further elucidate the Figure 36. Change of O(KL2sL2s) intensity with behavior of the Cu deposits, dcu from growth of Cu on ZnO(1120) at 300 K.. synchrotronInsert shows submonolayer range (ref. 127).

t/

1

518

1.2 ~Efi 0.8 D9

.._.,.

,. 0.4

-..

.... ----.-

& I

,

I

.

I

4

I

I

I

I

--_.~ I

12

8

__

I

16

dcu (~.1

Figure 37. Cu/ZnO(1120). Changes with dcu in initial- and final-state contributions to AEB(Cu 2psi) (ref. 127).

6.8 ~

,

,

,

i

,

6.4

6.0 ...-..

56 t,~

~. 4.8 "

B

Ill

4.4

9

2.2

A 9

1.8

3

~

4

~

9

5

.

6

dcu (A)

Figure 38. EELS energy shifts with dcu from a 6-A Cu/ZnO(1120) when heated to 875 K. Ep= 98 eV. (Ref. 127).

radiation based investigations were carried out [127] for depositions over the 0 < dcu < 18 A r a n g e . Core-level spectroscopy showed EB(Cu 2Psi) reaching the bulk Cu value at dcu= 17.8/k., and the Auger Ekin(Cu LsM4,fM4,5) reached similarly the bulk Cu value after a monotonic change of about 4 eV, and we noted a narrowing of this peak with increasing dcu, in agreement with an increasing cluster diameter. This observed shift in Auger kinetic energy is given by the difference between the energy shifts of the one-hole initial state and the two-hole final state. A r a t h e r strong increase in E Bwith dcu may be expected from very small supported metal clusters, however, because of additional finalstate effects AE~ compared to the bulk metal reference state [129]. If we assume t h a t the levels involved in the transition are localized and t h a t the screenings of the first and second hole are similar [129], then the change in the Auger p a r a m e t e r Aa= AF~n(]kl) + AEB0) is related to the shift in the final state by which also holds for metallic clusters. Figure 37 shows the changes AEi, (/) and AE~(]) with dcu, and we see, perhaps with exception of the lowest end of the coverage range, t h a t the observed shift entirely is of final-state nature. For this surface we therefore conclude t h a t the combined stability of the initial-state shift and the plasmon

519 frequency indicate t h a t the shift is due primarily to a size effect, and t h a t the combined use of the two effects is very valuable in elucidating the behavior of ultrathin film electronic properties. If we anneal the Cu-deposited ZnO(ll2) surface, however, a quite different behavior occurs in the electronic structure in the Cu-deposited film with increasing dcu. By annealing a 6-A Cu/ZnO(1120) surface to 875 K for 45 rain at 4• s Pa we obtain [127] a sharp Z n O ( l l 2 0 ) - ( l • LEED p a t t e r n appears, and AES analysis shows a signal corresponding to only 0.2 ML, consistent with a strong agglomeration of the copper particles. In the corresponding EELS we observe [ 127] clear copper features with a strong energy loss at 5.7 eV and absence of both the Cu plasmon at 6.7 eV and the oxygen dangling-bond surface loss at 7.5 eV. Upon deposition of Cu onto this surface, a shift of the 5.7-eV peak (Figure 38, curve C) toward higher energy is observed, and for dcu > 5 ,~ the loss energy is equal to the COpvalue of the non-annealed surface. The free-electron plasmon relation applied to this shift gives a charge transfer of 0.28 e per adatom which is exactly the numerical Cu --~ O charge transfer obtained theoretically for Cu20 [130], and it is therefore reasonable to consider the 5.7-eV loss peak as originating from Cu(I) species at the surface or at interstitial or substitutional positions in the top layer. The considerable agglomeration of Cu and the absence of the O dangling bond cause us to suggest substitutional O-diffusion in the surface at elevated temperature, leading to vacancies in the substrate and oxydation of Cu. This is somewhat in contradiction to the work of Didziulis et al. [116] and Ernst et al. [119] who both reported metallic Cu at elevated temperatures. 8.5. Ni on ZnO(0001) a n d ZnO(0001) Contrary to the Cu/ZnO systems, the Ni/ZnO systems have not been controversial. The deposition of nickel onto ZnO surfaces have been much less studied, perhaps because of the smaller range in applications. An investigation by AES and UPS was carried out already early [1_31] on the deposition over the 0.05 to 5 ML of Ni on the ZnO(0001) and ZnO(0001) surfaces at RT. With the sample at RT, the Ni film is found to grow in layer form on both surfaces. The width of the Ni 3d band as seen in He(I) and He(II) UPS has been developed when d~i has reached 1 ML, indicating that the dispersion of the 3d bands for a 2-D Ni film is as great as for 3-D bulk Ni. Peaks at 6 and 4.3 eV is tentatively interpreted as due to Ni chemisorbed to oxygen from the substrate, and peaks at 2.1 and 3.1 ev are attributed to Ni atoms and 2-D Ni clusters, implying a relaxation shift of 2.5 eV for the Ni atom on the Zn-terminated surface. On the oxygen-terminated ZnO(0001) surface an upward band bending of 0.6 eV with respect to the fiat band case is found after Ni deposition to compensate for the contact potential difference. On the zinc-terminated ZnO(0001) surface, nearly no band bending was found.. This is explained by a dipole layer between the surface Zn atoms and the Ni atoms of the first layer having the Ni atoms negatively charged. On the oxygen-terminated face, and not on the zinc-terminated face, oxygen diffusion into the Ni layer is observed, which could be facilitated and perhaps becomes possible only by this band bending.

520 9. COPPER ON FLUORITE METAI~OXIDE S T R U C r U R K S Last we consider the class of fluorite (CaF2) structures. Among the fluoritestructured metal-oxide crystal surfaces only U02 and yttria-stabilized Zr02 have been investigated with regard to metal adsorption., and so far none with Cu or Ni. We have recently investigated Cu adsorption onto yttria-stabilized Zr02 [132]. A pure Zr02 crystal is monoclinic at RT. At very high t e m p e r a t u r e s (2640-2950 K) a bulk cubic phase exist. This cubic phase can, however, be stabilized to lower t e m p e r a t u r e s by addition of certain oxides, such as Y20s (yttria), CaO and MgO, to the pure ZrO2. Yttria-stabilized cubic Zr02, YSZ, has been found to be a suitable substrate for silicon-on-insulator device technology [133] where it may well replace silicon-onsapphire (SOS) devices. It is a defect solid (Y20s)m(ZrO2)l.m solution which maintains a cubic structure over the 0.08 < m < 0.4 range, and the lattice constant increases almost linearly with m from 5.13 to 5.18 A [134]. The defects are oxygen vacancies created to preserve lattice neutrality when y+s ions are substituted for Zr 4§ ions in the CaF2-type structure. These vacancies give rise to high oxygen ion mobilities, and various oxygen sensors and solid electrolytes made of stabilized ZrO2 are based upon this ion-conducting property. F u r t h e r m o r e it is used as a substrate or as bufferlayer for YBa2CusOT. ~ superconductor thin films and as a support for Rh in heterogeneous catalysis such as in synthesis of ethanol from syngas. There are contradictory assignments of the electronic bandstructure scheme in the literature. There has been studies on polycrystalline surfaces, most recently by Wiemhffer et al. [135], and the reported bandgap varies from 3.7 eV [136] to 7.2 eV [137]. We have studied the clean yttriastabilized ZrO2(100) surface with the common composition m = 0.10, using AES, EELS and LEED, and followed the growth of Cu onto t h a t surface over the 0 _< dcu < 50 /k deposition range [132]. The surface was initially cleaned by heat t r e a t m e n t to 1370 K, causing Zr some loss of oxygen from the surface layers (without heati i I 1 t r e a t m e n t to high t e m p e r a t u r e it 220 200 180 160 140 was not possible to obtain LEED Energy loss (eV) p a t t e r n s due to charging). Due to strong overlap of the Zr and Y Figure 39. Core-level EELS from MNN-transitions, core-level EELS yttria-stabilized cubic ZrO2(100). is preferred (Figure 39) for Ep= 514 eV (ref. 132). identification for YSZ [138], although further refinement is ..-....

e~

u..I

...-... LLI

-...... Z

521 1.o . 9

0.8

9-= 0.6

O-KLL Zr-MNN

OlD~ ~

~ o

g

.,-

0.4

O

z

0.2

0

"

0

10

"

20

30

-

40

50

Depositionthicknessdcu (,~,) Figure 40. Changes in O(KLL) and Zr(MNN) normalized Auger intensities with dcu for yttria-stabilized ZrO2(100) (ref. 132).

needed to extend it into a quantitative method for surface composition analysis. As seen from the normalized AES Zr(MNN) and O(KLL) exponential attenuations with dcu (Figure 40), the growth at RT follows the MSM mode (due to low mobility of the deposited Cu the mode here basically belongs to the Volmer-Weber mode, i. e. pure "islanding"). The LEED pictures show that the surface is of a p ( l • simple bulk-truncation structure.The EELS from the clean surface exhibited strong loss features at 7, 14.5, 26, 34 and 40 eV, respectively. The assignment of the 14.5-eV loss is controversial in the literature since it has been interpreted either to a volume plasmon [139,140], to an interband of 2p to 3s states of oxygen in the conduction band [141] or to a collective excitation of more ligand np electrons in the valence band towards ligand (n+l)s states in the conduction band [142], coinciding somewhat with ref. 138. We therefore performed a structure calculation for an YSZ(100) cubic crystal, with Y doped uniformly in the lattice, and obtained a plasmon energy cop of 21.0 eV (assuming t h a t the valence electrons include 2p electrons of O and 4d and 5s electrons of Zr and Y). Analyzing the EEL spectra we find t h a t the energy loss at 9.5 eV decreases with increasing Ep, and disappears for E, > 250 eV, and this is believed to be one of the characteristics of a surface plasmon. This point supports an assignment of the 14.5-eV loss to a volume plasmon. Furthermore, an analysis of the dielectric function for Zr02 shows t h a t the 14.5-eV loss has character of a collective excitation of valence electrons, hence the assignment to a single-electron transition is not reasonable. Upon Cu deposition, the 14.5-eV loss shifts downwards in energy, for dcu > 6 A, and the relative intensity of the 7.8- and 10.0-eV losses changes drastically. The losses at 7.0 and 10.0 eV are clearly identified as due to the well-known Cu surface and bulk plasmons, respectively [132].

522 The electronic structure of the clean yttria-stabilized Zr02(100) hence has been more clarified, a LEED pattern has been obtained from this surface, the changes caused by Cu deposition have been elucidated, and the mode for the growth of copper clusters on the surface at room temperature has been determined.

Acknowledgements The cooperation with my coworkers in this work, I. Alstrup, J.E.T. Andersen, B. Ealet, Q. Ge, E. Gillet, L. Gui, Q. Guo, J.-W. He, S.A. Komolov, E.F. Lazneva, F. Matthiesen, J. Nerlov, H.N. Waltenburg and M.-C. Wu, the skillful drawings by H. Vib~k, and the support from the Danish Natural Science Research Council through Center for Surface Reactions and through a synchrotron-radiation research grant, the Carlsberg Foundation, the Thomas B. Thrige Foundation and the Elkraft Corp. are gratefully acknowledged. I thank the Research Institute of Electronics for providing excellent conditions and for fruitful comments by Y. Fukuda during a sabbatical stay at Shizuoka University, Japan.

Referenoe~ *) Present address: Research Institute of Electronics, Shizuoka University, H a m a m a t s u 432, Japan. 1 V.E. Henrich, Rep. Progr. Phys. 48 (1985) 1481. 2 L.C. Dufour and M. Pedereau, in: Surface and Near-Surface Chemistry of Oxide Materials, eds. J. Nowotny and L.-C. Dufour, Elsevier, Amsterdam, 1987. 3 P.A. Cox, Transition Metal Oxides: An Introduction to Their Electronic Structure and Properties, Clarendon Press, Oxford, 1992. 4 J.E.T. Andersen and P.J. Moller, Surf. Sci. 258 (1991) 247. 5 G. Ertl and J. Kiippers, Low Energy Electrons and Surface Chemistry, VCH Publ., Weinheim, 1985. 6 (a) S.A. Komolov, Total Current Spectroscopy of Surfaces, Gordon and Breach, Philadelphia, 1992. (b) P.J. Moller and M.H. Mohamed, Vacuum 35 (1985) 29. 7 T.M. French and G.A. Somorjai, J. Phys. Chem. (1970) 2489. 8 P.J. Moller and J.-W. He, Nucl. Instrum. Meth. Phys. Res. B 17 (1986) 137. 9 J.E.T. Andersen and P.J. Moller, J. Vac. Sci. Technol. A10 (1992) 497. 10 W. Wei, J. Vac. Sci. Technol. A6 (1988) 2576. 11 J.-W. He and P.J. Moller, Surf. Sci. 178 (1986) 934. 12 I. Alstrup and P.J. Moller, Appl. Surf. Sci. 33/34 (1988) 143. 13 J.-W. He and P.J. Moller, Surf. Sci. 180 (1987) 411. 14 H.M. Neergaard, MSc thesis (in Danish), Univ. Copenhagen, 1992. 15 P.J. Moiler and H.N. Waltenburg, to be publ. 16 M.-C. Wu and P.J. Moller, Surf. Sci. 279 (1992) 23. 17 D.G. Lord and M. Prutton, Thin Solid Films 21 (1974) 341. 18 K. Takayanagi, K. Yaki and G. Honjo, Thin Solid Films 48 (1978) 137. 19 T. Conrad, J.M. Vohs, P.A. Thiry and R. Caudino, Surf. Interface Anal. 16 (1990) 446.

523 20 C. Noguera, J. Goniakowski and S. Bouette-Russo, Surf. Sci. 287/288 (1993) 188. 21 J.-W. He and P.J. MoUer, Chem. Phys. Lett. 129 (1986) 13. 22 N.C. Bacalis and A.B. Kunz, Phys. Rev. B 32 (1985) 4857. 23 C. Benndorf, H. Caus, B. Egert, H. Seidel and F. Thieme, J. Electron Spectrosc. Related Phenomena 19 (1980) 77. 24 (a) M. Prutton, J.A. Ramsey, J.A. Walker and M.R. Welton-Cook, J. Phys. C: Solid State Phys. 12 (1979) 5271. (b) M. Prutton, J.A. Walker, M.R. Welton-Cook, R.C. Felton and J.A. Ramsey, Surf. Sci. 89 (1979) 95. 25 A.R. Protheroe, A. Steinbrunn and T.E. Gallon, J. Phys. C: Solid State Phys. 15 (1982) 4951. 26 E.V. Stephanova, V.S. Stepanyuk, M.N. Rogaleva, O.V. Farberovich, A.A. Grigorenko and V.V. Mikhailin, Sov. Phys. Solid State 30 (1988) 1329. 27 R.P. Furstenau, G. McDougall and M.A. Langell, Surf. Sci. 150 (1985) 55. 28 Q.-G. Zhu, A.-D. Zhang, E.D. Williams and R.L. Park, Surf. Sci. 172 (1986) 433. 29 Y.W. Chung, W.J. Lo and G.A. Somorjai, Surf. Sci. 64 (1977) 588. 30 W. GSpel, J.A. Anderson, D. Frankel, M.Jaehnig, K. Phillips, J.A. Schfifer and G. Rocker, Surf. Sci. 139 (1984) 333. 31 M.H. Mohamed, H.R. Sadeghi and V.E. Henrich, Phys. Rev. B 37 (1988) 8417. 32 G.N. Raikar, P.J. Hardman, C.A. Muryn, G. van der Laan, P.L. Wincott, G. Thornton and D.W. Bullett, Solid State Commun. 80 (1991) 423. 33 A.K. See, M. Thayer and R.A. Bartynski, Phys. Rev. B 47 (1993) 13722. 34 P.J. Moller and M.-C. Wu, Surf. Sci. 224 (1989) 265. 35 C.C. Kao, S.C. Tsai, M.K. Bahl, Y.W. Chung and W.J. Lo, Surf. Sci. 95 (1980) I. 36 J. Nerlov, Q. Ge and P.J. Moiler, to be published. 37 Z. Zhang, S.-P. Jeng and V.E. Henrich, Phys. Rev. B 43 (1991) 12004. 38 S.C. Wang and G. Ehrlich, Phys. Rev. Letters 62 (1989) 2297. 39 U. Diebold, J.-M. Pan and T.E. Madey, Surf. Sci. 287/288 (1993) 896. 40 M.-C. Wu and P.J. Moller, Surf. Sci. 224 (1989) 250. 41 S.A. Lindgren and L. Walld~n, Phys. Rev. B 22 (1980) 5967. 42 M.-C. Wu and P.J. Moller, Phys. Rev. B. 40 (1989) 6063. 43 M.-C. Wu, Q.-L. Guo and P.J. Moiler, Vacuum 41 (1990) 1418. 44 M.-C. Wu and P.J. Moller, Chem. Phys. Lett. 171 (1990) 136. 45 R.G. Egdell, S. Eriksen and W.R. Flavell, Solid State Commun. 60 (1986) 835 46 M.-C. Wu and P.J. Moiler, Surf. Sci. 235 (1990) 228. 47 N.R. Avery, Surf. Sci. 111 (1981) 338. 48 K.K. Kleinherbers and A. Goldman, Surf. Sci. 133 (1983) 38. 49 M.A. Vannice, J. Catal. 44 (1976) 152. 50 H. Onishi, T. Aruga, C. Egawa and Y. lwasawa, Surf. Sci. 233 (1990) 261. 51 M.-C. Wu and P.J. Moiler, in The Structures of Surfaces Ill (Springer Ser. Surf. Sci. vol 24), eds. S.Y. Tong, M.A. Van Hove, K. Takayanagi and X.D. Xie, Springer-Verlag Berlin, Heidelberg 1991, p. 652. 52. M.-C. Wu and P.J. Moiler, Surf. Sci. 250 (1991) 179. 53 (a) H.H. Madden, J. Kiippers and G. Ertl, J. Chem. Phys. 58 (1973) 3401.

524 (b) K. Christmann, 0. Schober and G. Ertl, J. Chem. Phys. 60 (1974) 4719. 54 C.-C. Kao, S.-C. Tsai and Y.-W. Chung, J. Catal. 73 (1982) 136. 55 V.E. Henrich, G. Dresselhaus and H.J. Zeiger, Phys. Rev. Lett. 36 (1976) 1335. 56 S. Bourgeois, D. Diakit~, F. Jomard, M. Perdereau and R. Poirault, Surf. Sci. 217 (1989) 78. 57 A. Benninghoven, Z. Phys. 230 (1970) 403. 58 S. Bourgeois, F. Jomard and M. Perdereau, Surf. Sci. 249 (1991) 194. 59 S. Ciraci and I.P. Batra, Phys. Rev. B 28 (1983) 982. 60 P.J. Moller and Q. Guo, Thin Solid Films 201 (1991) 267. 61 C.C. Chang, J. Appl. Phys. 39 (1968) 5570. 62 S. Xia, C. Guo, L. Lin and D.E. Ellis, Phys. Rev. B 35 (1987) 7671. 63 J. Olivier and R. Poirier, Surf. Sci. 105 (1981) 347. 64 E. Gillet, B. Ealet and J.L. Berlioz, Surf. Interface Anal. 16 (1990) 461. 65 Q. Guo and P.J. Moller, Surf. Sci. 244 (1991) 228. 66 C.A.M. Mulder and J.T. Klomp, J. Phys. (Paris) 46 (Suppl. C4) (1985) 111. 67 G. Katz, Appl. Phys. Lett. 12 (1968) 161. 68 G. Katz, J. Mater. Sci. 5 (1970) 736. 69 A.G. Schrott, R.D. Thompson and K.N. Tu, MRS Symposia Proceedings, vol. 60, Materials Research Society, Pittsburgh, 1986, p. 331. 70 J.E.E. Baglin, A.G. Schrott, R.D. Thompson, K.N. Tu and A. Segmiiller, Nucl. Instrum. Methods Phys. Res. B 19/20 (1987) 782. 71 K.H. Johnson and S.V. Pepper, J. Appl. Phys. 53 (1982) 6634. 72 K. Nath and A.B. Anderson, Phys. Rev. B 39 (1989) 1013. 73 Q. Guo, P.J. Meller and L. Gui, Acta Phys. Polon. 81 (1992) 647. 74 Q. Guo and P.J. Moller, Vacuum 41 (1990) 1114. 75 M. Gautier, J.P. Dureaud and L. Pham Van, Surf. Sci. 249 (1991) L327. 76 S. Varma, G.S. Chottiner and M. Arhab, J. Vac. Sci. Tech. A 10 (1992) 2857. 77. J.G. Chen, M.L. Colaianni, W.H. Weinberg and J.T. Yates, Jr., Surf. Sci. 279 (1992) 223. 78 J.T. Klomp, in: Ceramic Microstructures '86, Role of Interfaces, eds. J.A. Pask and A.G. Evans, Plenum, New York, 1988, p.307. 79 Q. Zhong and F.S. Ohuchi, J. Vac. Sci. Technol. A 8 (1990) 2107. 80 V. Vijayakrishan and C.N.R. Rao, Surf. Sci. Lett. 255 (1991) L516. 81 A.D. Zdetsis and A.B. Kunz, Phys. Rev. B. 32 (1985) 6358. 82 D.W. Goodman, R.D. Kelley, T.E. Madey and J.T. Yates, Jr., J. Catal. 63 (1980) 226. 83 J.G. Chen, J.E. Crowell and J.T. Yates, Jr., Surf. Sci. 185 (1987) 373. 84 B. Ealet, E. Gillet, V. Nehasil and P.J. Moller, to be published. 85 M. Ricci, thesis, Univ. d'Aix-Marseille III, 1991. 86 C.C. Kao, S.C. Tsai, M.K. Bahl, Y.W. Chung and W.J. Lo, Surf. Sci. 95 (1980) 1. 87 R.L. Kurtz and V.E. Henrich, Surf. Sci. 129 (1983) 345. 88 C. Sanchez, M. Hendewerk, K.D. Sieber and G.A. Somorjai, J. Solid State Chem. 61 (1986) 47. 89 R.J. Lad and V.E. Henrich, Surf. Sci. 193 (1988) 81. 90 P.J. Moller and F. Matthiesen, to be published. 91 Q. Guo and P.J. Moiler, to be published.

525 92 M. Liehr, P.A. Thiry, J.J. Pireaux and R. Caudano, J. Vac. Sci. Technol. A 2 (1984) 1079. 93 P.A. Thiry, M. Liehr, J.J. Pireaux and R. Caudano, Phys. Rev. B 29 (1984) 4824. 94 L.H. Dubois, Surf. Sci. 119 (1982) 399. 95 T. Conrad, J. Ghijsen, J.M. Vohs, P.A. Thiry, R. Caudano and R.L. Johnson, Surf. Sci. 265 (1992) 31. 96 A.P. Baddorf and J.F. Wendelken, Surf. Sci. 256 (1991) 264. 97 V.E. Henrich, G. Dresselhaus and H.J. Zeiger, Phys. Rev. B 17 (1978) 4908. 98 K.C. Mishra, K.H. Johnson and P.C. Schmidt, J. Phys. Chem. Solids 54 (1993) 237. 99 R. Courths, B. Cord and H. Saalfeld, Solid State Commun. 70 (1989) 1047. 100 T. Hikita, T. Hanada, M. Kudo and M. Kawai, Surf. Sci. 287/288 (1993) 377. 101 A.D. Baden, P.A. Cox, R.G. Egdell, A.F. Orchard and R.J.D. Willmer, J. Phys. C: Solid State Phys. 14 (1981) L1081. 102 J.E.T. Andersen and P.J. Moller, Thin Solid Films 186 (1990) 137. 103 D.M. Hill, H.M. Meyer III and J.H. Weaver, J. Appl. Phys. 65 (1989) 4943. 104 J.E.T. Andersen and P.J. Moller, Appl. Phys. Lett. 56 (1990) 1847. 105 R.C. McCune and P. Wynblatt, J. Am. Ceram. Soc. 66 (1983) 111. 106 C. Webb, S.-L. Weng, J.N. Eckstein, N. Missert, K. Char, D.G. Schlom, E.S. Hellman, M.R. Beasley, A. Kapitulnik and J.S. Harris, Jr., Appl. Phys. Lett. 51 (1987) 1191. 107 J. HSlzl and F.K. Schulte, Solid Surface Physics, Springer, Berlin 1979, p. 87. 108 Y.W. Chung and W.B. Weissbard, Phys. Rev. B 20 (1979) 3456. 109 R. Dorn, H. Liith and M. Biichel, Phys. Rev. B 16 (1977) 4675. 110 W. GSpel, J. Pollmann, I. Ivanov and B. Reihl, Phys. Rev. B 26 (1982) 3144. 111 P.J. Moller and J.-W. He, Surf. Sci. 162 (1985) 209. 112 W. GSpel and U. Lampe, Phys. Rev. B 22 (1980) 6447. 113 G. Heiland and H. Liith, in: The Chemical Physics of Solid Surfaces and Heterogeneous Catalysis, vol. 3, eds. D.A. King and D.P. Woodruff, Elsevier, Amsterdam, 1984. 114 K. Jacobi, G. Zwicker and A. Gutmann, Surf. Sci. 141 (1984) 109. 115 P.J. Moller, S.A. Komolov and E.F. Lazneva, Surf. Sci. in press. 116 S.V. Didziulis, K.D. Butcher, S.L. Cohen and E.I. Solomon, J. Am. Chem. Soc. 111 (1989) 7110. 117 A.R. Raju and C.N.R. Rao, Sensors and Actuators B 3 (1991) 305. 118 C.T. Campbell, K.A. Daube and J.M. White, Surf. Sci. 182 (1987) 458. 119 K.H. Ernst, A. Ludviksson, R. Zhang, J. Yoshihara and C.T. Campbell, Phys. Rev. B 47 (1993) 13782. 120 P.J. Moller, S.A. Komolov and E.F. Lazneva, to be published. 121 S.A. Komolov and V.N. Strocov, Phys. Stat. Sol. B 167 (1991) 605. 122 I. Sch~ifer, M. Schliiter and M. Skibowski, Phys. Rev. B 35 (1987) 7663. 123 S. Bloom and I. Ortenburger, Phys. Stat. Sol. B 58 (1973) 561. 124 A. Ludviksson, K.H. Ernst, R. Zhang and C.T. Campbell, J. Catal. 141 (1993) 380.

526 125 Q. Ge and P.J. Moller, to be published. 126 P.J. Moller, S.A. Komolov and E.F. Lazneva, Surf. Sci. Lett. 290 (1993) L677. 127 P.J. Moller and J. Nerlov, Surf. Sci. in press. 128 C. Argile and G.E. Rhead, Surf. Sci. Rep. 10 (1989) 277. 129 G.K. Wertheim, Phys. Rev. B 36 (1987) 9559. 130 S. Evans, D.E. Parry and J.M. Thomas, Farad. Disc. Chem. Soc. 60 (1975) 102. 131 D. Schmeisser and K. Jacobi, Surf. Sci. 88 (1979) 138. 132 Q. Ge and P.J. Moller, Thin Solid Films, in press. 133 H.M. Mannsevit, I. Golecki, L.A. Moudy, J.J. Yang and J.E. Mee, J. Electrochem. Soc. 130 (1983) 1752. 134 V.S. Stubican, R.C. Hink and S.P. Ray, J. Am. Ceram. Soc. 61 (1978) 17. 135 H.D. WiemhSfer, S. Harke and U. Vohrer, Solid State Ionics 40/41 (1990) 433. 136 Z.M. Bluvshtein, G.P. Nizhnikova and U.V. Farberovich, Sov. Phys. Solid State 32 (1990) 548. 137 M. Morinaga, H. Adachi and M. Tsukada, J. Phys. Chem. Solids 44 (1983) 301. 138 J.S Solomon and J.T. Grant, J. Vac. Sci. Technol. A 3 (1985) 373. 139 K.-O. Axellsson, K.-E. Keck and B. Kasemo, Surf. Sci. 164 (1985) 109. 140 C. Palacio, J.M. Sanz and J.M. Martinez-Duart, Surf. Sci. 191 (1987) 385. 141 G.R. Corallo, D.A. Asbury, R.E. Gilbert and G.B. Hoflund, Phys. Rev. B 35 (1987) 9451. 149 G. Marletta and S. Pignataro, Chem. Phys. Lett. 124 (1986) 414.