Progress in Surface Science 73 (2003) 117–165
www.elsevier.com/locate/progsurf
Review
Vanadium oxide surface studies
S. Surnev, M.G. Ramsey, F.P. Netzer
*
Institut f€
ur Experimentalphysik, Karl-Franzens-Universit€at Graz, Universit€atsplatz 5, Graz A-8010, Austria
Abstract
The vanadium oxides can exist in a range of single and mixed valencies with a large variety
of structures. The large diversity of physical and chemical properties that they can thus possess
make them technologically important and a rich ground for basic research. Here we assess the
present status of the microscopic understanding of the physico-chemical properties of vanadium oxide surfaces. The discussion is restricted to atomically well-defined systems as probed
by surface techniques. Following a brief review of the properties of the bulk oxides the
electronic and geometric structure of their clean single crystal surfaces and adsorption studies,
probing their chemical reactivity, are considered. The review then focuses on the growth and
the surface properties of vanadium oxide thin films. This is partitioned into films grown on
oxide substrates and those on metal substrates. The interest in the former derives from their
importance as supported metal oxide catalysts and the need to understand the two-dimensional overlayer of the so-called ‘‘monolayer’’ catalyst. On the single crystal metal substrates
thin oxide layers with high structural order and interesting properties can be prepared. Particular attention is given to ultrathin vanadium oxide layers, so-called nano-layers, where
novel phases, stabilised by the substrate, form.
2003 Elsevier Ltd. All rights reserved.
Contents
1.
Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 118
2.
Some properties of vanadium oxide bulk systems . . . . . . . . . . . . . . . . . . . . . . 119
3.
Vanadium oxide bulk crystal surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 122
3.1. Clean surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 122
3.2. Adsorption studies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 128
*
Corresponding author. Tel.: +43-316-380-5189; fax: +43-316-380-9816.
E-mail address: [email protected] (F.P. Netzer).
0079-6816/$ - see front matter 2003 Elsevier Ltd. All rights reserved.
doi:10.1016/j.progsurf.2003.09.001
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4.
Vanadium oxide thin film surfaces. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.1. Vanadium oxide films on oxide substrates. . . . . . . . . . . . . . . . . . . . . . .
4.1.1. Vanadia on alumina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.1.2. Vanadia on tin-oxide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.1.3. Vanadia on ceria . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.1.4. Vanadia on titania . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.2. Vanadium oxide films on metal substrates . . . . . . . . . . . . . . . . . . . . . .
4.2.1. VOx (x 1) layers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.2.2. V2 O3 layers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.2.3. VO2 and V2 O5 layers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
4.2.4. Vanadium oxide nano-layers . . . . . . . . . . . . . . . . . . . . . . . . . .
134
134
135
137
139
139
144
145
148
154
154
5.
Synopsis and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 161
References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 163
1. Introduction
Vanadium oxides constitute a fascinating class of materials with outstanding
physical and chemical properties [1]. They are used in many technological applications, such as in electrical and optical switching devices, light detectors, critical
temperature sensors, write–erase media, and in heterogeneous catalysis [2,3]. The
interesting solid-state physics of vanadium oxides is centred around phase transitions, in particular metal–insulator transitions as a function of temperature, which
display peculiar structural, electronic, and magnetic behaviour and which still pose
open questions concerning their theoretical description [4–12]. For example, V2 O3
with a insulator-to-metal transition at 160–170 K has been initially regarded as a
prototype of a spin 1/2 Mott–Hubbard compound [4], where electron correlations
dominate the inter-atomic overlap leading to bands, but more recently the concept of
3d-ligand orbital hybridisation and charge transfer systems has been considered
more appropriate [4,8]. However, the electronic structure and the nature of the
magnetic properties of V2 O3 remain controversial and of interest [9–12].
From a chemical viewpoint, vanadium oxides are excellent catalysts used in the
manufacture of important chemicals and in the reduction of environmental pollution, in fact vanadium is the most important metal used in metal oxide catalysis [13].
The rich and diverse chemistry and the catalytic performance of vanadium oxides is
based on two factors, namely the variety of vanadium oxidation states, ranging from
2þ to 5þ , and the variability of oxygen coordination geometries. The latter comprise
octrahedra, pentagonal bypyramids, square pyramids, and tetrahedra, which can be
combined by shared corners, edges and faces, yielding an impressive variety of
structural arrangements. This structural richness is the source for the existence of
differently coordinated oxygen ions, which provide an important ingredient for the
physical and chemical surface properties.
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119
While the physical properties, i.e. the electronic and magnetic structure and their
phase transitions, of vanadium oxides in two dimensions remain yet to be explored,
their chemical performance, namely the catalytic relevance, has steered the scientific
interest towards surface studies of these oxides. Some surface properties of vanadium
oxides have been reviewed previously in the general treatment on the surface science
of metal oxides by Henrich and Cox [1]. More recently, Hermann and Witko [2] have
discussed the physical and chemical behaviour of vanadium oxide surfaces concentrating on theoretical concepts and results. In the present review we will adopt a more
experimental approach, and aim to assess the present status of the microscopic understanding of the physico-chemical properties of vanadium oxide surfaces from the
experimentalists’ viewpoint. We will restrict the discussion to atomically well-defined
surfaces and will neglect the vast literature on the catalytically relevant, but at the
atomic level less well characterised, surfaces of powder samples. In addition to single
crystal surfaces of bulk oxides the surfaces of thin films of vanadium oxides will be
included in the present treatment. Epitaxial thin films of oxides in general are obtaining ever growing importance both in the fundamental science of oxides and in the
high technologies of the upcoming decades. The growth and the surface properties of
thin films of vanadium oxides are therefore discussed prominently in this paper.
The review is organised as follows. In Section 2 the properties of vanadium oxide
bulk systems will be briefly introduced, and in Section 3 the single crystal surfaces of
bulk oxide samples will be investigated. Here, the geometry and the electronic
structure, as probed by various spectroscopies, of the clean surfaces will be discussed
first, and then the chemical reactivity of these surfaces as implied by adsorption
studies will be assessed. Since great difficulties are encountered in preparing single
crystals of vanadium oxides, substantial efforts have been undertaken to prepare
ordered oxide surfaces in the form of epitaxial thin films on suitable substrates.
Section 4 is therefore devoted to the presentation of results of surface studies of
vanadium oxide thin films on oxide and metal substrates. In view of the present-day
trends towards nano-science and -technology a particular focus of this review is
dedicated to the discussion of ultrathin films with thickness of only a few unit cells––
the so-called nano-layers––where novel physical and chemical properties as compared to the respective bulk materials may be encountered. Finally, an outlook with
possible future perspectives in Section 5 concludes this review.
2. Some properties of vanadium oxide bulk systems
The principal oxides of vanadium occur as single valency oxides in the oxidation
states from V2þ to V5þ , that is in form of VO, V2 O3 , VO2 , and V2 O5 . However, the
vanadium–oxygen phase diagram also includes mixed valency oxides containing two
oxidation states, such as V6 O13 with V5þ and V4þ and a series of oxides between VO2
and V2 O3 (e.g. V8 O15 , V7 O13 , V6 O11 , etc.), which contain V4þ and V3þ species [14].
These mixed valency oxides are formed by introducing oxygen vacancy defects into
the respective higher oxides. If the number of oxygen vacancies exceeds a certain
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Table 1
Basic properties of vanadium oxides
Oxide
Oxidation
state
(3d count)
Crystal structure
V2 O 5
VO2
V5þ (3d0 )
V4þ (3d1 )
V2 O 3
V3þ (3d2 )
Orthorhombic (layered)
Rutile (T > TC ), monoclinic
Corundum (T > TC ),
monoclinic
VOx
V2þ (3d3 )
Cubic (rocksalt)
TC (K)
340
168
Magnetic
structure
(T < TC )
Optical band gap
Diamagnetic
Diamagnetic
2 eV
Metallic (T > TC ),
0.7 eV
Metallic (T > TC ),
0.2 eV
Antiferromagnetic
(TN ¼ 168 K)
Magnetic
Metallic to semiconductor (depending on x)
value, the vacancies tend to correlate and form the so-called crystallographic shear
planes, i.e. the vacancies associate along a lattice plane and become subsequently
eliminated by reorganisation of VAO coordination units [1]. The result is a series of
oxides with related stoichiometries, such as the Mgneli phases with Vn O2n1 or the
Wadsley phases with V2n O5n2 formulas.
Some basic physical properties of the principal vanadium oxides are listed in
Table 1. The structure of V2 O5 can be approximated by zigzag ribbons of square
pyramidal VO5 units, which share edges thus building double chains along the bdirection. The chains are connected by their corners and the resulting layers are
stacked along the c-direction ([0 0 1] in Fig. 1). The result are octahedrally coordinated (distorted from regular shape) VO6 units with three principal VAO distances:
(vanadyl oxygen along the c-direction); VAO(2,3) ¼ 1.77–2.02 A
VAO(1) ¼ 1.58 A
Fig. 1. Perspective view of three layers of V2 O5 with am exposed (0 0 1) surface (V atoms dark balls, O
atoms light balls: weak van der Waals bonds omitted for clarity). In the third layer the corner- and edgesharing VO5 space pyramids are shown. From Ref. [51].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
121
(weak bonds in between the
(bridging oxygens in the basal plane); VAO ¼ 2.79 A
layers) (Fig. 1).
VO2 has a tetragonal rutile structure above and a monoclinic structure below 340
K, at which temperature the oxide undergoes an insulator–metal transition. In the
tetragonal lattice the VO6 octahedra are elongated slightly along the fourfold axis,
whereas in the monoclinic lattice the octahedra are more severely distorted yielding
(see insert of Fig. 2).
three VAO distances in the range 1.76–1.87 and 2.01–2.05 A
The adjacent VO6 units are connected by edges or corners in both structures. The
metal–insulator transition produces an abrupt change of the physical properties of
the VO2 system such as a large change in the resistivity or in the transmittance of
infrared radiation. This offers interesting possibilities for practical applications, e.g.
as thermistors and thermal switches or in ‘‘smart thermochromic windows’’ [16,17].
The sesquioxide V2 O3 occurs in the corundum structure above 168 K and in a
monoclinic phase below 168 K. The nature of the metal–insulator phase transition in
V2 O3 and of the associated magnetically ordered insulating phase is still a matter of
intensive experimental consideration and theoretical dispute [1,9,11,12,18,19]. The
vanadium monoxide VO crystallises in the cubic rocksalt structure, where all adjacent VO6 units are connected by their edges. The VO is less stable than the other
Fig. 2. Crystal structure of monoclinic VO2 with netplane stacking along the [0 1 1] direction. All atoms of
the VO2 (0 1 1) surface layer are emphasised by dark shaded balls. V(O) atoms are shown in large (small)
balls. Inequivallent V and O atoms of different coordination are marked. The insert shows a VO6 building
unit with the various VAO distances. From Ref. [15].
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principal vanadium oxides [20] and is prone to non-stoichiometry yielding a defective
rocksalt structure.
The electronic structure of vanadium oxides has been investigated experimentally
by various electron spectroscopic techniques. Since these techniques are in general
surface specific, this topic will be treated in connection with the surface studies on
bulk vanadium oxides in the following section.
3. Vanadium oxide bulk crystal surfaces
3.1. Clean surfaces
Clean surfaces of vanadium oxide bulk crystals can be prepared by cleaving, ion
bombardment followed by annealing in oxygen, or by scraping with a suitable tool,
e.g. a diamond file. Only the two former methods can produce ordered surfaces of
sufficiently long-range order to allow the characterisation by a LEED pattern. V2 O5
can be easily cleaved parallel to the (0 0 1) plane (see Fig. 1), whereas V2 O3 cleaves
along the (1 0 1 2) plane. VO2 cleavage provides a (0 1 1) surface (see Fig. 2), but does
not produce a surface of sufficient quality to show a LEED pattern [5]. However,
single crystals grown by the chemical transport technique have many specular surfaces, which after annealing in situ under oxygen atmosphere (e.g. 106 mbar O2 at
600 C [21]) are ordered displaying sharp LEED patterns.
The LEED patterns observed for V2 O5 (0 0 1), VO2 (0 1 1), and V2 O3 (0 0 0 1) surfaces showed the symmetries as expected from the respective bulk terminated surfaces [21–23]. Quantitative LEED investigations have not yet been performed on
these surfaces and possible surface relaxation parameters have not been determined.
Czekaj et al. [24,25] have addressed theoretically the question of the surface termination and geometric relaxation of the (0 0 0 1) surface of rhombohedral V2 O3 (i.e.
the basal plane of the corundum structure) by ab initio density functional theory
(DFT). They found the smallest ionic charging and the smallest relaxation effects for
a half metal layer V0 AOAV termination and suggested that this termination is the
most stable of the three relaxed bulk-type terminations (VAV0 AO and OAVAV0
being less favourable). Fig. 3 shows models of the (0 0 0 1) surface of V2 O3 with the
various terminations as adapted from Ref. [25]. However, on V2 O3 (0 0 0 1) type
surfaces of vacuum deposited thin films of vanadium sesquioxide a surface layer of
vanadyl V@O groups has been established both experimentally and theoretically as
the most stable surface configuration in a broad oxygen pressure range [26], as
discussed below in Section 4.
The electronic structure of vanadium oxide single crystal samples has been investigated by X-ray absorption spectroscopy using the total electron yield technique
[6,22,28,29], electron energy loss spectroscopy in transmission mode [27], and vacuum-ultraviolet reflectance spectroscopy [5]. These methods give mostly information
on the bulk properties. Photoemission measurements are more surface sensitive and
it is therefore appropriate to discuss them in more detail in this review. Vanadium 2p
core level X-ray photoelectron spectra (XPS) from the cleaved surfaces of the
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123
Fig. 3. Surface terminations of the rhombohedral V2 O3 (0 0 0 1) surface with different surface sites indicated. From Ref. [24].
Fig. 4. V 2p core level spectra of the vanadium oxides V2 O5 , V6 O13 , V4 O7 , V3 O5 and V2 O3 . From Ref.
[30].
principal vanadium oxides from V2 O5 to V2 O3 and of several mixed valent compounds are collected in Fig. 4 [30]. In V2 O5 the V 2p3=2 line is narrow (FWHM ¼ 1.2
eV) and is located at 517.2 eV binding energy (BE). For the other oxides (3d
count >0) this line is significantly broader due to the multiplet splitting of the
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Fig. 5. Theoretical (full line) and experimental (dots) V 2p core level XPS spectra of V2 O3 . Adapted from
Ref. [32].
photoemission final states, which is caused by the 2p core hole–3d electron interaction [31]. For VO2 the V 2p3=2 peak maximum occurs at 516.0 eV (FWHM ¼ 1.95
eV), and for V2 O3 at 515.85 eV (FWHM > 3 eV). The latter width, however, might be
caused to some extent by the occurrence of defects at the surface, as indicated by the
tail of the spectral structure to lower binding energy. The experimental V 2p XPS
spectrum of V2 O3 has been analysed with the help of cluster model calculations by
Uozumi et al. [32], and Fig. 5 compares their experimental and calculated data: the
spin-orbit split 2p3=2 and 2p1=2 peaks are observed together with the O 1s line in the
experimental spectrum. In addition to the main peaks, satellite structures labelled CT
are seen in the theoretical spectrum, which have been interpreted in terms of a charge
transfer (CT) mechanism: both main peak and satellite are strongly mixed states
between 2p1 3dn and 2p1 3dnþ1 L1 final state configurations (2p1 and L1 denote
holes in the V core and oxygen ligand states, respectively). The satellite associated
with the 2p3=2 peak can be seen in the experimental data (S), whereas the corresponding satellite of the 2p1=2 peak is obscured by the O 1s structure.
Early UV photoemission work of the valence bands of vanadium oxide single
crystals have been reviewed by Henrich and Cox [1]. Here, the more recent XPS
valence band spectra as recorded by Zimmermann et al. [8] are shown in Fig. 6. The
spectra have been corrected for an integrated background and display two spectral
regions, which may be vaguely associated with the V 3d derived states between EF
and 2–3 eV below EF and the O 2p valence states from 3–10 eV. V2 O5 has no intensity up to 2 eV below EF , consistent with its 3d0 configuration and the existence
of an optical gap of 2 eV, whereas V2 O3 is metallic with a noticeable intensity at EF .
The spectrum of VO2 has been shifted in energy to place the Fermi level at the top of
the known small energy gap, since the VO2 sample showed strong charging effects
during the measurements at room temperature [8]. The apparent separation of the
photoemission intensity into two spectral regions does not imply, however, that a
simple ionic picture of the vanadium oxides is appropriate; as shown below there is
evidence for considerable hybridisation which mixes the wavefunction character of
these spectral features.
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
125
Fig. 6. Valence band XPS spectra of V2 O3 , VO2 and V2 O5 , taken at room temperature and normalised to
equal height. From Ref. [8].
Fig. 7 compares the He II (hm ¼ 40:8 eV) spectrum of the V2 O5 (0 1 0) surface and
the total and atom projected density of states (DOS) from a DFT cluster calculation
[33], where the energy scale of the theoretical data has been shifted appropriately.
The total DOS reproduces the overall shape of the UPS spectrum quite well, and the
origin of the photoemission peaks can be identified from the partial DOS curves,
which allow to specify the emission from particular subunits of the surface. Accordingly, the most prominent peak in the experimental spectrum is associated with
emission from the terminal oxygen O(1) (see Fig. 1), whereas the peak at 7 eV BE is
derived from a mixture of vanadium with the bridging oxygen O(2) and O(3)
emission intensities (Fig. 1). The structure at 4 eV has been interpreted as a
combination of all three types of surface oxygen emissions. The ability to assign the
different structures in the photoemission spectrum to localised atom centres at the
surface is an important step to specify possible reaction sites at the surface, which
will be modified by the exposure to gas phase molecules and can thus be identified in
the experimental spectra.
The question of the nature of bonding is of interest for both physical and chemical
properties of vanadium oxide surfaces. It is related to the amount of hybridisation
between the O 2p and V 3d valence states, i.e. to the degree of ionic versus covalent
bonding character. Based on theoretical considerations Hermann and Witko [2]
have concluded that the bonding in vanadium oxide compounds has to be described
by a mixed ionic and covalent character of the vanadium–oxygen bonds. The
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Fig. 7. Comparison of the total and atom projected density of states calculated for a V16 O49 H18 cluster
with a photoelectron spectrum (He II radiation, normal emission). From Ref. [33].
experimental verification of V 3dAO 2p hybridisation has been provided by resonant
photoemission and by soft X-ray emission experiments. Resonant photoemission
spectra of V2 O3 excited at photon energies near the V 2p threshold have been
measured by Zimmermann et al. [8] and are reproduced in Fig. 8. The spectra have
been corrected by a Shirley-type background and have been normalised to the intensity of the O 2s peak at 22 eV BE. The V 3d structure at 0–2 eV BE shows the
expected pronounced enhancement at photon energies above the V 2p threshold (see
spectrum recorded with hm ¼ 517:2 eV) due to the resonant photoemission process,
but also the region of the O 2p band at 4–8 eV BE displays enhanced intensity as
compared to the off-resonance spectra (the two bottom curves in Fig. 8), reflecting a
strong admixture of V 3d states in this spectral region. Unfortunately, the spectra
also contain contributions from incoherent LVV Auger processes in this energy
region (marked V-LMM on the figure), which prevent a more quantitative analysis
of the data [8]. Nevertheless, the spectra demonstrate clearly the resonant photoemission enhancement of the O 2p band and confirm the V 3dAO 2p wave function
mixing due to covalent bonding contributions.
Soft X-ray emission spectroscopy (SXES) allows to study the partial components
of the density of valence states localised at an atom with a core hole. The process has
a clear selection rule of the angular momentum, because it is a dipole transition, and
it probes the local electronic structure at a particular atomic species, because the
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127
Fig. 8. Resonant valence band phototemission spectra at the V 2p threshold for V2 O3 . The upper and the
lower spectra correspond to on- and off-resonance conditions, respectively. From Ref. [8].
respective core hole is strongly localised. Shin et al. [34] have applied SXES to early
transition metal compounds, amongst them VO2 . Fig. 9 shows a comparison of
valence band photoemission and V L2;3 (2p) and O 1s SXES spectra of VO2 , as taken
from Ref. [34]. The V 2p X-ray emission reflects the V 3d partial density of states and
the O 1s emission reflects the O 2p partial density of states, as illustrated by the
schematic diagram at the right hand side of Fig. 9. The V 2p SXES spectra show two
structures, which coincide with the photoemission feature at 0–2 eV BE and the
higher BE part of the O 2p valence band (marked B on the figure). This indicates that
the photoemission peak close to the Fermi level is V 3d derived, as well known, but
also that the higher binding energy component of the O 2p band contains V 3d
contributions due to the VAO bonding. The energy position of the O 1s X-ray
emission coincides with the O 2p band at lower BE as well as the V 3d band at EF .
These observations clearly corroborate that O 2p and V 3d states are strongly hybridised with each other.
Whereas the electronic structure of vanadium oxides has been addressed both
experimentally and theoretically as mentioned above, the vibrational structure of
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Fig. 9. Comparison of the valence band photoemission spectrum and the V 2p and O 1s SXES spectra of
VO2 . From Ref. [34].
these oxides has been much less studied on controlled samples, although the phonons
structures of oxide powders as probed by Raman and infrared spectroscopy have
been used successfully in catalysis studies to provide specific fingerprints for catalyst
characterisation [35]. The surface vibrational properties of V2 O5 (0 0 1) have been
investigated in the study of Poelman et al. [36] by high-resolution electron energy loss
spectroscopy (HREELS) and have been analysed within the framework of the dielectric theory. An HREELS spectrum of a cleaved V2 O5 sample, recorded with a
primary energy Ep ¼ 12 eV and a resolution of 4 meV, is displayed in Fig. 10 [36]. To
the left of the elastic peak a group of prominent loss features is seen, which are
marked N1 –N9 on the spectrum, while to the right hand side of the elastic peak
several gain features are recognised. Table 2 contains the vibrational energies of the
loss peaks together with an assignment in terms of their principal vibrational modes,
as derived from an extension of the dielectric loss theory to anisotropic materials [37]
and from a comparison with experimental infrared data [38]. All loss peaks can be
attributed to surface phonon modes of V2 O5 and the loss frequencies correspond
well to infrared absorption bands of V2 O5 powder samples [39], supporting the diagnostic value of the phonon structure in recognising particular oxide phases. The
phonon spectra of vanadium oxides have also been used to characterise the oxides in
ultrathin film form and to determine the surface termination of oxide layers, as it will
be further discussed in Section 4.
3.2. Adsorption studies
Experimental adsorption studies on bulk single crystal surfaces of vanadium
oxides are rather scarse, but several theoretical simulations of adsorption processes
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129
Fig. 10. High-resolution electron energy loss spectrum of a cleaved V2 O5 (0 0 1) surface (Ep ¼ 12 eV,
resolution 4 meV, Hin ¼ Hout ¼ 45). From Ref. [36].
Table 2
Experimental loss peaks and their assignment to principal vibrational modes from Ref. [36]
Loss
N1
N2
N3
Energy (meV)
9.0 ± 1.0
28.0 ± 1.0
34.0 ± 3.0
Assignment
chain translation
VAO(1)a deformation
deformation
N4
N5
44.0 ± 1.0
55.5 ± 3.0
VAO(1) deformation
stretching + deformation
N6
N7
63.0 ± 3.0
72.5 ± 1.0
VAO(2)AV deformation
stretching
N8
N9
a
106.0 ± 1.0
126.0 ± 1.0
VAO(2)AV stretching
VAO(1) stretching
O(1), O(2), O(3) refer to the oxygen atoms specified in Fig. 1.
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have been published in the literature. This reflects the experimental difficulty in
preparing good-quality surfaces of vanadium oxide single crystals.
The adsorption of hydrogen is of particular interest, since hydrogen is ubiquituously present in hydrocarbon reactions, for which vanadium oxides are used as
catalysts [3,13]. An important question regards the formation of surface hydroxyl
groups and of hydrogen induced defects. In this context it is of interest, which type of
geometrically different oxygen atoms, for example on the V2 O5 surface, is removed
by the interaction with hydrogen. Tepper et al. [40] have investigated the adsorption
of molecular and atomic hydrogen on a vacuum-cleaved V2 O5 surface by HREELS,
angle-resolved UPS, and XPS. They found that molecular H2 interacts weakly with
V2 O5 (0 0 1), whereas atomic H strongly reduces the surface. Fig. 11 shows HREELS
spectra of V2 O5 (0 0 1) taken after exposure to increasing amounts of atomic H [40].
While the losses due to the vanadyl group and the triply coordinated oxygen atoms
O(3) (see Table 2) are little affected by the hydrogen, the loss at 105 meV, due to
vibrations involving the bridge-bonded oxygen atom O(2), disappears already after
small H doses. This indicates that atomic H affects the bridging surface O(2) atoms
more strongly than the vanadyl O(1) and bridging O(3) atoms. The loss at 365 meV,
Fig. 11. HREELS spectra of V2 O5 (0 0 1) after adsorption of increasing doses of atomic hydrogen
(Ep ¼ 7:5 eV, Hin ¼ Hout ¼ 65, T ¼ 300 K). From Ref. [40].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
131
visible after large H doses, is not due to surface OH, but has been attributed by the
authors to the CAH stretching mode of organic surface contaminations [40].
XPS core level spectra of V2 O5 (0 0 1) before and after exposure to atomic H are
displayed in Fig. 12 [40]. The V 2p XPS peaks develop a structure at the lower
binding energy (516.0 eV, see insert), which reveals reduction and the formation of
V4þ species. Contrary to expectation the spectroscopic results from the H exposed
V2 O5 surfaces showed no indications of the formation of OH groups. The UPS
spectra revealed no intensity in the region around 9–11 eV, which could be associated
with the OAH r bond, the HREELS data showed no loss near 450 meV due to the
OAH stretching vibration, and the O 1s XPS spectra showed no peaks at 531–532 eV
BE. However, for thin films of VO2 (1 1 0) on TiO2 (1 1 0) and V2 O3 (0 0 0 1) on
Au(1 1 1) substrates stable surface OH groups have been detected by the same group
[40]. Tepper et al. have proposed that stable OH groups do not form on the
V2 O5 (0 0 1) surface for geometric reasons, but the observed reduction of the surface
indicated that there must be an interaction of H atoms with the surface. It has
Fig. 12. XPS spectrum of the O 1s and V 2p core levels for cleaved V2 O5 (0 0 1) before and after exposure
to atomic hydrogen (hm ¼ 625 eV, T ¼ 300 K). From Ref. [40].
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therefore been suggested by the authors that OH groups are formed intermediary,
which are mobile and combine to form water which desorbs from the surface [40].
In contrast to the experimental observations the ab initio DFT calculations of
embedded clusters, simulating the V2 O5 (0 1 0) surface, of Hermann et al. [41] have
predicted the existence of stable OH groups at all oxygen sites, with the strongest
adsorptive binding at the vanadyl O(1) site (3.05 eV), followed by the O(2) (2.76 eV)
and the O(3) (2.50 eV) sites. Yin et al. [42] have performed first-principles DFT
calculations under periodic boundary conditions to rationalise the active oxygen
sites of V2 O5 (0 1 0). These calculations found also that the vanadyl oxygen is the
most active site for hydrogen adsorption among the three kinds of surface oxygen
sites. The desorption ability of OH species from the surface has also been calculated,
and the removal of O(1)AH species, formed by H adsorption at the vanadyl oxygen,
was found to be energetically preferable [42]. Hermann et al. [41] have calculated the
desorption energies to remove surface O, OH, or H2 O, the latter formed by approach
of a second H atom to surface OH, leaving a surface vacancy behind after the desorption process. These authors reported that the removal of surface O always requires more energy than for surface OH, and more for surface OH than for surface
H2 O. At the time of this writing it appears that the question of the most active
surface O site for H adsorption on the V2 O5 (0 1 0) surface is unresolved, and that
further experimental studies are required to settle this discrepancy between experimental and theoretical results.
The adsorption of H2 O on cleaved V2 O3 (1 0 1 2) surfaces has been studied by
Henrich et al. using UPS [1,43]. The formation of OH groups has been reported as
indicated by the OAH r bond signature in UPS at 10–11 eV BE, but since cleaved
V2 O3 (1 0 1 2) surfaces are highly defective, the dissociative adsorption of H2 O almost
certainly occurred at defect sites. Recently, the Henrich group has investigated the
adsorption of H2 O on pure and Cr-doped V2 O3 (0 0 0 1) surfaces by UPS and XPS
[44]. At 180 K adsorption temperature molecular H2 O adsorption has been detected
with a sticking coefficient close to unity for water doses up to 1000 L (1 Langmuir
(L) ¼ 1 · 106 torr s). At 273 K adsorption temperature only dissociative adsorption
of H2 O has been found with a low sticking coefficient of <103 . However, it is likely
that the dissociation of the water molecules occurred also at surface defects such as
steps and kinks in this case. This is corroborated by very recent XPS H2 O adsorption
studies on almost defect-free V2 O3 (0 0 0 1) thin film surfaces on a Rh(1 1 1) substrate
[45], where only molecular water species have been identified.
The molecular and dissociative adsorption of H2 O on V2 O5 (0 1 0) has been investigated theoretically using DFT theory with periodic boundary models by Yin
et al. [46]. According to these calculations molecular adsorption of H2 O on
V2 O5 (0 1 0) occurs favourably both at the exposed pentacoordinated V atoms and at
all three types of surface O sites, with the vanadyl oxygen acting as the most favourable site. Water dissociation hardly occurs on the stoichiometric surface, because of the significant Coubomb repulsion of the lattice oxygen around the exposed
V centre to the approaching hydroxyl species. The water surface bonding has been
described by a combination of both coordination interaction and hydrogen bonding
effects. The chemisorption of water on V2 O5 (0 0 1) has also been addressed by
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
133
semiempirical quantum chemical methods, which are based on modifications of the
INDO (intermediate neglect of differential orbitals) technique [47,48]. Qualitatively
similar results as in the ab initio DFT calculations were reported for the relative
energetic order of adsorption sites, but the absolute energies were different. In the
latter two studies the calculated adsorption energies were compared with experimental data derived from thermal desorption measurements on powder samples.
This comparison is however questionable, since the theoretical energies refer to ideal
single crystalline surfaces whereas the thermal desorption data are dominated by the
desorption from low coordination sites, that are abundantly present on powder
samples.
The adsorption states of ammonia on both Br€
onsted and Lewis acid sites of the
V2 O5 (0 1 0) surface have been investigated by periodic boundary DFT calculations
[49]. The Lewis acid sites are present as the pentacoordinated exposed V atoms on
the bare V2 O5 (0 1 0) surface, whereas the Br€
onsted sites have to be created by hydroxylation of the surface. It was found that NH3 adsorption can take place energetically on both types of surfaces, but that the Br€
onsted sites are favoured. On the
Br€
onsted acid sites ammonium NHþ
4 species are formed when NH3 molecules adsorb
at the vanadyl derived O(1)AH groups, which are the most reactive sites toward
ammonia adsorption. On the non-hydroxylated V2 O5 surface on-top oriented NH3
species on the exposed V sites, the Lewis acid sites, are reported to be most stable
[49], which is in agreement with the conclusions of an earlier semiempirical extended
H€
uckel calculation [50].
The physisorption of ethane and propane has been considered on a V2 O5 (0 0 1)
surface with oxygen vacancies within the framework of a molecular mechanics
simulation by K€
amper et al. [51]. These physisorption sites may be the starting point
for subsequent catalytic reactions. The energetically most favourable adsorption site
is on top of a vacancy of twofold coordinated O(2). It has been proposed that the O
vacancy creates a so-called ‘‘van der Waals cage’’, that may be responsible for the
trapping of the hydrocarbon molecules.
In view of the significance of defects for the reactivity of vanadium oxide surfaces,
the surface reduction and the comcomitant formation of oxygen vacancies, e.g. by
reactive chemisorption and subsequent desorption, is an important topic. Hermann
et al. [52] have studied theoretically the hydrogen assisted oxygen desorption by DFT
from the V2 O5 (0 1 0) surface modelled by embedded clusters, and have examined the
structural and energetic consequences of hydrogen interacting with different oxygen
sites. The presence of hydrogen at the oxide surface facilitates the oxygen removal,
because the hydrogen interaction weakens the binding of the surface oxygen atoms
with the neighbouring V atoms; this weakening is larger for surface H2 O than for
surface OH species, the desorption energies being lowest for the bridging O(3) sites.
Toledano et al. [53] have investigated the adsorption of CO on Cr-doped
V2 O3 (0 0 0 1) by UPS, and have observed a reduction of the surface after large CO
exposures, suggesting that the adsorbed CO reacts with lattice oxygen and desorbs as
CO2 at T > 400 K.
The relationship between the semiconductor–metal phase transition at the surface
of cleaved VO2 single crystals and the adsorption of atomic hydrogen has been
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investigated, using UPS, by Bermudez et al. [54]. Evidence has been found for a
partial reduction of the surface, leading to the formation of adsorbed OH and an
increase in the V 3d electron density; the adsorbed hydrogen seems to impede the
phase transition within the photoemission sampling depth ( 6 two lattice constants).
The thermal reduction by heating air-cleaved V2 O5 (0 0 1) surfaces under an oxygen
atmosphere to 500 C in a vacuum chamber has been characterised by X-ray
photoelectron diffraction (XPD) by Devriendt et al. [55]. Comparison of the V 2p
azimuthal XPD scans with single scattering calculations indicated that the
V2 O5 (0 0 1) has been transformed into V6 O13 (0 0 1) by the heating procedure; this
transformation of V2 O5 into V6 O13 occurs through ordering of oxygen vacancy
defects and the creation of crystallographic shear planes, via a change of the sharing
of VAO coordination units from corner-linked to edge-linked octahedra. The
changes of the surface structure and morphology of vanadium oxides as a result of
reduction processes have not been investigated as yet on bulk single crystal surfaces,
but atomic-level studies performed on thin films of vanadium oxides have revealed
that dramatic morphology changes can occur. These will be reported in Section 4.
4. Vanadium oxide thin film surfaces
Vanadium oxide thin film investigations are undertaken for numerous reasons.
They can provide model surfaces, which are important as the surfaces of the bulk
oxides are difficult to prepare. Moreover, such models on metallic substrates can be
used to avoid charging problems of the single crystal oxides and thus allow the
application of the many electron based surface science techniques. Thin vanadium
oxide films are also interesting in their own right with technological applications
ranging from devices through sensors to protective coatings and catalysis. The nanostructures that arise can be quite novel with properties very different from the bulk
oxides.
4.1. Vanadium oxide films on oxide substrates
Technologically the greatest interest in V-oxides on oxides springs from the importance of such systems in heterogeneous catalysis. They form the most important
supported metal oxide catalysts and can be found used on SiO2 , Al2 O3 , TiO2 and
ZrO2 supports. Catalyst scientists have shown that the support plays a much more
active role than the name implies with both the catalytic activity and selectivity being
affected by the support oxide material. For instance, depending on the support, the
catalytic activity for methanol oxidation can vary over three orders of magnitude.
The precise origin of the mechanism of this metal oxide-support effect is as yet
unknown. As the specific interaction between the support oxide and the vanadium
oxide seems critical, insight into the preparation of the vanadium oxides at the
atomic level is necessary. At present the lack of knowledge of their structure and
composition prevents a rational design of supported vanadium oxide catalysts [13].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
135
Supported vanadia catalysts consist of two-dimensional vanadia overlayers as
well as V2 O5 crystallites above a monolayer coverage. Numerous fundamental
questions are still open including the critical question of the exact nature of the twodimensional overlayer––its structure, oxidation state and the thickness of this socalled ‘‘monolayer’’. The physico-chemical techniques generally employed by the
catalytic community, such as X-ray diffraction or infrared spectroscopy, are not truly
monolayer sensitive. Thus for the catalytic community this ‘‘monolayer’’ is defined
as the amorphous interfacial region as determined by X-ray diffraction [56,57], and
the term ‘‘monolayer catalyst’’ is also used for oxide multilayers. The numerous
surface science techniques are able to address these important issues with wellcontrolled model studies. It must be noted that due to charging problems with
electron-based techniques the choice of the substrate oxide is limited either to conducting bulk oxides such as TiO2 or to epitaxial oxides grown on metals such as
Al2 O3 . Also it is possible that the UHV requirements of most surface sensitive
techniques may result in in-situ grown vanadia being different to those in real catalysts.
4.1.1. Vanadia on alumina
Studies of vanadia deposition on alumina have been performed both on single
crystal a-Al2 O3 (0 0 0 1) [58] and also on thin, well ordered alumina formed by the
oxidation of the surface of NiAl(1 1 0) [59]. Both have involved a number of techniques that gives a degree of confidence in the conclusions drawn. Moreover, the use
of identical reactive evaporation conditions (107 mbar O2 at room temperature)
along with an overlap in techniques used allows a comparison of the two model
alumina substrates. In the synchrotron radiation XPS and near-edge X-ray absorption fine structure (NEXAFS) study of the growth of vanadium oxide from
submonolayer to multilayers on Al2 O3 (0 0 0 1) it was concluded that a stable, conducting V2 O3 layer forms with short range order with the vanadium cations in an
octahedral oxygen coordination. Both the lineshape and energy positions in NEXAFS are characteristic of the vanadium oxidation state [60] and, unlike XPS, do not
suffer uncertainties with interpretation due to charging. As can be seen in Fig. 13 for
increasing coverages the V L2;3 near-edge features grow without significant energy
shifts and have very similar fine structure to those of single crystal V2 O3 . In the very
low coverage regime the NEXAFS lineshape lacks definition, which the authors
suggested may be a result of poor short-range order at the vanadia/alumina interface. In contrast, the V XPS features grow but display a discontinuous shift from low
to high binding energy for increasing coverages. The final multilayer value for the
V 2p3=2 binding energy reached a value of 515.5 eV (charging was compensated for
by a flood gun and the Al 2p level used as a reference), very close to that of single
crystal V2 O3 . Although the low coverage binding energy of 517 eV is very close to
that of bulk V2 O5 , the authors considered this to be a coincidence due to a final state
effect arising from the very small clusters in the low coverage regime. Such a final
state interpretation is supported by the lack of shift in the NEXAFS and has been
elaborated upon in the recent growth study on an alumina thin film support [59]. In
this latter work the conclusions of the previous single crystal Al2 O3 work is
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Fig. 13. Vanadium L-edge NEXAFS and XPS of O 1s and V 2p of a-Al2 O3 (0 0 0 1) for increasing vanadia
coverages from clean to approximately 10 ML. From Ref. [58].
supported and extended with in situ STM and infrared reflection-absorption spectroscopy (IRAS). In the low coverage regime the vanadia overlayer appears as ho,
mogeneously distributed small particles with diameters in the range of 20–30 A
which are coverage independent. Bias dependent apparent height measurements
indicated that the vanadia is incorporated into the alumina surface. The XPS behaviour with growth is essentially identical to that observed in the single crystal
study suggesting that the conclusions from these two model substrates are transferable. Interestingly, the Al-oxide XPS signal also grows with increasing vanadia at
the expense of the metallic Al signal from the NiAl substrate. This growth of the
alumina film during the reactive evaporation of vanadia indicates that the vanadia
particles act as oxidation catalysts.
Particularly telling in this study are the IRAS results as a function of vanadia
coverage as displayed in Fig. 14a [56]. In the low coverage regime the substrate
phonon feature at 866 cm1 is quickly suppressed while the two features which grow
at 1046 cm1 (m1 ) and 945 cm1 (m2 ) saturate in intensity by 1 monolayer. At high
coverages a third feature appears at 715 cm1 (m3 ). These features do not appear on
exposure to metallic V (see Fig. 14b) and must be associated with the vanadia
particles, with m1 and m2 being VAO species localised either at the particle surface or
its interface with the alumina. The bands observed have been assigned from an
empirical correlation between Raman stretching frequencies and crystallographically
determined bond lengths. Band m1 (bond order 2) represents a vanadyl species
(V@O) on the surface of the particles. Although vanadyl groups are not structural
elements of truncated bulk V2 O3 surfaces, CO adsorption proves that they are a
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
137
Fig. 14. (a) IR spectras taken at 300 K for vanadia coverage series. (b) Comparison for particles prepared
under UHV conditions (V, thin line) and in an oxygen ambient (VOx , thick line). From Ref. [59].
surface species. Such species have also been observed on vanadium oxides on metals
(see Section 4.2). Band m2 (bond order 1.6) has a vibrational frequency close to that
of AlVO4 powder samples and is associated with an interface species involving Al, V
and O. Clearly, as it is not observed on vanadium deposition (i.e. without oxygen,
Fig. 14b), it is not simply a AlAOAV termination configuration at the alumina
surface. Band m3 (bond order 1) is indicative of asymmetric stretching vibrations
of a single bridge species (VAOAV) and appears at a frequency where strong
absorption bands of V2 O3 are observed (see Section 4.2).
4.1.2. Vanadia on tin-oxide
SnO2 has a rutile-type structure like TiO2 , albeit with significantly different lattice
constants, making the VOx /SnO2 system interesting for comparison with the much
studied TiO2 system. Moreover, tin dioxide is the active material in gas sensors
whose sensitivity and selectivity is enhanced by the addition of vanadium oxides.
Atrei et al. [61,62] have investigated vanadium and vanadia deposited on the oxygen
deficient (4 · 1) reconstructed surface of SnO2 (1 1 0) that forms from sputter annealing in UHV. Both with reactive evaporation of vanadia or through oxidation via
lattice oxygen from the substrate an epitaxial rutile like VO2 film is formed with a
two-dimensional tin oxide layer on top of it. For low exposures, metallic vanadium
reacts with the substrate to form vanadium oxide and metallic tin. High coverages
yield metallic vanadium and metallic tin in the XPS spectrum, both metallic signals
disappear on annealing (Fig. 15) with the vanadia feature appearing at a binding
energy typical for V2 O3 . Effectively the same oxide forms via reactive evaporation
with XPD results indicating that the vanadium oxide grows epitaxially even at room
temperature and low energy ion scattering (LEIS) proving the presence of a Sn-oxide
layer encapsulating it. The authors considered that a mixed oxide is formed with a
local structure around the V atoms being a distorted rutile VO2 . The experimental
XRD patterns compare well with the multiple scattering simulations for the V 2p and
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S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
Fig. 15. O 1s, V 2p and Sn 3d XPS spectra of the clean SnO2 surface, after deposition of 6 MLE of vanadium and after annealing at 800 K for 10 min. The spectrum of the clean surface is divided by a factor 2.
From Ref. [61].
Fig. 16. Comparison of the V 2p and O 1s XPD patterns calculated for the bulk truncated rutile structure
of VO2 (1 1 0) with the experimental patterns for 6 MLE thick vanadium oxide on SnO2 . From Ref. [61].
O 1s patterns bulk truncated rutile VO2 (1 1 0) as shown in Fig. 16. A similar structure
result has been obtained on TiO2 , however, the presence of a disordered two dimensional Sn-oxide on top of the vanadia has no counterpart on the TiO2 substrate.
The authors considered that it may arise from diffusion of the coordinately unsaturated Sn of the (4 · 1) reconstruction of the SnO2 substrate.
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
139
4.1.3. Vanadia on ceria
The possibility to produce V2 O5 in UHV by oxygen exposure has recently been
demonstrated on CeO2 [63]. For low exposures of vanadium the characteristic V2 O3
V 2p3=2 XPS binding energy of 515.7 eV is observed along with a reduction of the
ceria surface. ÔNormal’ post-oxidation treatments of heating and oxygen exposures
in the 107 Torr range produced no change in the oxidation state of vanadia. In
contrast, post-oxidation at relatively high oxygen pressures (103 Torr for 1 h at 400
K) caused a sharpening and shift of the V 2p feature to 516.9 eV consistent with the
formation of V2 O5 . As illustrated in Fig. 17 the higher oxidation state could only be
achieved for (sub)monolayer vanadia coverages with thicker layers remaining as
V2 O3 even under extreme conditions such as exposure to an oxygen plasma. This
suggests that bulk V2 O5 films are not stable in UHV. Recently similar results have
been achieved on TiO2 [64] indicating that V2 O5 is not necessarily more thermodynamically stable on CeO2 as originally suggested.
4.1.4. Vanadia on titania
This system is relatively well represented by in situ, controlled UHV studies. On
the one hand this reflects the industrial importance of the system but probably also
reflects the fact that non-charging single crystal TiO2 surfaces are relatively easy to
Fig. 17. V 2p and O 1s XPS of vanadia films on CeO2 (1 1 1) as a function of vanadium coverage. Each film
was oxidised in 103 Torr of O2 for 1 h at 400 K. From Ref. [63].
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S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
prepare in situ. Methanol oxidation to formaldehyde has been observed in temperature programmed desorption studies for monolayer V2 O3 on TiO2 (1 1 0), but not
for the clean substrate or for multilayer V2 O3 on TiO2 (1 1 0), suggesting that UHV
models for the real catalysts are good [65]. More recently the reactivity of
(sub)monolayer V2 O5 on TiO2 produced by post oxidation at 103 Torr was investigated. For this system the CH2 O was produced at considerably lower temperature
(485 vs. 615 K for V2 O3 ) [64] suggesting the catalytic importance of V2 O5 .
Wang and Madix [66,67] have produced V2 O5 films on TiO2 (1 1 0) by co-dosing
VOCl3 and water. The oxidation state was monitored by both XPS and UPS and the
reactivity was also followed by the oxidation of methanol. The results are similar to
those of Wong and Vohs for the V2 O5 produced at high O2 pressure [65] but with
some differences suggesting that the production method/exact structure of the film is
important. From the hydrolysis of VOCl3 V2 O5 films above a monolayer could also
be produced. As evident in the XPS uptake series of Fig. 18 the first layer is exclusively V5þ while for higher layers V4þ also becomes evident as a minority species.
With increasing coverage the rate of formaldehyde production from methanol increases with a maximum in reactivity at one monolayer. In the reaction the vanadia
Fig. 18. V 2p XPS for increasing vanadia coverage on TiO2 (1 1 0). From Ref. [67].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
141
is reduced from V5þ to a mixture of V4þ and V3þ , the latter being dominant. The
methanol reaction on this reduced surface also produces formaldehyde albeit at a
higher temperature, similar to that for monolayer V2 O3 [67]. The lower reactivity of
the 1 ML reduced V2 O5 again clearly demonstrates the importance of V5þ to the
reactivity. That the multilayer V2 O5 on the other hand produces no formaldehyde
highlights the importance of VAOATi linkages of the monolayer for the reaction.
The structure of the V2 O5 monolayer has as yet not been investigated, on consideration of the results for powder samples Wang and Madix have suggested a
O@V(AOATi)3 structure. Another interesting aspect of the Wang and Madix results
is the relative ease with which their reduced monolayer could be reoxidised to V2 O5 ,
requiring 2 · 106 as opposed to 103 Torr of O2 of Wong and Vohs. Clearly, investigations of the detailed structure of the V2 O5 monolayers produced by the two
methods and that after reduction would be interesting.
Prior to the relatively recent discovery of the means to produce V2 O5 in UHV
there has been many studies of vanadia on TiO2 (1 1 0) where either V2 O3 , VO2 , and
VO have been reported to form. Films produced by reactive evaporation have been
investigated by a number of different groups. Invariably beyond one monolayer
electron spectroscopic investigations, including XPS, NEXAFS and Auger spectroscopy, indicate films of V2 O3 stoichiometry with no long range order indicated by
LEED [68–70]. The XPS of Guo et al. [68] indicates VO2 in the first monolayer and
V2 O3 for thicker films while Biener et al. [69] indicate V2 O3 in both monolayer and
multilayers and show with both NEXAFS and XPS that the TiO2 substrate is unperturbed on vanadia growth. Interestingly, the HREELS spectra of Guo et al. do
not display the feature of a vanadyl termination that is prominent in the spectra of
V2 O3 films grown on metal substrates (see Section 4.2) or on alumina [69] under
similar conditions. Sambi et al. [70] have explicitly reproduced the growth conditions
of V2 O3 and the XPS and Auger fingerprint of Refs. [68,69] in order to investigate
its structure with XPD. The resulting 2p patterns (see Fig. 19) are remarkably
similar to those of the substrate. This indicates that the vanadia films do not have
the corundum structure of bulk V2 O3 but a rutile like lattice isomorphic to the
substrate. It was suggested that this arises because of the good lattice match of
rutile VO2 with TiO2 and that such a structure can accommodate a large amount
of oxygen defects in so-called M
agneli phases (Vn O2n1 with n P 4). Recently,
however, this structure has been shown not to be unique to films grown on the TiO2 .
On SnO2 , where there is a large lattice mismatch, the same rutile type VO2 XPD
pattern was also obtained, with a V2 O3 stoichiometry indicated by XPS [61] (see
Section 4.1.2).
In contrast, post-oxidation of vanadium deposited on TiO2 has been shown to
yield a variety of vanadium oxide films. In a combined XPS and NEXAFS study
Price et al. [71] investigated vanadium deposited on TiO2 followed by exposure to
oxygen at room temperature. Sub-monolayer vanadium was seen to oxidise to V2 O3
while reducing the TiO2 surface to Ti3þ . Post-oxidation did not change the oxidation
state and thick vanadium covered surfaces also oxidised to V2 O3 .
The group of Granozzi et al. have been extensively studying vanadia produced by
post-oxidation on TiO2 and have been able to produce films of either VO2 or VO.
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Fig. 19. Ti 2p3=2 and O 1s 2p patterns for clean, stoichiometric and ordered TiO2 (1 1 0) compared to the
V 2p and O 1s patterns for vanadia films on TiO2 (1 1 0) obtained for coverages of 5.3, 18 and 18 ML
annealed for 10 min at 473 K. From Ref. [70].
For submonolayer coverages of V the XPD and LEED results are consistent with
epitaxial vanadium oxide and suggests that the V atoms occupy Ti substitutional
sites in the topmost layer without interdiffusion [72]. Ordered vanadia films of up to
5 ML thickness were grown by stepwise vanadium deposition and annealing in
oxygen (0.5 ML V steps, 106 mbar O2 at 423 K). The (1 · 1) LEED structure together with the XPD indicate a crystalline epitaxial VO2 layer with a rutile structure.
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
143
Beyond monolayer coverages XPS displayed a bulk VO2 V 2p3=2 binding energy of
516.5 eV. The electronic structure as determined by UPS was not consistent with
neither V2 O5 nor V2 O3 but rather VO2 in a semiconducting phase [73]. For such VO2
films a reversible semiconductor-to-metal phase transition was observed on heating
above 360 K [74]. The ability to grow such epitaxial VO2 is particularly important
because of the difficulty of growing bulk VO2 . In (sub)monolayer coverages XPS of
the V 2p indicated V2 O3 [73] while in a later synchrotron radiation study two distinct
V 3p emissions at 39.8 and 41.9 eV of approximately equal weight were observed
[75]. The low binding energy emission showed photoelectron diffraction intensity
modulations consistent with a pseudomorphic VO2 monolayer, the latter in contrast
gave no indications of local order. The authors suggested this 41.9 eV binding energy
feature was due to V2 O5 , however the energy separation to the VO2 feature of 2.1 eV
is far greater than would be expected from V 2p results (1.3 eV) and this assignment
must be questioned.
By careful optimisation of growth conditions the Granozzi group could also grow
films (4 ML) of VOx (x 1) on TiO2 [76] and were able to determine its structure
and orientation [77]. Here the films were also formed by stepwise vanadium evaporation but the annealing was in UHV (1–10 min at 400–500 K) with the oxidation
arising from bulk to surface oxygen migration from the substrate. XPS showed a 3þ
oxidation state at the interface while for thicker films the V 2p3=2 binding energy was
513.6 eV, indicative of VO. No long range order was obtained but strong XPD
modulations similar for both the O 1s and V 2p lead to the conclusion of a rocksalt
structure consistent with bulk VOx (x 1). Comparison of the XPD to multiple
scattering theoretical simulations of both VO(1 1 0) and VO(1 0 0) orientations of the
NaCl structure were made. Fig. 20 shows the experimental V 2p 2p plot in comparison with the undistorted and distorted VO(1 0 0)/TiO2 simulations, that gave the
best agreement. The model illustrating the structure and expitaxial relationship of
VO(1 0 0) to TiO2 (1 1 0) that was concluded is shown in Fig. 21. The [0 0 1] azimuth
of the overlayer is aligned with the [
1 1 2] direction of the substrate. To achieve a
match to the substrate the overlayer experiences a strain, which results in a 12–16%
vertical interlayer contraction and a buckling in the [1 1 0] substrate direction of
.
0.5 A
The XPD results can give the impression that well ordered layers are being
formed. However, recent STM investigations of vanadia on titania, produced by
both reactive evaporation and also by stepwise oxidation, revealed that the surfaces
in general were poorly ordered and only local order is achieved at best [78]. In order
to avoid misinterpretation it should be remembered that XPD is only a probe for the
local coordination.
Clearly a plethora of ultrathin V-oxide structures and stoichiometries are possible
on oxide substrates. Also evident is that the fine details of preparation are crucial,
making a scientific discourse on the subject challenging. Apart from the need for
sample reproducibility there is also a need for more adsorption and reactivity studies
particularly given the importance of such systems to catalysis. As yet there have been
no atomically resolved scanning probe investigations published which could be very
fruitful.
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Fig. 20. (a) Experimental V 2p 2p plot. (b) Best fit MSC-SW V 2p pattern simulation from VO(1 0 0)/TiO2
without buckling and (c) with buckling along the [0 1 1] direction of the overlayer. (d) The O 1s pattern
simulation with buckling along [0 1 1]. From Ref. [77].
4.2. Vanadium oxide films on metal substrates
In this section the growth of thin V-oxide layers on single-crystal metal surfaces
will be reviewed. The surfaces of thin oxide layers can be prepared with a structural
order and cleanliness exceeding that of corresponding bulk oxide crystals. Furthermore, by controlling the oxide thickness, oxygen pressure, substrate temperature
and structure, new oxide phases of varying stoichiometry and structure can be
fabricated which can exhibit novel, non-bulk-like properties. In the following, dif-
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
145
Fig. 21. Model illustrating the matching of a VO(1 0 0) overlayer to the TiO2 (1 1 0) substrate. Surface
bridging oxygen atoms of the substrate are shown as white circles. The buckling of the overlayer along its
[0 1 1] azimuth is shown. From Ref. [77].
ferent V-oxide film structures with a formal stoichimetry of VOx (x 1) and V2 O3
will be examined first. These oxide materials have structures, which are more or less
similar to their bulk counterparts. A special class of ultrathin V-oxides, so-called
vanadium oxide nano-layers, will be reviewed next, which are only stable in the
ultrathin layer (up to few monolayers) limit and exhibit rich and complex phase
diagrams with many novel structures. It is appropriate at this point to introduce two
definitions for the oxide coverage, which will be extensively used throughout this
section. In cases when the stoichiometry and the morphology of the oxide film are
well defined, the term monolayer (ML) will be used to indicate the relative portion of
the substrate’s surface covered by the oxide layer. In the more general case, one may
define the V-oxide thickness as corresponding to the evaporated amount of V, as
determined by a quartz monitor and referred to the density of the metal substrate
surface. The term monolayer equivalents (MLE) has been adopted to indicate this
relative oxide coverage.
4.2.1. VOx (x 1) layers
The monoxide of vanadium VOx occurs over a large composition range of
0:8 < x < 1:3 [79]. VOx crystallises in a defective rocksalt structure, which contains a
large number of O and V vacancies. Although the origin of the stability of such large
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a number of vacancies is not understood at present, it is believed that the creation of
vacancies reduces the internal energy more than increasing the entropy. Because of
this non-stoichiometric nature the lattice parameter of VOx can vary in a broad
range as a function of the oxygen content.
VO(1 1 1) films with the rocksalt structure have been grown on Cu(1 0 0) and
Ni(1 1 0) substrates by Kishi et al. [80,81] in the oxide coverage range up to 1 ML.
for both
The LEED patterns are hexagonal and reveal a surface unit cell of 2.88 A
Cu(1 0 0) and Ni(1 1 0) substrates, which value is virtually the same as the lattice
parameter for the bulk VO(1 1 1) surface. V 2p3=2 XPS spectra show a chemical shift
of 0.6 eV to lower binding energy relative to that of bulk V2 O3 , implying a valence
state of V 6 3þ for the oxide films. On the basis of this Kishi et al. have suggested a
schematic model for the VO(1 1 1) surface, which consists of a bilayer of V2þ and O2
ions, similar to that proposed by Galloway et al. for FeO layers on Pt(1 1 1)
[82]. Because the VO bilayer represents a polar surface, it is prone to instability
upon further oxide growth. In fact, Kishi et al. [80,81] have observed the development of a (2 · 2) overlayer on top of the VO(1 1 1) surface at oxide coverage of 2
ML. The authors suggested that the latter structure may have a termination similar
to that of Fe3 O4 (1 1 1) layers on Pt(1 1 1) [83]. In the light of the results discussed
below in Section 4.2.4, it is however more likely that the small negative shift in
the V 2p spectra is due to the formation of an interfacial oxide, rather than being
due to oxide overlayers with a VO stoichiometry, which are predicted to be unstable
[20].
Granozzi et al. have shown that thicker VOx layers (up to 14 ML) can be grown
on Pt(1 1 1) by reactive evaporation of V in a controlled water background (1 · 109
mbar) [84,85]. The VOx stoichiometry has been confirmed from the V 2p3=2 binding
energy (BE) of 513.3 eV, which is between the typical BE values for V metal (512.2–
512.7 eV) and bulk V2 O3 (515.1–515.9 eV) samples. The V/O ratio was estimated to
vary between 0.8 and 1.3, depending on the history of the deposition procedure.
Although the VOx films on Pt (1 1 1) were not particularly well ordered in the long
range, the LEED images suggested that the films are grown epitaxially to the substrate. In contrast, X-ray photoelectron diffraction (XPD) oscillations are quite
strong, revealing a high degree of short-range order. Fig. 22a shows 2p plots for
V 2p3=2 (at kinetic energy (KE) of 973.6 eV) and O 1s (at KE of 955.4 eV) photoemission intensities for a 7 ML thick film. Both diffraction patterns are similar and
exhibit threefold symmetry, as expected for a (1 1 1) rocksalt structure. By comparing
the experimental XPD 2p patterns with those obtained from multiple scattering
calculations from clusters of VO (1 1 1) bi-layers with an O- and V-termination
Granozzi et al. [84] found a good agreement only for the O-terminated surface. A
schematic model for the VO(1 1 1) structure is presented in Fig. 22b, with the arrows
showing possible forward scattering directions. The lattice constant was determined
, with the top-most interlayer distance
to be within the range of 2.77 and 2.92 A
contracted by 7% with respect to the bulk value [84]. Although the reduction of the
interlayer distance could to some extent decrease the high electrostatic energy of the
polar VO(1 1 1) surface, it remains unclear how an unreconstructed (1 1 1) surface
with the rocksalt structure could be maintained for such thick VO films.
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
147
Fig. 22. (a) XPD 2p plots for V 2p3=2 and O 1s photoemission lines of a 7 ML thick VOx film. (b) Structure
model (side view) of the VO(1 1 1) structure. From Ref. [85].
An original approach for growing V-oxide overlayers on metallic substrates has
been recently used by Niehus et al. [86,87]. These authors took advantage of the
ability of the Cu3 Au(1 0 0) surface to accumulate subsurface oxygen after a proper
treatment. The subsurface oxygen has served as an oxidising medium for the growing
oxide film. Depending on the oxygen content, V-oxide films of different stoichiometry and structure have been obtained [86]. For example, a VO(1 0 0) layer with a
rocksalt structure has been proposed to form at a low oxygen content, as suggested
by SPA-LEED (Fig. 23a) and STM (Fig 23b) measurements [86] which reveal a
square lattice structure with a lattice parameter of 0.28 nm. Unfortunately, the authors did not specify the thickness of the vanadium oxide overlayer, but the STM
image in Fig. 23b suggests a heterogeneous surface, consisting of small, ill-defined,
VOx patches, rather than a well developed and atomically ordered VO(1 0 0) layer.
Summarising, ultrathin films with VOx bulk-like rocksalt structures have been
proposed to grow on several metal surfaces, however whether VOx stoichiometry can
indeed be grown by physical vapour deposition must await further confirmation.
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Fig. 23. (a) SPA-LEED of the Cu3 Au(1 0 0)AO surface and a VOx layer. (b) STM image showing different
CuO and VO domains. From Ref. [86].
4.2.2. V2 O3 layers
The growth of ordered V2 O3 films by evaporation of metallic vanadium on
Au(1 1 1) and psubsequent
oxidation has been first reported by the Somorjai
ffiffiffi pffiffiffi
group [88]. A ( 3 3)R30 structure has been observed in LEED in a broad oxide
þ
coverage
pffiffiffi pffiffiffi range and the 3 oxidation state has been identified by AES. The
( 3 3)R30 structure of the vanadium oxide film is consistent with that of the
(0 0 0 1) face of the V2 O3 corundum-type structure. p
The
ffiffiffi in-plane lattice parameter of
) differs by only 2% from the 3 times the lattice parameter of
the V2 O3 film (5.1 A
), which is favourable for the epitaxial oxide growth. Since
Au(1 1 1) (2.88 A
then, ordered V2 O3 (0 0 0 1) layers have been reported to grow on many other single
crystal metal surfaces, such as Cu(1 0 0) [89,90], Pd(1 1 1) [26,91,92], Rh(1 1 1) [93],
Re(0 0 0 1) [94], Cu3 Au(1 0 0) [86,87] and W(1 1 0) [95].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
149
In a series of studies of our group vanadium oxide films have been grown by
reactive evaporation of V metal in an oxygen atmosphere of 2 · 107 mbar on
Pd(1 1 1), with the substrate kept at 250 C [26,91,92]. For oxide coverages exceeding
2 MLE a bulk-like V2 O3 phase has been identified in XPS and NEXAFS [91]. V 2p3=2
XPS spectra displayed a single peak with a BE of 515.4 eV, which is well within the
range of BE values reported for the bulk V2 O3 [31]. The V2 O3 stoichiometry of the
V-oxide films on Pd(1 1 1) in this coverage range has been further confirmed by
the vanadium L- and the oxygen K-edge NEXAFS spectra, which are presented in
Fig. 24. The V L-edge spectra for 5 MLE (Fig. 24a) show two broad structures at
517 and 523.5 eV, which are related to excitations of electrons from the V 2p3=2
(LIII -edge) and V 2p1=2 (LII -edge) levels, respectively, into unoccupied V 3d states.
Fig. 24. NEXAFS spectra of vanadium oxide on Pd(1 1 1) for various oxide thicknesses: (a) vanadium
L-edge features, (b) oxygen K-edge features. From Ref. [91].
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2 , )0.5 V; 0.1 nA) and (b) high-resolution (200 · 200 A
2 , )0.3 V; 0.1
Fig. 25. (a) Large-scale (1000 · 1000 A
nA) STM images of 3 MLE V oxide on Pd(1 1 1). The areas denoted A (A0 ) and B on the image in (b)
correspond to different terminations of the V2 O3 (0 0 0 1) surface. From Ref. [92].
The O K-edge spectra (Fig. 24b) correspond to transitions from the O 1s level to
unoccupied O 2p levels and show a doublet structure near the edge at 530 eV,
which can be attributed to the crystal-filed-splitting of the V 3d bands. Both V L- and
O K-edge spectra at 5 ML are very similar to those from single crystals of V2 O3 ,
reported by Abbate et al. [60] and Chen [96], which confirms the V2 O3 identification
of the thick V-oxide films on Pd(1 1 1).
Large-scale STM images of thick (>3 MLE) V2 O3 films on Pd(1 1 1) (Fig. 25a)
have shown that the bulk-type V2 O3 phase grows in the form of random-shaped
[92].
three-dimensional crystallites with lateral dimensions between 50 and 200 A
Atomically resolved STM images (Fig. 25b) have revealed a hexagonal surface
along
structure on the top facets of the islands with a lattice parameter of 4.9 ± 0.1 A
the Æ1 2 1æ directions, which is in remarkable agreement with the in-plane lattice
parameter of the V2 O3 (0 0 0 1) face. The unit cell
by 30 with respect to the
pffiffiffiis rotated
pffiffiffi
Pd(1 1 1) unit cell vectors and LEED shows
a
(
3
3
)R30
pattern. The mismatch
pffiffiffi
between the V2 O3 (0 0 0 1) lattice and the 3 direction of the Pd(1 1 1) surface is 3.7%,
i.e. slightly higher than that on Au(1 1 1) [88].
An important issue and interesting question is how the V2 O3 (0 0 0 1) surface is
terminated. Atomically resolved STM images of V2 O3 layers on Pd(1 1 1) (Fig. 26a)
[92] have demonstrated that two different terminations (indicated A and B) coexist
on the V2 O3 (0 0 0 1) surface. They both exhibit the hexagonal structure and the periodicity of the V2 O3 (0 0 0 1) surface, but are well distinguished by the detailed shapes
of the atomic features in STM (Fig. 26b). The atomic features of the A-type terminations have a circular form, while those of the B-type surface exhibit a triangular
shape in the STM images. For the oxide preparation conditions employed for the
V-oxides on Pd(1 1 1) two stable terminations of V2 O3 (0 0 0 1) have been found in the
DFT calculations [92]. One corresponds to a V2 O3 (0 0 0 1) surface terminated by
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
151
2 , 1.0 V; 1.0 nA) of the V2 O3 (0 0 0 1) surface revealing two different
Fig. 26. (a) STM image (200 · 200 A
2 ) of the A and B terminations. (c) Structural
terminations A and B. (b) Magnified portions (27 · 18 A
models (left-hand side) and STM simulations (right-hand side) of the vanadyl-oxygen (A) and bulk-type
oxygen (B) terminated V2 O3 (0 0 0 1) surface. (d) HREELS phonon spectrum of the V2 O3 (0 0 0 1) surface.
From Ref. [92].
V@O dimers, the so-called vanadyl group, i.e. OAVAO3 AV2 AO3 AV in terms of
layers, as schematically shown in the top panel of Fig. 26c. The other V2 O3 (0 0 0 1)
termination consists of three oxygen atoms per unit cell, as in the bulk, i.e.
O3 AV2 AO3 AV2 (bottom panel of Fig. 26c). As seen from the comparison of
experimental and simulated STM images the A-type surface is terminated by the
V@O dimers, whereas the B-type surface is terminated by bulk oxygen planes. The
existence of terminal oxygen species on the V2 O3 (0 0 0 1) surface has been also
supported by the HREELS spectra shown in Fig. 26d, where the loss peak at 130
meV has been assigned to the excitation of V@O stretching vibrations. Similar peaks
have been found in vibrational spectra of V2 O5 (0 0 1), where terminal oxygen atoms
[36]. The bond
are connected to a V atom via a short double bond of length 1.58 A
[92], i.e. very
length of the V@O dimer in model A of Fig. 26c is calculated to 1.61 A
close to that of the vanadyl group on the V2 O5 (0 0 1) surface.
Recently, V@O terminated V2 O3 (0 0 0 1) films have been also reported by the
Freund’ group to grow on Au(1 1 1) and W(1 1 0) surfaces, as revealed by the
HREELS and IRAS spectra [95]. The preparation process included the evaporation of metallic vanadium in an oxygen atmosphere, followed by annealing at 700 K
in 5 · 108 mbar of oxygen. The UPS and NEXAFS spectra have shown the
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appearance of a small gap (0.3 eV), which is induced by the formation of vanadyl
groups at the surface. In addition, a V-termination was found to form upon short
annealing to 600 K with electron bombardment. The V-termination has been identified in the UPS, core-level XPS and NEXAFS spectra as well as from the different
adsorption behaviour of probe molecules, such as CO, CO2 and H2 O, when compared to the V@O terminated surface [95]. It has been suggested that the V-termination is formed by the electron beam induced removal of vanadyl oxygen atoms
from the V@O termination, with the resulting surface similar to that of
Cr2 O3 (0 0 0 1), which is terminated by a half Cr layer [97].
Other different terminations of the V2 O3 (0 0 0 1) surface have been recently proposed by Niehus et al., based on STM and low-energy ion scattering spectroscopy
(ISS) results [87]. In this study a V2 O3 film has been fabricated by evaporation of 5
ML of vanadium at 300 K on a Cu3 Au(1 0 0) surface, followed by subsequent oxidation at 650 K in 5 · 107 mbar oxygen atmosphere. As already mentioned in
Section 4.2.1, the oxygen pre-treatment of the Cu3 Au(1 0 0) substrate appears to be
decisive for the growth of atomically flat V-oxide layers, as convincingly demonstrated in the large-scale STM image of Fig. 27a [87]. The corresponding LEED
pattern is due to the superposition of the diffraction patterns of two hexagonal
domains, rotated by 90. The atomically resolved STM image (Fig. 27b) revealed a
, which together with the measured
hexagonal structure with a periodicity of 5 A
terrace height of 3 A is consistent with the geometry of the V2 O3 (0 0 0 1) surface.
Concerning the question of how this surface is terminated, Niehus et al. have argued
that two V2 O3 (0 0 0 1) terminations, which are different from those considered above,
are possible. Their ISS data were interpreted with a model consisting of a full vanadium layer termination, although such a surface has a polar character and would
be quite unstable.
Bias-dependent STM imaging has shown that by varying the bias voltage between
)0.3 and )1.3 V much sharper features appear in the STM images (Fig. 28a), which
consist of bright spots arranged in distorted hexagons. The average distance between
Fig. 27. V2 O3 film on Cu3 Au(1 0 0)AO: (a) survey STM image. Inset: LEED pattern of the thick V2 O3
film. (b) Atomically resolved STM image, z-corrugation 0.15 nm, tunnelling parameters: )0.2 V, 1.0 nA.
From Ref. [87].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
153
Fig. 28. High-resolution STM image of V2 O3 (0 0 0 1) on Cu3 Au(1 0 0)AO. (a) On terraces; the links between vanadium atoms as possible positions of oxygen atoms are marked in the lower left. (b) Close to
steps or larger defects an undisturbed hexagon can be also obtained. (c) Schematic structure model
consisting of the full vanadium layer termination, stabilised by 1/3 ML of additional oxygen in pseudobridge positions. From Ref. [87].
, which is close to the distance between the V atoms (2.89
the protrusions was 3 A
A) in the full layer of V2 O3 (0 0 0 1). To explain the distortions of the hexagons
Niehus et al. have assumed that it is driven by the tendency of the polar full metal
layer terminated (0 0 0 1) corundum surface to reduce its energy either by reconstruction or by an adsorbate of 1/3 ML coverage. Most naturally oxygen or hydroxyl
groups could play the role of such adsorbates. The ISS data suggested the pseudobridge sites as the most likely locations of the oxygen atoms (or OH species), as
schematically shown in the STM image (Fig. 28a) and in the atomic model (Fig. 28c).
Interestingly, the on-top position of the O atoms above the V atoms, which would
correspond to the vanadyl termination discussed above, has been rejected by the
analysis of the ISS data. Occasionally, near defects, a regular hexagonal structure has
been observed in the STM images (Fig. 28b). The latter has been interpreted by
Niehus et al. as an unperturbed hexagonal full layer of vanadium atoms, which is
stabilised by the presence of defects [87].
In summary, irrespectively of the type and symmetry of the metal substrate, or of
the preparation method, all V2 O3 films grow with the (0 0 0 1) face of the corundum
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structure parallel to the substrate plane. This is an interesting result, having in mind
that the cleavage plane of the bulk-type V2 O3 crystals is not the basal (0 0 0 1), but
the (1 0 1 2) surface. Concerning the termination of the V2 O3 (0 0 0 1) surface, it appears that different surface terminations could exist, depending on the oxide film
fabrication procedure.
4.2.3. VO2 and V2 O5 layers
To the best of our knowledge, there are no studies reporting on the growth of
ordered and stoichiometric VO2 and V2 O5 films with a bulk-type structure on metal
substrates. This could be partly due to the reducing role of the metal surface, or to
kinetic limitations. However, such films have been successfully grown on oxide
supports, as discussed in Section 4.1. In the next paragraph it will be shown, however, that V-oxide phases with a formal oxidation state of equal or even higher than
4+ may form in the ultrathin oxide film limit on metal substrates.
4.2.4. Vanadium oxide nano-layers
The V-oxide thin-film phases discussed above have more or less a bulk-type
character, i.e. their physical and chemical properties are similar to those of bulk oxide
samples. However, when ultrathin oxide films with a thickness of the order of one to
several layers are grown on metals, their stoichiometry may differ from those of
known bulk phases, because the proximity to the metal substrate perturbs the electronic structure of the oxide. Interfacial oxide layers specific to and stabilised by the
interface may thus form and these are the subject of this section. Since these layers are
prototypical for the properties of matter in ultrasmall dimensions, we will use the
term nano-layers for these phases. Here, as an example, the peculiar growth mechanism and atomic structure of vanadium oxide nano-layers will be discussed, which
form in the early stage (up to 2–3 ML) of vanadium oxide film growth on Pd(1 1 1).
The layer-dependent evolution of vanadium oxides structures on Pd(1 1 1), grown
by reactive evaporation of V metal in an oxygen atmosphere of 2 · 107 mbar, has
been extensively investigated in Graz, using a variety of techniques including LEED,
STM, high-resolution XPS, NEXAFS, HREELS, XPD and I–V LEED combined
with first-principles DFT calculations [26,91,92,98–104]. At submonolayer oxide
coverages (<0.5 MLE), two well-ordered V-oxide phases form after the reactive
evaporation. They exhibit (4 · 4) (Fig. 29a) and ‘‘zigzag’’ structures (Fig. 29c), as
observed in the STM images [92,101,102]. The DFT calculations have revealed that
the (4 · 4) and the ‘‘zigzag’’ structures can be rationalised in terms of models with
formal V5 O14 and V6 O14 stoichiometries, as presented in Fig. 29b and d, respectively
[102]. The (4 · 4)-V5 O14 and the ‘‘zigzag’’-V6 O14 phases are thermodynamically stable only at submonolayer V coverages and at higher oxygen pressures (e.g. high O
chemical potentials 1) as seen in the phase stability diagram (Fig. 30a), obtained
1
The chemical potential lO is related to the oxygen pressure p and the temperature T through
lO ðT ; pÞ ¼ lO ðT ; pO Þ þ 12kT lnðp=pO Þ. lO ¼ 0 corresponds to a pressure at which oxygen condenses at the
surface; lO ¼ 1 eV corresponds to p 107 mbar at 520 K.
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
155
Fig. 29. Atomically resolved STM images (left panels) and DFT models (right panels) of the (4 · 4)-V5 O14
(a and b), zigzag-V6 O14 (c and d), and (2 · 2)-surface-V2 O3 (e and f) oxide phases. The insets in (a–c)
display the STM simulations for the respective models. From Ref. [102].
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S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
Fig. 30. (a) Surface phase diagram as a function of the oxygen pressure lO and the vanadium coverage xV .
Stable phases are indicated with black bars. From Ref. [102]. (b) Change of surface energy on deposition
of vanadium and oxygen on Pd(1 1 1) versus the vanadium coverage. From Ref. [26].
from the DFT calculations [101,102]. Note that the oxidation state of the (4 · 4))V5 O14 phase is only apparently higher than 5þ , since the oxide layer is oxygen-terminated at both oxide-metal and oxide–vacuum interfaces. Under more reducing
conditions (annealing in UHV or in a H2 background), corresponding to a lowering
of the O chemical potential, the (4 · 4) and ‘‘zigzag’’ phases transform into another
interface-stabilised phase, which exhibits a (2 · 2) honeycomb structure in the STM
(Fig. 29e). The DFT analysis [20,98] has shown that the (2 · 2) phase has a formal
V2 O3 stoichiometry, and the structure has thus been designated as surface-V2 O3 (Fig.
29f). The latter consists of two V atoms per unit cell located in fcc and hcp threefold
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
157
hollow Pd(1 1 1) sites and three O atoms above VAV bridge sites. The simulated
STM image (inset of Fig. 29e) is consistent with the experimental one and reveals
that the V atoms are imaged as bright protrusions. The surface-V2 O3 structure has
been also recently confirmed by XPD [103] and I–V LEED [104] measurements. The
average VAO interlayer spacing, as estimated from the XPD data, is 0.72 ± 0.07 A
[103], which is very close to the DFT derived value of 0.723 A [98]. An important
feature of the s-V2 O3 structure is the V-termination at the Pd interface, which is due
to the strong V-Pd bond [20]. This causes an XPS V 2p core level shift of 1 eV to
lower BE with respect to the bulk-type V2 O3 , which was also predicted by DFT [98],
but which has lead previously to an erroneous assignment of this phase to a VO [91].
This means that the XPS binding energy is not necessarily a good measure for the
oxide stoichiometry in the thin film limit and has to be interpreted with care. The
strong V-Pd bonding has also been suggested to be the reason for the observed
wetting behaviour of the s-V2 O3 layer during oxidation–reduction cycles [101], as
discussed below.
Growth of additional V2 O3 layers on top of the s-V2 O3 is energetically quite
unfavourable, as seen in the DFT phase stability diagram (Fig. 30b) [20,26,102]. For
V-oxide growth on Pd(1 1 1) in the range of 0.6 ML a sudden transformation from
s-V2 O3 to VO2 -type phases takes place [26]. Two different VO2 -type thin film phases
have been experimentally observed in STM, with rectangular (Fig. 31a) and hexagonal (Fig. 31c) lattice structures, which initially coexist with the s-V2 O3 layer
[26,92]. The VO2 -stoichiometry has been confirmed by XPS [91], where V 2p core
level spectra show a component with a binding energy of 516.3 eV, typical for VO2
compounds [31]. The rectangular VO2 phase (VO2 -rect) grows in the form of 2D
(Fig.
islands of rectangular shape and lattice spacings of 3.0 ± 0.2 and 3.8 ± 0.2 A
31a). The rows parallel to the shorter lattice spacing are misaligned by 7.5 with
respect to the Æ1 1 0æ substrate directions, which results in the appearance of two
rotationally different rect-VO2 domains, displaying an angle of 15 with respect to
each other in the STM (Fig. 31a) and LEED (Fig. 31b) images. The hexagonal phase
(VO2 -hex) displays a (1 · 1) lattice, which is commensurate with the Pd(1 1 1) substrate (Fig. 31c). The DFT calculations [20,26] have confirmed that two unsupported
thin-film V oxide phases with formal VO2 stoichiometry are indeed conceivable. One
of them has a rectangular structure with relaxed surface lattice parameters of 2.96
and consists of alternating three oxygen and two vanadium layers (Fig.
and 3.71 A
31d). The other VO2 phase has a hexagonal structure with relaxed surface lattice
, i.e. very close to the Pd(1 1 1) substrate, and is composed of a
parameters of 2.87 A
(1 · 1) vanadium layer sandwiched between two oxygen layers (Fig. 31e). It is important to stress, that these ultrathin VO2 layers are again particular to the interface
and bear no resemblance to the bulk-type rutile VO2 phase. As shown in the DFT
analysis [20,26], the VO2 layers on Pd(1 1 1) derive their stability from an improved
lattice matching to the substrate, which lowers the interfacial energy. By means of
the ±7.5 rotation of the overlayer with respect to the substrate, the oxygen atoms
at the PdAVO2 -rect interface avoid the Pd-hollow sites, which are less favourable
than the on-top and bridge sites (Fig. 31f). A similar situation could also be invoked
to rationalise the growth of thin VO2 -like layers on Cu3 Au(1 0 0), which exhibit a
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S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
· 275 A
; 0.5 V, 1.0 nA) of 1 MLE V oxide on Pd(1 1 1), showing 3 domains
Fig. 31. (a) STM image (275 A
of the VO2 -rect phase. The upper-right insert is a high-resolution STM image, the centre insert a theoretical simulation. The theoretical calculations reveal that the V atoms, although below the O atoms, are
· 133 A
; 0.6 V, 0.5 nA) of the VO2 -hex phase. (c)
imaged as protrusion in the STM. (b) STM image (133 A
LEED pattern of the 1 MLE V-oxide/Pd(1 1 1) surface. (d) Structural model of the VO2 -rect phase. (e)
Structural model of the VO2 -hex phase. (f) Model of the orientation of the VO2 -rect unit cell on the
Pd(1 1 1) surface, revealing row matching. From Ref. [26].
rectangular structure, but with unit cell dimensions different from that of the rutile
VO2 (1 1 0) surface [86].
On further V-oxide deposition the VO2 -rect phase collapses, since the growth of
the next oxide layer on top of the VO2 -rect is again energetically unfavourable, as
seen in the DFT phase diagram (Fig. 30b). This collapse can be also understood
from simple electrostatic arguments, since the VO2 -rect layer has a non-zero dipole
moment and the addition of further layers would lead to an increase of the electrostatic energy. The further oxide growth commences on top of the VO2 -hex phase
with the formation of a (2 · 2) overlayer, which is a precursor phase to the bulk-type
V2 O3 [26,92]. As demonstrated in the DFT phase diagram (Fig. 30), a (2 · 2) overlayer with a similar structure as s-V2 O3 stabilises the VO2 -hex phase [20,26]. Eventually, for oxide coverages above 3–4 ML the oxide layer converges to the bulk-type
V2 O3 phase, which is the stable oxide phase under the given experimental conditions.
The detailed knowledge of the phase diagram of the V-oxide-on-Pd(1 1 1) system
in the ultrathin film limit has allowed us to follow the mutual transformation of the
various interface-stabilised oxide phases upon cycling reducing and oxidising conditions, which is of relevance for heterogeneous catalysis. In this respect, oxide nanoparticles on metal surfaces are useful models to study the morphology changes and
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
159
the mass transport processes of catalytic systems under reaction conditions. Oxidation and reduction treatments are important industrial practice for catalyst activation and regeneration and the detailed information on the mass transport leading
to morphology changes is a relevant issue. Using a variable-temperature STM, the
morphology and structure oscillations of the oxide phase on the metal surface have
been observed in situ as a result of alternating exposures to H2 and O2 at elevated
temperature (top panel of Fig. 32) [101]. The STM image of Fig. 32A presents a
Pd(1 1 1) surface, decorated with 0.7 ML V-oxide and contains clean Pd(1 1 1)
areas, islands of the VO2 -rect phase and a region of the s-V2 O3 layer. Upon reduction (exposure to 1 · 109 mbar H2 at 250 C) the VO2 -rect islands shrink in size
and transform into the s-V2 O3 phase, which eventually spreads out and wets the
whole Pd surface (Fig. 32A–C). Upon oxidation (exposure to 5 · 108 mbar O2 at
250 C) the s-V2 O3 phase transforms first into the ‘‘zigzag’’-V6 O14 phase while
dewetting the Pd surface (Fig. 32D) and then into the (4 · 4)-V5 O14 phase (Fig. 32E).
A repetitive reduction in H2 leads again to a surface completely covered by the
s-V2 O3 layer (Fig. 32F). The whole sequence of STM images collected during the
reduction–oxidation–reduction cycle has been evaluated quantitatively, yielding
the percentage areas of the bare Pd patches and the various V-oxide phases, which is
displayed on the bottom panel of Fig. 32. Note that after the first reduction step
there is an induction period (2200–4000 s), which is presumably due to the necessity
of generating free Pd sites for the dissociation of O2 .
The physical origin of the reduction/oxidation behaviour of V-oxides on Pd(1 1 1)
can be explained with the phase stability diagram as derived from the DFT calculations [101,102] (Fig. 30a). At a given V coverage of say 0.5 ML and an oxygen
chemical potential of lO ¼ 1:0 eV, the ‘‘zigzag’’-V6 O14 and (4 · 4)-V5 O14 phases
coexist at the surface. On lowering the oxygen potential, e.g. to lO ¼ 2:0 eV, the
‘‘zigzag’’ and the (4 · 4) phases become unstable and decompose into VO2 -rect,
s-V2 O3 and some bare Pd patches. Under more reducing conditions, e.g. at
lO 6 2:25 eV, the s-V2 O3 remains the only stable oxide phase. The observed reversible wetting–dewetting behaviour of the V-oxide layers on Pd(1 1 1) is related to
the well-known SMSI (strong metal support interaction) effect [105], where it is
thought that the metal particles of a supported metal catalyst become encapsulated
by a thin oxide layer under hydrogen treatment [106].
To conclude this section the growth behaviour of V-oxides on a Rh(1 1 1) substrate is worth mentioning. Rh(1 1 1) is structurally very similar to Pd(1 1 1), but
exhibits a higher affinity towards oxygen. As a result, a completely different oxide
phase diagram has been established by LEED and STM in the early stages of
V-oxide growth [93]. Without going into the details of the complex growth behaviour
of vanadium oxides on Rh(1 1 1), we would like to mention briefly one particular
V-oxide phase,
pffiffiffi which grows in the form of well-ordered 2D islands with a rectangular (5 3 3) structure, as shown on the STM image of Fig. 33. In between the
islands individual flat hexagonal star-like clusters are visible on the Rh(1 1 1) terrace
with lateral dimensions of approximately
pffiffiffi12 A. These clusters are mobile and are the
building units of the rectangular (5 3 3) oxide islands. The study of their diffusion
behaviour and of their condensation into the rectangular oxide structure allows to
160
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
Fig. 32. Top: Sequence of STM images of a vanadium oxide decorated Pd(1 1 1) surface exposed to reducing and oxidising ambient conditions. (A) The Pd(1 1 1) surface, annealed at 300 C in vacuum after
oxide deposition, 65% covered with two coexisting V-oxide phases: s-V2 O3 and r-VO2 . (B) After 1000 s
exposure to 1 · 109 mbar H2 at 250 C. (C) After 2000 s exposure to 1 · 109 mbar H2 at 250 C. The
surface is now reduced and completely covered by s-V2 O3 . (D) After 2000 s exposure of surface (C) to
5 · 108 mbar O2 at 250 C. The zigzag V6 O14 oxide appears. (E) After 400 s exposure of surface (D) to
5 · 108 mbar O2 at 250 C, followed by pump-down of the O2 . The surface contains zigzag V6 O14 and
(4 · 4) V5 O14 oxide phases along with some free Pd areas. (F) Reduction of surface (E) by 1250 s exposure to 3 · 109 mbar H2 at 250 C. The surface is again wetted by the reduced s-V2 O3 phase. Bottom:
Percentage areas of bare Pd and various V-oxide covered parts of the surface plotted versus the time
during a reduction–oxidation–reduction cycle, as evaluated from STM images. From Ref. [101].
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
161
pffiffiffi
2 , +1.5 V; 0.1 nA) of a rectangular (5 3 3) vanadium oxide islands
Fig. 33. STM image (200 · 200 A
on Rh(1 1 1). Between the islands star-like clusters are seen. From Ref. [93].
follow the growth kinetics of oxide nano-structures at the atomic level, and such
experiments are presently under way in our laboratories in Graz.
5. Synopsis and outlook
The surface science of vanadium oxides has been reviewed in this article to assess
the present status of knowledge on the atomic structure, the electronic properties,
and the adsorption behaviour of well-defined surfaces of these oxides. Two types of
surfaces have been considered: surfaces of bulk single crystals and surfaces of thin
films grown on other oxide or metal substrates. While the former can be prepared by
cleavage of crystals in UHV or by ex situ cutting and polishing followed by various
in situ cleaning procedures, the latter are created by epitaxial in situ growth using
methods of physical vapour deposition. The electronic structure of vanadium oxide
surfaces has been investigated by electron spectroscopic techniques mostly on bulk
single crystal surfaces, whereas the atomic-level characterisation of surface structures
has been performed predominantly on thin film surfaces. This latter fact is mainly
due to some technical advantages in the preparation of thin film surfaces and in the
application of scanning probe microscopies to these samples. In the area of thin film
162
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
studies the mechanism of epitaxial growth and the morphology of the resulting thin
films of vanadium oxides are active areas of research.
The interest in adsorption studies on well-defined vanadium oxide surfaces is
driven by the relevance of these oxides in the field of heterogeneous catalysis. Although vanadium oxides are amongst the most important materials in oxide catalysis, the molecular level characterisation of elementary reaction steps of many
industrially relevant reactions involving vanadium oxides are not or only poorly
understood. In this context the value of model studies on atomically defined surfaces
is well recognised and an increase in the scientific activity of adsorption and reaction
studies on ordered vanadium oxide surfaces is expected. For example, the role of
defects on the surface reactivity of vanadium oxides with the related presence of V
surface species in different oxidation states is an interesting and important question.
Moreover, the characterisation of the so-called ‘‘monolayer catalyst’’, which appears
to exhibit superior catalytic properties as compared to a conventional mixture of
oxides [107], is a fascinating subject. In our present day knowledge these monolayers
catalysts are a few layers of e.g. vanadium oxides on another oxide substrate, but the
morphology, atomic structure, and chemical composition of these oxide-on-oxide
overlayers are largely unknown. The detailed elucidation of their atomic-level
properties is a necessity, if the full catalytic potentials of these systems are to be
exploit in a rational knowledge-based way. Studies on thin film systems under UHV
conditions are a way to approach this problem.
The progress in experimental techniques of surface science during the last decades
[108] has been paralleled by the progress in the development of theoretical methodologies. DFT-based ab initio methods have now the ability to tackle complex
oxide surfaces and examples of theory assisted structure determinations of novel
ultrathin vanadium oxide phases have been given in this review. Also, the simulation
of reaction steps on particular surface centres on vanadium oxides has been successfully attempted and has helped the interpretation of experimental results. There
can be no doubt that theoretical simulations will play an important role in advancing
the field of model catalytic studies, in particular on complex oxide surfaces such as
those of vanadium oxides.
In the field of nano-science vanadium oxide nano-layers have interesting potential
as model oxide systems with a high flexibility in structural and electronic behaviour.
As mentioned in Section 4.2.4 novel vanadium oxide phases have been detected in
the ultrathin limit. The investigation of oxide properties in reduced dimensional
systems are only just beginning, and potential applications of ultrathin oxide layers
in new areas of high technology can be envisioned. For example, the preparation of
nano-tubes based on mixed-valent vanadium oxides has recently been described
[109], and a field-effect transitor made of individual V2 O5 nano-fibres has been reported [110]. In a different area, namely solid-state electrochemistry, an interesting
application of vanadium oxides has been reported recently, where V2 O5 thin films
have been tested as electrodes with intercalating properties for rechargeable Li
micro-batteries [111,112]. Thus, the field of vanadium oxide surface research appears
to be open into many different directions, and we expect an active scientific development along many different lines.
S. Surnev et al. / Progress in Surface Science 73 (2003) 117–165
163
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