Fu-2007

Surface Science Reports 62 (2007) 431–498www.elsevier.com/locate/surfrep

Interaction of nanostructured metal overlayers with oxide surfaces

Qiang Fua,b,1, Thomas Wagnera,∗

a Max-Planck-Institut fur Metallforschung, Heisenbergstrasse 3, D-70569 Stuttgart, Germanyb State Key Laboratory of Catalysis, Dalian Institute of Chemical Physics, The Chinese Academy of Sciences, Zhongshan Road 457, Dalian 116023, PR China

Accepted 29 July 2007

Abstract

Interactions between metals and oxides are key factors to determine the performance of metal/oxide heterojunctions, particularly innanotechnology, where the miniaturization of devices down to the nanoregime leads to an enormous increase in the density of interfaces. Onecentral issue of concern in engineering metal/oxide interfaces is to understand and control the interactions which consist of two fundamentalaspects: (i) interfacial charge redistribution — electronic interaction, and (ii) interfacial atom transport — chemical interaction. The present paperfocuses on recent advances in both electronic and atomic level understanding of the metal–oxide interactions at temperatures below 1000 ◦C,with special emphasis on model systems like ultrathin metal overlayers or metal nanoclusters supported on well-defined oxide surfaces. Theimportant factors determining the metal–oxide interactions are provided. Guidelines are given in order to predict the interactions in such systems,and methods to desirably tune them are suggested.

The review starts with a brief summary of the physics and chemistry of heterophase interface contacts. Basic concepts for quantifying theelectronic interaction at metal/oxide interfaces are compared to well-developed contact theories and calculation methods. The chemical interactionbetween metals and oxides, i.e., the interface chemical reaction, is described in terms of its thermodynamics and kinetics. We review the differentchemical driving forces and the influence of kinetics on interface reactions, proposing a strong interplay between the chemical interaction andelectronic interaction, which is decisive for the final interfacial reactivity. In addition, a brief review of solid–gas interface reactions (oxidation ofmetal surfaces and etching of semiconductor surfaces) is given, in addition to a comparison of a similar mechanism dominating in solid–solid andsolid–gas interface reactions.

The main body of the paper reviews experimental and theoretical results from the literature concerning the interactions between metals andoxides (TiO2, SrTiO3, Al2O3, MgO, SiO2, etc.). Chemical reactions, e.g., redox reactions, encapsulation reactions, and alloy formation reactions,are highlighted for metals in contact with mixed conducting oxides of TiO2 and SrTiO3. The dependence of the chemical interactions on theelectronic structure of the contacting metal and oxide phases is demonstrated. This dependence originates from the interplay between interfacialspace charge transfer and diffusion of ionic defects across interfaces. Interactions between metals and insulating oxides, such as Al2O3, MgO, andSiO2, are strongly confined to the interfaces. Literature results are cited which discuss how the metal/oxide interactions vary with oxide surfaceproperties (surface defects, surface termination, surface hydroxylation, etc.). However, on the surfaces of thin oxide films grown on conductingsupports, the effect of the conducting substrates on metal–oxide interactions should be carefully considered.

In the summary, we conclude how variations in the electronic structure of the metal/oxide junctions enable one to tune the interfacialreactivity and, furthermore, control the macroscopic properties of the interfaces. This includes strong metal–support interactions (SMSI), catalyticperformance, electrical, and mechanical properties.c© 2007 Elsevier B.V. All rights reserved.

Keywords: Oxide surfaces; Metal films; Oxide films; Model systems; Metal–support interaction; Interface reaction; Charge transfer; Titanium oxide; Strontiumtitanate; Aluminum oxide; Magnesium oxide; Silica; Catalysis; Growth; Epitaxy

∗ Corresponding author. Tel.: +49 711 6891470; fax: +49 711 6891472.E-mail addresses: [email protected] (Q. Fu), [email protected] (T. Wagner).

1 Current address: Dalian Institute of Chemical Physics, The Chinese Academy of Sciences, Zhongshan Road 457, Dalian 116023, PR China. Tel.: +86 41184379976; fax: +86 411 84694447.

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432 Q. Fu, T. Wagner / Surface Science Reports 62 (2007) 431–498

Contents

1. Introduction............................................................................................................................................................................ 4322. Fundamental aspects of metal–oxide interaction ......................................................................................................................... 434

2.1. Physics and chemistry of heterophase interface contacts .................................................................................................... 4342.1.1. Physics at metal–semiconductor interfaces .......................................................................................................... 4342.1.2. Chemistry at metal–semiconductor interfaces....................................................................................................... 435

2.2. Electronic interaction of metals with oxides — charge redistribution ................................................................................... 4362.2.1. Local charge redistribution ................................................................................................................................ 4362.2.2. Space charge transfer ........................................................................................................................................ 437

2.3. Chemical interaction of metals with oxides — mass transport............................................................................................. 4382.3.1. Chemical interactions at metal/oxide interfaces .................................................................................................... 4382.3.2. Thermodynamic considerations .......................................................................................................................... 4402.3.3. Kinetic consideration ........................................................................................................................................ 441

2.4. Interplay between electronic and chemical interactions ...................................................................................................... 4422.4.1. Cabrera–Mott theory in solid–gas interface reactions ............................................................................................ 4422.4.2. Cabrera–Mott theory in solid–solid interface reactions .......................................................................................... 4432.4.3. Generalized Cabrera–Mott theory in interface reactions ........................................................................................ 445

3. Experimental methods ............................................................................................................................................................. 4453.1. Preparation of model systems ......................................................................................................................................... 446

3.1.1. Oxide surfaces ................................................................................................................................................. 4463.1.2. Metal overlayers on oxide surfaces ..................................................................................................................... 4473.1.3. Inverse model systems....................................................................................................................................... 449

3.2. Characterization techniques............................................................................................................................................ 4493.2.1. Electron-based spectroscopy .............................................................................................................................. 4493.2.2. Scanning probe techniques ................................................................................................................................ 4503.2.3. Transmission electron microscopy ...................................................................................................................... 4513.2.4. Other techniques .............................................................................................................................................. 452

4. Interaction of metals with mixed conducting oxides .................................................................................................................... 4534.1. Metals on TiO2 ............................................................................................................................................................. 453

4.1.1. TiO2 surfaces ................................................................................................................................................... 4534.1.2. TiO2 bulk defect chemistry ................................................................................................................................ 4554.1.3. Metal–TiO2 interactions .................................................................................................................................... 457

4.2. Metals on SrTiO3 .......................................................................................................................................................... 4634.2.1. SrTiO3 surfaces................................................................................................................................................ 4634.2.2. SrTiO3 bulk defect chemistry............................................................................................................................. 4654.2.3. Metal–SrTiO3 interactions................................................................................................................................. 467

5. Interaction of metals with insulating oxides ................................................................................................................................ 4715.1. Metals on Al2O3 ........................................................................................................................................................... 472

5.1.1. α-Al2O3 surfaces ............................................................................................................................................. 4725.1.2. Metal interactions with bulk Al2O3 .................................................................................................................... 4755.1.3. Metal interactions with alumina films.................................................................................................................. 476

5.2. Metals on MgO............................................................................................................................................................. 4775.2.1. Metal interactions with bulk MgO ...................................................................................................................... 4775.2.2. Metal interactions with MgO films ..................................................................................................................... 480

5.3. Metals on SiO2 ............................................................................................................................................................. 4805.3.1. Metal interactions with bulk SiO2 ....................................................................................................................... 4805.3.2. Metal interactions with silica films ..................................................................................................................... 482

6. Summary ............................................................................................................................................................................... 484Acknowledgements ................................................................................................................................................................. 485References ............................................................................................................................................................................. 485

1. Introduction

Metal/oxide interfaces play critical roles in many applica-tions including materials science, microelectronics, and chemi-cal applications. Metal–oxide interactions, which consist of in-terfacial charge redistribution and/or mass transport upon in-terface formation, are important factors determining propertiesand performances of the heterojunctions.

For materials science, metal/oxide interfaces can be found inmany technological materials, such as functional ceramics withmetals, oxide dispersion-strengthened alloys, oxide coatingson metals functioning as thermal barriers or natural corrosionprotection layers, etc. Fundamental problems with thesesystems include the adhesion strength, mechanical stability,and fracture behavior of the interfaces, which are all closelyrelated to metal–oxide interactions [1,2]. To further understand

Q. Fu, T. Wagner / Surface Science Reports 62 (2007) 431–498 433

Fig. 1. Schematic of a metal island on a flat oxide support in thermodynamicequilibrium.

this, consider Fig. 1 which shows a schematic of a metalisland supported on a flat oxide surface. In the simplifiedcase of absence of surface energy anisotropies of the metal,a mechanical equilibrium for this system can be written byYoung’s equation:

γoxide = γmetal cos θ + γinterface. (1)

Here, γoxide is the surface free energy of the oxide substrate,γmetal is the surface free energy of the metal overlayer,γinterface is the interface free energy including any metal–oxideinteraction, and θ is the contact angle between γmetal andγinterface. Experimentally, the work of adhesion (Wad) ratherthan the interface energy can be measured. Wad is expressedby the following formula:

Wad = γmetal + γoxide − γinterface. (2)

Thus, the mechanical strength at the interfaces is closelyrelated to the metal–oxide interactions. Furthermore, accordingto (1), γinterface strongly influences the growth behavior of metaloverlayers, determining wetting or non-wetting of the metal,layer growth or island growth of the metal, etc.

In the field of microelectronics, metal–oxide–semiconductor(MOS) field-effect transistor (FET) devices are core compo-nents. As shown in Fig. 2, two interfaces are very importantin the devices: the upper interface of gate electrode with di-electric oxide layer and the lower interface of Si with dielectricoxide layer [3]. In the near future, the actual gate lengths of thedevices will be scaled down to 10 nm and the thicknesses ofdielectric oxide layers will be a few atomic layers [4]. The per-formance of the devices sensitively depends on properties of thetwo interfaces, and they have to be stable in order to be success-ful. Chemical reactions and atom diffusion across interfacesmust be avoided to prevent any failure of the devices, and thedensity of interface electronic states should be well-controlledto obtain stable transport properties at the junctions [5].

Finally, chemical applications will be highlighted. Manycatalytic systems consist of nanosized metal catalysts supportedon oxides. It has been found that the interaction between metalsand oxide supports, so-called metal–support interactions, areof great importance in heterogeneous catalysis [6–13]. Inparticular, the strong metal–support interaction (SMSI) wasfirst suggested by Tauster et al. to explain the suppressionof both H2 and CO chemisorption capacity of metal clusters

Fig. 2. Schematic of important regions of a MOS field-effect transistor gatestack. From [3].

supported on TiO2 which are reduced at high temperatures [14,15]. Later, SMSI was widely observed in many metal/oxidecatalytic systems [8–13]. Two major factors contribute tothe SMSI states, an electronic factor and a geometric factor.The electronic factor is determined by a perturbation of theelectronic structure of the metal catalyst, which originatesfrom charge transfer between the metal and the oxide, whilethe geometric factor results from a thin layer of reducedoxide support physically covering the metal particles (calledthe encapsulation or decoration model), which blocks activecatalytic sites at the metal surface. Fig. 3 is used as anexample that illustrates the encapsulation of a Ru particle byits amorphous titania support [16].

Presently, the sizes of many devices and technologicalmaterials are rapidly decreasing to the nanoregime. As weapproach the nano limit, the density of interfaces increasessubstantially such that the effect of metal–oxide interactionsbecomes more and more significant. In past decades, muchprogress has been made thanks to modern surface sciencetechniques and advanced calculation methods. Model systemswere developed, which consist of nanometer scale metalclusters or ultrathin metal films supported on well-defined oxidesurfaces under the conditions of ultrahigh vacuum (UHV).In this manner, studies of charge redistribution and/or masstransport at metal/oxide interfaces have been significantlysimplified and many of the modern surface science techniquescan be applied to study the interfacial atomic and electronicstructures. Moreover, many sophisticated calculation methodsallow for the derivation of the atomic structure and chargedensity distribution at metal/oxide interfaces. This theoreticalwork contributes to the understanding of the nature ofmetal/oxide interfaces in many aspects, which helps spurfurther experimental development to catch up with the advancesin the theory.


Fig. 3. (a) A Ru particle partially covered by amorphous titania; (b) A Ru particle completely covered by amorphous titania. From [16].

The knowledge of metal/oxide interfaces has been discussedin some excellent books, conference proceedings, andreviews [1,2,17–29]. Many of the fundamental questionsabout these systems have been addressed. In the presentpaper, we concentrate on the understanding of metal–oxideinteractions at low temperatures (below 1000 ◦C), with specialemphasis on experimental and theoretical results obtainedin the past decades using the model systems consisting ofnanostructured metal overlayers on well-defined oxide surfaces.Two aspects of the interaction are particularly examined: thecharge redistribution and the mass transport. It is demonstratedhow these two aspects are influenced by different factors.In addition, generalized guidelines are given to predict theinteraction in the metal/oxide systems and the procedures totune them in a desirable way.

2. Fundamental aspects of metal–oxide interaction

2.1. Physics and chemistry of heterophase interface contacts

Metal–semiconductor (MS) interfaces are the most well-studied heterophase junctions. The knowledge of both thephysical and chemical aspects of the interface contributesmuch to the understanding of other heterophase interfaces,e.g., metal/oxide interfaces. A central issue of concern for theMS interface is the electronic properties of the device, whichcan be characterized by the Schottky barrier height (SBH). Inpast decades, various theories have been developed to describethe basic mechanism of the barrier formation [30–32], which isnow discussed below.

2.1.1. Physics at metal–semiconductor interfacesSchottky [33] and Mott [34] suggested that energy bands

of a metal and a semiconductor in contact adjust so as toalign both vacuum levels. They considered long-range chargetransfer, but ignored any local interaction at the interface. Dueto the readjustment of the vacuum levels, it was necessary forthe Fermi energy levels (EF ) to be “bent” to higher (or lower)energy. This forms a barrier at the interface that originates froma space charge layer at the semiconductor surface. The barrierheight, φB,n , is simply the difference between the metal workfunction φM and the electron affinity of the semiconductor χS :

φB,n = φM − χS . (3)

Fig. 4. Energy band diagram of a metal–n-type semiconductor interface at theSchottky limit. VBB, band bending at semiconductor surface.

This is the Schottky limit for the band bending (Fig. 4). In theSchottky approximation, the charge transfer process and theform of space charge layers are governed by Poisson’s equation

d2Vd2r= −

ρ(r)

εε0, (4)

where ρ(r) is the space charge density, ε is the staticdielectric constant of the semiconductor, and ε0 is thevacuum permittivity in free space. This has been extensivelydocumented and can be found in the classic Schottky contacttheory [35,36].

Bardeen [37] first considered that there may be interfacestates located in the semiconductor band gap. In caseof sufficiently high density of the interface states, thesemiconductor is completely screened by the interface statesand the SBH is given by

φB,n = Eg − φo (5)

where Eg is the band gap energy of the semiconductor. φois defined as the energy measured from the valence bandmaximum (VBM) at the semiconductor surface to the levelbelow which all interface states must be filled for chargeneutrality at the surface. This phenomenon is known as Fermilevel pinning. It should be noted that Eq. (5) only considers thelocal charge transfer between the metal and the interface states.

As shown in Fig. 5, Cowley and Sze [38] assumed acontinuum of interface states with a constant density of states,Dis , and proposed a hybrid approach, where they combined the


Fig. 5. Energy band diagram of a metal–n-type semiconductor contact with aninterfacial layer after Cowley and Sze [38]. φ0: energy level measured fromthe VBM at the semiconductor surface to the level below which all interfacestates must be filled for charge neutrality at the surface, 1φn = image forcebarrier lowering; 1: potential across interfacial layer; εi t : dielectric constantof interfacial layer; δ: thickness of interfacial layer; VBB: band bending atsemiconductor surface.

Schottky limit and the Bardeen limit in order to analyze chargebalance at the interface. They derived the following relation:

φB,n = γ (φm − χ)+ (1− γ )(Eg − φo

)(6)

where γ , the interface parameter, is defined by

γ =(

1+ e2 Disδ/εi t

)−1(7)

with δ defined as the thickness of the interfacial layer, and εi t isthe permittivity of the interface layer. γ can be compared to theexperimentally determined slope parameter Sφ = ∂φB,n/∂φm .For a high density of interface states (Dis → ∞, γ → 0), thebarrier height becomes independent of the metal work function,i.e. the Bardeen limit; if Dis = 0 and γ = 1, it results inthe Schottky limit. Covalent solids, e.g. Si, present smaller γ ,indicating the dominant role of the local charge transfer andthe strong Fermi level pinning; while ionic semiconductorswith wide band gap, e.g., SrTiO3 and SiO2, have larger γ ,which suggests that interface states do not affect the interfaceformation significantly [39–41].

Concerning the physical origin of the interface states, animportant theoretical approach is based on the concept of metal-induced gap states (MIGS) [42–44]. Heine [42] suggestedthat at a MS interface, the metal electronic wave functionstail into the semiconductor in the energy range in which theconduction band (CB) of the metal overlaps the band gap of thesemiconductor. These tails give rise to a continuum of MIGS inthe semiconductor. The charge neutrality level (CNL), φCNL, ofthe MIGS is defined such that the tail of the metal wave functionpenetrating into the semiconductor will carry no charge when itcoincides with EF . The charge transfer is thus governed by theposition of EF relative to φCNL.

Tung [32,45,46] has used another way to account forthe experimentally observed strength of Fermi level pinningon different semiconductors. His bond polarization theoryassumes that the most significant charge rearrangement, when

a metal comes into contact with a semiconductor, is due tothe formation of polarized chemical bonds at the interface.The transferred charge and the interface dipole can be treatedby the electrochemical potential equalization [47,48] methodemployed in molecular physics. It has been shown that thepolarization of interface bonds leads to a weak dependence ofthe SBH on the metal work function.

2.1.2. Chemistry at metal–semiconductor interfacesBesides the physical explanations, there is much experimen-

tal evidence showing the importance of chemical mechanismsin the charge transfer and Schottky barrier formation at MS in-terfaces.

Andrews and Phillips [49] attempted to correlate the SBHswith the heats of formation of transition metal silicides, 1H f(tmSi). They showed a linear relation between the experimentalφB,n and 1H f (tmSi) indicating the strong effect of interfacechemical bonding on the barrier formation. At transitionmetal–Si interfaces, Ottaviani et al. [50] also identified thestrong relationship between SBHs and interface reactions.The reaction between the two contacting solids produces aninterfacial layer, which dominates the SBH. The relationshipis manifested through the correlation between the SBHsand eutectic temperatures for the transition metal–silicide–Sisystems. Brillson [51] revealed a systematic dependence ofSBHs on interface reactivity. The SBH of metals on individualcompound semiconductors, such as ZnO, ZnS, CdS, and GaP,demonstrates a sharp transition as a function of heat ofreaction 1HR (from the most stable known metal–anion bulkcompound). The local charge redistribution, which is madeevident by new interface compounds for reactive metals and bydiscrete interface states for non-reactive metals, determines thebarrier formation [52].

On the other hand, the chemical trends in SBHs weremade to manifest with the concept of electronegativity.Originally, Kurtin [39] used the electronegativity difference,1X , between the constituents of the compound semiconductore.g. 1XGaAs = XAs − XGa to describe the semiconductorionicity, and showed a correlation between the slope parameter,S, and 1X . Monch [31,53] applied Pauling’s electronegativitymodel to explain the charge transfer across the MS interface.It is proposed that the transferred charge varies proportionallyto Xm − X S , which is simply the difference between theelectronegativities of the metal, Xm , and the semiconductor,X S . In case of Xm − X S = 0, the metal EF will coincide withthe CNL of MIGS and no charge transfer occurs.

Thus, the Schottky barrier formation should be attributedto a two-fold charge transfer upon the interface formation: thelocal charge transfer and the long-range charge transfer [54].The long-range charge transfer between the metal band andsemiconductor band occurs over the semiconductor spacecharge region with a scale of screening length L [35]. Theprocess can be described by the Schottky contact theory.The local charge rearrangement happens at the interface withdimensions of A. The local interaction may be characterizedin the band picture using the MIGS concept [42–44,55] or in


the bond picture which considers interface chemical bondingbetween the metal and semiconductor atoms [45,54,56].

It should be mentioned that there is a strong interplaybetween the physical and chemical aspects of the interface. Onthe one hand, charge transfer, formation of interface dipole,and SBHs are strongly correlated with the chemical interaction(e.g., chemical reaction and atom diffusion) at the interfaces.On the other hand, charge transfer produces strong electricfields at the interface, which further affects the interfacereaction or diffusion dynamics [30].

2.2. Electronic interaction of metals with oxides — chargeredistribution

Interfacial contact between a metal and an oxide can resultin charge redistribution at the junction similar to the situation atthe MS interfaces discussed above. The electronic interactionsare simply driven by the principles of the system energyminimization and the continuity of the electric potential in asolid [32].

Electron transfers always occur at reactive interfaces, wherealready existing chemical bonds are broken and new ones areformed resulting in interdiffusion over length scales of morethan a single monolayer and formation of new phases with athickness of at least one monolayer [29]. These systems willbe considered in Section 2.3. Here, the discussion of electronicinteractions is limited only to non-reactive interfaces, wherethe localized charge redistribution occurs among atoms at theinterface and/or the delocalized charge transfers between themetal and the space charge region of the oxide. Many theoriesand methodologies have been developed to reveal the variouscontributions to the electronic interactions.

2.2.1. Local charge redistributionThe local charge redistribution is the electron rearrangement

involving a few atomic layers at the interface. Depending on themetal/oxide system, different mechanisms may be dominant inthe electronic interaction [22,57]. The following sections detailthe relevant mechanisms.

2.2.1.1. Empirical correlations from dispersion force. At ametal–insulator interface, the mutual polarization of thetwo media gives rise to dispersion forces. The dispersioncontribution, i.e. the van der Waals interaction, is quiteweak compared to other interactions, such as the electrostaticinteraction [22,58]. Therefore, it is believed that this interactionis important only in the systems consisting of a noble metaldeposited on a wide band gap oxide [59]. Under this condition,the van der Waals contribution has been shown to qualitativelyaccount for some of the interaction energy experimental results.From this, an empirical correlation has been derived showingthat the adhesion energies at metal/oxide interfaces increase asthe plasmon energy of the metal increases and/or as the bandgap of the oxide narrows [59,60]. Stoneham [61] found anotherempirical correlation where the optic dielectric constant, ε∞, ofthe oxide provides a simple classifying rule. An ε∞ higher thana critical value leads to wetting; indeed oxides with sufficientlyhigh ε∞, such as TiO2 and Nb2O5, exhibit SMSI effects incatalysis.

Fig. 6. Image charge interaction. (a) Interaction between a point charge (q) anda metal surface. (b) Image charges at a metal/oxide interface. From [2].

2.2.1.2. Image charge theory. Stoneham and Tasker [62]first considered that many phenomena associated withmetal/nonmetal interfaces with a large dielectric constantmismatch can be understood in terms of image interactions dueto charges in the nonmetal. Based on classical electrodynamics,they suggested an “image charge theory” which describes theinteraction between a metal and an ionic crystal by electrostaticforces [62,63].

Positioning a charge, q, at distance, z0, from a metal surfacewill cause rearrangement of the metal charge and induce aquasi-free image charge on the metal surface. Such a chargedistribution results in a potential of the form

V (r, z) =−q[

r2 + (z0 + z)2]1/2 (8)

where r is the distance parallel to the surface. At a metal–ionicoxide interface, anions and cations in the oxide induce imagecharges in the metal. The image theory assumes that themetal–oxide interaction energy originates from the attractiveCoulomb force between ions in the oxide and their images inthe metal (Fig. 6). The interaction may give adhesion energy ofthe order of joules per square meter [22].

It is known that the valence band (VB) of metal atomssupported on an ionic oxide can be distorted by the electrostaticfield of the oxide. This phenomenon is called the polarizationeffect, and results in electron redistribution among the differentcomponents of the metal band [64–67]. As an example, uponadsorption of Pd or Ru on the surface oxygen of sapphire,electrons are depleted from the d2

z orbital and move to thelateral d orbitals. The polarization of these atoms makes theadsorbate positive above the O ions and negative between theO ions [66]. In the case of ultrathin metal overlayers, e.g. onemonolayer (ML) of adsorbate, the binding of metals to oxidesis dominated by the polarization effect, which mimics themacroscopic “image charge” effect.

2.2.1.3. Metal-induced gap states. Noguera and Bordier [57]tried to understand the nature of bonding between a metal andan oxide using the concept of MIGS, which was introducedby Heine to model MS junctions [42]. They treat the metalas a jellium, whose eigenstates are plane waves, and theoxide as a tight-binding system with local orbitals. MIGS areproduced by matching the delocalized metal wave functionswith exponentially decaying oxide states in the band gap energyrange at a defectless interface. The effect of the metal jellium


seems to allow some states originally in the valence andconduction bands of an oxide to be pulled into the band gapof the oxide. The MIGS are associated with a smooth densityof states in the band gap and their density decays from theinterface into the oxide over a length of several A [68]. Thislength correlates with that of the ionicity of the oxide wherelarger band gap energy leads to smaller length.

The gap states are derived from the bulk bands of the oxide.Thus, they may have donor-like or acceptor-like characterdepending on the position of the states in the band gap. Likethe φCNL of MIGS at MS interfaces, Noguera and Bordier statethat charge neutrality occurs if the MIGS are filled up to EZCP,the zero-charge point. The significance of EZCP is that it can beused to predict if charge transfer will occur when a metal is incontact with an oxide. If, for example, EF of a metal is aboveEZCP, a transfer of electrons from the metal to the oxide wouldtake place, creating an interface dipole.

2.2.1.4. Interface bonding. Formation of chemical bonds isan important mechanism for charge transfer at metal/oxideinterfaces. It is known that either covalent bond or ionic bondcan be established between metal atoms and surface oxygenions or cations of oxides.

Polar covalent bonds, in which the metal and oxideelectronic orbitals are strongly hybridized, are partly ionic.As proposed by Pauling [69], the ionic character of a polarcovalent bond between atoms A and B is determined by theirelectronegativities, XA and XB. The bonding charge shifted tothe more electronegative atom is expressed by the relation [70]

1q/e = 0.16 |XA − XB| + 0.035 |XA − XB|2 . (9)

In case of a compound solid, the electronegativity is given bythe geometric mean of the atomic values of the constituents,e.g. XAB = (XA XB)1/2 [31]. Using the simple conceptof Pauling’s electronegativity, one can predict the directionand amount of charge transfer at metal/oxide interfaces.An alternative way to describe the charge transfer throughmetal–oxide bonding can be adopted from the bond polarizationtheory of MS interfaces suggested by Tung [45]. Theconglomerate of the entire metal/oxide interface is regardedas a giant “molecule” and the charge redistribution in the“molecule” can be treated using methods applied in molecularphysics. The transferred charge per interface bond is estimatedto be

q =∣∣∣∣φM − χM Ox − Eg/2

Eg + κ

∣∣∣∣ (10)

where φM is the metal work function, χM Ox the electron affinityof the oxide, Eg is the band gap energy of the oxide, and κ theparameter related to interface characters [45].

An ionic bond is formed via electron donation from oneatom to another without orbital mixing. At a metal/oxideinterface, the ionic bonding is accompanied by charge transferbetween interfacial atoms, leading to a redox reaction.Experimentally, ionic bonds are often observed during theinitial stage of deposition of reactive metals on oxides. An

example of this is the oxidation of metal adatoms and reductionof topmost cations which take place via metal-to-oxygen andoxygen-to-cation charge transfer [71,72]. The bonding and theresultant reaction concern atoms at the interface, and, therefore,may not necessitate atomic diffusion over more than one latticedistance.

It is worth mentioning that formation of interface bondingstrongly depends on oxide surface properties including, but notlimited to, surface termination, surface reconstruction, surfaceimpurities, and surface point defects. A good example canbe found for the Ag/MgO interface where chemical bondformation is demonstrated to be not important [73,74]. Ona defective surface, however, a Ag atom sitting on a neutralMg vacancy, V 0

s , donates two valence electrons to the four Oatoms surrounding the vacancy which forms strong ionic bondsbetween the Ag and O atoms [74]. As we will show below,defects on oxide surfaces, such as Al2O3(0001), MgO(100),SiO2, and TiO2(110), greatly affect metal adsorption and theatomic interaction energy.

2.2.1.5. Theoretical calculations. Many sophisticated calcula-tion methods allow for the derivation of the charge density dis-tribution at metal/oxide interfaces. Interface electronic effects,such as polarization, bonding, charge transfer, etc., can be ex-plicitly derived from these computational results.

To model metal/oxide interfaces, two main approaches areused in electronic structure calculations, a cluster model and aslab modeling method [75–78]. In the cluster model the solidis replaced by a finite cluster of atoms and is based on theassumption that all interactions are locally well-described. It isimportant to carefully choose the cluster such that all importantaspects of the electronic structure of the infinite system canbe reproduced. The slab model is to take advantage of thetranslational symmetry of the system and model the interfaceby slabs. In this model, a slab consisting of a finite number oflayers mimics the semi-infinite system, with a 2D translationalperiodicity. In this model, a band structure calculation isusually performed in the reciprocal space using delocalizedplane waves (PW) for expanding the wave functions. Thecluster approaches often use the real-space molecular orbitals,e.g. the linear combination of atomic orbitals (LCAO) of thesurface atoms. These calculations can be performed using semi-empirical methods, density functional theory (DFT), or from abinitio principles. The advantage of these quantum mechanicalcalculations is that they contribute to the understanding ofthe nature of metal/oxide interfaces in many systems wheretechnology has not advanced to the point where the necessaryexperiments can be done.

2.2.2. Space charge transferOxides, such as TiO2, SrTiO3, ZnO, etc., have relatively

small band gap energies, (usually <3.5 eV) [17]. Intrinsic orextrinsic dopants may produce free electronic carriers inside thesolids causing them to become equivalent to n-type or p-typesemiconductors. Thus, for these systems, MS contact theoryis applicable to describe interface contact phenomena betweenmetals and the semiconducting oxides.


Recall from 2.1.1 that ionic semiconductors, e.g., SrTiO3,ZnO, and SiO2, present a larger γ and Sφ [39–41], whichsuggests that interface states on their surfaces do not controlthe interface formation. It is therefore reasonable to assumethat charge transfer happens primarily between metals and thespace charge regions at oxide surfaces. The relative positionof EF of the metal and the oxide determines the direction andmagnitude of the charge transfer. This process can be describedusing the Schottky contact theory (see Fig. 4). It is now usefulto discuss two cases describing supported metal overlayers onoxides, those with continuous metal overlayers on oxides andthose with metal clusters or discontinuous overlayers on oxides.

(1) For an infinite interface, (continuous metal overlayerson oxides) under the condition of strong depletion of ann-type semiconducting oxide, the space charge density isapproximated by a step function with

ρ(z) = eND 0 ≤ z ≤ d, (11)

where ND is the donor concentration in the bulk semiconductor,and d is the thickness of the depletion layer. Poisson’s equation(4) becomes

d2Vd2z= −

eND

εε0, (12)

with the relevant boundary conditions, V (z = 0) = VBB =

(φm − φsc)/e the band bending at the surface, and V (z = d) =

0. The thickness of the space charge layer is calculated to be

d =(

2εε0 |VBB|

eND

)1/2

. (13)

The transferred charge per unit area from the semiconductor tothe metal, Q, is

Q = (2εε0eND |VBB|)1/2 . (14)

(2) The finite interface (metal clusters or discontinuousoverlayers on oxides) model is used to describe the situationof small clusters on semiconducting oxides. Verykios andcoworker [79] suggested a physical model consisting ofspherical metal particles of radius, rM , embedded in thesemiconductor bulk (Fig. 7). Within approximation (11),Poisson’s equation in 3D form can be analytically solved. Thethickness of the depletion region, d , and band bending, VBBhave a relation of

VBB =eNd

ε0ε

(d2

2−

r2M6−

d3

3rM

). (15)

The number of electrons transferred at the interface is nowgiven by

Q =4π Nd

3

(d3− r3

M

). (16)

As seen in the above equations, both d and Q are functions ofthe cluster size, rM . For example, on doped TiO2, the amount ofcharge transferred to the metal is dependent on the metal clustersize. They report 0.5 electrons per metal atom for a cluster size

Fig. 7. Physical model used to simulate the contact of a metal cluster with asemiconducting oxide support. From [79].

of 1.5 nm and only 0.01 for crystallites larger than 10 nm [79].Eq. (16) well illustrates the effect of size on charging the metalclusters.

The long-range charge transfer between a metal and an oxideproduces space charges and band bending at the oxide surface.Experimentally, band bending with a few tenths eV has beenobserved when various metals are deposited on ZnO [80,81],TiO2 [82–84], and SrTiO3 [85,86]. In these systems, there existinterfacial dipole fields due to the space charges in the devices.Typical space layers extend for 10–100 nm, and lead to anenergy difference of 0.2 eV and fields of 106–107 V/m [87].The built-in electric fields clearly have a strong effect onreaction kinetics as we will discuss below.

2.3. Chemical interaction of metals with oxides — masstransport

In this case chemical interaction between metals and oxidesis regarded as an interaction which results in chemical reactionsat the interfaces. The term chemical reaction is used forthose cases where interdiffusion occurs over length scales ofmore than a single monolayer or where new phases with athickness of at least one monolayer are formed [29]. Thechemical interaction involves mass transport over more thanone lattice distance. The driving force behind the mass transportin the solids is a gradient in the electrochemical potential,consisting of a gradient in the electrical potential and/orchemical potential.

2.3.1. Chemical interactions at metal/oxide interfacesDue to the nature of the products formed at metal/oxide

interfaces, the chemical interaction can be generally classifiedinto four different groups, redox reaction, alloy formation,encapsulation, and interdiffusion [20,23]. Each case is nowdiscussed and the mass transport processes during everyreaction are to be demonstrated.

Case 1: Redox reaction. A redox reaction at a metal/oxideinterface (MeI

‖ MeIIOx ) will occur by oxidizing the metaloverlayer and reducing the oxide substrate. The reaction can beschematically written as:

MeI‖ MeIIOx → MeIO

y ‖ MeIIOx−y .

Of all the reactions occurring at metal/oxide interfaces,these are the most frequently observed, especially for reactivemetals on oxides such as TiO2 and SrTiO3 [18,29,88,89].For example, thin epitaxial Nb films on TiO2(110) have


Fig. 8. (a) HRTEM image of Nb/TiO2 (110) interface formed at room temperature. A distorted interface between Nb and TiO2 crystal is present [90]. (b) HRTEMimage obtained on the 5% Pt/CeTbOx catalyst reduced at 1173 K, resulting in the formation of CePt5 alloy [93]. (c) HRTEM image of a 0.5% Rh/Ce0.8Tb0.2O2−xcatalyst reduced at 1173 K showing the encapsulation of Rh particle by support [94]. (d) HRTEM image of Cr/TiO2 (110) interface formed at 400 ◦C. The interfacelayer of TixCryOz is from the interdiffusion between Cr overlayers and TiO2 substrate.

been prepared by molecular beam epitaxy (MBE) at roomtemperature. Investigation of the interfaces of these systemsby high resolution transmission electron microscopy (HRTEM)and electron energy loss spectroscopy (EELS) confirmed thatthe first two monolayers of Nb were oxidized, resulting inthe partial reduction of a ∼2 nm thick TiO2 layer near theinterface [90] (Fig. 8(a)). Mass transport in the form of oxygen(i.e. oxygen vacancy) diffusion, commonly occurs during thereaction process (e.g., [91,92]).

Case 2: Alloy formation. At some interfaces, stableintermetallic compounds may be formed according to

MeI‖ MeIIOx → MeIMeII

y ‖ MeIIOz,

or

MeI‖ MeIIOx → MeI Oy ‖ MeIMeII

z ‖ MeIIOx .

The former case, where MeII is reduced during alloy formation,occurs for noble metals supported on oxides of CeO2 [95–97],SiO2 [98–100], Al2O3 [96], and SnO2 [101]. Pt nanoparticlessupported by thin films of silica, alumina, and ceria weresubjected to hydrogen reduction at temperatures up to 1073 K.Formation of Pt-rich Pt3Me (Me = Si, Al, Ce) alloy phaseswere identified in all three systems after the reactions [96].Similar reactions of the alloy formation were observed for othernoble metals, such as Pd, Rh, Ni, and Cu, on these oxides.Fig. 8(b) gives one example of formation of an intermetalliccompound at a Pt/CeTbOx interface.

In cases of reactive metals deposited onto Al2O3 and SiO2substrates, the second case may happen. The reaction leads toformation of a thin layer of aluminide or silicide sandwichedbetween the substrate and a top layer of metal oxide [102,103].


Table 1Summary of interface reactions at metal/oxide interfaces

Chemical interaction at MeI‖ MeIIOx Interfacial products Main mass transport process Typical systems

Redox reaction MeIOy ‖ MeIIOx−y Oxygen RMa on TiO2 & SrTiO3

Alloy formation MeIMeIIy ‖ MeIIOz Cations NMb on CeO2 & SiO2

MeIOy ‖ MeIMeIIz ‖ MeIIOx RMa on Al2O3 & SiO2

Encapsulation MeIIOx−δ ‖ MeI‖ MeIIOx Cations NMb on TiO2 & CeO2

Interdiffusion MeIMeIIOy Metal atoms and/ Ni/Al2O3or support atoms Al/spinel

a RM: reactive metals.b NM: noble metals.

In alloy formation processes, it is mostly expected thatcations are physically extracted from oxide substrates andincorporated in metal overlayers.

Case 3: Encapsulation. Also called decoration, theencapsulation reaction involves mass transport from the oxidesupport onto the surface of metal particles. It necessitatesthe cations’ outward diffusion from the bulk to the interface(e.g.[13]). Encapsulation is a special interface process wherebymetal particles are physically covered by a thin layer ofreduced oxide support (Fig. 8(c)) according to the followingrelationship:

MeI‖ MeIIOx → MeIIOx−δ ‖ MeI

‖ MeIIOx .

This reaction results in blocking active catalytic sites on metalsurfaces and contributes to the SMSI state [8,11,13,14]. Noblemetals, including Pt, Pd, and Rh on TiO2 are good examples ofencapsulation reactions. Ultrathin layers of TiO2−x (0 < x <

1), which have a thickness of several atomic layers, have beenconfirmed to emigrate onto surfaces of metal particles (see [13,94] and references therein). Similar results can be found fornoble metals on CeO2 [95].

Case 4: Interdiffusion. It is known that metals may diffuseinto their oxide supports and/or substrate atoms may diffuse tothe metal surface. Such interdiffusion leads to the formation ofinterdiffusion zones or mixed oxides (e.g. ternary oxides andoxide solid solutions) at the interfaces (Fig. 8(d)). They can beexpressed as

MeI‖ MeIIOx → MeIMeIIOy .

One example is that NiAl2O4 has been confirmed to format the Ni/Al2O3 interface [104]. Raj et al. have also shownthat an ion exchange reaction between Al and spinel ceramicof (MgO)·(1.25Al2O3) produces interdiffusion zones at theinterface [105,106].

The four reactions at metal/oxide interfaces are summarizedin Table 1 and show the typical reaction systems and masstransport processes. We now turn to a discussion of thethermodynamic and kinetic aspects of these interface reactionsin the following sections.

2.3.2. Thermodynamic considerationsThermodynamics can be used to find out if a chemical

reaction is favorable. The classic thermodynamics may consider

the bulk thermodynamic data. For interface reactions, theinfluences of interface and surface should be taken intoconsideration. These two cases are being discussed here.

(1) Bulk thermodynamics: As a first approximation, simplebulk thermodynamic calculations for solid state reactions areoften used to predict interface reactions in metal/oxide systems.For example, a redox reaction at a metal/oxide interface may beexpressed as follows:

MeI+MeIIOx → MeIOy +MeIIOx−y .

The Gibbs free energy change for this reaction, 1G R = 1HR−

T 1SR , will suggest the feasibility of the reaction. For solidstate reactions, changes in entropy are negligible such thatthe enthalpy changes for the reactions can be simply used forthermodynamic criteria. As discussed by Campbell [23], theheat of formation of oxides (1Ho

f ) is a decisive thermodynamicparameter for description of the interface reactions. The heatsof oxide formation per mole of oxygen (1Ho

f in kJ/mol O)were compiled in his review paper. He showed that 1Ho

f can besuccessfully applied to explain the experimental results in manymetal/oxide systems. The trend in metal overlayers’ reactivityto a specific oxide, e.g. TiO2, can be explained very well by thevariation in 1Ho

f of the metals [28].An alternative way to describe reactivity of metals to

oxides is to introduce the concept of oxygen affinity of metals(pO) [107]. pO is defined as

pO = − log pO2 = −1G f /RT, (17)

at 1000 K for the equilibrium between the metal and its lowestoxide:

1G f is given for the reaction of a metal with one mole ofO2:

2n

Me+ O2 →2n

MeOn .

pO is often used to compare the reactivity of metals tooxides [20,108].

To obtain a rough estimate if a reaction occurs, the Gibbsfree-energy-changes for reactions, such as alloy formation, canbe calculated using the corresponding reaction formula.

(2) Thermodynamics including interface terms: In the abovethermodynamic calculations, energies associated with surfacesand interfaces were not included. In fact, the contribution from


Fig. 9. Schematics showing the mass migration of TiOx (x < 2) onto metalclusters driven by the minimization of surface energy of the whole system [13].

the surface and interface energies to the total free energy couldbe significant, particularly for nanostructured metal overlayerson oxides where the ratio of interface to bulk is very large. At ametal/oxide interface, interface energies include metal surfaceenergy γmetal, oxide surface energy γoxide, and interface energyγinterface (recall Fig. 1). These factors may play a critical role insome reaction processes, such as oxidation and encapsulationreactions.

Oxidation reaction: Metals generally have a larger surfaceenergy than an oxide such that 3D metal islands are favoredto grow on oxide surfaces due to γoxide < γmetal + γinterface.If the metal islands are transformed into oxides, the decreasein surface energy may reverse the above inequality and drivethe conversion of 3D islands to 2D structures. This can beseen in the Cr/SrTiO3(100) system, where we have observedthe flattening of Cr oxide islands after Cr oxidation above600 ◦C [91]. Also, STM investigations show that V clusterssupported on a thin alumina film become flatter in case ofoxidation at 800 K in UHV [109]. Finally, room temperatureadsorption of oxygen to Cu islands supported on TiO2(110)allows 2D Cu islands to form on the surface at the expense ofthe existing 3D islands [110].

Encapsulation: It has been proposed that minimization ofthe surface energy of a system is one of the main driving forcesfor an encapsulation reaction [13,111,112]. Only in systemswhere metals have high surface energies and oxides have lowsurface energies, can supported metal clusters be encapsulatedby oxide support layers (Fig. 9). The surface energetic factorexplains why Pt and Pd but not Au and Ag (both metals havelower surface energies) have been subjected to the reactions.Similarly, oxides with low surface energy, e.g. TiO2 and V2O5,undergo encapsulation reactions more easily than oxides withrelatively high surface energies, such as SiO2 and Al2O3.

Restrictions of thermodynamic calculations are now beingaddressed. Thermodynamic data of interfacial reactants andproducts are often unavailable so, as an approximation,bulk thermodynamic data are typically used to calculate thedynamics. These bulk data are frequently different from thatof the interface phase. Secondly, the chemical composition ofthe interface phases is often unknown. This can be seen in the

case of a redox reaction at a metal/oxide interface, where anoxidized metal (MeIOy) and reduced oxide (MeIIOx−y) thatpossess different stoichiometry, may form complex phases atthe interfaces. Accurate phase determination is complicated,since the resulting phases are often approximately a fewatomic layers or unit cells thick. Thus, calculations can onlybe performed assuming specific reaction phases. Moreover,equilibrium conditions are assumed in these calculations, whichmay be very wrong since solid state reactions are oftenkinetically limited, especially at relatively low temperatures(e.g. <1000 ◦C). Therefore, nonequilibrium thermodynamicsand reaction kinetics must be considered as discussed below.

2.3.3. Kinetic considerationInterface reaction kinetics can be controlled by the reaction

itself, the mass transport process, or both. If the interfacereaction controls the kinetics, the growth rate of newlyformed layer is linearly dependent on the reaction time.Whereas parabolic reaction kinetics, where the reaction rate isproportional to the square root of time, is observed in case ofdiffusion-controlled reactions (e.g. [113]). At many metal/oxidesystems, as we will see below, interface reactions at relativelylow temperatures are often limited by mass transport, e.g. thetransport processes shown in Table 1.

Using the hopping model for ionic diffusion in solids,Fromhold and Cook [114] derived an equation for the ionicdiffusion current J in the steady-state approximation

J = 4aν exp(−W/kB T ) sinh(ZeE0a/kB T )

×

[C(L)− C(0) exp(ZeE0L/kB T )

1− exp(ZeE0L/kB T )

]. (18)

W is the activation energy; ν is the ionic attempt frequency;2a is the ionic jump distance; E0 is the electric field appliedexternally or from space charges in solids; Ze is the effectivecharge per particle of the ionic species undergoing transportthrough the lattice; C(L) and C(0) are the bulk defectconcentrations of the diffusing ionic species at x = L andx = 0. Eq. (18) is the nonlinear diffusion equation in thehomogeneous field limit, which can be simplified by variousapproximations.

Approximation (1). In case of E0 → 0, which would occurif the effect of electric field is negligible or it is diffusion ofneutral particles in a solid, Eq. (18) changes to

J = −4a2ν exp(−W/kB T ) [C(L)− C(0)] /L , (19)

which can be written as

J = −D [C(L)− C(0)] /L (20)

with a diffusion coefficient D = 4a2ν exp(−W/kB T ). This isthe form of Fick’s first law showing that the particle current islinearly related to the spatial concentration gradient ∂C/∂x .

Approximation (2). Eq. (18) under a low electric field and forthe limiting case of a high temperature, i.e., E0 � |kB T/Zea|,results in

J = µE0

[C(L)− C(0) exp(ZeE0L/kB T )

1− exp(ZeE0L/kB T )

]. (21)


Fig. 10. Schematic diagrams showing the potential energy of an interstitialmetal “ion” as a function of position near the metal/oxide interface. The electricfield at oxide surface, which is generated by the interfacial charge transfer,lowers the energy barrier for ions moving away from the metal/oxide interface.After [116,117].

Here, µ = (Ze/kB T )4a2ν exp(−W/kB T ) is the mobilityof the diffusing particle, which is related to the diffusioncoefficient, D, through the well-known Einstein relationµkB T = ZeD. The above equation is the integral form ofthe ordinary linear diffusion equation in the homogeneous fieldlimit

J = −D∂C(x)

∂x+ µE0C(x), (22)

where the ionic diffusion current is linearly dependent on theelectric field.

Approximation (3). In the case where the driving forcedue to the concentration gradient is insignificant compared tothat from the electric field, we have C(L)/C(0) → 1 andexp(ZeE0L/kB T ) � 1. This is the high electric field and lowtemperature limiting case. Eq. (18) becomes

J = 4aCν exp(−W/kB T ) sinh(ZeE0a/kB T ), (23)

or

J = 2nν exp(−W/kB T ) sinh(−ZeaVM/kB T L), (24)

where n is the number of ions per unit area which are in aposition to jump the barrier W , n = 2aC . This is the equationgiven by Mott, which is based on the assumption that nonlineardiffusion under the aiding potential VM (Mott potential) is rate-limiting [115,116]. It can be seen from Eq. (23) that an electricfield E0 modifies the activation energy for the ionic motionfrom W to W − ZeE0a as shown by Fig. 10. The additionalterm describes either an extra driving or retarding force for thetransport of ionic defects. The ionic diffusion current showsan exponential dependence on the electric field such that themacroscopic electric field contributes largely to ionic diffusion.At relatively low temperatures, kB T � W , ordinary diffusiongiven by (20) is negligible and the diffusion of ionic defects isoften thermally limited. Therefore, electric fields from surfaceor space charge layers may play an important role in the defectdiffusion and the reaction kinetics. Such a field-driven effectis quite general and critical in many interface reactions. Wediscuss this in more detail in the next section.

2.4. Interplay between electronic and chemical interactions

At low temperatures, e.g. room temperature, defect diffusionin solids is often thermally limited, whereas a large electric field

could lower the energy barrier thereby influencing the ionicdiffusion and chemical reaction. As discussed in Section 2.2,the electronic interaction, in particular, the long-range chargetransfer at metal/oxide interfaces, produces space charges andconsequently, induces an electric field. It is therefore expectedthat there is interplay between the electronic interaction and thechemical interaction in some cases. The coupling between thedefect diffusion and charge transfer enables interface reactionsto be dependent on the electronic structure of the interfaces.The process is explained in the framework of a GeneralizedCabrera–Mott theory [13,85,89], which is now discussed indetail.

2.4.1. Cabrera–Mott theory in solid–gas interface reactionsThe Cabrera–Mott theory was first suggested to explain the

low temperature oxidation of metal surfaces. Later, it was foundthat the etching of semiconductor surfaces can be understoodusing the same theory. Here, we give a brief overview of the twodifferent surface reactions. Special attention is given to chargetransfer and mass transport processes at the surfaces.

2.4.1.1. Low temperature oxidation of metal surfaces. Manytheories have been developed to describe the oxidation of metalsurfaces ([117,118] and references therein). Cabrera and Mottsuggested a theory of formation of thin films by metal oxidationat low temperatures [115,116]. In this so-called Cabrera–Motttheory, they considered two important points, (1) electrontransfer occurs between surface adsorbed oxygen and metal and(2) the resulting electric field in the oxide layer can lower thebarrier for ionic mobility. Both of these cases are discussed insome detail.

To begin the discussion of the Cabrera–Mott theory, weconsider the fact that a layer of atomic oxygen is adsorbedat the gas–solid interface during the oxidation reaction. Thecentral assumption of Cabrera and Mott is that electrons canpass through the oxide layer from the metal to the oxygen atom.This charge transfer is facilitated by thermionic excitation fromthe metal into the conduction levels of the oxide or by electrontunneling through the oxide barrier [119,120]. This electronicmotion is considered to be rapid compared to the ionic motionsuch that an electric field across the oxide layer forms anda quasi-equilibrium state is set up between the metal and theadsorbed oxygen. A contact potential VM , known as the “Mottpotential”, forms and is defined by the initial difference in EFof the metal and O2p level. The Mott potential VM is defined as(see Fig. 11)

VM = (φ0 − φL)/e, (25)

where φ0 is the work function of the metal, and φL is the energydifference between the vacuum energy level and O2p energylevel.

As stated above, an electric field, E0 = −VM/L(t), mayoccur both in the oxide layer with thickness of L(t) and atthe metal/oxide interface. It may lower the energy barriers forthe initiation of ionic motion and enable the metal cations oroxygen anions to move through the oxide layer without muchhelp from temperature. For very thin films, the field is large


Fig. 11. Electronic levels in the metal, oxide, and adsorbed oxygen (after Cabrera and Mott [116]): (a) before electrons have passed through the oxide, (b) whenequilibrium is reached. The thickness of grown oxide layer is L(t). VM is the “Mott potential” producing an electric field E0 at the metal surface. φ0 is the workfunction of the metal, and φL is the energy difference between the vacuum energy level and O2p energy level. χ0 is the energy difference between the conductionband minimum (CBM) in the oxide and EF of the metal and χL the energy difference between the CBM in the oxide and the O2p energy level.

such that ion motion may show an exponential dependence onthe electric field. The ion diffusion current can be expressedusing the formula (24).

The Cabrera–Mott theory predicts a very rapid initial growthrate followed by a sharp leveling off to a very slow growth stageat thickness of 50–150 A. This reaction mechanism has beenfound to be active in many oxidation experiments [116–118,121–124].

2.4.1.2. Etching of semiconductor surfaces. Reactions ofsemiconductor surfaces with reactive gases, which producevolatile products, are used for surface etching. One importantand well-studied example is the etching of Si via exposure ofa Si surface to atomic F generated in fluorocarbon plasma orfrom dissociation of XeF2 molecules. The interaction of theSi surface with F atoms leads to growth of a fluorosilyl SiFx(x < 4) layer on the surface. The layer is composed of SiF,SiF2, and mostly SiF3. In a quasi-steady etching condition,the thickness of the layer (l) is nearly constant with valuesbetween 10 and 30 A. Etching occurs from the surface of thislayer, where volatile SiF4 is formed. The structure of the Sisurface is schematically shown in Fig. 12. The etching rate hasbeen found to be strongly influenced by Si-doping [125–132].The following facts were well-established in those papers: (1)heavily doped n-Si (n+-Si) etches much faster than undoped Siand lightly n- or p-doped Si; (2) lightly doped p-type and n-typeSi have similar etching rates; (3) heavily doped p-Si (p+-Si)exhibits the lowest etching rate.

Winters et al. [125,126] suggested that the Si etchingreactions can be described on the basis of the Cabrera–Motttheory (Fig. 12):

(1) Atomic F adsorbed on a Si surface readily becomesnegatively charged. This is because the electron affinity levelof F lies below the EF of Si when the atom is close tothe surface [126,131,133]. An electron is transferred from Sisubstrate to a surface F-atom in order to equilibrate the energylevels of Si and F. The process proceeds via electron tunnelingthrough the SiFx layer, which is less than 3 nm thick, and thusfeasible for this process to occur.

(2) There is also an electric field similar to what wasmentioned in Section 2.4.1.1. Now, the presence of surfaceanions F− forms an electric field, E0, and bends the surfaceSi bands upwards [126,131]. The concentration of surface F−

(n) and E0 at the SiFx layer are related to Si-doping by

E0 ∝ n = εSiFx

(EF − Ea)surf

el. (26)

Here, εSiFx is the dielectric constant of the SiFx layer, EFthe Fermi level of the Si surface, Ea is the affinity level ofF at the surface, and l is the thickness of the SiFx layer.The expression (26) indicates that an increase in n and E0 isexpected if Si changes from p+-doping to to n+-doping. As wewill discuss below, it is n and E0 which influence the etchingrate significantly.

(3) The reaction rate is determined by F− diffusion throughthe SiFx layer toward the SiFx–Si interface as well as thereaction of F− with Si and SiFx. The large electric field E0 candrive the diffusion of F− from the gas–SiFx interface towardthe SiFx–Si interface; moreover, the high concentration of F−

adsorbed on the surface (n) favors the reaction of F + SiFn →

SiFn+1 (n < 4). Accordingly, n+-Si has a higher etchingrate than lightly doped Si and p+-Si. The doping effect in theetching of semiconductor surfaces can be mostly attributed tothe differences in charge transfer process between surface F andSi as well as the field-assisted anion diffusion process.

2.4.2. Cabrera–Mott theory in solid–solid interface reactionsIn previous papers, we have systematically studied reactions

at solid–solid interfaces using metal/oxide model systems. Thereaction kinetics of metal oxidation and metal encapsulationon oxide surfaces was found to be controlled, besides othermaterials parameters, by EF of both solid phases, and thereaction results can be explained in the framework of theCabrera–Mott theory [13,85,89]. We now discuss the metaloxidation and metal encapsulation reactions on oxide surfaces,respectively.

2.4.2.1. Oxidation of metals on oxide surfaces. The redoxreaction at metal/oxide interfaces was investigated, with an


Fig. 12. (a) Schematic diagram to illustrate the spatial structure of Si surface during etching reactions. The thickness of the interface reaction layer SiFx is l. (b)Schematic diagram to illustrate the band structure at F–Si interfaces. The charge transfer between F adatoms and Si crystals produces electric field E0 at Si surface.After Winters and Haarer [126].

emphasis on the dependence of the reaction kinetics onelectronic structure of both metal and oxide. We studiedoxidation of ultrathin metal overlayers on SrTiO3(100) andTiO2(110) surfaces. The experimentally determined metaloxidation rate was found to closely correlate with EF of boththe metal and oxide phases. We concluded that (1) p-type dopedoxide favors metal oxidation, (2) metals with low surface workfunction have high oxidation rate, and 3) the coupling betweeninterface charge transfer and O2− diffusion at the interface isa critical factor for the reaction kinetics, which is given below[85,89].

To begin with, for the oxidation of identical metals, e.g. Cron different SrTiO3(100) and TiO2(110) crystals, the oxidationrate is dependent on the doping state, i.e., free electron carrierconcentration [e′] of SrTiO3 crystals or TiO2 crystals, but isvirtually independent on the oxygen vacancy concentration[V ••O ] in the crystals. The reaction rate monotonically decreaseswith the increase of the free electron carriers in the oxidecrystals. For a p-type oxide, this means that the system favorsmetal oxidation.

Second, for the oxidation of different metals on identicaloxide crystals, it is found that the reactions show a systematicdependence on the work function of the metal films. Theoxidation rate here increases with decreasing surface workfunction of the metal.

Third, long-range charge transfer after contact betweenmetals and SrTiO3 (or TiO2) crystals produces space chargesat the oxide surfaces and induces band bending. The electronicinteraction is controlled by the relative position of EF of themetal and that of the oxide (See Section 2.2.2). The oxidationof reactive metals on SrTiO3 or TiO2 substrates requires O2−

transport from the oxide bulk to the interface or equivalently aninward diffusion of V ••O . The electric field resulting from spacecharge transfer dominates the O2− diffusion at the interface,and the corresponding interface reactivity and oxidation rates.In case that EF of the metal is greater than EF of the oxide,charge is transferred from the metal to the oxide and negativespace charges form in the oxide. The resulting electric fieldat the interface promotes outward diffusion of O2− and metaloxidation. This is schematically shown in Fig. 13.

2.4.2.2. Encapsulation of metals on oxide surfaces. In aprevious work, the model system Pd/TiO2(110) was usedto evaluate the correlation between metal encapsulation and

Fig. 13. Upper panel: schematic diagram showing the energy bands of ametal/oxide interface in the case of EF (MeI) > EF (MeIIOx ). Lower panel:negative space charges at oxide surface regions and the electric field E0produced by the interfacial charge transfer process, promoting the outwarddiffusion of O2−. The arrangement of EF of metals and oxides as EF (MeI) >

EF (MeIIOx ) favors oxidation reactions. See [85,89].

electronic structure of TiO2 crystals. Some previous results ofthe encapsulation reactions at metal/oxide interfaces were alsocarefully revisited therein. The following points were derivedfrom such investigations [13]:

(1) The studies in the interaction of Pd clusters withdifferently doped TiO2(110) crystals show that n-type dopedTiO2 crystals favor the encapsulation of Pd clusters. Moregenerally, we proved that encapsulation reactions occur only formetals with large work function, e.g. Pt, Pd, and Rh, supportedon n-type doped oxides. The prerequisite to the onset of theencapsulation is that EF of the metal phase should be muchlower than that of the oxide, EF (metal) < EF (oxide).

(2) The encapsulation reactions at metal/TiO2 interfacesinvolve the dominant outward diffusion of interstitial titaniumcations, Tin+i (n = 3 or 4), from the bulk to the TiO2 surface.At relatively low temperature, the outward diffusion of Tin+imust be facilitated by the space charges as shown in Fig. 14.This requires EF (metal) < EF (oxide). In other words, thesystem must have a metal with a large φM and a stronglyn-type doped TiO2. The equilibration of both EF results ina charge transfer from the occupied donor states in TiO2 tothe metal and, consequently, an upward bending of the TiO2bands. Positive space charges form at TiO2 surface such that


Fig. 14. Upper panel: Schematic diagram showing the energy bands of ametal–TiO2 interface in the case of EF (Metal) < EF (TiO2). Lower panel:Positive space charges at TiO2 surface regions and the electric field E0produced by the interfacial charge transfer process, promoting the outwarddiffusion of Ti interstitial ions, Tin+ (n < 4). The arrangement of EF ofmetals and oxides as EF (MeI) < EF (MeIIOx ) favors encapsulation reactions.See [13].

interstitial titanium cations are driven to diffuse toward theinterface (Fig. 14).

2.4.3. Generalized Cabrera–Mott theory in interface reactionsWe have shown that the four reactions at both solid–gas and

solid–solid interfaces demonstrate strong interplay between theinterfacial electronic interaction and the chemical interaction.Such interplay originates from the strong coupling between theinterfacial mass transport and charge transfer. Table 2 lists thetransport processes in the reaction types discussed above. Thecoupling between the mass transport and charge transfer mayrequire different electronic configurations of the two contactingphases in order to favor the interface reaction (see Table 2). Theclassic Cabrera–Mott theory can be generalized to apply to bothsolid–gas and solid–solid interface reactions.

The results introduced above demonstrate that interfacereactions can be uniquely controlled by the relative positionsof EF in both phases before establishing contact. A variationof these positions is directly coupled to changes in the electricfield at the interface and modifies the corresponding reactionrates. These results can be combined in a unique picture bycalculating the surface electric field ES , i.e. electric field atthe interface, as a function of the EF positions in both phasesat the interface. An analytic expression of ES was derived byKingston and Neustadter [134]:

ES = ±

(kT

q L D

)F(φs, φb), (27)

where

F(φs, φb) =√

2[

sinhqφb

kT

(qφb

kT−

qφs

kT

)−

(cosh

qφb

kT− cosh

qφs

kT

)]1/2

(28)

(+ for φs > φb, − for φs < φb).

Table 2Summary of reactions at solid–solid and solid–gas interfaces

Interfacereactions

Contactingphases A ‖ B

Diffusing species& direction

EF arrangementsfavoring reactions

Redoxreaction

Metal ‖ Oxide O2−, A← B EF (A) > EF (B)

Encapsulation Metal ‖ Oxide Tin+i , A← B EF (A) < EF (B)

Surfaceoxidation

Metal ‖ Gas O2−, A← B orMn+

i , A→ BEF (A) > EF (B)

Surfaceetching

Gas ‖ Si F−, A→ B EF (A) < EF (B)

This equation shows how the surface electric field varieswith EF of two contacting interface phases. At a metal/oxideinterface, φb is the energy difference between EF and Ei(intrinsic energy level) of the bulk oxide, which reflects theinfluence of the oxide and can be changed by varying the EFof the oxide (see schematic inset in Fig. 15(a)), φs is the energydifference between EF of the metal and Ei of the oxide, varyingwith the EF of the metal (schematic inset in Fig. 15(b)), andL D is the Debye length [36]. From a plot of ES as a functionof φb and φs one can predict the condition under which metalswill react with oxides. Fig. 15(a) shows ES as a function ofφb for the case where φs = 1 eV, Eg = 3 eV (band gapenergy) and T = 290 K. The upward bending (φb > 1 eV),flat band (φb = 1 eV), and downward bending (φb < 1 eV)will be experienced, respectively, as the oxide’s EF shifts downfrom the VB to the CB. This illustrates that a change of theoxide EF leads to a distinct modification of the direction andmagnitude of the electric field. A plot of ES as a functionof φs with φb = 1 eV is depicted in Fig. 15(b). Here, φsdecreases with increasing φM , which shows that ES changesfrom positive values (downward bending) to negative values(upward bending).

The calculation demonstrates that ES changes with EF ofthe metal (Fig. 15(b)) or the oxide (Fig. 15(a)) monotonicallyand this is one way to predict the thermal stability ofthe metal/oxide interface semi-quantitatively. After contactbetween the metal and the oxide, ES can be established bycharge transfer across the interface, subsequently affectingion transport in the vicinity of the interface as illustratedby Eq. (23). At relatively low temperatures, ES is a criticalparameter that determines transport processes at the interface incases including the initiation of metal oxidation at metal/oxideinterfaces, the surface reactions in reactive Si etching, theoxidation of metal surfaces in gas, and the encapsulationof metal nanoclusters by oxide supports. The systematicdependence of these interface reactions on EF is consistentwith the monotonic change of ES with EF in the materials.

3. Experimental methods

As we have discussed in the introduction, a deeperunderstanding of the nature of metal–oxide interactions relieson the development of better theoretical models and advancesin experimental techniques. In the model systems, ultrathin


Fig. 15. Dependence of the surface electric field ES on φb and φs . (a): φs isconstant (1 eV) for a distinct metal; EF of the oxide shifts from CB towardVB, varying φb from positive to negative (small inset). (b): For a distinct oxide,φb is regarded as constant (1 eV); EF of the metal moves down, resulting indifferent values for φs (small inset). The constant K is a function of the oxide’smaterial parameters and T . For a given oxide and constant temperature, K isconstant.

metal overlayers or metal nanoclusters are grown on well-defined oxide surfaces. These simplified systems significantlyreduce the complexity in real systems in order to studyfundamental interface processes. In this chapter, we presentimportant methods and strategies to prepare these systems. Themost frequently used techniques to characterize metal/oxideinterfaces are subsequently introduced, with an emphasis on thebenefits and limitations of these experiments.

3.1. Preparation of model systems

As stated above, model systems consist of nanostructuredmetal overlayers supported on well-defined oxide surfaces.These systems necessitate the preparation of well-defined oxidesurfaces and well-controlled deposition of metal overlayers,which is discussed in detail.

3.1.1. Oxide surfaces

3.1.1.1. Single crystal oxides. The simplest way to prepareclean and well-ordered single crystal surfaces is by cleaving

in UHV [36]. The surfaces obtained in this way are generallystoichiometric and exhibit a low defect density. There is alimitation to the cleaving of a single crystal surface. First,it may only be used for brittle materials, such as MgO andZnO crystals. Second, cleavage is only possible along certaincrystallographic directions, which results in the formation oflow surface energy surfaces with distinct cleavage planes. Forexample, MgO with the rock salt structure cleaves along the{001} surfaces, and crystals with the wurtzite structure such asZnO cleave along the non-polar faces {1010}. Therefore, only afew surfaces can be studied in this way. MgO(100) is the mostused cleaved surface [135–140]. The cleavage of some polarsurfaces, such as MgO(111), ZnO(0001), and ZnO(0001), hasbeen also reported ([24,76] and references therein).

It is also possible to prepare oxide surfaces by mechanicalcutting and polishing in air. This method is applied to the mostcurrently used oxide crystals, including TiO2, Al2O3, SrTiO3,SiO2, as well as the cleavable MgO and ZnO crystals. Differentorientation surfaces (of a polar or non-polar nature) can beobtained by the process. Ion sputtering is then used to removesurface impurities, which may originate from the polishingprocess, contaminations from exposure to air, segregation ofbulk impurities to the surface, or from previous experimentscarried out on the surface. Surface order, which may have beendisrupted by this harsh treatment, can be achieved by annealingin oxygen or vacuum. Sometimes, cycles of sputtering andannealing are required to get clean and well-ordered surfaces.

When preparing oxide surfaces one should note thefollowing problems.

(1) Surface stoichiometry: For compound surfaces, inparticular oxides, adequate attention must be paid to thechange in surface stoichiometry. On many oxide surfaces, ionsputtering preferentially removes oxygen from the surfaces,resulting in formation of oxygen vacancies and oxygen-deficient surfaces. Also, vacuum heating may cause desorptionof surface oxygen or metal atoms and, therefore, deviation fromthe surface stoichiometry.

(2) Surface reconstruction: During the surface preparationprocesses, oxide surfaces may undergo surface reconstructions,which result from rearrangement of surface atoms, desorptionof surface atoms, or the addition of atoms to the surface fromdeposition. The most commonly observed reconstruction onTiO2(110) surfaces has a (1 × 2) symmetry with a doublingof the periodicity along the [110] direction [28]. SrTiO3(100)has various surface structures, such as (1 × 1), (2 × 1), (2 ×2), c(4 × 2), c(6 × 2), (

√5 ×√

5)R26.6◦, etc [141–146].(√

3 ×√

3)R30◦, (3√

3 × 3√

3)R30◦, and (√

31 ×√

31)R ±9◦ reconstructions have been found on the α-Al2O3(0001)surface [147–150]. ZnO(0001)-(1× 1) can be transformed intoa (2× 2) surface structure [151].

(3) Surface termination: Another important character ofoxide surfaces is the termination. For example, SrTiO3(100)surfaces have two kinds of terminations: SrO- and TiO2-terminated surfaces [152,153] depending on how they areprepared. This is similar to the case of cutting a ZnO crystalperpendicular to the c axis creating a Zn-terminated (0001)surface on one side and an O-terminated (0001) surface on


the other side [154,155]. For the bulk truncated α-Al2O3(0001)surface there are three different terminations: O layertermination, single Al layer termination, and double Al layerstermination [25,156]. Oxide surfaces can be hydroxylated undercertain conditions, e.g. exposure to water, which leads toa fully (or partially) covered hydroxyl surface with surfaceOH-termination. These surface terminations vary sensitivelywith the surface preparation conditions, for example heatingtemperature, heating atmosphere, and surface sputtering.

(4) Bulk defect chemistry: For a variety of oxide crystals,prolonged and/or high temperature heating may cause massexchange between the surface and bulk [157,158]. Theconcentration of defects in the bulk will change during thesurface preparation process, which is particularly important forthe mixed conducting oxides, such as TiO2 and SrTiO3, sinceboth the ionic and electronic conductivities in the bulk solidswill change during the surface treatment process. As we willdiscuss in Section 4, defect chemistry of the oxides plays acritical role in metal–oxide interactions. The variation of theoxide defect chemistry as a function of surface preparationconditions has to be known in detail.

The above discussion suggests that the chemical stoichiom-etry, surface composition, surface atomic structure, and mor-phology sensitively depend on the preparation conditions. It is,therefore, important to be aware of the dependence of surfaceproperties on the treatment conditions, such as the sputteringenergy and sputtering time, heating temperature and time it isapplied, and oxygen partial pressure during preparation of thesurfaces.

3.1.1.2. Supported thin oxide films. Another important methodto prepare model oxide surfaces is to grow thin films supportedon other solids. The main advantage of the method is thatadequate conductivity can be obtained by preparation of anultrathin oxide film on a conducting substrate or by growinga doped thin oxide film while these thin films could mimic thesituation of the bulk oxide materials [159]. On the surfaces ofthe thin oxide films, electron/ion spectroscopy (photoemissionelectron spectroscopy (PES), Auger electron spectroscopy(AES), and ion scattering spectroscopy (ISS)) and scanningtunneling microscopy (STM) measurements become feasible.

Oxide films can be synthesized using various advanced thinfilm growth techniques, in particular MBE. These surfacescan be prepared by direct vapor evaporation of bulk oxidesusing a high energy laser or, alternatively, thermal heating.Most often, metals are evaporated onto clean well-orderedsingle crystal surfaces under molecular O2 atmosphere or usingatomic oxygen [160–162]. In some cases, metal depositionis separated from the oxidation process: metal evaporation isconducted in UHV to form epitaxial metal films, which aresubsequently introduced to an oxygen atmosphere for growthof oxide films. This technique has been applied to the growthof oxides of Ti [163–165], Ce [166], and Fe [167,168]. Forexample, Fig. 16 shows atomically flat iron oxide films grownon Pt(111) using different preparation processes [168]. Thanksto the well-developed MBE technique, a wide range of thin

Fig. 16. Low energy electron diffraction (LEED) patterns, STM images, andstructural schematics of iron oxide films grown on Pt(111). From [168].

oxide films have been deposited on refractory metals or singlecrystal oxides and the oxide surfaces can be well-controlled.

An alternative method to produce thin oxide films isto oxidize a metal or an alloy single crystal surface. Thethickness and structure of the oxide layers depend on oxygenpartial pressure, oxidation temperature, and oxidation time.This method has been successfully used for preparation ofsystems of SiO2/Si(111) [169,170], Al2O3/Al(111) [171,172],and Al2O3/NiAl(110) [26,27]. For example, epitaxial Al2O3-films grown on NiAl(110) show a high degree of crystallinity,very low surface roughness, and good reproducibility.Experimentally, Al2O3/NiAl(110) has been extensively usedfor model oxide surfaces [26,173].

3.1.2. Metal overlayers on oxide surfacesThere are many methods to prepare metal films, such as

physical vapor deposition (PVD), chemical vapor deposition(CVD), and chemical solution deposition (CSD), which havebeen described extensively, see [174]. Here, we give a shortdescription of the most commonly used preparation routesfor ultrathin metal films and metal nanoclusters supported onmodel oxide surfaces.

3.1.2.1. Metal deposition by evaporation. Vapor evaporationunder UHV is by far the most applied method to depositmetal overlayers on oxide surfaces. When a metal is vapordeposited onto an oxide support, various atomic processestake place at the surface. Fig. 17 shows a schematic diagramof the elementary atomic processes and the correspondingcharacteristic energies for metal adatoms on an oxidesurface [175,176].

Atoms which arrive at the surface may reside on the surfacefor a certain length of time before returning to vacuum. Thisis known as the residence time τa . τa is determined by the


Fig. 17. Schematic of elementary steps and characteristic energies taking placefor deposition of a metal onto an oxide surface. Z represents the dimensionperpendicular to the surface. Three basic atomic processes are critical in theinterface formation: the in-plane surface atom diffusion (Z = 0), the out-of-plane surface atom diffusion (Z > 0) including the down-step diffusion andthe up-step diffusion, and the atom interdiffusion (Z < 0). Ea is the adsorptionenergy; Ed is the surface diffusion energy;E1 is the up-step diffusion energy;E−1 is the down-step diffusion energy.

adsorption energy Ea ,

τ−1a = νa exp(−Ea/kT ), (29)

where νa is the atomic vibration frequency. It is expected thathigher surface temperature results in a shorter residence timeand induces re-evaporation (i.e. desorption) of some adatoms.This makes the net sticking coefficient smaller than one. Formany transition metals on oxides at room temperature, thesticking coefficient is near unity ([26] and references therein),indicating that this temperature is not sufficiently large to causedesorption of the particles or the particles are caught at defectsduring their residence on the oxide surface.

Adsorbed atoms can also move on the surface. This diffusionprocess depends on the diffusion activation energy, Ed . Forheterogeneous nucleation, the diffusing atoms may find surfacedefect sites, e.g., steps and vacancy sites, and get trapped thereforming nuclei for subsequent growth process. In the processof homogeneous nucleation, a stable nucleus is formed byaggregation of n (n > 1) adatoms on a regular surface site.After the saturation density of islands has been reached, it isin the stage of island growth. Upon further metal deposition,coalescence of islands happens and finally leads to continuousfilm formation. During the growth and coalescence processes,surface diffusion between the island and the oxide substratelevel becomes possible, as shown in the right side of Fig. 17.Up-step (which occurs for 3D growth) or down-step diffusion(which occurs for 2D growth) of adatoms influences the metalfilm growth mode and the island shape [177,178]. The diffusionis controlled by the activation energies for up-step and down-step diffusion, E1 and E−1 [81].

The energies shown in Fig. 17 vary with the metal/oxidesystem, and thus are strongly dependent on the metal–oxideinteraction. For example, a strong interaction between a metaland an oxide results in large activation energies of Ea , Ed , andE1. Weak desorption, homogeneous nucleation, and layer-likegrowth are thus expected. Studies of the surface processes inthe initial stage of metal deposition allow for the determinationof these activation energies and yield insight into the strengthof the interaction.

3.1.2.2. Size-selected deposition of metal clusters. An elegantmethod to deposit metal nanoclusters on oxide surfaces isthe use of size-selected molecular beams and softlanding ontosubstrates [179–184]. Clusters generated by laser evaporation,magnetron sputtering, or thermal evaporation sources arepassed through an ion optics system, where mass selection andcontrolling of cluster energy are conducted by quadrupole massspectrometer and electrostatic lenses. Subsequently, the mass-selected clusters are deposited under UHV with low kineticenergy onto the oxide surfaces. The softlanding of the clusteronto the substrate avoids the fragmentation or fission of thecluster. The method allows for the preparation of supported,monodispersed metal nanoclusters with a very narrow sizedistribution. Since the bonding between mass-selected clustersand the support should not be strong, thermal stability of themodel systems is quite low. Reliable experiments may only beperformed below room temperature.

3.1.2.3. Chemical vapor deposition. The CVD route to metaloverlayers is carried out by exposing organometallic precursorsto a substrate in vacuum. Heat is then applied to the systems toremove the ligands and convert the metal into the zero-valencestate. This method is sometimes referred as metal-organic vaporphase epitaxy (MOVPE).

A special modification of CVD is the atomic layer deposition(ALD) or atomic layer epitaxy (ALE) [185–190]. In thismethod, film growth takes place in a cyclic manner. One growthcycle normally produces a single monolayer or a fraction ofit. These cycles can be repeated until thicker films are formed.The basic characteristic of an ALD process is the self-limitingnature of the surface reactions. Because of this, ALD enablessimple and accurate control of film growth process at an atomiclayer level. The ALD technique has been widely used todeposit metal oxide films. Nevertheless, Ta [189,190], Ru [191,192], Pt [193], and other transition metals [194] have beensuccessfully synthesized by the ALD process. The thicknessof metal overlayers or the size of metal particles can beprecisely controlled by the number of cycles used. The abilityto control the growth of metal overlayers at the atomic levelmakes ALD a promising method for preparation of metal modelsystems.

3.1.2.4. Solution chemical deposition. Various solution chem-ical methods have been utilized to synthesize metal overlay-ers. Sol–Gel is one of the most used solution methods. In thisprocess, metal precursors are dispersed in aqueous or organicsolutions. The metal containing solutions are spin-coated ordip-coated onto substrate surfaces, which are subsequently sub-jected to further treatments to get pure metal overlayers. Some-times, colloidal chemistry is applied to control the size of col-loids in the precursor solutions and, thus, the size of the de-posited metal particles. Wet chemical impregnation is anothercommon method of metal catalyst preparation in real catalyticsystems. The method is occasionally used for the model sys-tems. Further details about the synthesis of metal on model sup-ports can be found in Ref. [195].


3.1.3. Inverse model systemsMost studies of metal–oxide interactions involve the model

systems of metal overlayers supported on well-defined oxidesurfaces. An alternative route towards the fabrication ofmodel systems is the deposition of nanostructured oxidelayers on well-defined metal surfaces. This process forms aninverse model system, which has been useful to explore themetal–oxide interactions [7,196,197]. In case of electronicinteraction between metals and oxide supports, the ratio of theamount of metal to that of semiconducting oxide support mustbe very small (e.g., very small metal clusters on large oxide)such that the transferred charge to or from the metal is sufficientto affect the electronic structure of the metal, such as themetal Fermi Energy. In an inverse model system, the very largenumber of electrons in the metal support can easily modifythe electronic structure of the supported semiconducting oxidelayer. The electronic interaction between the two phasesbecomes more significant such that the interaction can be easilyprobed [7].

The oxide-on-metal system also allows one to model thepromotion effect of oxides on metal catalysts [198,199] becauseit has been found that the promotion of some catalytic reactionsmay happen near the metal/oxide phase boundary. In inversemodel systems, the perimeter of the oxide islands offers anactive catalytic site, which can be well-studied by surfacescience techniques.

An inverse model system consists of an oxide layersupported on a well-defined metal surface. The growth of oxidefilms has been described in Section 3.1.1.2. Systems of titaniaand vanadia on Pt and Rh [197], ceria on Rh [200–202], vanadiaon Pd(111) [198,203–205], and vanadia on Rh(111) [199,206] have been synthesized and studied in detail. It has beenshown that the metal–oxide interactions lead to the formationof reduced oxide phases with lower oxidation state in thevicinity of the metal interface. The lower oxidation state derivesfrom the chemical interactions with the substrate atoms at theinterface and/or from the geometry-related effects, i.e., latticematching.

3.2. Characterization techniques

3.2.1. Electron-based spectroscopyElectron-based spectroscopy, which includes X-ray photo-

electron spectroscopy (XPS), ultraviolet photoelectron spec-troscopy (UPS), AES, and EELS, are extensively used tostudy metal–oxide interactions. They yield information aboutinterface electronic structures and enable the understandingof charge transfer processes at metal/oxide interfaces. Thesetechniques are also surface-sensitive. As a consequence, smallchanges in surface chemical composition, e.g., surface diffusionand interdiffusion, can be detected such that atom movementsat interfaces are able to be in situ monitored under UHV.

3.2.1.1. Photoemission spectroscopy. XPS has been appliedto investigate the changes in core level binding energy (BE).For supported metal overlayers, in particular metal clusters,

core level BE shifts contain two contributions: the initialstate [207–212] and final state effects [211–217].

The initial state effects arise from the electronic structurefactors which are present in the neutral atoms beforephotoemission. The mechanisms for this state include thefollowing factors [211]:

(i) Charge transfer toward or from the core-ionized atom dueto metal–oxide interaction (interatomic charge transfer).This can be seen in the oxidation of transition metals onoxides which often produces large positive shifts of metalcore level BEs, typically >1.5 eV [18].

(ii) Electric field that arises from the effective charges insupports or metal overlayers, for example, charging effectson insulating supports and electric field effects frominterface space charges.

(iii) Surface core level shifts including contributions of reducedmetal–metal coordination number and rehybridization ofvalence levels (intra-atomic charge transfer). The twocontributions have opposite signs, such that the surfacecore levels typically present small negative BE shifts(≤300 meV) [208,210]. In case that the initial stateeffects are dominant in BE shifts, the shifts can be usedto determine interfacial charge transfer, bonding state ofinterface metal atoms, and interface reactions.

The final state effects arise from the charge rearrangementor relaxation which screens the core holes created after photoe-mission. The screening effects depend on the surrounding en-vironment, i.e., coordination number, ligands from atmosphere,and substrate. In small metal clusters supported on weakly in-teracting substrates, the screening of core holes in the parti-cles is limited compared to the bulk metal. The extra energy ofthe unscreened charge should be proportional to the Coulombenergy e2/2R (R: radius of clusters), which results in a posi-tive shift of the core level BE by a similar amount [214–217].If the final state effects are dominant, the shift of the metalcore level BEs is inversely dependent on the cluster diame-ters. This can be seen in Fig. 18 which shows BE positionof Cr2p spectra as a function of Cr nominal thickness for Cron a SrTiO3(100) grown at room temperature [218]. PositiveBE shifts originating from final state effects can also be foundin other systems with magnitudes up to 1 eV, e.g., 0.7 eV inPt/TiO2(110) [219], 1.0 eV in Ag/TiO2(110) [220], and 0.8 eVin Pt/SrTiO3(100) [221].

In most systems, both the initial state and final state effectscontribute to the metal core level BE shifts. The Augerparameter has been introduced which helps us to separate thetwo effects. If we reference all the BE values with respect toVBM the BE shifts can be given by

1BE = 1E−1R (30)

where 1E is energy shift due to the initial state effects and 1Ris the relaxation energy due to the final state effects. The Augerelectron energy is influenced by the same effects as core levelBE. The shift of the Auger electron kinetic energy (1KE) iswritten as

1KE = −1E+ 31R. (31)


Fig. 18. Cr2p3/2 core level BE as a function of the thickness of Cr overlayerson SrTiO3(100) grown at room temperature. The insert figure is the dependenceof BE shift (1E) on the reciprocal of Cr cluster radius (1/R). From [218].

The Auger parameter α′ is defined as [222–224]:

α′ = KE+ BE. (32)

Based on certain assumptions and the three equations above,the relaxation energy, 1R is simply estimated to be half of theshift of the Auger parameter (1α′) which means that the initialstate contribution is equal to 1BE+ 1

21α′ [225,226]:

1R =121α′; 1E = 1BE+

121α′. (33)

Therefore, the initial state and final state effects can bedistinguished in this way [221,227].

For core level spectra recorded from oxide supports, changesin the spectra can be simply associated with the initialstate effects. The charge transfer between metal adatoms andsubstrate surface atoms causes BE shifts or line shape changesin the corresponding substrate core levels. For example,redox and encapsulation reactions at metal–TiO2 interfacesoften result in reduction of surface Ti. Ti2p spectra becomeasymmetric due to the presence of shoulder peaks from Tin+

(n < 4), e.g. see Fig. 40 in Ref. [28]. On the other hand, spacecharge transfer between the two contacting phases leads to theband bending and the rigid shifts of all substrate core levels.Bending of oxide bands can be monitored as a function of metalcoverage by measuring the rigid shift.

UPS provides information on the VB structure. The VB ofsupported metal clusters varies with the metal cluster size orthe metal coverage. With a decreasing cluster size, the metalVB shows a similar shift to that of the core levels. The stronglocalization of electrons in small clusters causes a narrowing ofthe VB. The shift of oxide VB spectra, however, reflects theband bending at the oxide surface. In particular, the changeof oxide VBM positions after metal adsorption can be usedto determine direction of space charge transfer between thesupport and the metal. UPS can often be used to measure thesurface work function as a function of metal coverage [228].

3.2.1.2. Electron-excited spectroscopy. AES is widely used tomonitor the metal film growth process due to the simplicity

of the technique and the speed at which a measurement canbe taken. Quantitative or qualitative analysis of AES intensity-coverage data allows extracting film growth modes and thestrength of the interaction [177,178,229]. Combining AESdata acquisition with ion beam sputtering, one can conducta depth profile analysis [230]. The depth profiling of ametal/oxide interface gives chemical composition distributionfrom the surface to the bulk, which is powerful in investigationof interface reactions, especially encapsulation or decorationreactions [231].

As an Auger transition includes three ionization processes,AES measures a convolution of three densities of stateswhich makes it hard to get direct information about thesurface electronic structure. Nevertheless, Auger peak shiftsoccasionally allow one to identify the change in chemical states.Bernath et al. [232] studied Ti deposition onto Al2O3(0001)surface by AES. With increasing Ti coverage, the initial Al LVVAuger peak position at 61 eV (from Al3+ in Al2O3) shifts to66 eV, which is characteristic of metallic Al, i.e., metallic AlLVV transition. This shift is due to the reduction of the sapphiresurface by absorbed Ti.

EELS is another electron promoted technique that consistsof inelastic scattering of low energy electrons by surfacespecies. The application of EELS in metal/oxide interfacesmakes use of the interaction of incoming electrons with surfacephonons, bulk and surface plasmons, and interband transitions.The features related to electronic structures of oxide surfacesand metal overlayers are in the range 1–50 eV. The primaryenergies (E0) are in the 100 eV range.

HREELS, a high resolution form of EELS, is performed atlow primary energies, with E0 < 10 eV. It is used to revealthe surface optical phonons or Fuchs–Kliewer modes [233]on oxide surfaces. The intensities of oxide surface phononscan be attenuated by formation of a metallic free carrier layeratop an oxide surface or by charge transfer from the metaloverlayers to the substrate. Additionally, free carriers in themetallic overlayers would give rise to a plasmon-like moderanging from zero to higher energies. The elastic peak canbe broadened asymmetrically toward the positive side [234].Changes in intensity of surface phonons, the full width at halfmaximum (FWHM) of the elastic peak, and loss-peak positionscan give information about electronic interaction betweenmetal overlayers and oxide surfaces. Petrie and Vohs appliedHREELS to study Pt films on ZnO surfaces. The HREELSresults indicate that there are only weak interactions at thePt/ZnO (0001) interface, while charge transfer and Schottkybarrier formation occurs at the Pt/ZnO(0001) interface [235].

3.2.2. Scanning probe techniquesSTM can deliver images of solid surfaces down to atomic

scales, which helps us to identify metal–oxide interactions, suchas the SMSI state, adsorption sites, and metal interaction withsurface defects. The tunneling characteristics of the techniquelimits STM to oxide crystals with relatively small band gaps,e.g., TiO2, SrTiO3, and ZnO. The nucleation of Pd on aTiO2(110) surface was investigated through STM. The imagesshow the presence of dimer and tetramer Pd clusters but


Fig. 19. (a) Low energy ion scattering (LEIS) spectra of (bottom) the clean TiO2(110) surface, (center) after evaporation of 25 monolayers Pt at room temperature,and (top) after the high temperature treatment causing encapsulation. (b)–(e) STM results after the high temperature treatment. (b) Overview (2000 A×2000 A). Mostclusters show hexagonal shape elongated along the substrate [1] direction (type-A). A few square clusters (type-B) are seen. (c) Small-scale image (500 A×500 A),filtered to show the structure of the encapsulation layer on type-A clusters. (d) Atomic-resolution image of an encapsulated “type-A” cluster. (e) Atomic-resolutionimage of a square “type-B” cluster, showing an amorphous overlayer. (f) STS of the different surfaces. From [11].

no single Pd atoms at the nucleation stage. The nucleationof Pd clusters is preferred at step edges. On terraces, Pdatoms adsorb on the five-fold coordinated Ti cations betweentwo bridging oxygen rows [236]. Dulub et al. succeeded inrecording atomically resolved STM images on Pt clustersgrown on a TiO2(110) surface, which were in the SMSI stateafter high temperature annealing in UHV. The encapsulationlayer with a striped zigzag structure on (111)-oriented Ptclusters was proposed to be a slightly oxygen-rich, oxygenterminated TiO1.1(111) double layer [11] (see Fig. 19). Moreoften, STM is used for in situ morphological investigations ofmetal overlayers during growth process, thermal treatment, orgas adsorption. Surface processes, such as nucleation, growth,coalescence, sintering, gas-induced island restructuring, etc.,have been studied for various metals supported on SrTiO3(100)and TiO2(110) surfaces [91,220,237–242].

STM can also be applied to image wide band gap oxidesurfaces provided that a well-ordered ultrathin oxide filmis supported on a conductive support. Metal interactionswith oxide films have been extensively explored by STM,for example, in metal/Al2O3/NiAl(110) systems by Freund’sgroup [26,243] and metal/SiO2/Mo systems by Goodman’sgroup [100,161,244].

Scanning tunneling spectroscopy (STS) provides informa-tion on the local electronic structure of solid surfaces. Both theoccupied and unoccupied energy states are probed by rampingthe bias voltage in positive and negative directions. The magni-tude of tunneling current (I ) at one bias voltage (V ) is related tothe density of states of the sample at that energy. The obtainedI –V curves reflect the electronic character of the tunnelingsites. In supported model systems, STS spectra can be recorded

on an individual metal particle, which allows for the exploringof the electronic structure of individual metal clusters. Valdenet al. acquired STS spectra from various Au clusters supportedon a TiO2(110) surface. They observed a metallic to nonmetal-lic transition of the clusters with a decrease in size [245]. Theappearance of band gap on small clusters is caused by the deple-tion of the density of states near EF and probably due to quan-tum size effects [246]. In Pt/TiO2(110) systems, I –V curveswere acquired from clean Pt clusters and Pt clusters in the SMSIstate. The spectra showed that the electronic structure of clus-ters in the SMSI state is more semiconductor-like due to deco-ration with TiOx layers [11,247] (e.g. Fig. 19(f)).

Atomic force microscopy (AFM) or scanning forcemicroscope (SFM) is not limited by the sample conductivity,but does not present a high spatial resolution as STM.Recently, non-contact AFM and dynamic-mode SFM havebeen developed to yield atomic-resolution images of someoxide surfaces, such as α-Al2O3(0001), TiO2(110), andMgO(100) [147,248,249].

3.2.3. Transmission electron microscopyTransmission electron microscopy (TEM) is one of the most

powerful tools to study metal/oxide interfaces. ConventionalTEM (CTEM) is the most common technique to investigate thenucleation, growth, and coalescence of metals on oxides [250].The direct particle size and shape determination is possiblefor particles larger than 3 nm. The orientation relationship ofmetal islands with support can be acquired when employingtransmission electron diffraction (TED). Polli et al. [251]have characterized Pt growth on SrO-terminated SrTiO3(100)surface via TEM. Fig. 20 gives plan-view and cross-sectional


Fig. 20. Plan-view and cross-section bright field (BF) images recorded from20 nm Pt films supported on SrTiO3 (100) surface covered by 2 ML SrO.From [251].

images recorded from a 20 nm thick Pt film grown onSrTiO3(100). Both images clearly show the presence of twokinds of Pt islands, as shown by the different morphologiesof Pt islands. The pyramids and truncated pyramids are from(100)-oriented Pt grains while the flatter and hexagonal islandspresent (111) orientation. The standard bright field (BF) imagesessentially show 2D projections of the shapes of the particles.Selected zone dark field (DF) and weak beam dark field(WBDF) methods have been developed to obtain the 3D shapesand internal structures of small particles [24,252].

The power of HRTEM in direct imaging of atomicstructures of solids is of great value in investigations ofmetal–oxide interactions. Interface reactions between metalsand oxides, e.g., redox reaction, encapsulation, alloy formation,and interdiffusion, can be directly observed in HRTEM(see Fig. 8). Bernal et al. [94,95,253,254] have shown thatelectron microscopy can make an outstanding contributionto understanding of the effects of SMSI. HRTEM revealedthe actual nature of the metal–support interactions in severalmetal/oxide systems, e.g. noble metal (NM)/CeO2 andNM/TiO2 systems. More often, HRTEM is used to distinguishbetween the reactive and non-reactive metal/oxide interfaces.At reactive interfaces, the newly formed interface phases canbe identified. In the case of a non-reactive interface, whichis atomically sharp, experimental HRTEM results may becompared with the results of image simulation and electronicstructure calculation. This can give information about interfacebonding, translation state, and defects (e.g. [29,255]).

Electrons can be used not only for imaging in TEM butalso for chemical analysis in the analytic electron microscope(AEM). EELS and the related electron energy-loss near-edge structure (ELNES) analysis are based on the inelasticscattering processes due to interactions of transmission electronbeams with samples. EELS can provide chemical analysisat high spatial resolution. AEM offers a spatial resolutionclose to 1 A and an energy resolution close to 0.1 eV [256].That means that electronic structures at atomic sites canbe probed. Experimentally, the spatial difference techniquesas well as EELS line scan across the interfaces are usedto study metal/oxide interfaces. Three spectra need to be

acquired: one from a region containing the interface andtwo from the nearby bulk materials on each side of theinterface. The interface-specific EELS components are obtainedusing the spatial difference method. The difference spectrumrepresenting the EELS of interface atoms gives informationabout oxidation state and chemical environment of these atoms.Such a technique has been successfully applied to derive theinterface electronic structure at Cu/α-Al2O3, Cr/SrTiO3, andNi/SrTiO3 systems [257,258].

3.2.4. Other techniquesISS, including low energy ion scattering (LEIS) and medium

energy ion scattering (MEIS) spectroscopy, may be used forsurface analysis. Among them, LEIS is extremely sensitive tothe composition of the topmost layer. Each surface elementgives a distinct peak in the ion scattering spectrum. Thearea of one peak is proportional to the coverage of thecorresponding element in the top layer. Interface reactions,in particular encapsulation, can be effectively studied by thismethod (e.g., see Fig. 19(a) recorded from Pt/TiO2(110) in theSMSI state) [11,259,260].

X-ray scattering, in particular, the grazing incidence X-ray scattering (GIXS) using synchrotron radiation sources,is another powerful technique for characterization of singlecrystal oxide surfaces and metal/oxide interfaces. Renaud hasgiven a detailed review of this topic [25]. Compared to theelectron-based surface techniques, this X-ray-based surfaceanalytical technique possesses a few unique advantages. First, itis not subject to charging effects such that the surface analysiswill not be limited by the insulating character of most oxidecrystals. Secondly, this technique can be applied in situ, e.g., onan oxide surface at high temperature, at high pressure, orduring the growth of metal overlayers (e.g. [261]). Furthermore,X-ray interacts weakly with solids and, thus, quantitativeanalysis based on a single scattering calculation is possible.Therefore, the surface X-ray scattering technique allows forthe experimental determination of the atomic structure of oxideinterfaces with high accuracy (e.g. [173,262]).

Electron diffraction techniques including low electronenergy diffraction (LEED) and reflection high energy electrondiffraction (RHEED) are applied to investigate the surfacestructure. On oxide surfaces, either LEED or RHEED deliversquantitative information about surface reconstruction andrelaxation. Upon metal deposition, the interfacial structure,such as orientation of metal overlayers with substrate, newinterface phases formed by reactions, and metal phasetransformation, can be effectively studied by the surfaceelectron diffraction methods [91,263,264].

Molecular beam experiments are frequently applied to studythe metal–oxide interactions [243]. In the experiments, variousmolecular probes were used to test the surface reactivity. Forexample, metal catalysts in SMSI states present much lowerH2 and CO chemisorption capacity [14]. The charging state ofmetal clusters supported on oxide surface can be detected byfrequency shifts of the intramolecular vibration of adsorbed COmeasured by infrared (IR) or EELS [265,266].


Fig. 21. (a) Structural model of the rutile TiO2(110)-(1× 1) surface. Large white balls represent oxygen, and small black balls represent titanium atoms. 6c-Ti: six-fold coordinated Ti atoms; 5c-Ti: five-fold coordinated Ti atoms; 2c-O: two-fold coordinated protruding O atoms (bridging O); 3c-O: fully three-fold coordinated Oatoms. (b) STM image (200 A×200 A) showing surface defects. It is accepted that the contrast in empty-states STM images of TiO2(110) is dominated by electroniceffects. The bright rows correspond to the five-fold coordinated Ti atoms, while the dark troughs are from the bridging oxygen rows. Bright spots centered on darkrows (defect A) are from missing single oxygen vacancies. The dark spots on bright rows (defect B) are attributed to subsurface oxygen vacancies. From [270].

4. Interaction of metals with mixed conducting oxides

In electronic solids, atomic constituents exhibit negligiblemobility while electronic free carriers can be easily excited andtransferred. In ionic solids, e.g. solid electrolytes, at least oneionic component is highly mobile whereas another (at least)ionic component is virtually immobile. For mixed conductingsolids, both electrons (holes) and ion(s) are mobile such thatthese solids exhibit electronic and ionic conductivities. Thesematerials have the ability to transform chemical energy orinformation into electric energy or information (and vice versa),and therefore possess many important applications [267].Oxides, such as TiO2, SrTiO3, CeO2, ZrO2, SnO, and ZnO,are typical mixed conductors. Inside these oxides, mobileelectronic and ionic defects can be generated according todistinct defect reactions. The defect concentration is a functionof the corresponding mass-action laws, temperature, oxygenpartial pressure, and extrinsic dopants (e.g., see [268,269]).

Due to the mobility of both electronic and ionic defectsin a mixed conducting oxide, the interaction of a metal withthe oxide depends on both ionic and electronic conductivities.Here, the electronic interaction is mainly determined by theelectronic character of the oxide. As discussed in Section 2.2.2,the interfacial charge transfer between a metal and a mixedconducting oxide can be treated similar to the case ofa metal–semiconductor junction. In this system, the oxidebehaves like a semiconductor. On the other hand, the chemicalinteraction involves atom diffusion at interfaces such that thetransport of ionic defects in the oxide should be considered. Insuch a case, the oxide is regarded as an ionic solid. Interactionsbetween metals and mixed conducting oxides (TiO2 andSrTiO3) are reviewed here. We show how the interactions varywith oxide surface properties, defect chemistry of oxides, andmetal overlayers.

4.1. Metals on TiO2

There is a large body of literature concerning the surface andinterface studies of metal–TiO2 systems, principal among theman article by Diebold which gives an excellent overview of thesurface science of TiO2 [28]. Here, we focus on the interactionsof metals on well-defined rutile TiO2(110) surfaces under UHVconditions. TiO2 surface defects, TiO2 bulk defects, and metaloverlayers are found to be the most important factors whendetermining the metal–TiO2 interactions. These factors are nowdetailed below.

4.1.1. TiO2 surfacesThe interactions of metals with TiO2(110) surfaces at the

initial stage of metal growth (mainly at submonolayer coverage)is discussed in this section. We show that surface defects ofa TiO2(110) surface play a critical role in adsorption andnucleation of metal adatoms. As expected, the metal interactionon a defect-free surface is quite different than that on a defectivesurface. It is therefore essential to discuss both cases separately.

4.1.1.1. Defect-free surfaces. Bulk truncated TiO2(110) sur-faces present two kinds of termination: the polar surface termi-nated with either Ti or O, and the non-polar surface containingboth under-coordinated Ti and O. Many experiments have con-firmed that only the non-polar surface is stable [28]. On this sur-face there exist six-fold coordinated Ti atoms (6c-Ti), five-foldcoordinated Ti atoms (5c-Ti), two-fold coordinated protrudingO atoms (bridging O) (2c-O), and fully three-fold coordinatedO atoms (3c-O). These surface atoms are schematically shownin Fig. 21.

In the case of metal adsorption on the TiO2(110)-(1 × 1)

surface, the adatoms may register at different sites [271]: atthe monocoordinated site, atop of 5c-Ti or protruding 2c-O;at the dicoordinated site, bridging two 2c-O; at the three-fold


coordinated site, “between” site formed by two 2c-O and one3c-O or “adjacent” site by one 2c-O and two 3c-O; at thetetracoordinated site, the hollow site surrounded by two 5c-Tiand two 3c-O. Generally, one adsorption site is preferred overthe others. Such a preferable geometric arrangement results inthe strongest interface bonding between the metal and substrateatoms. As discussed in Section 2.2.1.4, the bonding strengthis strongly correlated with electronegativity of two bondedelements. As a rule of thumb, Pauling electronegativity (X M ) ofmetals can be used to find the local interaction of metal adatomswith the TiO2(110)-(1× 1) surfaces. As detailed below, metalswith X M < 1.9 form interface bonds with surface O whereasmetals with X M > 1.9 tend to register on surface Ti cations.

For metals with X M < 1.9 including alkali metals, earlytransition metals, and Al, bonding of these metals with theTiO2 surfaces mainly occurs via surface O. There is an indirectcharge transfer from the metal adatoms to surface Ti ions whichis mediated by the surface O. The adatoms try to maximize theirO coordination numbers and, thus, should sit on the three-foldO coordination sites or sites bridging two 2c-O.

For example, in the case of Na-adsorption on TiO2(110)Onishi et al. [272] proposed a “Na2O-dimer” model in whicha Na atom bonds to two bridging O atoms. Lagarde et al. [273]applied extended X-ray absorption fine structure (EXAFS) todetermine the adsorption site of Na. They, however, foundthat for Na coverage ranging between 0.25 and 0.5 ML themetal is “between” a three-fold coordinated site where it isbonded to two bridging O atoms and one in-plane O atom.The same geometric configuration of Na adatoms on TiO2(110)has been experimentally observed by Murray et al. [274] andNerlov et al. [275]. In order to determine the adsorption siteof K on the TiO2(110)-(1 × 1) surface, surface extended X-ray absorption fine structure (SEXAFS) combined with STMand non-contact AFM were used for the investigation of 0.15ML K deposited on the TiO2 surface [276]. The experimentaldata suggest that K also occupies a three-fold coordinated sitein which K bonds to two bridging O atoms and one in-plane Oatom. A similar three-fold hollow site seems plausible for Ca-adsorption on TiO2(110) surface [277]. Ab initio calculationsand molecular dynamics (MD) simulations further confirmedthat the preferred adsorption positions for alkalies on TiO2(110)are the three-fold O coordination sites. Both “between” and“adjacent” sites should be considered and the adsorption sitesmay change with coverage [278,279].

Diebold et al. studied electronic structure of ultrathin Fefilms on TiO2(110) with soft X-ray photoelectron spectroscopy(SXPS) and resonance photoemission [18]. The lack ofhybridization of the “Fe” and “Ti” states suggests that thebonding of Fe occurs predominantly via surface O and that atopbonding above a Ti atom appears unlikely. Atomic-resolutionSTM images recorded on a V/TiO2(110) surface is proof that Vadatoms adsorb in the three-fold hollow sites. On these systems,the V atom bonds to two bridging oxygen atoms and one basaloxygen atom [280].

The formation of metal–O interface bonds facilitates thelocal charge transfer from metal to TiO2. Since surface O atomsare in their closed-shell configuration (O2−) the additional

electrons transferred from the metal adatoms will populateempty orbitals of Ti atoms. The transferred charge localizes atthe 3d orbital of the five-fold coordinated Ti ion resulting inthe appearance of occupied Ti3d-derived gap states [71]. Thiselectronic interaction can be observed by XPS measurements,where shoulder peaks at low BE in XPS Ti2p spectra appeardue to the electron transfer from metal to surface Ti. The resultshave been found in many metal/TiO2 interfaces [281–286].

Metals with X M > 1.9, e.g. noble metals, try to bond withsurface Ti cations in most cases. There is a very weak, if any,charge transfer from the substrate to the metal adatoms whichmay be facilitated via formation of metal–Ti bonding.

Au-adsorption on the TiO2(110)-(1 × 1) surface has beenextensively studied. Theoretical investigations show that Auatom may bind to either a 5c-Ti atom on the basal planeor atop a bridging O atom. The most favorable site dependson Au coverage and calculation methods. Nevertheless, thetwo sites have quite similar BEs and are comparable to eachother [271,287,288]. The electronic interaction between Au andTiO2 is very weak. For example, MIGS and metal polarizationeffects are found to contribute to the interaction [271,288].Experimentally, STM images show that many of small Auclusters are located on top of the bright rows that arise fromsurface five-fold coordinated Ti atoms. Thus, it appears that Auatoms nucleate on top of the Ti cations [240]. Another veryimpressive STM result from Tong et al. also confirms that smallAu clusters, which are deposited on TiO2(110)-(1× 1) by size-selected Au2 ∼ Au4 beams, are located atop of 5c-Ti [289].XPS and LEIS studies in Ag deposition on stoichiometricTiO2(110) surface lead to a similar conclusion that Ag atomsare preferentially bonded to Ti rather than O [290].

Horsley has conducted a molecular orbital study of Ptinteraction with TiO2. The calculations favor a model in whichthe nearest neighbor of Pt atom is the Ti ion and covalentmixing between Ti3d and Pt5d orbitals occurs. There is someionic contribution to the Pt–Ti bond and the Pt atom isnegatively charged by 0.11 e/atom [291]. A later ab initiostudy by Xu et al. used the same geometric structure andconcluded that Pt is also negatively charged [292]. Adsorptionof Pt on a 5c-Ti atom on TiO2(110) has been directly imagedby STM [293]. XPS investigations of Pt-adsorption on an idealTiO2 surface indicate that the preferable Pt-adsorption site ison top of 5c-Ti [294,295]. For Pd on TiO2(110), STM imagingshows that Pd atoms also adsorb on the 5c-Ti ions [236,240].

4.1.1.2. Defective surfaces. Various defects can be introducedon the ideal TiO2(110)-(1×1) surface and they are very criticalfor metal–TiO2 interactions. Oxygen vacancy and hydroxylgroup are the most important point defects on TiO2(110)surface. Surface oxygen vacancies can be created by UHVheating, ion sputtering, and irradiation by electron beams orUV light. The oxygen vacancies introduce Ti3d-derived gapstates ∼0.8 eV below the CBM and partially reduce the surfaceTi4+ ions. In STM images, these vacancies appear as brightfeatures centered on dark rows and connect two neighboringbright rows, which suggest that they are from the bridging Orows (see Fig. 21). The concentration of surface O vacancies


can be well-controlled by the surface treatments [28,296].Dissociation of water on many defective oxide surfaces can beobserved [297]. On defective TiO2(110) surfaces, it has beenfound that water reacts very efficiently with oxygen vacanciesin a wide temperature range, which results in formation ofhydroxyl groups on the surface. The oxygen vacancy is imagedas a faint bright spot in the dark rows while the feature fromthe OH group appears brighter in the STM images than thevacancy [298–300]. Steps are always present on single crystalsurfaces due to a miscut of the crystal. On TiO2(110) surfaces,step edges run predominantly parallel to the 〈001〉 and 〈111〉-type directions. At the step edges, under-coordinated O andTi atoms can be found [270]. Finally, the (1 × 1) surfacemay be subjected to reconstruction during heating the surfacein UHV or under oxidizing conditions. The most commonlyobserved reconstruction is the (1× 2) structure. Both missing-row and added-row models were proposed to explain the (1×2)

reconstructed surface structure [301–304]. Presently, the added-row reconstruction, in particular “added Ti2O3 rows” model, isgenerally considered to be a viable model [28].

Again, as discussed in Section 4.1.1.1, metal adsorption onTiO2(110) relies on X M of the metal. Reactive metals withX M < 1.9 prefer to adsorb onto surface O via formation ofmetal–O bonds and the interface bonding strength is strong.There is no tendency for the metals to decorate step edgesor oxygen vacancies because of an oxygen deficiency inthese sites. The metal atoms cover the surface evenly withoutpreferable decoration at any defective sites. It is also consistentwith the large adsorption energy and small diffusion lengthwhich are expected for the highly reactive metals on thestoichiometric TiO2 surface. Thornton and coworkers havesystematically studied the adsorption of Na [274], K [276],and Ca [277] on TiO2(110) surfaces with scanning probemicroscopy (SPM). The metal-induced surface structures areobserved to distribute evenly on the surfaces. Al and V interactstrongly with the TiO2(110) surface. No obvious preferentialnucleation of Al [305] and V [280,306] clusters at the surfacedefect sites has been observed.

For the noble metals with X M > 1.9, the surface defectshave a significant effect on the metal adsorption on TiO2(110).The role of steps, reconstructions, and oxygen vacancies inthe surface metal adsorption, which are all quite important, isdiscussed.

Steps: Steps behave as trapping sites for noble metaladsorption such that preferential nucleation of noble metalclusters at the edge sites is often observed. Studies on Ag [220,239,307], Cu [238], Au [245,308–311], and Pd [236,312]growth on the stoichiometric TiO2(110) surfaces demonstratethat most of metal clusters predominantly cover steps. Theweak interaction of the metals with the stoichiometric surfacesites enables that the adatoms have enough mobility on thesurfaces and they can diffuse to the step edges which containunder-coordinated Ti atoms [248,270] where they act as sinksfor metal formation.

Reconstructions: On the TiO2(110)-(1 × 2) surface, defectsites, such as added rows and cross-link structures, act as trapsor diffusion barriers for metal atoms. This causes the mobility

of the noble metal atoms on the reconstructed surface to bestrongly reduced [311–313]. The metal particles show highdispersion on terraces and, therefore, have small size and highisland density. It has been found that the Ag cluster densityon the reconstructed surface is larger than that on the (1 × 1)

surface [220,313]. A similar phenomenon was observed forCu [314], Au [311,315], and Pd [312] on the (1 × 2) surface.Berko et al. found that on the reconstructed surfaces Ir, Rh, andPt particles evenly covered both terraces and step edges and thesteps are not decorated preferentially [241,316,317].

Oxygen vacancies: On both (1 × 1) and (1 × 2)TiO2(110)surfaces, oxygen vacancies exert a strong influence on theinteraction of metal adatoms with the rutile surfaces. Tong et al.confirmed that Ag clusters tend to nucleate at oxygen vacanciesdue to the strong bonding of Ag atoms with the vacancies.Their calculations show that the bonding of an Ag atom toan oxygen vacancy site is stronger by about 0.5 eV than thatof any other site on the terrace [313]. High resolution STMimages reveal that Au clusters nucleate primarily at bridgingoxygen vacancies at 130 K. Calculations show that for a singleAu atom the most stable configuration is adsorption in anoxygen vacancy site. The Au bond on the stoichiometric surfaceoriginates from bond polarization while the Au vacancy bondis covalent with very little charge transfer [309]. The result isfurther confirmed by recent calculations [318]. They found thatAu remains neutral on regular sites of TiO2(110) surface butbecomes negatively charged (Au−) when trapped at an oxygenvacancy.

Surface hydroxyl groups: The recent research results suggestthat surface oxygen vacancies react quickly with residualwater even under UHV conditions such that the TiO2(110)surfaces are mostly hydroxylated [298]. Besenbacher andcoworkers found that the interaction between Au clusters andthe hydroxylated TiO2(110) surface is quite weak while strongAu–TiO2 interaction occurs for the reduced TiO2(110) surfacewith oxygen vacancies and the oxidized TiO2(110) surface withadsorbed oxygen atoms [319].

4.1.2. TiO2 bulk defect chemistryChemical reactions of metal overlayers on TiO2 surfaces

involve the diffusion of ionic defects between the bulk andsurface of the oxide. Therefore, the metal–TiO2 interfacereactions are closely related to TiO2 bulk defects. We show howthe metal–TiO2 interface reactions including metal oxidationand metal encapsulation depend on the bulk defect chemistryof TiO2 crystals.

4.1.2.1. Defect chemistry of TiO2. The defect chemistry andelectronic structure of TiO2 solids have been extensivelystudied (see [28] and references therein). It is now well-accepted that the main ionic defects in TiO2 are interstitialtitanium ions (Ti••••i or Ti•••i ) and oxygen vacancies (V ••O orV •O) (Kroger–Vink notation). The diffusion mechanism of Tiand O is that Ti diffuses in the solids as an interstitial atom andO is transported via vacancy diffusion [320–322].


Typical defect reactions in the solids occur according to thefollowing equations:

OXO ↔ V ••O + 2e′ + 1/2O2 (34)

TiXTi + 2OX

O ↔ Ti••••i + 4e′ + O2 (35)

and

NbXTi ↔ Nb•Ti + e′, for Nb donors on Ti sites. (36)

Here, OXO and TiX

Ti are lattice O and Ti sites. Temperature and O2partial pressure (PO2 ) determine the concentration of the ionicand electronic defects in TiO2. Furthermore, extrinsic dopantssuch as Nb and Fe also affect the defect equilibria. Thus, theionic and electronic defects in TiO2 can be controlled by hightemperature heating, extrinsic doping, and ion sputtering [296,323–325]. We show how these treatments change the TiO2defects and electronic structure.

High temperature heating: Heating TiO2 in an oxygen-deficient atmosphere (e.g. UHV and H2) is known as hightemperature reduction (HTR). As shown in Eqs. (34) and (35),this leads to the loss of O and contributes to the formationof Ti and O defects, acting as intrinsic donors. Stoichiometricrutile TiO2 is an insulator with band gap energy of 3.05 eV atroom temperature. In non-stoichiometric TiO2−x (0 < x < 1),defect states (essentially Ti3+), which originate from oxygenvacancies, are present in the band gap at ≥2.3 eV aboveVBM [296]. If the density of the defect states is sufficientlyhigh they can develop into a shallow CB, which produces freeelectrons in the CB and thereby shifts EF towards the CB edge.

Extrinsic doping: Nb-doping in the parts-per-thousand rangegives rise to the formation of shallow donor states 0.02 –0.03 eVbelow the CBM. The donor states cause n-type semiconductingbehavior of TiO2 [323–325].

Ion sputtering: Ar+ sputtering can preferentially remove Oatoms from the solid surfaces [296] by introducing O vacanciesin the topmost surface layer (e.g., Ar+ energy with 200 eV).However, heavy ion sputtering (e.g., Ar+ energy with 3 keV)produces O vacancies down to subsurface regions and maychange the stoichiometry of the bulk region near to surface [13,92,326].

4.1.2.2. Effect of TiO2 bulk defect on metal oxidation on TiO2.Metal oxidation is the most common reaction at metal–TiO2interfaces. We show that the TiO2 bulk defect chemistry has astrong effect on the metal oxidation reaction. Recently, Fu et al.studied the oxidation of Cr overlayers (6 A thickness) on threedifferently prepared TiO2(110) crystals [89]. The TiO2 defectchemistry was controlled by HTR, doping, and sputtering.

In Case (A), TiO2 is doped with Nb (donor, 0.01 at.%) whichwas Ar+-sputtered (200 eV, 10 min) and UHV annealed at800 ◦C for 1 h. Case (B) is undoped TiO2(110) which was Ar+-sputtered (200 eV, 10 min) and UHV annealed at 800 ◦C for 1h. Case C, which is also undoped TiO2(110), was oxidized inair (800 ◦C, 6 h) and lightly Ar+-sputtered (200 eV, 1 min). Thefree electron concentration ([e′]) in Case A is higher than thatfor Case B and C. The oxidized TiO2 (Case C) has the lowest[e′] because annealing in an oxygen-rich atmosphere tends to

remove oxygen vacancies, i.e. the intrinsic donors, in the crystal(Eq. (34)). Accordingly, we can conclude that [e′] in TiO2decreases from the Case A to Case C. It should be mentionedthat introduction of donors, e.g. Nb, into TiO2 may decrease[V ••O ] according to a compensation reaction [327]. Among thethree crystals, Case B possesses the largest [V ••O ].

To investigate the bulk defect effect on metal oxidation, Croverlayers were grown on the different TiO2 crystals and allwere stepwise heated in UHV to 600 ◦C. In situ XPS wasperformed to monitor the oxidation process after heating. Thefollowing details the results. For Case A, Cr oxidation beganat 370 ◦C, but metallic Cr was still observed on the surface at550 ◦C. With Case B, Cr was almost fully oxidized at 460 ◦C,strikingly different than Case A. Case C was already lightlyoxidized at 280 ◦C and nearly converted to Cr oxide at 370 ◦C.These results indicate that the reaction rate of Cr oxidationon TiO2 is inversely proportional to [e′] in TiO2 rather thancorrelated with

[V ••O

]in TiO2. Lower [e′] in TiO2 favors faster

oxidation of Cr overlayers.Domenichini et al. have systematically studied the role of

the TiO2 bulk stoichiometry in the interface reaction betweenMo and TiO2(110) [92]. In that work, three different TiO2(110)crystals were prepared. The first (#1 TiO2) was annealed at925 K for 72 h in air. This crystal then presents the bulk andsurface stoichiometry of TiO2. From the first crystal, a secondone (#2 TiO2) was further prepared by Ar+ sputtering (20 min,3 keV). This procedure produces defects in the upmost severalatomic layers, as discussed above. Thus, the #2 TiO2 crystal isstoichiometric in the bulk but its surface is reduced. The thirdone (#3 TiO2) was subjected to cycles of Ar+ sputtering (20min, 3 keV) and UHV annealing (873 K, 30 min). The #3 TiO2crystal is both bulk reduced and slightly surface reduced. Again,one can conclude that [e′] in #3 TiO2 crystal is higher than thatin either #1 or #2 TiO2 crystals.

To compare these substrates, a varying amount ofMo (1–3 eq ML) was deposited on surface #1 through #3,which were then stepwise annealed in UHV up to about 750 ◦C.XPS investigations show that the deposited Mo becomesoxidized for temperatures greater than 400 ◦C on the #1 and#2 TiO2 crystals, but Mo/#3 TiO2 was still metallic at 750 ◦C(see Fig. 22). Based on these results, three conclusions canbe made: (1) Mo oxidation is not influenced by the substratesurface defects; (2) Mo oxidation relies much more on the bulkstoichiometry; and (3) Faster Mo oxidation occurs on a TiO2crystal with a lower concentration of electron defects, [e′].

The reaction results taken from both Cr/TiO2(110) andMo/TiO2(110) interfaces are quite consistent with each other,which suggest that metal on a TiO2 crystal with a high [e′]is more resistant against the oxidation reaction. These resultsfurther support the concept suggested in Section 2.4 andFig. 13. For n-type doped TiO2 (Case A in [89]) or the TiO2crystal with bulk non-stoichiometry (#3 TiO2 in [92]), EF isclose to the CBM. In case of contact between Cr (or Mo) and theTiO2 crystals, EF (Cr or Mo) < EF (TiO2) results in electrontransfer from TiO2 to the metal. Positive space charges format the TiO2 surfaces which hinder O2− outward diffusion tothe interface. Therefore, metal oxidation is kinetically limited


Table 3Treatments for the six TiO2 crystals

Crystals Surface preparation Color — n Encapsulation

#A, undoped Standard procedure Pale grey — Low No#B, undoped Standard procedure plus UHV heating at 800 ◦C for 9 h Blue — High Yes#C, doped Standard procedure Dark blue — High Yes#D, doped Standard procedure plus UHV heating at 800 ◦C for 9 h Dark blue — High Yes#E, undoped Standard procedure plus light sputtering (200 eV, 1 min) Pale grey — Low No#F, undoped Standard procedure plus heavy sputtering (1000 eV, 10 min) Light blue — High Yes

A qualitative evaluation of the electron density (n) in CB of TiO2 was performed considering the defect chemistry of TiO2 and the crystal’s color after treatment.Theintensity of this blue color corresponds to different n in the CB of TiO2: a darker color, a higher n [321,322]. Results on the Pd encapsulation are also included.After [13].

Fig. 22. Evolution of the BE of the Mo3d5/2 peak main component duringannealing under UHV for 3.3 eqML of molybdenum deposited on bulkand surface stoichiometric TiO2, #1 TiO2, (�); 3.3 eqML of molybdenumdeposited on stoichiometric bulk and surface reduced TiO2, #2 TiO2, (�); 3.05eqML of molybdenum deposited on bulk and surface slightly reduced TiO2, #3TiO2, (•). From [92].

even at high temperature at Cr/TiO2 (Case A) and Mo/#3 TiO2interfaces.

4.1.2.3. Effect of TiO2 bulk defect on metal encapsulation onTiO2. The encapsulation reaction is another important reactionoccurring at metal–TiO2 systems. Like the metal oxidation onTiO2, metal encapsulation shows a strong dependence on thebulk defect of the TiO2 substrate.

Fu et al. have systematically investigated Pd interactionon six different TiO2(110) crystals [13]. Table 3 presents thedata of these experiments. Generally speaking, electron-rich,or n-type doped, crystals were obtained by HTR in UHV, Nb-doping, or by strongly sputtering the surface. Pd clusters withnominal coverage of∼1.5 nm were deposited on the differentlytreated TiO2 crystals at 200 ◦C and all were stepwise heatedin UHV to 720 ◦C. Characterization by XPS, AES, AFM, andhigh resolution Rutherford backscattering spectroscopy (RBS)confirms that the encapsulation of Pd clusters on TiO2 mainlydepends on [e′] in the TiO2 crystals. Pd encapsulation onlyoccurs on the electron-rich or n-type doped TiO2 crystals.

Bourgeois et al. reached a similar conclusion, which wasderived from XPS studies of the Ni/TiO2(100) system [328].They found that the covering-up of nickel by a thin layer ofTiO2 occurs more easily in the case of bulk non-stoichiometricTiO2−x crystals. The TiO2−x substrates were prepared by HTR

in Ar–H2 mixture at 1000 K for 10 h and they are n-typedoped due to the introduction of oxygen vacancies. The n-type conductivity of the TiO2−x substrates contributes to thecovering-up reaction of Ni atop.

Furthermore, Berko et al. studied the thermal behaviorsof Rh supported on three TiO2(110) surfaces [326]. Noencapsulation of Rh particles was observed on a well-orderedTiO2(110) surface and a slightly Ar+-sputtered TiO2(110)surface. On the other hand, clear evidence is found forencapsulation of Rh crystallites supported on a stronglyAr+-sputtered surface. The heavy ion sputtering producessubsurface non-stoichiometry which leads to the presence ofan electron-rich near-surface region that contributes to theencapsulation of Rh particles.

Experiments from other groups also show that encapsulationof some noble metals, e.g. Pt, Pd, and Rh, takes place onstrongly reduced TiO2 crystals which are subjected to HTRbut does not occur on bulk stoichiometric crystals [8,11,231,259,329,330]. In these cases, it known that the HTR treatmentleads to n-type doping in TiO2 crystals, which favors theencapsulation reactions.

Based on these results, we come to a general conclusion thatmetal encapsulation reactions are only possible on TiO2 crystalsthat are electron-rich in the bulk. The n-type conductivity ofTiO2 can be obtained by HTR, extrinsic doping via donors,or heavy ion sputtering. These results are consistent withthe concept described in Section 2.4 and Fig. 14. Since theencapsulation process requires the amplified outward diffusionof Tin+i to TiO2 surfaces and an electronic configuration ofEF (TiO2) > EF (metal), it can be understood that n-type TiO2crystals and noble metals with large work function (e.g. Pt, Pd,and Rh) are the prerequisites for the encapsulation reaction.Interface contact between the metals and the n-type TiO2 resultsin formation of positive space charges at the TiO2 surfaceswhich drives the outward diffusion of Ti cations and, thus, theencapsulation reaction (Fig. 14).

4.1.3. Metal–TiO2 interactionsIn this section, we consider the effect of metal overlayers

on metal–TiO2 interactions, in which both the thermodynamicand kinetic factors is discussed. The interaction data of variousmetals on TiO2(110) surfaces are to be revisited. Four differentmetal interactions with the TiO2(110) surface can be observed.The metal work function is the critical parameter whichinfluences the metal–TiO2 interactions.


Fig. 23. Relationship between surface energy γ and work function φ ofdifferent transition metals. If available, φ of fcc(111) and bcc(100) surfacesare displayed because of the orientation of the metals on TiO2(110). Otherwiseφ values are taken from polycrystalline samples [331]. The surface energydata were taken from calculations of Mezey and Giber [332]; these data arequalitatively consistent with experimental data [333,334], providing larger γ

values for Pt, Pd, Rh, and Ir and smaller γ values for Au, Ag, and Cu. In regionI, encapsulation is expected (φ > 5.3 eV and γ > 2 J m−2) while in region II(φ < 4.7 eV) oxidation of metals on TiO2 is possible. After [13].

4.1.3.1. Thermodynamic aspect. Using thermodynamics, it ispossible to predict if a reaction between a metal and TiO2 isfeasible. A good example for a typical reaction at metal–TiO2interfaces is the redox reaction. The corresponding reactionequation can be written as follows:

M+ TiO2 → MOx + TiO2−x . (37)

As discussed in Section 2.3.2, the heats of formation of oxides(1Ho

f ) can be simply used to find if the reaction occurs.Diebold has made a general conclusion that redox reactionsare favored by the condition of 1Ho

f < −250 kJ/mol O. Thissimple rule has been confirmed by many redox reaction resultsat metal/TiO2 interfaces [28].

The thermodynamic criteria for encapsulation at metal–TiO2interfaces mainly rely on the metal surface energy, γ . It hasbeen shown that the metals with γ > 2 J m−2 are likely to

experience encapsulation reactions [13,111]. The dependenceof the encapsulation reactions on the surface energy ofmetals was illustrated by Fig. 23 which was reproducedfrom [13].

4.1.3.2. Kinetic aspect. Reaction kinetics at metal–TiO2interfaces is determined by mass transport processes at TiO2surface regions. In TiO2, the mass transport involves thediffusion of defects of Tin+i (n ≤ 4) and/or V x+

O (x =1 or 2). Depending on the interface reaction, one has toconsider the different defect diffusion processes. For example,the oxidation of metals on TiO2 is controlled by the outwarddiffusion of oxygen anions (O2−), i.e., the diffusion of oxygenvacancies (Vx+

O ) in the reverse direction, in the vicinityof the interfaces. The outward diffusion of O2− has beenobserved at many metal/TiO2 interfaces, e.g., Al/TiO2 [335],V/TiO2 [336], Cr/TiO2 [284], and Mo/TiO2 [337]. Theencapsulation of metals on TiO2 is limited by the outwarddiffusion of titanium interstitial cations Tin+i from TiO2 bulkto the interface [13]. Many previous results have establishedthat Tin+i ions possess a high diffusivity in TiO2 at elevatedtemperatures (e.g. 500–800 K) and the Ti ion diffusion iscritical to other TiO2 surface processes including oxygen-induced surface restructuring and bulk-assisted reoxidation onTiO2 surfaces [301,329,338–341].

Many experiments have shown that the metal–TiO2interface reactions are often thermally limited at relativelylow temperatures [18,20,29,90]. Therefore, it is importantto consider reaction kinetics. The generalized Cabrera–Motttheory suggests that the diffusion of ionic defects at TiO2surfaces should be strongly coupled with the charge transfer atmetal–TiO2 interfaces. We now discuss two different electronicconfigurations at metal–TiO2 interfaces: EF (metal) > EF(TiO2) and EF (metal) < EF (TiO2). Different reaction kineticscan be seen in the two cases.

In the case of EF (metal) > EF (TiO2), the interface contactcauses electron transfer from the metal to TiO2, which resultsin formation of negative space charges in TiO2 and a downwardbending of TiO2 bands. This case favors the outward diffusion

Fig. 24. Schematics showing the electronic interaction between metals and TiO2. (a) EF (metal) > EF (TiO2) and downward bending after contact, favoring metaloxidation; (b) EF (metal) < EF (TiO2) and upward bending after contact, favoring metal encapsulation.


Fig. 25. Dependence of the heat of formation of the metals’ oxide, 1Hof , on the

work function φ of the metals [331]. 1Hof was taken for the values of the most

stable oxide of the metals [331]. The solid horizontal depicts the borderline at1Ho

f = −250 kJ/(mol O). All metals can be classified into four regions I, II,III, and IV (for details see text).

of O2− in TiO2 and, thus, promotes the metal oxidation(Figs. 13 and 24(a)).

The electronic configuration of EF (metal) < EF (TiO2)

causes electron flow from TiO2 to the metal upon interfaceformation. Positive space charges form at the TiO2 surfaceand the TiO2 bands bend upwards, which drives the outwarddiffusion of Tin+i and favors metal encapsulation (Figs. 14 and24(b)).

Therefore, it is expected that there is a strong dependenceof metal–TiO2 reaction kinetics on the electronic structuresof the metals and TiO2. In the following we present previousexperimental data to support the above working hypothesis.

4.1.3.3. Metal–TiO2 interactions. Table 4 summarizes thesurface and interface studies of metals on TiO2(110) whichhave been investigated by PES, AES, TEM, and othertechniques. Metal oxidation is the reaction most commonlyobserved at the metal–TiO2 interfaces. The reaction strengthcan be characterized experimentally by the thickness of themetal layer which may undergo oxidation (see Table 4).Electronic interactions, including the interface bonding andspace charge transfer, were studied primarily by PES. Theinterface bonding results in reduction of surface Ti4+ oroxidation of surface Ti3+ ions. The space charge transfer causesband bending, 1(EF − EV ) (eV), which was determined fromthe BE shifts in PES spectra, such as Ti2p, O1s, O2p, etc. Thevariation of surface work function, 1φ, reflects the relativeposition of EF of the metal and TiO2 before contact betweenthe phases and was measured by UPS. All the electronicinteraction data including interface bonding, 1(EF − EV ), and1φ are also listed in the table.

The metals can be classified into four groups accordingto the different interactions. This is shown in Fig. 25, whichillustrates the category of the metals based on their workfunctions. As shown below, each group will interact withTiO2(110) quite differently.

Fig. 26. Variation of φ, Ebend, and hII (height of UPS peak around 1.0 eVrelated with Na-adsorption), represented by full, half filled, and open circles,respectively, with INa on TiO2(110). The broken line indicates the completionof the first layer. From [272].

(1) Alkali and alkaline earth metals (Cs, K, Na and Ca, Ba)(φ < 3.0 eV, region I in Fig. 25).

These metals possess the highest reactivity to TiO2(110)surfaces concerning their relatively high oxygen affinities.Within one monolayer coverage, fully ionized metals formon TiO2. Further exposure of the metals leads to formationof multilayers of oxidized metals. This was observed by Ladand Dake [342] where they reported that K2O multilayers,more than 9 monolayers thick, grew at 300 K by extractingoxygen from the TiO2 bulk to the surface. The K2O layersare stable against annealing at 900 K. Similar K-adsorption-induced O diffusion to the surface was reported by Heiseet al. [282]. For the Ca/TiO2 interface, XPS data showed that3 nm thick Ca overlayers were oxidized during depositing Caon TiO2 surfaces [283]. For other alkali metals, oxidation wasalso observed at multilayers coverage and the oxidation wasaccompanied by a significant rearrangement of oxygen acrossthe interfaces [72,272,343].

Since the work function of these metals (<3 eV) is muchsmaller than that of TiO2(110) (ca. 5.2 eV [17]), electronsflow from the metals to TiO2 after interface contact formation.The transferred electrons may reside on surface Ti ions andreduce them from Ti4+ to Ti3+ [272,282,342,343]. Some ofthe electrons are delocalized into the CB of TiO2 within adepth comparable to the Debye screening length L D [367]. Theaccumulation of electrons in the TiO2 surface regions causes adownward band bending as shown in Fig. 24(a) and leads toan observed 1(EF − EV ), ranging from 0.3 to 1 eV. In thesesystems, the total decrease in the surface work function is in therange of −2 to − 4 eV, which is consistent with work functiondifferences between the metals and TiO2. Fig. 26 displaysthe variation of surface work function and band bending asa function of Na coverage in the Na/TiO2 system [272]. Itcan be seen that the band bending is as large as 1 eV with acorresponding work function decrease of 3.4 eV. This effect hasalso been seen in other alkali metal–TiO2 systems [84,342,343].

The strong downward band bending and high densitynegative space charges at the TiO2 surfaces accelerate O2−


Table 4Interaction of metal overlayers on TiO2(110)

Metals Interaction@RT Thermal stability Epitaxy Band bending1(EF − EV )

(eV)

1φmax (eV)

Cs [72,84] CsOx & Cs:1 ML < θ < 7 ML [84]Reduction of Ti4+

Desorption 0.3 eV [84] −2.2 [84];−3.5 [72]

K [282,342] K2O: θ > 9 ML [342]Reduction of Ti4+

KOx is stable at 900 K 1 eV [342] −2.2 [342];−3.5 [282]

Na [272,275] Na0+ & Na2O: θ > 3 MLReduction of Ti4+ [272]

1 eV [272] −3.4 [272];−2.5 and−2.9 [275]

Ba [343] Ba0+ & Bax+: θ > 1 MLReduction of Ti4+

0.4 eV [343] −3.3 [343]

Ca [283] CaO: θ > 3 nmReduction of Ti4+

Complete oxidation at900 K

Mg [344] Formation of Mg+

Reduction of Ti4+−2.0 [344]

Al [335] Al2O3: θ < 3 MLReduction of Ti4+

Oxidation of Al

Nb [29,90] NbOx: θ < 2 MLReduction of Ti4+

Oxidation of Nb (100)[001]Nb‖(110)[001]TiO2

Ti [345,346] TiOx: θ < 2 MLReduction of Ti4+

Oxidation of Ti,interdiffusion

0.44 eV [345]

V [29,347,348] VOx at θ < 2 MLReduction of Ti4+

Oxidation of Vinterdiffusion

(100)[010]V‖(110)[001]TiO2 [29]

0.2 eV [347] −1.2 [347]

Cr [88,284,349] CrOx at submonolayerReduction of Ti4+

Oxidation of Cr,interdiffusion

(100)[010]Cr‖(110)[001]TiO2 [88]; bccCr(100) [349]

−1.5 [284]

Fe [71,285,286,349,350] FeOx at submonolayerReduction of Ti4+

Oxidation bcc Fe(100) [349] −1.3 [284];−0.8 [350]

Mo [351,352] Metallic Mo Reduction ofTi4+

Oxidation (100)[001]Mo‖(110)[001]TiO2(reducedcrystal) [352]; bcc(110)texture on oxidizedTiO2(110) surface [351]

Ag [220,290,353] Metallic AgNo reduction of Ti4+

Sintering Downwardbending [290,353]

Cu [29,238,350,354] Metallic CuNo reduction of Ti4+

Sintering (111) [110]Cu‖(110)[001]TiO2 [29]; fccCu(111) [349]

0.5 [350]

Au [310,355,356] Metallic AuNo reduction of Ti4+

Sintering (111) [110]Au‖(110)[001]TiO2(RT);(112) [110]Au‖(110)[001]TiO2 (775K) [356]

0.1–0.15 eV [355];−0.2 [310]

Ni [183,227,328,357–359] Weak interactionNiδ+ [358]; Niδ− [227]

Sintering orEncapsulation

(110)[110]Ni‖(110)[001]TiO2 [359]

−0.5 [358]


Table 4 (continued)

Metals Interaction@RT Thermal stability Epitaxy Band bending1(EF − EV )

(eV)

1φmax (eV)

Rh [82,231,326] Rh negatively chargedOxidation of Ti3+ to Ti4+

Encapsulation fcc Rh(111) [316] −0.8 [82]

Pd [13,330,360,361] Metallic PdOxidation of Ti3+ to Ti4+

Encapsulation (111)[121]Pd‖(110)[001]TiO2 [360]

−0.3 eV [13] 0.5 [361]

Pt [11,83,259,294,362–365]

Metallic PtOxidation of Ti3+ to Ti4+

Encapsulation fcc Pt(111) [11,259] −0.9 [83];Upward bendingto formSchottkybarrier [363,365]

0.5 [294]

O2 [296,301,329,338,366] Oxidation of Ti3+ to Ti4+ Formation of TiO2 orTiaOb layers onsurfaces

−0.8 eV [296] 1.1 [296]

θ : thickness of the metal layer which undergoes oxidation; 1(EF − EV ), band bending: positive for the downward bending and negative for the upward bending;1φmax shows the maximum change of surface work function during the adsorption of metals on TiO2(110).

outward diffusion and promote oxidation at interfaces at roomtemperature (see Figs. 13 and 24(a)). Therefore, adsorptionof the metals on TiO2(110) surfaces results in growth ofmultilayers (>3 ML) of metal oxides at room temperature.

(2) Early transition metals (Mo, Fe, Cr, V, Ti, Nb, and Hf)and Al, 3.75 eV < φ < 5.0 eV (region II in Fig. 25).

These metals are thermodynamically favored to react withTiO2 (1Ho

f < −250 kJ/mol O). Redox reactions were oftenobserved right at the interfaces, i.e., reduction of surface Tiand oxidation of the first metal layer. Further growth of metaloxides necessitates the extraction of O from TiO2. In manycases the O diffusion and growth of oxidized metals is limitedat room temperature. The thickness of the oxidized metal layersranges from submonolayer coverage (Cr [284], Fe [71,285,286,368], and Mo [351,352]), through ∼2 ML (Nb [29,90],Ti [345,346], and V [29,347,348]), up to ∼3 ML (Hf [28]and Al [335]). Mostefa-Sba et al. studied Fe deposition onTiO2(110) by XPS and AES. They found that the redox reactiontakes place only at the interface between the metal and the oxidesurface. i.e., during the completion of the first Fe layer or atthe periphery of the Fe islands [286]. The same conclusion wasreached by Nakajima et al. [368]. Marien et al. used HRTEMto image the Nb–TiO2(110) interface showing that only thefirst two monolayers of Nb were oxidized at room temperature(Fig. 8(a)) [90]. It should be mentioned that annealing theseinterfaces at higher temperature in UHV may facilitate theoxidation via bulk diffusion of oxygen in TiO2 [90,284,337]. Atsome systems, interdiffusion occurs at elevated temperatures.This further confirms that oxidation of the metals on TiO2surfaces is thermally limited at room temperature due to thelow mobility of the diffusing atomic/ionic species.

The decrease in the surface work function upon adsorptionof these metals on TiO2(110) surfaces was observed to be upto ∼1.0 eV (Table 4). Such an electronic configuration withEF (metal) > EF (TiO2) results in long-range charge flow frommetal bands to TiO2 bands and a downward band bending after

establishing interface contact (Fig. 24(a)). However, 1(EF −

EV ) is observed to be below 0.5 eV, which is much weaker thanthat at alkali metal/TiO2 interfaces. The local charge transfer isfacilitated by formation of interface bonding between metal andsurface oxygen. Fig. 27 shows XPS and UPS results recorded atthe V/TiO2 interface by Zhang and Henrich [347]. In the Ti2pspectra (Fig. 27(a)), shoulder peaks at low BE appear after 2ML V deposition. The change originated from a local electrontransfer from V atoms to surface Ti ions. On the other hand,the long-range transfer of electrons from V to TiO2 bends theTiO2 surface bands downwards such that the O2p VB spectrumand O1s core level peak are shifted to higher BE by ∼0.2 eV(Fig. 27(b)).

Table 4 shows that the charge transfer and downward bandbending at these metal/TiO2 interfaces is not as strong as thatin alkali metal/TiO2 systems. Therefore, it is well-expected thatthe driving force behind the O2− diffusion is also smaller thanthat at alkali metal/TiO2 interfaces such that oxidation of themetals is generally limited within three ML.

(3) Mid-to-late transition and noble metals (Ag, Au, and Cu),4.6 eV < φ < 5.4 eV (region III in Fig. 25).

The heats of oxide formation of these metals are abovethe borderline of −250 kJ/mol O, and as a consequence,neither metal oxidation nor TiO2 reduction was observed atthe metal/TiO2 interfaces over a wide range of temperatures.In one example, the Ti2p spectra recorded from the TiO2(110)surface which was covered by Au [369] (Fig. 28) did not showany changes in line shape of the spectra, which indicates thatno reactions occur at the Au/TiO2 interface. Due to the weakinteraction, the metal adatoms have quite high mobility on theTiO2 surfaces, even at very low temperatures, and tend to form3D clusters [220,238,239,354,369].

Work function of these metals is comparable to that ofTiO2. The change in the surface work function during interfaceformation is quite small, normally below ±0.5 eV. Thereis no significant band bending because of the similarities of


Fig. 27. (a) XPS spectra of Ti2p core levels for a stoichiometric TiO2(110)surface before (solid curve) and after (dashed curve) deposition of∼2ML V. (b)UPS spectra for a stoichiometric TiO2(110) surface with different V coverage.From [347].

Fig. 28. XPS spectra of Ti2p from TiO2(110) surface covered by Au with athickness from 1.1 to 52 A (nominal thickness). From [369].

EF (metal) and EF (TiO2). For example, upward bending of∼−0.2 eV was observed in case of 0.1 ML Au-adsorption onTiO2(110) [310] while 0.1–0.15 eV downward bending wasreported for Au on TiO2 with coverage less than 1 ML [355].Due to the small charge transfer between the metal and TiO2, itwas concluded that the interaction at the interfaces is very weak.

(4) VIII B metals (Pt, Pd, Rh, Ir, and Ni), φ > 5.4 eV (regionIV in Fig. 25).

There was no oxidation of the metals observed, even atelevated temperatures, for these surfaces. Prolonged UHVannealing at high temperatures, however, leads to theencapsulation reaction. This was shown for Pt in [11,111,247,

259,362,363,370], Pd in [13,329,330,360], Rh in [82,231,316,326,370], Ir in [184], and Ni in [183,328,357].

An increase in the surface work function during depositionof the metals on TiO2(110) was often observed, which indicatesan electronic configuration of EF (metal) < EF (TiO2) priorto interfacial contact. For example, Negra et al. [361] andSchierbaum et al. [83,294] reported a 0.5 eV increase in thesurface work function after Pd and Pt-adsorption on TiO2(110)or (100) surfaces. Because the metal Fermi level is lower thanthat of TiO2, contact between the metals and TiO2 causeselectron transfer from TiO2 to metals [13,82,231,294,326,361,364] and may lead to the formation of a Schottky diode [83,363,365]. This unique charge transfer process leads to upwardbending of TiO2 bands, formation of positive space charges atTiO2 surfaces, oxidation of surface Ti, and negatively chargedmetal clusters. For example, a detailed surface analysis ofRh/TiO2 interface was performed by Sadeghi and Henrich [82].They found that shoulder peaks originating from Tin+ (n <

4) on a sputtered TiO2 surface were largely weakened after1.5 ML Rh deposition (Fig. 29(a)). Such a change originatedfrom a local electron transfer from Tin+ (n < 4) ions to Rh,i.e., oxidation of surface Ti ions instead of surface Ti reductionat V/TiO2 interfaces as shown in Fig. 27(a). The long-rangecharge transfer from TiO2 to Rh resulted in an upward bendingof TiO2 bands, which was confirmed by the O2s shift to a lowerBE after 0.5 ML Rh deposition (Fig. 29(b)).

The unique electron transfer, which occurs from TiO2 to themetals (Pt, Pd, Rh, etc.), results in formation of positive spacecharges at TiO2 surfaces, which promotes the outward diffusionof Tin+i at the interfaces (see Figs. 14 and 24(b)). Therefore, itcan be concluded that encapsulation is uniquely observed formetals with a large work functions (see Fig. 23) on TiO2 [13].

(5) Oxygen-induced restructuring and bulk-assisted reoxida-tion reactions on TiO2 surfaces.

One well-known TiO2 surface reaction is the so-calledoxygen-induced restructuring of TiO2(110) surfaces. Thereaction describes the growth of surface phases whenheating TiO2(110) surfaces in O2 environment. The dominantmechanism of the reaction is the outward diffusion of Tin+iions from the bulk to surface, where they react with gaseousoxygen and form phases like TiO2 or TiaOb (e.g. Ti2O3) [301,329,338,366]. Due to the large electronegativity of oxygen it isreasonable to suggest that the electronic energy level of O2pis lower than EF of TiO2 (EO2p < EF (TiO2)) when oxygenis adsorbed on TiO2 surfaces. Thus, exposure of TiO2 to O2leads to transfer of electrons from TiO2 to surface adsorbedoxygen atoms. Such an electron transfer process causes upwardbending of the TiO2 surface energy bands and formation ofpositive space charges at the TiO2 surface. Henrich et al. haveshowed that exposure of reduced TiO2 surfaces to O2 at RTresults in 1(EF − EV ) of −0.8 eV and 1φ of 1.1 eV as shownin Fig. 30 [296]. Therefore, the outward diffusion of Tin+itowards the surface is driven by negatively charged oxygenatoms on the TiO2 surfaces and the positive space charges atthe near-surface regions in TiO2. The reaction process at theO2–TiO2 interfaces can be well-compared to the encapsulationat VIII B metals/TiO2 interfaces (see Figs. 14 and 24(b)). For


Fig. 29. (a) XPS spectra of Ti2p core levels for slightly reduced TiO2(110)before (solid curve) and after (dashed curve) deposition of 1.5 ML Rh. (b) UPSspectra for slightly reduced TiO2 before (solid curve) and after (dashed curve)deposition of about 0.5 ML Rh. Fom [82].

comparison, data concerning the O2–TiO2 interface reactions isshown at the end of Table 4.

Another TiO2 surface reaction is the bulk-assisted reoxida-tion of reduced TiO2 surfaces in which the surface reducedTi ions are reoxidized upon heating a reduced TiO2 surfacein UHV [339,340]. The mechanism of the reaction is the out-ward diffusion of O2− or inward diffusion of Tin+i . To drivethe O2− outward diffusion or Tin+i inward diffusion, it isclear that a downward band bending and formation of negativespace charges at TiO2 surfaces are necessary. Actually, elec-tron accumulation at TiO2 surface and downward band bend-ing of 0.2–0.3 eV have been observed on a reduced TiO2 sur-face, for example see Fig. 35 in Ref. [28]). The schematicsshown in Figs. 13 and 24(a) can also explain the bulk-assistedreoxidation. Therefore, the oxygen-induced restructuring andbulk-assisted reoxidation reactions are dominated by the sameconcept.

In summary, we could show that the metal–TiO2 interactionsdepend strongly on φ of the metals. Redox reactions, observedfor metals with small work functions (e.g. Na, K, Al), showa strong inverse correlation with work function: the smallerthe φ, the stronger the oxidation of the metal. On the otherhand, metals with very large work function (from Ni to Pt)undergo encapsulation reactions on TiO2(110). The interfacereactivity between metals and TiO2 can be categorized into thefour groups according to their work function values (Fig. 25).

Fig. 30. Work function change (1φ), Fermi level (EF − EV ), position ofsurface-state UPS peak (ES − EV ), and normalized amplitude of the ELSpeak A versus normalized intensity of UPS surface-state peak (αS) for Ar-ionbombardment (solid points) and subsequent oxygen exposure (open points) ofthe TiO2(110) surface. Arrows indicate the sequence in which data were taken.From [296].

Such a correlation cannot be drawn between the reactivity andthe thermodynamic data, e.g. 1Ho

f (see Fig. 25). All the aboveexperimental results can be understood within the generalizedCabrera–Mott theory, as suggested in Section 2.4.

4.2. Metals on SrTiO3

SrTiO3 is a model material of the perovskite-type oxides(ABO3). It has many technological applications, e.g., insensors, photocatalysts, and gate dielectrics. In addition,SrTiO3 has been widely used as a substrate for epitaxialgrowth of films of high-Tc superconductors, metals, andoxides. SrTiO3 surfaces and metal interactions on thesesurfaces were extensively studied. SrTiO3 surfaces, SrTiO3bulk defect chemistry, and metal overlayers strongly affect themetal–SrTiO3 interactions.

4.2.1. SrTiO3 surfacesThe (100), (110), and (111) facets are the most commonly

studied SrTiO3 surfaces. Among them, the (110) and (111)surfaces are polar, while the (100) surface is non-polar.Most research focused on the more stable SrTiO3(100)surface. Important surface characteristics, including surfaceterminations, surface defects, and surface reconstructions, are


addressed in the following, and the influence of these surfaceproperties on deposited metal overlayers are discussed in detail.

4.2.1.1. Surface terminations. A SrTiO3 crystal consists ofa stack of alternating TiO2 and SrO layers. Therefore, aSrTiO3(100) surface may have two distinct terminations: TiO2-termination and SrO termination. The surface terminationcan be experimentally controlled. Kawasaki et al. [152] havedeveloped a wet chemical etching method to prepare the TiO2-terminated surface. It has been found that etching of pre-polished SrTiO3(100) substrates in NH4F-buffered HF (BHF)solutions selectively removes surface SrO groups as well ascalcium impurities leaving the surface terminated with Ti andO [152,153,371]. The reproducibility of the surface etchingprocess has been improved through hydroxylation of thetopmost SrO layer, which enhances the etching-selectivity ofSrO relative to TiO2 in a BHF solution [91,372]. Atomicallyflat surfaces with a TiO2-termination can be obtained by thechemical etching followed by heating in O2 and UHV. Thesurface termination can be switched to the SrO-terminationby depositing a monolayer of SrO on the TiO2-terminatedsurface [251,373–375].

EELS measurements indicate that the electronic structuresof the two differently terminated surfaces are different. TiO2-terminated surfaces show intrinsic surface states while SrO-terminated surface does not exhibit such states [373]. It hasbeen seen that the photocatalytic activity on the SrO-terminatedsurface is lower than that on the TiO2-terminated surface [373,375]. Furthermore, the surface termination has strong effects onmetal growth on the SrTiO3(100) surface.

Polli et al. [251] have investigated Pt growth on the twodifferently terminated SrTiO3(100) surfaces. They showed thatthe SrO-terminated surface has a mixture of Pt(100) andPt(111). Compare this to a TiO2-terminated surface wherePt islands form with a (100) orientation. The differencebetween orientations was explained as being caused by theweaker bonding between Pt and the SrO-terminated surfaceas compared to Pt on the TiO2-terminated surface. Thisconclusion has been tested by a DFT calculation from Sholland coworker [376]. They found that Pt atoms prefer tosit on top of O atoms on both surfaces, but the bondingstrength at Pt/TiO2–SrTiO3 interface is stronger than that atPt/SrO–SrTiO3 interface in cases of a Pt coverage above 1ML. A similar conclusion was made based on DFT studiesfor Pd/SrTiO3 and Co/SrTiO3 interfaces. The TiO2-terminatedsubstrate is energetically favorable for adhesion of Pd filmswith Pd atoms bonded on top of O atoms [377,378]. At theCo/SrTiO3 interface, the strongest cohesion between Co andSrTiO3 is also found for the Co/TiO2–SrTiO3 interface wherethe interfacial Co atoms sit on top of the interfacial O atoms.Indirect electronic coupling was observed between interfacialTi and Co atoms mediated by the O atoms [379].

4.2.1.2. Surface point defects. Surface oxygen vacancies andTi3+ ions can be generated by ion bombardment, whichproduces electronic states with predominant d electroncharacter in the band gap [380]. The gap states are located

several tenths of an eV below the CBM [380–382]. Comparedto stoichiometric SrTiO3(100) surfaces, defective SrTiO3(100)surfaces exhibit a higher reactivity to gases such as NO [382,383], methanol, [384], and acetaldehyde [385]. It has beendemonstrated that the surface defects influence the metalinteraction as well. Conard et al. [386] studied Cu growth onboth stoichiometric and reduced SrTiO3(100). They found thatthe number of Cu islands is higher on the reduced surface thanon the stoichiometric surface. This was attributed to a strongerinteraction of Cu with the defective surface. We applied in situAES to monitor Cr growth on different SrTiO3(100) surfaces.The decay of substrate signals and increase of Cr signals for Cron Ar+-sputtered SrTiO3(100) are much faster than those onstoichiometric SrTiO3(100), which indicates a more layer-likegrowth and, thus, stronger interaction of Cr on the reduced andhighly defective surface [387].

4.2.1.3. Surface reconstructions. Many reconstructions of theSrTiO3(100) surface have been experimentally observed. Theyinclude 2 × 1 [145,388–395], 2 × 2 [142,381,388,396–398], c(4 × 2) [142–144,146,390,393–395,399,400], c(4 ×4) [142,390,398], c(6 × 2) [143,144,389,395], (

√5 ×√

5) −

R 26.6◦ [141,142,401,402], c(√

13×√

13) [403], and (√

13×√

13)− R 33.7◦ [142,389]. A summary of SrTiO3(100) surfacereconstructions is given in Table 5. It can be seen that thereconstruction depends sensitively on the annealing conditions,e.g., atmosphere (oxidation or reduction), temperature, andtime, as well as surface conditions prior annealing. As oneexample, Jiang and Zegenhagen [144] found that a c(4 ×2) surface is obtained by annealing a pristine surface withassistance of hydrogen but a c(6×2) reconstruction is observedupon heating O2 annealed surfaces. Castell [390,399] reportedthat the creation of 2 × 1 and c(4 × 4) was achieved throughBHF etching and UHV annealing while a c(4× 2) surface wasfound by Ar+ ion sputtering and subsequent annealing in UHV.

Two kinds of surface models have been suggested to explainthe SrTiO3 surface reconstructions, Ti-rich surface modelsand Sr-rich surface models. Oxygen vacancies on the topmostTiO2 layer can be produced by vacuum annealing. Therefore,the surface superstructure is attributed to an ordering of Ovacancies [144,381,401,402]. Formation of TiOx (x ≤ 2)-typesurface phases was suggested to explain some experimentalresults, e.g., a Ti2O3 surface phase for the (2× 1) surface [390]and a TiO2 overlayer atop a bulk-like TiO2 layer for (2×1) andc(4× 2) structures [145,146]. In contrast to the Ti-rich surfacemodels, Kubo et al. [141,142] found that a structural modelconsisting of an ordered Sr adatom at the O four-fold site ofa TiO2-terminated layer can explain previous experimental andtheoretical results. Liang and Bonnell [404] also explained theintergrowth of a SrO-rich surface phase on SrTiO3(100) with alamellae structure (Srn+1TinO3n+1).

In brief, the surface reconstruction is closely relatedto the change in surface composition. Explanation of thereconstructions can be complicated by a number of factorswhich include surface non-stoichiometry, presence of TiO2and SrO terminations, diffusion of defects between bulk andsurface, and surface segregation of impurities. In order to


Table 5Surface reconstructions on SrTiO3(100)

Reconstructions Sample preparations Techniques

2× 1 UHV, 800–900 ◦C, 20–120 min LEED, STM [388]

Vacuum, 750–800 ◦C, 1 h RHEED [389]

UHV, 600–800 ◦C, 30 min–9 h STM [390–392]

O2, 1050–1100 ◦C, 0.5–5 h, after ion milling TEM [145,395]

2× 2 UHV, 1000–1200 ◦C, 2–20 min LEED, STM [388]

Vacuum, T < 700 ◦ C LEED, AES [396]

UHV, 650 & 730 ◦C RHEED [381]

UHV, 800 ◦C, 5 h, after Ar+ sputtering STM [397,398]

UHV, 1000–1250 ◦C STM [142]

c(4× 2) H2, 950 ◦C, 2 h LEED, STM, AES [143,144]

UHV, 900–1150 ◦C, only after sputtering STM [390,393,394,398–400]

UHV, 1000–1250 ◦C STM [142]

O2, 830–930 ◦C, after ion milling TEM [146,395]

c(4× 4) UHV, 950 ◦C, 2 h, only after O2 heating LEED, STM, AES [143,144]

O2, 800–1000 ◦C, 15 h RHEED [389]

c(6× 2) UHV, up to 1200 ◦C, some minutes STM [402]

UHV, 1200 ◦C STM [141,142]

O2, 1050–1100 ◦C, after ion milling TEM [395]

(√

5×√

5)-R26.6◦ O2, 800–1000 ◦C, 15 h RHEED [389]

UHV, 1200 ◦C, several minutes STM [142,401]

c(√

13×√

13) UHV, 950 ◦C, 2 h STM [403]

(√

13×√

13)-R33.7◦ O2, 800–1000 ◦C, 15 h RHEED [389]

UHV, 1250 ◦C STM [142]

clarify the atomic structure of all the different SrTiO3 surfacereconstructions further detailed studies are necessary.

The reconstructed SrTiO3 surfaces are significant intechnological applications of SrTiO3, e.g., in the fields of filmgrowth, catalysis, and other surface chemical reactions. Forexample, Silly and Castell [393,394] performed an elegantexperiment of metal growth on SrTiO3(100) surfaces wherethey demonstrated that the shape and orientation of supportedmetal nanocrystals can be selected by controlling the oxidesubstrate reconstruction. On a SrTiO3(100)-(2 × 1) surface Pdforms hut-shaped nanocrystals; while on a SrTiO3(100)-c(4 ×2) surface hexagonal and truncated pyramid nanocrystals areobserved (see Fig. 31). It was found that the two reconstructedSrTiO3(100) surfaces have different surface energies and theinterface energies between the metal and the substrates varieswith the substrate crystallography. As a consequence, thesupported nanocrystals can have different equilibrium shapesand orientations.

4.2.2. SrTiO3 bulk defect chemistryLike TiO2, chemical reactions occurring on SrTiO3 surfaces

are closely related with SrTiO3 bulk defects. In this section, thedefect chemistry of SrTiO3 will be discussed first. Then, wewill show how the reactions on SrTiO3 surfaces are dependenton the defect chemistry of SrTiO3. The reactions include

SrTiO3 surface reactions upon high temperature annealing andmetal–SrTiO3 interface reactions.

4.2.2.1. Defect chemistry of SrTiO3. SrTiO3 is the prototypeof mixed conductors whose defect chemistry has been studiedin detail. Schottky-type defects are dominant in SrTiO3 solids,which include oxygen vacancies, V ••O , and strontium vacancies,V ′′Sr [405]. Assuming two-fold ionization of the O vacancies, thedefect reaction of O vacancies (V ••O ), free electrons (e′), andlattice oxygen sites (OX

O) is described by

O XO ↔ V ••O + 2e′ + 1/2O2. (38)

The defect reaction of lattice Sr sites (SrXSr) and OX

O can bewritten as

OXO + SrX

Sr ↔ V ′′Sr + V ••O + SrORP-phase. (39)

The formation of V ••O and V ′′Sr has been suggested to beaccompanied by growth of Ruddlesden–Popper (RP) phases[SrO · (SrTiO3)n] [406] in the solids which act as a sink for theexcess Sr. However, direct experimental proof of an RP phaseformation in the SrTiO3 solid has not been given as yet. Instead,recent results indicate that Sr vacancies can be created only atthe crystal surfaces and all excess Sr migrates to the surfaceswhere secondary SrOx phases grow on top of the surfaces of


Fig. 31. (a) STM image (140× 140 nm2) of ca. two monolayers Pd deposited onto SrTiO3(100)-(2× 1) at room temperature followed by 650 ◦C anneals showinghut-shaped nanocrystals; (b) STM image (140 × 140 nm2) of Pd deposition of around one monolayer onto SrTiO3(100)-c(4 × 2) at 175 ◦C followed by 650 ◦Canneals showing two kinds of nanocrystal shapes, hexagons and pyramids. From [394].

the single crystals [407–409]. The defect reaction of the cationsub-lattice can be written as

SrXSr + 2e′ + 1/2O2 ↔ V ′′Sr + SrOsecond-phase. (40)

For doped SrTiO3, the following equilibria may exist in thesolids:

LaXSr ↔ La•Sr + e′, for donor of La on Sr sites,

NbXTi ↔ Nb•Ti + e′, for donor of Nb on Ti sites,

and

FeXTi ↔ Fe′T i + h•, for acceptor of Fe on Ti sites. (41)

The above defect reactions suggest that the concentrations ofV ••O , V ′′Sr, e′, and hole (h•) are a function of temperature, O2partial pressure (pO2), and extrinsic dopant concentrations.Qualitatively speaking, annealing in an oxygen-deficientatmosphere (e.g. UHV) increases V ••O and e′ concentrations([V ••O ] and [e′]

)while heating in an oxygen-rich environment

increases V ′′Sr and h• concentrations([V ′′Sr] and [h•]

).

In SrTiO3, all cations move via Sr vacancies and anionsdiffuse via O vacancies. It is found that the diffusivity of Tiis much lower compared to that of Sr and the Sr diffusivityis orders of magnitude lower than the diffusivity of oxygenvacancies [409,410].

In the following, we show that SrTiO3 defect chemistryplays an important role in SrTiO3 surface reactions and inmetal–SrTiO3 interface reactions.

4.2.2.2. Effect of SrTiO3 bulk defects on SrTiO3 surfacereactions. Surface phases often form on SrTiO3 surfaceswhen heating the crystals at elevated temperatures. Generally,annealing in an oxidizing atmosphere produces SrO-richsurface phases (Eq. (40)) while heating in a reducingenvironment results in growth of TiOx-rich surface phases (Eq.(38)). The SrTiO3 surface reactions are similar to the surfacerestructuring of TiO2 surfaces. As discussed in Section 4.1.3.3,

oxidation of reduced TiO2 crystals drives the outward diffusionof titanium interstitial ions to the surfaces, where they react withgaseous oxygen and additional TiO2 or TiaOb structures formon the surfaces [301,329,338,366]. The formation of surfacephases on SrTiO3 surfaces is also related to defect chemistryin SrTiO3. In particular, the production and diffusion of Srvacancies in SrTiO3 dominate the SrTiO3 surface reactionsoccurring at high temperatures in both oxidizing and reducingatmosphere.

Surface reactions in oxidizing atmosphere: According tothe defect reactions in SrTiO3, heating SrTiO3 in oxidizingatmosphere increases

[V ′′Sr]

at the SrTiO3 surface (Eq. (40)).The gradient in

[V ′′Sr]

drives the inward diffusion of V ′′Sr. Thismeans that Sr ions will diffuse outwards from the bulk tothe surface, where they can react with gaseous oxygen andform SrO-rich surface phases. Szot et al. have observed acontinuous accumulation of SrOx on SrTiO3(100) surfacesin the case of oxidizing conditions [411,412]. Sr surfacesegregation and growth of SrOx secondary phase on SrTiO3surfaces under oxidizing conditions have been also confirmedby many groups [407,413–418]. The amount of the formedsurface SrOx layers increases with the doping level of La orNb in SrTiO3 since the extrinsic donors can create more Srvacancies and enhance Sr diffusion in the solids [413–415].Rahmati et al. also showed that the density of surface SrO-rich islands depends on SrTiO3 surface orientations and islanddensity decreases in the sequence of (100), (110), and (111).The result was attributed to that the Sr segregation rate isrelated with SrTiO3 surface energy which shows the order ofγ (100) < γ (110) < γ (111) [416,417].

Surface reactions in reducing atmosphere: In reducingconditions, oxygen may desorb from the surface (see Eq. (38)),leading to an increase in [V ••O ] at the surface. Accordingto Schottky equilibrium, [V ′′Sr] subsequently decreases. Thegradient in Sr drives inward diffusion of Sr ions from thesurface to the bulk. This results in Ti enrichment as well asO- and Sr-deficiency at the surface. It has been demonstrated


that reducing SrTiO3 surfaces in vacuum above 900 ◦C causesformation of various Ti-rich surface phases, including Ti-richphases of TiO and Ti2O [411], TiO2 and Magnelli-type Ti-richphases [419], TixOy Magneli phases [413], Ti2O3 islands [420],anatase TiO2(100) islands [421], and TiO islands [422]. Anexception is the result reported by Liang and Bonnell [404].They found that SrO-rich islands formed on SrTiO3(100) afterannealing the surface in UHV at 1300 ◦C.

4.2.2.3. Effect of SrTiO3 bulk defects on metal oxidation onSrTiO3. In order to study the influence of SrTiO3 defectchemistry on metal–SrTiO3 interface reactions, Cr/SrTiO3interfaces were used as model systems and the reactivity of theCr/SrTiO3 interface was studied on various SrTiO3 crystals [85,89]. Five SrTiO3(100) crystals were prepared in different ways.SrTiO3#1: undoped and strongly reduced; SrTiO3#2: Nb-doped(n-type doping) and normally reduced; SrTiO3#3: undoped andnormally reduced; SrTiO3#4: Fe-doped (p-type doping) andnormally reduced; SrTiO3#5: undoped and oxidized. Accordingto the SrTiO3 defect chemistry discussed in Section 4.2.2.1and the defect model suggested by Moos and Hardtl [405],the defect concentrations

[V ••O

], [V ′′Sr], [e

′], and [h•], can be

calculated. The data is shown in Fig. 32. It can be seen that[e′]

in the five crystals decreased in the order #1 to #5 whereas [V ••O ]

decreased in the sequence #1, #4, #3, #5, #2. The decreasing [e′]in the crystals in the order #1 to #5 results in a shift of EF fromCB to VB, which transforms the n-type doping of #1 to p-typedoping of #5.

The thermal stability of Cr clusters (nominal film thickness6 A) on the five SrTiO3(100) crystals was studied underidentical experimental conditions. XPS was used to monitor Croxidation on the SrTiO3 surfaces and determine the reactiononset temperature (TRO, the temperature point where oxidationof Cr is first observed). We found that Cr on SrTiO3 #1 was notoxidized below 810 ◦C (TRO > 810 ◦C) while Cr oxidationoccurred on SrTiO3 #5 as low as 280 ◦C. The TRO of Croxidation on SrTiO3 #2, SrTiO3 #3, and SrTiO3#4 were 760 ◦C,640 ◦C, and 460 ◦C, respectively. The large difference in TROreflects significant variation in the Cr oxidation rate on SrTiO3,which decreases in the order Cr/SrTiO3#5 to Cr/SrTiO3#1.Plotting the data of TRO with [e′] reveals a systematic increaseof TRO with

[e′]

(Fig. 32(a)). A smaller [e′] in SrTiO3corresponds to faster oxidation of Cr on SrTiO3. In contrast,we did not find such a systematic dependence between TRO and[V ••O ] (Fig. 32(b)).

It has been shown that oxidation of metals on SrTiO3(100)necessitates outward diffusion of oxygen ions at the SrTiO3surfaces [91,237]. If the oxygen diffusion is controlled byFick’s law (Eq. (20)) the oxygen transport in SrTiO3 andthe rate of Cr oxidation on SrTiO3(100) should be stronglyrelated with [V ••O ]. However, this work demonstrated that theCr oxidation rate was controlled by [e′] rather than [V ••O ] asshown in Fig. 32. This behavior could be explained by assumingthat at relatively low temperatures (<900 ◦C) the Cr oxidationon SrTiO3(100) is determined by the interface electric fieldand [e′] in SrTiO3. In the case of interface contact between Crand the electron-rich SrTiO3, e.g. SrTiO3 #1 and SrTiO3 #2,

Fig. 32. TRO (the temperature point where oxidation of Cr is first observedby XPS) for Cr oxidation on STO as a function of [e′] (a) and as a functionof [V ••O ] (b) in the STO crystals #1 to #5. Reproduced from [89]. Thedefect concentrations were calculated according to the defect model suggestedin [405].

the interfacial electronic configuration is similar to that shownin Fig. 14. In that case, the positive space charge hinders theoxygen outward diffusion and Cr oxidation. On the other hand,the contact of Cr with the p-type doped SrTiO3, e.g. SrTiO3#5, leads to the interfacial electronic configuration shown inFig. 13. Then, the oxygen outward diffusion and Cr oxidationare promoted by a negative space charge (compare Fig. 9 inRef. [89]). Such a systematic investigation of Cr oxidationon SrTiO3(100) clearly shows that the metal/oxide interfacereactions are closely related to the defect chemistry in oxides,and the reactions are tunable by variation in the electronicstructure of the oxides. These interfacial reactions can beexplained in the framework of the Generalized Cabrera–Motttheory (see Chapter 2.4.3).

4.2.3. Metal–SrTiO3 interactionsVarious metal overlayers have been grown on SrTiO3

surfaces. The main results of the metal interactions onSrTiO3(100) surfaces are listed in Table 6. Here, we focus a fewimportant characteristics of metal–SrTiO3 interfaces, includingthe interface electronic structure, growth and epitaxy of themetal overlayers, and interface reactions.


Table 6Summary of metal interactions on SrTiO3(100) surfaces. OR I, OR II, OR III, OR IV, and OR V are different orientations of metals on SrTiO3(100) (see the text)

Metal/SrTiO3 systems Interaction Reaction Growth and epitaxy

Y/SrTiO3(100) [423] Oxidation of Y Layer growth

Y, Ba, Ti, /SrTiO3(100) [424]

Oxidation of metals; reduction of Ti4+

to Ti3+ and Ti2+Oxidation of metal overlayers at RT and500 ◦C; strong outward diffusion of O

Al/SrTiO3(100) [89,108] Oxidation at RT OR I

Cr/SrTiO3(100) [89,91,177,218,237,258]

Formation of interface bondingbetween Cr and O; SrTiO3 bandbending after Cr deposition;

Oxidation at elevated temperatures; Reactiontemperature depends on SrTiO3-doping andsurface orientation

Island growth; 100–600 ◦C: OR II;>700 ◦C: OR II & OR III

Mo/SrTiO3(100) [89,263,425,426]

Formation of Mo–O bonds at interfaces Oxidation of Mo above 900 ◦C Metastable fcc Mo at low coverage;OR II at low T ; OR III & OR IV athigher T

Fe/SrTiO3(100) [89,427] Oxidation of Fe above 800 ◦C TiO2-terminated: OR II; TiO2- andSrO-terminated surface: major OR II& minor OR IV

Fe/SrTiO3(100)-c(4× 2) [400]

Truncated pyramid islands with OR II

Co/SrTiO3(100)-(2× 2) [397]

Truncated pyramid islands with fccOR I

Ni/SrTiO3(100) [108,258,428–431]

Ni bonds with outmost surface O toform 2D NiO; work function increasesby 1.3 eV after Ni deposition

Simultaneous Multilayer (SM)growth [428]; OR I [108,258]

Cu/SrTiO3(100) [386,424,432]

Cu bonds to O: strong interaction No reaction but sintering of Cu duringannealing at 500 ◦C

OR I at 100 ◦C; Cu(100) andCu(111) with CuOx interface layer

Ag/SrTiO3(100)-(1× 2) [392]

Icosahedral nanocrystals

Au/SrTiO3(100)-(1× 2) [391]

Wetting at submonolayer Au;Dewetting as annealing

Truncated triangle island with smallamount of icosahedral islands; ORV

Pd/SrTiO3(100) [108,433,434]

600 ◦C: OR I T < 250 ◦C: {111}with ORV

Pd/nanoline−structured SrTiO3(100) [435]

Encapsulation of Pd with TiO1.36 or TiO1.37layers

Pd/SrTiO3(100)-(2× 1) &SrTiO3(100)-c(4× 2) [393,394]

On SrTiO3 (100)-(2× 1): hut islandswith OR III; On SrTiO3 (100)-(4×2)

@ RT: Hexagonal islands with OR V;On SrTiO3(100)-(4× 2) @ 460 ◦C:truncated pyramid islands with OR I

Pt(621)/SrTiO3(621) [436] 600 ◦C: (621)[012]Pt‖(621)[012]SrTiO3

Pt(620)/SrTiO3(620)&(622) [437]

Pt adsorbs more at steps; interfacestrength larger on low index surfaces

DFT: Possible step-flow growth fromDFT calculations

Pt/SrTiO3(100) [86,221,438]

Charge transfer from SrTiO3 to Pt:0.6e/Pt atom; band bending of −0.6 eVafter Pt; SBH ∼ 0.6 eV

Pt/SrTiO3(100) [439] XPS, band bending of −0.6 eV


Table 6 (continued)

Metal/SrTiO3 systems Interaction Reaction Growth and epitaxy

Pt/SrTiO3(100) [251,440] TiO2-terminated surface: pyramidalislands with OR I; SrO-terminatedsurface: pyramidal islands with OR Iand hexagonal islands with OR V

Ir/SrTiO3(100)-c(√

13×√

13) [403]Needle islands with OR I

Fig. 33. (a) Ti2p spectra from the ordered STO (100) surface as a function of Y coverage showing Ti reduction from 4+ to 3+ and 2+ nominal valences [424], and(b) Ti2p spectra from the sputtered STO (100) surface with different coverages of Pt [221].

4.2.3.1. Electronic interaction of metals with SrTiO3. Theadsorption of metals on SrTiO3(100) surfaces may cause theformation of interface bonding and charge transfer betweenmetal adatoms and SrTiO3 surface atoms. The local chargetransfer occurring via interfacial bonding depends on theelectronegativity of the metals. Reactive metals with smallelectronegativity, such as Y, Ba, and Ti [424], K [441], andCr [91,218], could reduce the surface Ti4+ ions where chargetransfer occurs from the metal atoms to SrTiO3. In contrast,adsorption of noble metals with large electronegativity, e.g. Pt,may result in electron transfer from SrTiO3 to the metal atomssuch that surface reduced Ti ions (e.g. Ti3+ or Ti2+) areoxidized by the metals [86,221]. The two different chargetransfer processes were exemplified by the evolution of Ti2pspectra from two SrTiO3(100) surfaces which were covered byY and Pt, respectively (Fig. 33). The electron transfer from Yto Ti resulted in an increase in the relative intensity of peaksfrom reduced Ti while the intensity of the reduced Ti peak wasweakened by electron transfer from surface Ti to Pt upon Ptdeposition.

The space charge transfer between metals and SrTiO3 isdetermined by the interface electronic configuration, i.e., EFof the metal overlayers and EF of the SrTiO3 crystal beforecontact. The work function of the clean and reduced SrTiO3surfaces is around 4.2 eV [86,380,442]. Metals with surfacework function above 4.2 eV form Schottky-type contactswith SrTiO3 single crystals. In case of interface contact,electrons flow from the SrTiO3 CB to the metal VB whichproduces positive space charges and upward band bendingat SrTiO3 surfaces. For example, Chung et al. [86] observeda band bending of −0.6 eV on SrTiO3(100) surface after1 ML Pt deposition. Copel et al. have performed XPS andUPS investigations of Pt–SrTiO3(100) and Au–SrTiO3(100)surfaces showing Pt-induced band bending of about −0.6 eVand little band bending for Au metallizations [439]. Parkand Kim showed that metal/Nb-doped SrTiO3(100) junctionwith a low work function metal of Ti exhibited linearcurrent–voltage (I –V ) characteristics while for Ni and Ptmetals with large work function the junctions showed rectifyingI –V characteristics and notable hysteresis due to the formation


Fig. 34. Plot of the lattice mismatch between the overlayer film and SrTiO3substrates, | f | as a function of the oxygen affinity of the overlayer metal, pOshowing the range of epitaxy of metals. Region A: epitaxial growth of metals;Region B: polycrystalline growth of metals; Region C: chemical reactions ofmetals. From [108].

of Schottky barrier. The results clearly indicate that metal workfunction influences the Schottky barrier height and the interfaceresistance switching [443]. Shimuzu et al. have measuredthe SBHs from various metal/SrTiO3 junctions and plottedthe SBHs as a function of Pauling’s electronegativity of themetals. They show that the determined slope parameter SX isabout 1, which suggests that the interfacial electronic statesare not dominant in the interface charge transfer process atmetal/SrTiO3 interfaces (see discussion in Section 2.1.1) [438].

4.2.3.2. Growth and epitaxy of metals on SrTiO3. SrTiO3 hasbeen often used as a substrate for epitaxial growth of metalfilms. Here, we review the results of metal growth and epitaxyon SrTiO3. It is shown that the metal epitaxial growth is closelydependent on the metal–SrTiO3 interactions.

Growth: SrTiO3(100) possesses a relatively low surfaceenergy about 0.9 J/m2 for the unreconstructed and TiO2-terminated surface ([444] and references therein). According toexpression (1), it is well-expected that most metals grow on theSrTiO3 surfaces with the island growth mode (Volmer–Webermode) [108]. Layer growth was observed only for a few cases,Y growth on SrTiO3(100) due to strong reaction between Y andSrTiO3 [423] and the formation of wetted Au monolayer islandson SrTiO3(100)-(2 × 1) because of the commensurate epitaxyat submonolayer coverage [391].

Metal islands on SrTiO3 surfaces tend to form in such amanner that the total energy 1E is minimized

1E =∑(hkl)

γMe(hkl) AMe(hkl) + γi Ai − γSTO Ai . (42)

Here,∑

(hkl) γMe(hkl) AMe(hkl) is the sum of the products ofsurface energy of every metal facet, γMe(hkl) and the metalfacet area, AMe(hkl), γi is the interface energy, γSTO is SrTiO3surface energy, and Ai is interface area. Obviously, the shapeand orientation of metal islands are determined by γMe(hkl),γSTO, and particularly γi . Wagner and coworkers [91,108,237,251,263,425,433,445] have systematically studied the epitaxialgrowth of various metals on SrTiO3(100). They suggested that

γi is closely related to lattice mismatch and interface reactivity.The lattice mismatch is described by f = (aSub − aMe)/aMe(aMe: lattice parameter of metal; aSub: lattice parameter ofsubstrate) and the interface reactivity is characterized by theoxygen affinity pO of the metals. γi increases monotonicallywith f 2 but decreases monotonically with pO. They found thatmetal epitaxy can be achieved only in a certain range of | f |and pO (see Fig. 34). For high | f | and low pO, polycrystallinefilms may be observed. In case of a very high pO, a strongchemical reaction between the metal and the substrate mayoccur, which also prevents metal epitaxy [108]. Silly andCastell have applied STM to image the growth of metalnanocrystals on reconstructed SrTiO3(100) surfaces [391–394,397,400]. They found that γ ∗, which is the energy differencebetween the interface energy and the substrate surface energy,plays a critical role in the shape and orientation of nanocrystals.The surface reconstruction leads to variation in SrTiO3 surfaceenergy and the interface energy and, thus, can be used tomanipulate the growth of metal nanocrystals [393,394].

Epitaxy: The most often observed epitaxial orientationrelationships for fcc and bcc metals on SrTiO3(100) are OR Iand OR II, respectively (see Table 6):

OR I, “cube-on-cube” : (100)SrTiO3 ‖ (100)Me,

[001]SrTiO3 ‖ [001]Me

OR II, “cube-on-cube with 45◦orientation” :(100)SrTiO3 ‖ (100)Me, [001]SrTiO3 ‖ [011]Me.

The OR I and OR II are mainly favored by their low interfaceenergies because the (100) surface of an fcc or bcc metal is notthe surface with the lowest surface energy. DFT calculationshave shown that Pt [376], Pd [377,378], and Co [379] atomstend to register on top of O atoms on the TiO2-terminatedSrTiO3(100) surface. Fig. 35 shows that OR I results in ahigh density of near coincident sites for fcc metal atoms onan O sub-lattice. Therefore, this orientation results in lowinterface energies. The same argument applies for bcc metals onSrTiO3(100). Here, metals such as Cr and Mo [258,425,426],try to bond with surface O. Formation of OR II is also promotedby a high density of near coincident sites for bcc metal atomson O (Fig. 35). In addition, OR I and OR II are favored by thesmall misfits.

As discussed above, the low interface energy in the casesof OR I and OR II relies on the registry of metal adatoms onthe O sub-lattice on the TiO2-terminated SrTiO3(100) surface.If SrTiO3(100) is SrO-terminated, or there exists an interfaceoxide layer between metal overlayers and the SrTiO3 surface,the interface energy may be increased to the point that the (100)orientations will not be dominant in the metal epitaxy. In sucha case, (110) orientations (OR III and OR IV) are observed forbcc metals and (111) orientations (OR V) for fcc metals:

OR III : (100)SrTiO3 ‖ (110)bcc-Me, [001]SrTiO3 ‖ [001]bcc-Me,

OR IV : (100)SrTiO3 ‖ (110)bcc-Me, [110]SrTiO3 ‖ [111]bcc-Me,

OR V : (100)SrTiO3 ‖ (111)fcc-Me, [001]SrTiO3 ‖ [011]fcc-Me.

For example, growth of Cr on SrTiO3(100) above 700 ◦Cresults in the formation of an interfacial CrOx layer, which


causes the appearance of OR II and (110) Cr with OR III [237].A similar result was observed by Francis et al. [432]. Theyfound that purely epitaxial Cu(100) films were obtained on aclean SrTiO3(100) while a mixture of Cu(100) and Cu(111)domains was observed in the presence of a small amountof copper oxide at the interface. The presence of surfaceterminated SrO layer could also lead to orientations other thanthe (100) orientations. Ono et al. found that Fe grows onthe flat TiO2-terminated SrTiO3(100) with OR II. However,both (100) Fe (OR II) and (110) Fe (OR IV) were observedon the as-polished SrTiO3 surface, which consists of bothTiO2- and SrO-terminations [427]. Polli et al. observed that Ptpresents OR I on purely TiO2-terminated SrTiO3(100) whilePt(100) (OR I) and Pt(111) (OR V) have been recorded on theSrO-terminated surface [251]. The epitaxial growth of metaloverlayers on SrTiO3(100) can be controlled by modificationof the SrTiO3(100) surface. This concept is applicable for mostoxide surfaces.

4.2.3.3. Chemical interaction of metals on SrTiO3. The mostimportant interface reaction between metals and SrTiO3 ismetal oxidation. Metals of Y, Ba, Ti, Al, Si, Cr, Fe, and Mo,can be oxidized on SrTiO3 surfaces [85,89,91,423,424]. Asimple thermodynamic criterion for metal oxidation on SrTiO3surfaces is 1Ho

f < −250 kJ/mol O. The heats of oxideformation per mole of oxygen (1Ho

f in kJ/mol O) were takenfrom Ref. [23].

Hill et al. found that Y, Ba, and Ti reacted with SrTiO3.At 300 K, the reaction was diffusion limited. At elevatedtemperatures, extended out-diffusion of O was activated andthe metal overlayers could be fully oxidized [424]. In order tostudy the reaction kinetics in more detail, Fu et al. performeda systematic study of the thermal stability of ultrathin films ofAl, Cr, Fe, and Mo on the SrTiO3(100) surface [89]. Al, Cr,Fe, and Mo films of a thickness of 6 A were deposited ontoidentical SrTiO3(100) surfaces. The interface reactivity wasstudied by stepwise UHV annealing combined with in situ XPSmeasurements. The determined reaction onset temperatures(TRO) were plotted as a function of the metal work function φM .The figure is reproduced in Fig. 36 showing close correlationwith EF of metals (see Section 2.4.3.).

Encapsulation reactions were seldom observed for metalson SrTiO3 surfaces. Silly and Castell recently reported theencapsulation of Pd nanocrystals on SrTiO3(100) [435] wherethe SrTiO3 substrate has been subjected to extended UHVannealing and the surface is covered by anatase TiO2(100)islands. On the strongly reduced SrTiO3 surface Ti and Oatoms are mobile enough at elevated temperatures such that Pdclusters were encapsulated by TiO1.36 or TiO1.37 layers uponUHV annealing. Therefore, the reaction can be regarded as Pdencapsulation on TiO2 islands supported on the SrTiO3(100)surface.

5. Interaction of metals with insulating oxides

Al2O3, MgO, and SiO2 crystals are excellent insulators.They have large band gap energies (Eg), i.e., Eg(Al2O3) =

Fig. 35. Schematics of the interface between a TiO2-terminated SrTiO3(100)and a bcc metal overlayer (a) and the interface between a TiO2-terminatedSrTiO3(100) and an fcc metal overlayer (b) with OR I and OR II, respectively(top view). From [108].

8.8 eV, Eg(MgO) = 7.7 eV, and Eg(SiO2) = 9 eV [17,41]. Electrons in these oxides are strongly localized andthe production and diffusion of ionic defects in the oxidesare limited. Therefore, long-range charge transfer and iontransport do not occur in these oxides at relatively lowtemperature (e.g., <1000 ◦C). Metal interactions with theseinsulating oxides are often confined to the interfaces, whichonly involve oxide surface atoms and metal adatoms incontact with the substrate surface. The interactions are stronglydependent on the surface properties of oxide substrates.Surface stoichiometry, surface terminations, and surface defectsare the most important factors influencing the metal–oxideinteractions.

In order to circumvent the charging problems encounteredon surfaces of single crystal insulating oxides, thin oxidefilms have been grown on conductive supports which were


Fig. 36. Plot of TRO (the temperature point where oxidation of metal overlayeris first observed by XPS) for Al, Cr, Fe, and Mo metal films on identicalSrTiO3(100) and TiO2(110) substrates as a function of φM [89].

widely used as model systems [159]. For these systems ofmetal overlayer/oxide film/conductive support, both the oxidefilm surface and the support are critical in the metal–oxideinteractions.

Interactions between metals and insulating oxides (Al2O3,MgO, and SiO2) are reviewed in this chapter. We show howinteractions vary with oxide surface properties and how thesupports of the thin oxide films influence the metal interactionson the top of the oxide film surfaces.

5.1. Metals on Al2O3

Alumina (Al2O3) is widely used as a catalyst support,for structural ceramics, and as a substrate for film growth.Among the various polymorphs of alumina, corundum, α-Al2O3, is the most stable phase and has been subjected toextensive studies [26,76]. In particular, the crystallographicallysimple and energetically stable α-Al2O3(0001) surface offersa good playground for understanding of insulating oxidesurfaces and metal–Al2O3 interactions. In the following,details of the surface properties of α-Al2O3(0001), whichinclude surface terminations, surface reconstructions, andsurface hydroxylation, are given. The effects of these Al2O3surface properties and metal overlayers on the metal–Al2O3interactions are discussed in detail. In addition to the singlecrystal Al2O3 surfaces, ultrathin alumina layers grown onconductive substrates have been widely used as model systemsfor supported metal overlayers. In the subsequent part, metalinteractions with alumina films are reviewed.

5.1.1. α-Al2O3 surfacesAlumina surfaces are often prepared by mechanical

polishing, ion sputtering, and annealing. Naturally, differentsurface preparation processes result in various surfaceproperties, which, in turn, may cause different behaviorsof metal growth on these surfaces. Here, we discuss thesurface terminations and reconstructions as well as the surfacehydroxylation of the α-Al2O3 surfaces.

5.1.1.1. Surface terminations and reconstructions. The unitcell of bulk α-Al2O3 can be described as a hexagonal unitcell containing six formula units of Al2O3 [446]. This unitcell consists of six close-packed hexagonal O layers. Allayers, which are not coplanar but buckled, are placed betweenthese O layers. All of the ions are stacked along the c axisof the unit cell in a sequence R-AlAlO3-R (R: continuingsequence in the bulk). For the bulk truncated and clean α-Al2O3(0001) surface there exist, from a geometrical standpoint,three different terminations [25,156]: O layer termination(O3AlAl-R), single Al layer termination (AlO3Al-R), anddouble Al layer termination (AlAlO3-R). Theoretical worksuggests that the 3 surfaces possess different thermodynamicstabilities [156,447–449]. The different surface terminationshave been experimentally observed by different groups(Table 7).

O layer termination: The O layer terminated surface hasa large surface dipole moment and surface dangling bondswhich cause the surface to be energetically unstable under mostenvironmental conditions [156]. The O layer termination wasobserved experimentally only by Toofan et al., who reporteda mixture of 2:1 O/Al-terminated surface domains [450].Presence of H or adsorption of a reactive metal can stabilizethe O layer terminated surface.

Single Al layer termination: The single Al layer terminatedsurface is generally accepted to be the most stable α-Al2O3(0001)-(1 × 1) surface [25,67,156,447–453,468–473]because this surface is non-polar. The surface Al atoms canrelax inwards so that they are almost coplanar with respect tothe second O layer (see Fig. 37). The relaxation is accompaniedby a rehybridization of surface Al atoms to an sp2 orbitalconfiguration, which significantly stabilizes the surfaces viacharge autocompensation [468]. The surface relaxation wascalculated to be about −85% [67,156,448,469,471–473]. Theexperimentally observed relaxations were smaller, rangingfrom −51% to −63% [25,451,452,474].

Double Al layer termination: The (1 × 1) surface may besubjected to reconstructions in cases of surface O desorptionor Al deposition onto the surface [147–150,466]. XRD andLEED investigations have revealed a (

√31 ×

√31)R ± 9◦

reconstruction on α-Al2O3(0001) surfaces which were heatedin UHV at T > 1200 ◦C [148–150] or covered by Al [150,466]. The Al-rich reconstructed surface has been confirmed tobe terminated by a double Al layer, which contains surfacedomains with a Al(111) structure separated by hexagonalnetwork of domain walls [147,148,448]. Such a structure hasbeen directly imaged using dynamic-mode SFM by Barth andReichling [147].

In addition, disordered Al-rich surface phases were obtainedthrough ion sputtering, which preferentially removes surfaceoxygen [464,465,475].

Many research results have shown that the sapphire surfaceterminations and reconstructions have a strong effect on thenature of bonds at metal/Al2O3 interfaces and the morphologyof the supported metal clusters.

An extensively investigated system is the Cu/α-Al2O3(0001)interface. Gautier et al. [476] found that the size of Cu clusters


Table 7Selected experimental data showing the surface terminations of α-Al2O3(0001)

Termination Surface treatments Surface structure Techniques

O layer termination Heating in O2 plasma T < 750 ◦C: (1× 1); T < 1000 ◦C :(√

3×√

3)-R30◦; T > 1100 ◦C :(√

13×√

13)-R ± 9◦; Mixture of O-and Al-terminated domains

LEED [450]

Single Al layer termination UHV annealing up to 1100 ◦C; (1× 1); 63% inward relaxation of topAl layer; H presence on surface evenafter annealing at 1100 ◦C

TOF-SARS, LEED [451]

Annealing in air at 1500 ◦C, 3 h (1× 1); 51% inward relaxation of topAl layer

CTR diffraction [25]

Heating at 650 ◦C using atomicdeuterium beam

(1× 1); 51% inward relaxation of topAl layer

LEED [452]

Heating in O2 at 1570 K for 10 hfollowed by UHV heating at 870 Kfor 10 min

(1× 1) CAICISS, RHEED [453]

Double Al layer termination UHV heating, O2 annealing, Aldeposition, or Si etching

UHV, T > 1250 ◦C : (√

13×√

13);Excess Al deposition at 800 ◦C:(√

13×√

13); O2 annealing at1000–1200 ◦C: (1× 1);

√13×

√13:

cubic layer with composition of Al2Oor AlO

LEED [150]

UHV, 1350 ◦C for 20 min (√

13×√

13)-R ± 9◦: two Al planesclose to metallic Al(111)

GIXD [148,149]

UHV, 1300 ◦C (√

13×√

13)-R ± 9◦: hexagonal(111) Al surface domains withhexagonal domain walls

SFM [147]

OH layer termination Exposure to 10−4 Pa H2O Formation of surface OH HREELS, XPS [454,455]

Exposure to H2O (1× 1); Chemisorption of H2O LITO, TPD, LEED [456,457]

Surface in ambient conditions Fully hydrated surface with Otermination; 2.3 A up disordered Olayer from adsorbed water

CTR diffraction [458]

Exposure of Al-terminated surfaceto >1 Torr water followed byoxygen plasma at RT

1 ML coverage of surface OH XPS [459]

Exposure of clean surface to waterdrops followed by oxygen plasmaat RT

0.5 ± 0.1 ML coverage of surface OH XPS [460]

Exposure to water vapor Formation of surface OH XPS [461–463]

Other surface phases 1 keV Ar bombardment 3 nm γ -Al2O3 layer with highdensity defects

TEM [464]

UHV heating or ion sputtering Surface Al-rich phases; Withincreasing T : change from (2× 2),(3√

3× 3√

3)R30◦, to(√

13×√

13)R ± 9◦

LEED, XPS, EELS [465]

Al deposition (√

13×√

13)R ± 9◦: Al-rich surfacephases

LEED, EELS, AES [466,467]

is larger on a reconstructed (√

31×√

31)R±9◦ surface than thaton a (1 × 1) surface. The same conclusion was made by Gotaet al. [477] based on SEXAFS studies during the initial stage ofCu growth on the two surfaces. It is concluded that the differentmorphologies of Cu clusters on the two surfaces originate fromthe different bonding strength of Cu with the Al2O3 surface.Stronger adhesion of Cu on the (1 × 1)Al2O3(0001) surfacefavors the formation of higher density of Cu clusters atop.

Scheu [475] has shown that bonding at a Cu/(0001)Al2O3interface depends on the substrate preparation and that thebonding can be manipulated by different substrate cleaningprocesses. As discussed above, the (1×1) Al2O3(0001) surfaceis terminated by a single Al layer. On this surface strong inwardrelaxation of the top Al layer results in that the second Olayer is almost coplanar with the topmost Al layer (Fig. 37).With surface O available, Cu-adsorption onto the surface results


Fig. 37. Atomic layer sequence and layer spacings along the [0001] directionfor single Al-terminated α-Al2O3 surface model determined by Guenardet al. [474]. Small grey balls represent Al, and big red balls represent O.Note that Al and O are almost coplanar at the topmost surface. Reproducedfrom [458].

in the formation of Cu–O bonds. A combined approach ofHRTEM and ELNES reveals the existence of Cu–O bonding atthe internal MBE-grown Cu–Al2O3 interface. The bonding isof mixed ionic–covalent character with substantial Cu3d–O2phybridization similar to that in Cu2O [257,475,478]. Cu-O bondformation for copper growth on the (1 × 1) surfaces has alsobeen observed by Guo et al. [479] and Varma et al. [480]. Incontrast to the (1 × 1)Al2O3(0001) surface, the reconstructedsurfaces or Ar+-sputtered surfaces are surface Al-rich. On thesesurfaces, Cu atoms tend to bond with surface Al atoms, whichresults in formation of interfacial metallic Cu–Al bonding.Scheu has observed Cu–Al bonds at the Cu/(0001)Al2O3interface that had been Ar+ sputter cleaned [475] and Gotaet al. confirmed that Cu clusters are bonded to the Ar+ sputter-cleaned (

√31×

√31)R ± 9◦ surface via Al atoms [477].

DFT calculations of Cu and Pd interactions with α-Al2O3(0001) demonstrate that the interaction mechanismdepends on the surface stoichiometry [481]. The interaction iscovalent-like for the metals on the Al-rich surface but ionic-like for the metals on the oxygen terminated surface. On thestoichiometric surface the interaction is very weak and mainlydue to polarization effects. The first principle studies fromZhang et al. [482,483] suggest that the interface energy ofmetal/Al2O3 is a function of the substrate surface stoichiometryand oxygen chemical potential. The most stable structure ofAl/Al2O3 and Nb/Al2O3 interfaces is reached for metals onthe oxygen terminated surface as long as the oxygen chemicalpotential is above a critical value. At the Ag/Al2O3 interface, Altermination results in the formation of a stable interface even atrelatively high oxygen potentials.

5.1.1.2. Surface hydroxylation. Many theoretical results showthat hydroxylation of the clean α-Al2O3(0001) surfaces mayresult in a further lowering of the energy of these surfaces [156,448,449,484,485]. Thus, the above-mentioned O- and Al-terminated surfaces are expected to be reactive to water. Abinitio calculations revealed that molecularly adsorbed wateron Al-terminated surfaces is metastable and can dissociatereadily. The H2O dissociative reactions produce two typesof surface OH groups: OadsH and OsH (Oads: water oxygen;Os: surface oxygen) [471,472,486]. The OH-terminated α-Al2O3(0001) surfaces have been experimentally observed byvarious techniques, e.g., SFM [147], XRD [458], thermaldesorption [456,457], EELS [454,487], and PES [455,459–463,488,489]. These OH-terminated surfaces were simply obtained

Fig. 38. O1s spectra measured at a 10◦ take-off angle for hydroxylated α-Al2O3 surface (a), and the same after 0.8 ML Co deposition (b). From theintensity ratio of OH to lattice oxygen O1s, the coverage of surface OH on thehydroxylated surface was calculated to be 1 ML. Co deposition removed 0.4ML surface OH groups. Reproduced from [459].

via exposure of the Al2O3 surfaces to water or air, see Table 7.The coverage of surface OH groups can be up to 1 ML.For example, Fig. 38 shows that the OH coverage on ahydroxylated Al2O3 surface is about 1 ML [459]. The surfaceOH groups are very sensitive to electron beam illumination,Ar+ sputtering, UHV heating, and adsorption of reactivemetals. For example, exposure of a hydroxylated Al2O3 surfaceto a large electron beam dose or high energy ion sputteringresulted in dehydroxylation of the surface [454,460,462,487].High temperature annealing or Al deposition can transform ahydroxylated surface to an Al-terminated surface [460,490].

Surface hydroxylation is of special importance for the α-Al2O3 surfaces. Alumina surfaces are quite often covered bywater or exposed to air during handling. Therefore, OH groupsare always present on the surfaces in case of no further specialsurface treatments. Furthermore, it is shown below that surfaceOH groups exert a significant influence on the metal interaction.

Cu/α-Al2O3(0001) interface is a well-studied system toillustrate the effect of OH on the interface formation. Kelberand coworkers applied XPS and DFT calculation to study Cuinteractions with hydroxylated α-Al2O3(0001) surfaces [461,462]. It was shown that the presence of surface OH leads to theformation of a Cu (I) monolayer up to 1/3 ML coverage. Inagreement with this result, a recent DFT calculation indicatesthat at Cu coverages below 1/3 ML Cu atom adsorbed onto aOH-terminated Al2O3 surface can remove the surface H andbind the surface through the surface oxygen [491]. The stronginteraction between Cu and the hydroxylated surface enhancesthe wetting of Cu which stabilizes Cu(I) adatoms in 2D islands.The strong interaction may originate from an exothermicreaction of surf-OH + Cu (gas) → 1

2 H2 + surf-O–Cu [462,491]. However, Wang et al. [492,493] and Lodziana et al. [481]


presented a quite different conclusion, suggesting weakinteraction between Cu and the surface OH. The interfacialOH can be stable even in the presence of two ML Cu [492,493]. Similar to the theoretical results, XPS experimentaldata do not support any strong chemical reactions betweensurface OH and Cu adatoms [460]. Lodziana et al. [481] havesuggested that surface defects which are introduced by thehydroxylation, rather than surface OH groups, are responsiblefor the experimentally observed enhancement of Cu interaction.

Experimental and theoretical work has shown that othersmetals, such as Al [455,460,494], Ti [455,488], Rh [489], andCo [459,460], can react with the surface OH to form oxidizedmetal overlayers at submonolayer coverage. For example,Chambers et al. [459] showed that Co deposition on fullyhydroxylated sapphire surfaces is accompanied by oxidation toCo2+ and removal of OH (Fig. 38). The interface reactions lead,instead of island growth, to technologically important 2D metalfilms growth on oxides.

5.1.2. Metal interactions with bulk Al2O3

Both electronic interactions and chemical reactions atmetal–Al2O3 interfaces are reviewed in detail here.

5.1.2.1. Electronic interaction of metals on Al2O3. Theinteraction of metals with Al2O3 surfaces presents a highdegree of complexity. Nevertheless, many theoretical andexperimental efforts have contributed to the understanding ofthe nature of interactions between metals and Al2O3 surfaces.Different interaction mechanisms have been proposed, such asvan der Waals force, interfacial bonding (covalent or/and ionic),and polarization. Both theoretical and experimental results arediscussed separately.

(1) Theoretical results: Johnson and Pepper [495] firstused a cluster model and molecular orbital theory to studythe metal–sapphire interfacial strength. They concluded that adirect chemical bond, which is mainly covalent, is establishedbetween metal atoms (Fe, Ni, Cu and Ag) and oxygen anions.The covalent interaction occurs through the hybridization ofmetal d orbitals and non-bonding O2p orbitals which producebonding states and antibonding states. An ionic componentis associated with metal-to-oxygen charge transfer at theinterface. The increasing occupation of antibonding orbitalsand decreasing metal-to-oxygen charge transfer explain thereduction in metal–sapphire contact strength through the seriesFe, Ni, Cu and Ag. This picture was confirmed by Nath andAnderson [496].

Alemany et al. [75] studied adhesion of 3d transition metalson α-Al2O3 surfaces using extended Huckel tight-binding bandstructure calculations. However, they found that two O–M andAl–M interactions are responsible for the adhesion strength.O–M repulsive closed-shell interactions are a destabilizingfactor while Al–M charge-transfer interactions favor interfaceformation.

Verdozzi et al. [67] have investigated Pt- and Ag-adsorptionon Al-terminated α-Al2O3(0001) using the local densityapproximation (LDA) and thick slabs (up to 18 O layers) intheir calculations. The nature of the oxide–metal bond was

Fig. 39. Adsorbate-induced charge density difference plot for Nb, Ru, andPd. Solid lines indicate charge accumulation, and dashed lines depletion, withlogarithmic increments. The 〈100〉 cut goes through the center of the adsorbatesand oxygen ions. For metals to the right of Mo in the periodic table, e.g., Ru andPd, charge transfer occurs from d2

Z orbitals to lateral d orbitals upon adsorptionon surface O. Metals to the left of Mo, e.g., Nb, present less than a half-full d shell, and the oxygen polarization goes from lateral d orbitals to d2

Zorbitals [66].

found to change drastically with the metal coverage. At 1ML both Pt and Ag prefer atop-Al sites and no evidence ofsignificant charge transfer occurs between metal overlayers andthe oxide. The BE is dominated by the polarization effect. Incontrast to the case of 1 ML, isolated metal atoms at 1/3 MLcoverage bind the surface up to 5 times stronger through largelyionic interactions. This result was generalized by Bogicevicand Jennison in a study in adsorption of various metals (Li,K, Y, Nb, Ru, Pd, Pt, Cu, Ag, Au, Al) atop 5 A Al2O3-films on Al(111) [66]. At 1/3 ML, the oxide–metal bond isionic, regardless of metal adsorbate. Metal atoms, attractedto the O2− ions, prefer to bind in the three-fold hollow site.The charge is transferred from metal atoms to the nearestneighboring oxygen ions. The degree of ionicity depends onthe metal Pauling electronegativity and metal ion radius. At 1ML coverage, the adhesion is almost purely electrostatic. Themetal overlayer is attracted to the O ions at the oxide surface bylateral polarization as illustrated in Fig. 39.

It has been shown that the theoretical results of metalinteractions on alumina surfaces rely on the choice ofthe surface termination (Al-terminated or O-terminated),calculation models (cluster or slab), metal coverage (isolatedatom or metal film), and choice of computational methodand approximations used. Nevertheless, a qualitative image ofmetal–Al2O3 interactions can be obtained based on the resultsavailable. The interactions can be qualitatively interpreted onthe basis of the metal’s Pauling electronegativity.

The bonding between the metals with small Paulingelectronegativity and Al2O3 is mainly ionic. As one example, atNb(111)/α-Al2O3(0001) interface, strong ionic bonds form byNb4d → O2p electron donation, which accounts for the highadhesive strength of O-terminated interfaces [497]. Ab initiocalculations of bonding at Al(111)/α-Al2O3(0001) interfacealso indicate that Al–O bonds constitute the primary interfacialinteraction. The bonds are very similar to the cation–anionbonds found in Al2O3 bulk and are mainly ionic [473].

The interactions of the metals with large Pauling elec-tronegativity with alumina surfaces are, however, dominatedby the polarization effect, in particular, on Al-terminated sur-faces and at high coverage, e.g., 1 ML. For example, at Pd/α-


Al2O3(0001) interfaces, the nature of the interaction has beeninvestigated by Gomes et al. [498,499]. They concluded that thePd-surface interaction comes from metal polarization, whichis caused by the surface electrostatic field. The same mecha-nism controls interactions of Pt and Au with α-Al2O3(0001)surfaces [66,67,500].

For metals with Pauling electronegativity in between e.g., Cuand Ag, either charge transfer or metal polarization is claimedto be responsible for the interactions [66,500,501].

(2) Experimental results: Adsorption of reactive metals ontoα-Al2O3(0001) is often accompanied by oxidation of the metalsin the submonolayer coverage regime. Oxidation of Nb [502],V [503], and Ti [504] on α-Al2O3(0001) was observed usingPES. The formation of M–O bonds is accompanied by chargetransfer from metal adatoms to the substrate surface. However,reduction of surface Al3+ was not often observed. Bieneret al. [503] argued that a non-localized charge redistribution onthe substrate surface instead of the localized charge transfer tothe substrate atoms, as observed at the V/TiO2 interface [348]seems to account for the oxidation of V atoms at the V/α-Al2O3(0001) interface.

In the case of non-reactive metals, such as Pd, on Al2O3surfaces charge transfer, if any, happens from the oxide tothe metals. For example, Ogawa and Ichikawa applied theKelvin probe technique to monitor changes in surface workfunction during Pd clusters growth on α-Al2O3(0001). Theirresults confirmed that the charge transfer occurs from Al2O3to Pd at low Pd coverage [505]. Gillet and coworkers studiedPd deposition on α-Al2O3 (1012). Analysis of the modifiedAuger parameter of Pd indicates that formation of interfacialPd–Al bonding promotes an electronic transfer from the oxideto Pd [506–508].

In conclusion, both theoretical and experimental resultsdemonstrate that the electronic interactions between metals andAl2O3 surfaces are dominated by either interfacial bonding ormetal polarization effect. The local electronic interactions arestrongly dependent on metal electronegativity. However, long-range interaction, such as space charge transfer, which wasobserved at metal/TiO2 and metal/SrTiO3 interfaces, has beenrarely observed.

5.1.2.2. Chemical interaction of metals on Al2O3 crystals.Various surface science studies have confirmed that oxidationreactions can happen between metals and α-Al2O3(0001)surfaces near room temperature. The metals include Al [455,460,466,467], Ti [232,455,488,504], Nb [502], V [503], andCu [462,480]. Almost exclusively, the oxidation reactionsare strictly limited to the interfaces. Subsequently, metallicoverlayers will develop with metal coverage above 1 ML.HRTEM was applied to investigate the interfaces of Al/α-Al2O3 [509], Ti/α-Al2O3 [510], Cr/α-Al2O3 [88,511], andCu/α-Al2O3 [257,510], which were all formed below 600 ◦C.The results clearly show that all the interfaces are atomicallysharp and no interface reaction phases thicker than a monolayerhave been observed.

The results indicate that only those metal adatoms rightat the interface, which are in contact with surface oxygen,

became oxidized. Alumina has a high thermodynamic stability(1Ho

f ∼ −600 kJ mol−1 O) and oxygen diffusion in the crystalis highly limited. The oxidation of metal multilayers may beeither thermodynamically impossible or/and kinetically limited.

5.1.3. Metal interactions with alumina filmsThe main problem of Al2O3 single crystals in surface studies

is the surface charging effect, which prevents the use of surfacetechniques that involve the emission or adsorption of chargedparticles. In order to circumvent this drawback ultrathinalumina layers have been grown on conductive substrates. Theoxide film-based model systems provide a good playground forsurface analysis of alumina surfaces and studies in interactionsbetween metals and alumina films.

5.1.3.1. Growth of alumina films. Well-ordered alumina thinfilms have been grown on different Ni/Al alloy surfaces.For example, epitaxial Al2O3-films on NiAl(110) show ahigh degree of crystallinity, very low surface roughness,and good preparation reproducibility [173,512,513]. Thus,the Al2O3/NiAl(110) system has been extensively used asmodel oxide surfaces [26,27]. Ultrathin alumina films can beobtained by oxidizing Ni3Al(111) [514,515] and NiAl(111)surfaces [516]. Oxidation of polycrystalline Al foils or singlecrystal Al surfaces also produces alumina films [171,172,517,518].

5.1.3.2. Electronic interaction of metals on alumina films. Inthe systems of metal-overlayer/oxide-film/conductive-support,both the oxide film surface and the support are critical inthe metal–oxide interactions. Here, we demonstrate how thesurfaces of the grown alumina films and the supports of thealumina films affect metal–alumina interactions.

Metal interactions with alumina films are strongly influencedby surface properties of the thin oxide layers. The orderedAl2O3-films grown on NiAl(110) have been extensively studiedby STM and surface X-ray diffraction (SXRD). The surface isfound to be terminated by a layer of O2−-ions [173,513,519].Therefore, adsorption of reactive metals, such as Al, V, and Cr,tend to build interfacial M–O bonds such that charge transferoccurs from the metal adatoms to surface oxygen, which resultsin the oxidation of the adsorbed metals [109,489,520]. Likebulk Al2O3 surfaces, surfaces of thin alumina layers can also behydroxylated by exposing the O-terminated surfaces to water.The surface OH groups tend to increase the dispersion of metaloverlayers on the hydroxylated surface. For example, STMexperiments on Rh and Pd deposition on thin alumina layersshowed that island density of the metals on the hydroxylatedsurfaces is higher than that on the clean surfaces [489,521,522].

In oxide film-based model systems the effect of the metalsupports has to be taken into consideration. If the thicknessof the supported oxide layer is very small, e.g. less than1 nm, electron tunneling happens between the underlying metalsubstrates and metal adsorbates [115,116,119]. In such a case,the support can have a strong effect on metal interactions onthe oxide surfaces. For example, a recent STM study in Au-adsorption on a thin alumina film (5 A) grown on NiAl(110)


Fig. 40. (a)–(b) STM images of Au pentamers on an alumina film on NiAl(110), which are adsorbed in A and B domains, respectively. The anglesbetween the chain (dotted lines) and the NiAl [001] direction (straight lines)are indicated. (c)–(d) Schemes showing the orientation of the hexagonal Alslattice with respect to the NiAl substrate. From [523].

demonstrated that 1D Au chains exhibit a preferentialorientation close to [001] direction of the underneath NiAl(100)(Fig. 40) [523]. The results unambiguously indicate theparticipation of the NiAl substrate in the Au binding on thealumina surface. Furthermore, many PES studies confirmed thatsupports may change the electronic structure of metal–aluminainterfaces. The support effects mainly originate from the chargetransfer between the metal overlayers and the underlying metalsupports. Depending on the EF of the metal overlayer and thesupport, two different charge transfer processes can occur in themetal-overlayer/oxide-film/conductive-supportsystem.

Deposition of Cs onto ultrathin alumina films supported onMo(110) induced large (0.9–1.1 eV) positive shifts in the BEof O KVV, O1s, Al2p feature of alumina film [524]. 1 A Aldeposition on an alumina layer grown on NiAl(110) resultedin a shift of the oxide component in Al2p, O1s, and O2p tohigher BE, 0.47 ± 0.03 eV [489]. A similar shift to higher BEin the range of 0.5 eV was detected during depositing V ontothe same oxide surface [109]. Considering that Cs, Al, and Vhave low work functions, a charge transfer is expected to takeplace from Mo to Cs [524], from NiAl to Al [489], and fromNiAl to V [109]. Consequently, downward band bending of theoxide layer may happen, which results in positive BE shiftsof the oxide components. The case of EF (metal overlayer) >

EF (support) is illustrated schematically in Fig. 41(a).Deposition of large work function metals on the supported

alumina films can reverse the charge transfer direction. Forexample, Sarapatka observed parallel negative shifts of oxideAl2p and O1s in the case of Ni or Pd deposition onto aluminathin layers formed by surface oxidation of polycrystallineAl foil [171,172]. Core level spectra and Auger parametersrecorded from Ni or Pd overlayers suggested negative chargingof the metal particles, which is consistent with a chargetransfer from the Al support. In another example, smallnegative BE shifts (around −0.1 eV) are also observed for

Rh-, Pd-, and Pt-adsorption on an alumina layer grown onNiAl(110) [489]. The electronic configuration in the systems ofNi/Al2O3/Al, Pd/Al2O3/Al, and Rh (Pd, Pt)/Al2O3/NiAl(110)can now be described by schematic shown in Fig. 41(b). Incontrast to Fig. 41(a), charge transfer takes place from metalsupports to metal overlayers because of EF (metal overlayer) <

EF (support), which results in downward band bending andnegative BE shifts in the oxides.

5.1.3.3. Chemical interaction of metals on alumina films. Inoxide film-based model systems, the defect structure of thinoxide layers provides some channels for metal migration intometal supports such that alloy formation may take place beneaththe oxide layers. Interdiffusion, mainly metal diffusion throughoxide layers into the metallic substrate, can limit the thermalstability and therefore is very critical at elevated temperatures.The thermal stability of Co, Rh, and Pd particles on a thinalumina layer supported on NiAl(110) was systematicallystudied as a function of annealing temperature [26,522].Annealing the surface up to 900 K almost completely removedall metal particles (see Fig. 42). STM experiments proved thatthe diffusion process is defect-mediated. Loss of material wasmainly observed in the vicinity of antiphase domain boundaries,which act as an important diffusion channel. Chen et al. [518]used EELS and AES to study the thermal behaviors of Nioverlayers deposited on Al2O3-films, which were preparedby oxidizing an atomically clean Al(111) surface. Annealingthe metal/oxide systems up to 700 K causes a strong inwarddiffusion of Ni overlayers. It was concluded that the metallicaluminum substrate provides a major driving force for theinward Ni diffusion by Ni–Al alloy formation. A similarresult was observed in the Cu/Al2O3/Al(111) system [525].In Pd/Al2O3/Al and Ni/Al2O3/Al systems, Sarapatka foundinward diffusion of Pd and Ni through alumina oxide layersand subsequent formation of PdAl and NiAl alloys under theoxides [171,172]. The interdiffusion of Pd and Ni can beretarded by increasing the thickness of the Al2O3 layer.

5.2. Metals on MgO

MgO has high chemical, thermal, and mechanical stabilitywhich allows it to be widely used as a substrate for epitaxialgrowth of metal films and as a catalytic support. Moreover,the MgO(100) surface is non-polar and experimentally well-characterized, making it a good model system for theoreticalstudies of insulating oxide surfaces.

5.2.1. Metal interactions with bulk MgOWe start with the most stable surface, MgO(100), which is

often subjected to metal adsorption or deposition. Two criticalfactors including MgO(100) surface properties and metaloverlayers in the metal–MgO(100) interactions are discussedin detail.

5.2.1.1. MgO(100) surfaces. The single crystal MgO canbe easily cleaved along the {100} planes, so well-formedMgO(100) surfaces are often obtained by cleavage in UHV or


Fig. 41. Flat energy band models of metal overlayer/ultrathin oxide layer/metal substrate systems. (a) metal overlayers have lower work function (φ) compared tothat of metal substrate; (b) metal overlayers have larger work function (φ) than that of metal substrate.

Fig. 42. Particle density as a function of the annealing temperature of Rh, Pd,and Co deposits on alumina layers grown on NiAl(110). The decrease in islanddensity is from metal migration into metal substrates. Co shows higher stabilitydue to its strong metal support interaction. Reproduced from [522].

air [526–528]. The surface reconstruction and point defects arestrongly dependent on the surface preparation condition.

Most of the well-prepared MgO(100) surfaces are unrecon-structed. However, surface defects may induce reconstructions.For example, a (

√2 ×√

2) surface structure was observedby helium atom scattering (HAS) measurements on cleavedMgO(100) single crystals [529]. The reconstruction is due tothe presence of a large compressive surface stress and a certainamount of surface defects. Ca segregation onto MgO(100) sur-faces leads to another (

√2 ×√

2)R45◦ reconstruction [528].On the well-defined (1 × 1)MgO(100) surfaces, both surfacerelaxation and surface rumpling are very small ([528] and ref-erences therein). For example, Robach et al. reported a relax-ation of −0.56%± 0.35% and a rumpling of 1.07%± 0.5% onMgO(100) which was determined by grazing incidence X-rayscattering (GIXS) [528].

Typically, MgO(100) surfaces exhibit defects, includingextended defects and point defects. The extended defects aresteps, line defects attributable to missing rows of Mg2+ andO2− ions, rectangular holes of nanometer size originatingfrom cleavage, complex adstructures due to adatoms, etc.All these surface features have been imaged by dynamic-

mode SFM [249] and NC-AFM [530]. The point defectscan be categorized into four groups: low coordinated sites,surface vacancies, divacancies, and impurity atoms. The lowcoordinated sites are four-coordinated ions located at step andedge sites and three-coordinated ions located at corners, kinks,etc. A surface oxygen vacancy is usually called as a surfaceF (Fs) center. Divacancies are created by removing a neutralMgO unit. MgO surface point defects have been reviewedby Pacchioni [77]. Single point defects on flat terraces ofMgO(100) were imaged by dynamic-mode SFM at atomicresolution (Fig. 43) [249]. The density of point defects isestimated to be roughly 1012–1013 defects/cm2. It has beenshown that surface defects strongly influence the nucleationand growth of metals on MgO(100) surface. Didier and Jupillefound that the growth mode of Ag on MgO(100) stronglydepends on the chemical purity of the oxide surface. On acarbon- and defect-free MgO(100) surface, Ag grows in a 2Dmode while the presence of any surface defects prevents sucha growth mode [531]. The nucleation of Au and Pd on cleavedMgO(100) surfaces has been studied by AFM, which indicatesthat nucleation is controlled by surface point defects [136,532].

5.2.1.2. Metal–MgO interactions. Many experiments havebeen performed to study the interface formation between metalsand MgO(100). Further theoretical work contributed much tothe understanding of these interactions. The following discussesthe interaction of various metals with MgO(100) surfaces,including both experimental and theoretical results.

(1) Experimental results: Using GIXS Renaud andcoworkers have systematically investigated the structure andmorphology of Ag [533,534], Pd [535], and Ni [536] filmsgrown on MgO(100) at RT. In all three cases, the epitaxialorientation of metals on MgO(100) is cube-on-cube and theadsorption site of metals is above oxygen ions on MgO(100).At submonolayer coverage, Ag overlayers have the bulk latticeparameter but Ni overlayers are strained to register with thesurface oxygen lattice of MgO(100). Pd with coverage less than1 ML has an average lattice parameter between those of bulkPd and MgO. At high coverage, e.g., >1 ML, Ag and Pd forma network of interfacial misfit dislocations to relax the strain,whereas the relaxation of Ni is facilitated by the growth of Ni


Fig. 43. Single point defects on a flat MgO(100) terrace imaged with atomicresolution by dynamic-mode SFM in two consecutive measurements (a) and (b)(9.1× 9.1 nm2). From [249].

clusters with an additional orientation of Ni(110) ‖ MgO(100).These authors concluded that structure and morphology atthe metal/MgO interfaces are mostly influenced by bondingstrength at the interface rather than the lattice parameter misfit.Trampert et al. [537] have investigated the atomistic structureof the Ag/MgO interface with HRTEM. These authors showedthat Ag islands grow cube-on-cube on the (100) surface of theMgO substrate. The interface was atomically abrupt, i.e. nochemical reactions like interdiffusion or formation of interfacialphases were observed. The misfit between Ag and MgOis accommodated by partial edge dislocations, leading to asituation where the Ag atoms between dislocations are sittingeither on Mg or O atoms/ions.

Other experiments suggest that there is a very weakinteraction between MgO surface and metals of Cu [140],Ag [533], and Au [532]. Metals often grow in form ofislands on MgO(100) at high temperature (above RT) and/orhigh coverages (above 1 ML). However, below a criticalsubmonolayer coverage, it was observed that Pd [138] andAg [538] grow in 2D islands on MgO(100) surfaces.

Chemical interaction of metals on MgO(100) was observedfor Ti- and Zr-adsorption on MgO. Ti and Zr adatoms diffuseinto the MgO bulk via Mg substitutional sites at RT. Noreactions were found at interfaces of Fe, Ni, Ge, and Ag with the

MgO surfaces. The trends of interface reaction at metal–MgOinterfaces are interpreted in terms of the reactivity of the metaladatoms to the surface oxygen atoms [135].

(2) Theoretical results: Many theoretical studies have beenconducted to study metal–MgO interactions. Several key issueson the interface formation have been addressed. They includethe adsorption sites of metal atoms on MgO, the nature ofelectronic interaction between metals and MgO, and the effectof surface defects on metal adsorption. All the three aspects arereviewed below.

For adsorption sites, most of the results indicate that ondefect-free MgO(100) surfaces metal atoms prefer to adsorb ontop of the surface O ions, which is consistent with experimentalresults [25]. Schonberger et al. calculated that the minimumenergy of the Ag/MgO interface occurs when the Ag atoms aresitting on top of the O atoms/ions [539]. Rosch and coworkersmade a systematic DFT study in the adsorption of variousmetal atoms on MgO(100). They show that metals of Cr,Mo, W, Ni, Pd, Pt, Cu, Ag, and Au all register on the oxidesurface anions [540–542]. The adsorption geometry of Pd [64,543], Cu [544], Ag [74,545,546], Au [543], and Al [546] onMgO(100) showed that the metal atoms sit on the surface Osites.

The electronic interaction between metals and MgO-surfaces with low defect densities is generally very weak [65,547]. The metal–MgO electronic interaction can be attributedto three mechanisms, metal polarization, chemical bonding,and MIGS. Goniakowski [65] suggested that Pd polarizationcaused by the MgO surface electrostatic field can be of greatimportance. There is a redistribution of electrons betweendifferent d components in the metal VB upon interfaceformation, which is similar to the case observed at metal/Al2O3interface (Fig. 39). In the case of metal adatom above surfaceO, electrons transfer from the d3z2−r2 orbital to the rest of the dband. Such an effect is similar to the macroscopic image chargeforce, which is observed at interfaces between a metal overlayerand an ionic crystal [62]. The polarization effect was found todominate in Cu, Ag, and Au interactions with MgO(100) [540,541]. Another example is the ab initio study in metal/MgO(100)systems [547]. They found that conventional MIGS produce asmooth DOS in the MgO band gap energy range. The interfacestates together with the polarization effect and bonding statesdetermine the interface characteristics, such as charge transfer,SBHs, etc. Finally, the interface bonding provides anotherimportant contribution to the metal–MgO interactions. In mostcases, the bonding is facilitated by the mixing of metal valenceelectron orbitals with O2p orbitals, which forms covalent bondsat the interface. Ni, Pd, Pt, and W [540,542] were found to formstrong bonding with the oxide anions and the bond has a polarcovalent nature with little charge transfer from the metal to theoxide.

Finally, we come to the effect of surface defects onmetal interactions. Extensive theoretical work demonstratesthat the presence of surface defects enhances the metal–MgOinteractions. Matveev et al. [548] calculated that the adsorptionof Cu, Ni, Ag, and Pd is stronger on Fs sites by 1–2.4 eVcompared with regular O2− sites. Zhukovskii et al. showed


that the nature of the interaction between Ag or Cu and MgOsubstrates with low defect density is of physisorptive character.Above the Fs sites, metal atoms are much more strongly bound.Adsorption onto the defect sites is accompanied by a substantialcharge transfer towards each adatom (δCu = 0.41 eV andδAg = 0.32 eV) as well as formation of partly covalent Me–Fsbonds at the interface [549]. In the Au/MgO(100) system, theadsorption of Au on the ideal MgO(100) surface is found to bevery weak with an adsorption energy of −0.13 eV/adatom. Onthe defective MgO(100) surface, Au adatoms prefer the vacancysites where the adsorption energy is −1.93 eV/adatom [550].The same role of oxygen vacancies in Pd-adsorption onMgO(100) was identified based on periodic ab initio electronicband structure investigation of the Pd–MgO interface [551].

5.2.2. Metal interactions with MgO filmsHighly ordered and stoichiometric MgO thin films can be

epitaxially grown on metallic substrates. These well-definedMgO-films are often used as model systems for supported metaloverlayers.

5.2.2.1. Growth of MgO films. MgO(100) thin films arefrequently prepared by evaporation of Mg onto metal substratesin the presence of O2. Occasionally, annealing in an oxygenatmosphere is necessary in order to crystallize the films.The supports include metal single crystals such as Mo(100)[552–554], Ag(100) [555], W(110) [556]. STM and electronspectroscopic investigations demonstrate that the grown MgO-films are smooth and possess the bulk properties of MgO [552,554]. Furthermore, the presence of surface defects on MgO-films has been identified and the density of the surfacedefects can be regulated by post-annealing or electronbombardment [557,558].

5.2.2.2. Metal interactions with MgO films. Like the Al2O3-film-based model systems, metal interactions with MgO-filmsdepend on both the MgO film surfaces and the supports,respectively.

(1) Surface effects: Similar to metal interactions on bulkMgO crystals, surface defects play an important role in theadsorption of metal atoms on MgO-films. Au-adsorption onMgO(100) films has been systematically investigated throughab initio calculations and laboratory experiments. Using mass-selected and softlanding techniques, Au8 clusters having thesmallest catalytically active size were deposited on MgO(100)films grown on Mo(100) [557,559]. Yoon et al. found that Au8clusters supported on defect-rich MgO(100) surfaces are activewhereas clusters deposited on virtually perfect MgO surfaceremain chemically inert. As illustrated in Fig. 44, the stretchvibration of CO adsorbed on mass-selected Au8 on MgO(100)with coadsorbed O2 shows a red shift on an F-center-richsurface with respect to the perfect surface. It was concludedthat there exists a larger degree of charge transfer from the F-center-rich MgO surface to the gold cluster than that on an F-center-free surface. Charging of the supported Au clusters bythe surface defects plays a key role in promoting their chemicalactivity [265]. Sterrer et al. applied electron paramagnetic

resonance (EPR), infrared spectroscopy (IR), and STM to studythe interaction of Au clusters with F centers on MgO surfaces.They also confirmed that Au particles adsorbed to color centersare indeed negatively charged while Au particles on regularterrace sites are neutral [558]. In addition to surface pointvacancies, hydroxyl groups on MgO surfaces also affect themetal adsorption process strongly [560], which is similar tothe cases observed in metal adsorption on hydroxylated TiO2,Al2O3, and SiO2 surfaces [319,461,462,489,491,561–563].

(2) Support effects: The other important factor, whichinfluences the metal–MgO-film interaction, is the metal supporteffect. Pacchioni and coworkers demonstrate that a metalsupport at the interface with a MgO-film results in charging ofadsorbed atoms with high electron affinity, like Ag and Au [543,564]. They concluded that an adsorbed Au or Ag atom isalmost neutral on a single crystal MgO(100) surface, while it isnegatively charged on a MgO(100) film supported on Mo(100).In Au/MgO(100) Au6s level is half filled and the configurationis atom-like, 5d106s1. However, in Au/MgO/Mo(100) the Au6slevel lies below the EF of Mo and both α and β componentsof the Au6s level are filled. Therefore, Au carries a net negativecharge and becomes Au− with an electronic configuration of5d106s2 (as shown in Fig. 45). The charge transfer, which maybe facilitated by tunneling effects, can occur from a metalsupport with low work function to adsorbed metal atoms withhigh electron affinity.

5.3. Metals on SiO2

SiO2 plays an important role in many technologicalapplications, for example as dielectric layer in microelectronicsand as catalyst support in heterogeneous catalysis. Over a broadrange of temperatures and pressures α-quartz is one of the moststable structures of SiO2, and the (0001) α-quartz surface canbe considered a model surface. However, the band gap of bulkSiO2 is very large, around 9.0 eV [565] such that it is difficultto study bulk SiO2 surfaces with many surface techniques.Studies of metal–SiO2 interactions often involve SiO2-filmsgrown on single crystal Si or refractory metal substrates, whichare then studied with various surface science techniques. Here,we discuss the interaction of metals with bulk SiO2 and SiO2-films.

5.3.1. Metal interactions with bulk SiO2

There have been a few experimental studies in single crystalquartz surfaces. The available results suggest two (0001) α-quartz surface structures. A LEED investigation of a (0001)surface, which was etched in HF solution, indicated a (1 ×1)-type surface [566]. A (

√84 ×

√84)-R11◦ reconstruction

was observed when heating the α-quartz surfaces in air above600 ◦C; the surface structure is related with the quartz α → β

phase transition which occurs at 573 ◦C [567]. Harte et al. havestudied the initial states of Cr and Ti growth on SiO2(0001).No large difference in the metal growth was observed on both(1×1) and (

√84×√

84)-R11◦ surfaces. Moreover, they foundthat deposition of Ti results in the formation of a Ti oxide layeron SiO2, while Cr forms metal clusters on the surfaces [568].


Fig. 44. Mass spectrometric signals pertaining to the formation of CO2 on Au8 deposited on (a) F-center-rich and (b) F-center-free MgO(100) thin films. (c) and(d) FTIR spectra measured for the same surfaces and with the same CO and O2 exposures as in (a) and (b), respectively, at various annealing temperatures. Massspectrometric data show that Au8 adsorbed on an F-center-free MgO(100) surface were essentially inactive for the combustion reaction (b); IR results indicate ared shift of the CO stretching frequency for the molecule adsorbed on Au8 supported on the defect-rich MgO thin film. From [265].

Fig. 45. DOS (density of states) curves for a Au atom adsorbed on top of O. (a) unsupported MgO(100) (3 layers); (b) supported MgO(100) (3 layers) on Mo(100).DOS of Au × 2. From [543].

Fig. 46. STM images of a SiO2/Mo (112) film. (a) 100× 100 nm2, Vs = 2 V, I = 0.2 nA, (b) 8× 8 nm2, Vs = 1.2 V, I = 0.35 nA, (c) 8× 8 nm2, Vs = 0.65 V,I = 0.75 nA. The arrow indicates an antiphase domain boundary and insets in (b) and (c) show the close up of the atomically resolved STM image (left) andsimulated image (right). From [580].

The intrinsic defects on bulk SiO2 surfaces may bedivided into two families. Oxygen-related defects of SiO2

include the peroxyl bridge ≡ Si–O–O–Si ≡, the peroxylradical ≡ Si–O–O·, the non-bridging oxygen ≡Si–O·, andovercoordinated oxygen O+3 (Si3) or O+3 (Si2O). Si-relateddefects include the two-fold coordinated silicon =Si, the

silicon dangling bond ≡ Si·, and the oxygen vacancy, whichmay be neutral VO (the weak ≡Si · · · Si ≡), the positivelycharged V •O corresponding to an E ′ center, or the wrongbond ≡ Si–Si ≡. All these defects have been identified byEELS [565] and ESR [569]. Experimental and theoretical datashowed that the interaction of metals, such as Au, Cu, Pd,


and Cs, on regular sites of SiO2 is weak, mainly throughdispersion interaction [570–572]. The surface defects, however,can stabilize the deposited metal atoms. Among the defect sites,the non-bridging oxygen forms the strongest bonds with metalatoms, which is followed by the silicon dangling bond, and theneutral oxygen vacancies [570,572].

Like TiO2, Al2O3, and MgO surfaces, hydroxylation mayoccur on SiO2 surfaces [573]. Silanol groups (−Si–OH)

can be introduced by reaction of SiO2 surfaces with water.The concentration, distribution, and nature of the silanolgroups clearly influence the properties for the technologicalapplication [561–563].

5.3.2. Metal interactions with silica filmsSiO2-films grown on conductive or semiconductive sub-

strates possess the chemical and physical properties of bulkSiO2 [574–576]. Furthermore, the thin SiO2-films can circum-vent the surface charging effect encountered in bulk SiO2 sys-tems. Therefore, the supported SiO2-films are good model sys-tems for studies in metal–SiO2 interactions.

5.3.2.1. Growth of silica films. The simplest route to prepareSiO2-films is to oxidize single crystal Si surfaces, eithervia thermal oxidation of Si in oxidizing atmospheres orby wet etching in solutions [5,577]. Other techniques, suchas ALD [187,188] and thermal evaporation of SiO in O2atmosphere [96,98], have been also used to grow thin SiO2-films. All the methods enable control over the thickness of theSiO2 layers, while the structure of the as-deposited layers isgenerally disordered.

From a structural point of view, it is highly desirablethat model systems consist of well-ordered films instead ofamorphous ones. Schroeder et al. [578,579] first reportedthe preparation of a thin crystalline SiO2-film on Mo(112).The experimental procedure consists of repeated cycles of Sideposition and subsequent oxidation, which is followed by afinal annealing step. The silica multilayer film is stoichiometricand fully covers the support surface. Recently, Freund andcoworkers have modified this process and prepared a monolayercrystalline SiO2-film on Mo(112) (Fig. 46) [580,581]. Theatomic structure of the epilayer was resolved by STM,infrared reflection-adsorption spectroscopy (IRAS), and DFTcalculation. The film consists of a 2D honeycomb-like networkof SiO4 tetrahedra with one oxygen of each tetrahedral unitbinding to the protruding Mo atoms of the metal surface,while the other three form Si–O–Si bonds with the neighboringtetrahedral unit. Goodman and coworkers [582,583] have alsogrown highly crystalline, well-ordered thin SiO2-films onMo(112) surfaces and characterized them by LEED, STM,and HREELS. The derived structural model of the SiO2-filmconsists of a layer of isolated [SiO4] clusters arranged in ac(2 × 2) structure on the Mo(112) surface with all oxygenatoms bonding to the Mo substrate. The physical propertiesof the SiO2-films at 1 ML coverage are influenced by the Mosubstrate, while films with coverage greater than 2 ML showproperties comparable to bulk-like SiO2 samples [584]. FlatSiO2-films have been also deposited on Mo(110) [574,585,

586] and Mo(100) [575] surfaces. In addition to Mo surfaces,Kundu and Murata [587] have prepared a single crystal SiO2film with the structure of β quartz on a Ni(111) surface andZhang et al. [588] deposited atomically flat silica films on aPd(100) surface.

5.3.2.2. Metal interactions with silica films. The effects of theSiO2film surfaces and the supports of the SiO2-films on metalinteractions with SiO2-films are given below.

(1) Surface effects: Surface defects are an important factorin the interaction between metals and SiO2. The presence ofsteps, antiphase domain boundaries, and oxygen vacancies onsurfaces of SiO2 epilayers grown on Mo has been confirmed byMIES, UPS, LEED, and STM [576,579,580,589]. Au clusterswere used to identify the role of the various defects in thenucleation and growth of metals on the SiO2 surfaces. Thestability of Au nanoclusters on defect sites follows the tendencyof oxygen vacancy complexes > step edges > line defects >

single oxygen vacancies [589]. In addition, the defects of Tiimpurities introduced onto SiO2 surfaces can strengthen themetal interaction with the surfaces. Metal clusters show amarked increase in island density and are sinter-resistant on theTiOx-modified SiO2 surface [590,591].

(2) Support effects: Depending on the thickness of thegrown SiO2 layer the supports may have a strong influenceon metal–SiO2 interactions. Toyoshima and coworkers [169,592] have studied the interaction of Ni with SiO2-filmsgrown on Si(111) and CO-adsorption at the Ni/SiO2/Si(111)system. CO-adsorption was suppressed on the Ni clusters thatwere deposited on SiOx/n-Si(111) but normal CO-adsorptionoccurred on Ni supported on SiOx/p-Si(111). It was concludedthat in the Ni/SiOx/n-Si(111) system the charge transfer takesplace from the donor level of n-Si to the Ni d orbital viaelectron tunneling through the thin SiOx interlayer (3 A). Thus,CO-adsorption at the Ni/SiO2/n-Si(111) system is inhibiteddue to retardation of σ -donation from CO to Ni. The chargetransfer between metal overlayers and Si(111) substrates hasbeen studied in metal/SiO2(∼15 A)/Si(111) systems by Ofneret al. [170]. They observed that the charge transports fromCs to the SiO2/Si interface which built dipole field acrossthe SiO2 layer and induced Si band bending. However, Indeposition did not cause any changes in SiO2/Si interfaceelectronic structure. The different electronic interactions inthe Cs/SiO2/Si and In/SiO2/Si systems are attributed to thedifferent electronegativities of Cs and In, which can be alsoexplained by the schematics shown in Fig. 41.

5.3.2.3. Chemical interaction of metals on silica films. Metaladatoms on SiO2 surfaces can be subjected to various surfaceand interface processes. Fig. 47 gives the schematic descriptionof the processes which could happen at metal/SiO2/Si modelsystems. The typical chemical interactions described inSection 2.3.1, which include redox reaction, alloy formation,encapsulation, and interdiffusion, have been observed atmetal–SiO2 interfaces [100,593].

Redox reaction: Pretorious et al. [103] found that Ti, Zr, Hf,V, and Nb react with SiO2 to produce oxides and silicides at


Fig. 47. Schematic of surface and interface processes in metal/SiO2/Si modelsystems. From [593].

the interfaces. They suggested a direct reaction between a silicafilm and a metal overlayer as follows:

Mx + SiO2 → MySi+Mx−yO2.

The reaction can be predicted based on thermodynamicconsideration as discussed in Section 2.3.2. Metals with anelectronegativity of less than 1.5 on the Pauling scale shouldreact with a SiO2 substrate. Similar to the above result, reactivemetals such as Al, Mg, Ti, Si, and Ge were found to reactwith SiO2 directly and cause a homogeneous decompositionof SiO2-films grown on Si(100) [594]. Annealing a CuMgalloy deposited on 650 nm SiO2 layers results in the formationof fcc MgO and the reduction of SiO2 [595]. In theTi/SiO2(50A)/Si(111) system, the metal oxidation and SiO2reduction occur even at room temperature [596].

Encapsulation: Powell and Whittington [597] suggested amechanism of encapsulation to explain the deactivation ofSiO2-supported Pt model catalysts. They observed that Ptparticles become partially immersed in the SiO2 surface witha concurrent formation of a SiO2 ridge around the base of thePt particles when annealing the catalysts at 1200 and 1375 K.The encapsulating process was driven by the minimization ofsurface free energy in the Pt/SiO2 system. Van den Oetelaaret al. studied the thermal stability of Cu particles supported ona thick SiO2 layer (400–500 nm) by LEIS, AFM, and RBS [593,598]. The disappearance of Cu from the outermost atomic layerof the UHV annealed Cu/SiO2 model catalysts was attributedto encapsulation of the Cu particles by silicides. The UHV-annealed Cu/SiO2 samples can be regenerated by exposure toair at room temperature for several hours. The reversibility ofthe surface process is very similar to the encapsulation reactionsobserved in metal/TiO2 systems [8].

Interdiffusion: Interdiffusion of metals into SiO2 layers isoften observed upon annealing metal overlayers supported onthin SiO2-films (∼10 nm). In metal/SiO2/Si systems, silicideformation between the metal overlayers and the Si substratedrives the metal diffusion through the SiO2 layer to theSiO2/Si interface [594]. It is believed that the interdiffusion isfacilitated by defects in the SiO2 layers, e.g., oxygen vacancies,pinholes, microchannels, or microvoids. The interdiffusion ofNi [599,600], Cu [593], Pd [601,602], and other transition

Fig. 48. HRTEM images of three particles after reduction showing formationof (a) Pt3Si with Cu3Au structure, (b) monoclinic Pt3Si, and (c) tetragonalPt12Si5. From [98].

metals [594] in SiO2 layers has been reported upon annealingof the metal/SiO2/Si model systems.


As we have discussed in Section 2.4, mass transport in anoxide can be assisted by an electric field therein. For example,Vogt and Drescher [603] demonstrated field-assisted Cu diffu-sion in SiO2. When annealing a 100 nm Cu/150 nm SiO2/Sisystem at 450–500 ◦C for 1 h, no Cu diffusion in the SiO2 layerwas observed. However, the thermal treatment of the interfacewith an external electric field results in strong enhancement ofCu transport. The applied electric field accelerates the trans-port process by assisting the ionized copper atoms to migratethrough the oxide towards the Si substrate.

Alloy formation: In metal/SiO2(thinlayer)/Si model systems,many metals can diffuse to SiO2/Si interface, where they reactwith Si substrates to form metal silicides [593,594,596,598–602,604].

Other studies also show that metal silicide formation cantake place via direct reaction between metal and SiO2. Van denOetelaar et al. [593] proposed a reaction mechanism to explainthe direct reaction between a noble metal and a thick SiO2-filmunder the condition of UHV annealing:

xCuδ+–O–Si–OUHV,HT−−−−−→ Cux –Si+ (x − 1)SiO2

+ “oxygen-compound” ↑ .

The “oxygen compound” represents oxygen that is releasedby the reduction of SiO2. The strong interaction between Cuand SiO2 supports results in the formation of Cuδ+ species,which could alter the surface and interface energy in the systemand make the reaction thermodynamically possible. A similarmechanism is found to be active for Rh3Si formation in the caseof thermal treatment of a Rh/SiO2/Mo model system above 850K in UHV [605] and for Pd silicide formation in a Pd/SiO2/Momodel system upon UHV annealing at 1000 K [100,606].

In the presence of H2, silicide formation takes place via adifferent mechanism [99,607]:

Mx + SiO2 + 2H2 → Mx Si+ 2H2O.

Thermodynamically, the driving force behind the process isthe formation of stable water molecules. For example, in aPt/SiO2 model system silicide formation was studied by variousTEM techniques [96,98,608]. Pt particles supported on free-standing amorphous SiO2-films (25 nm thick) were reducedin 1 bar H2 at 873 K. Pt3Si with Cu3Au structure, monoclinicPt3Si, and tetragonal Pt12Si5 were identified after the treatment(see Fig. 48). The reaction products result from reductionof SiO2 by atomic hydrogen, which involves dissociativeadsorption of H2 on Pt particles, reduction of SiO2 by atomichydrogen diffused from Pt, and migration of Si atoms intoPt to form Pt silicides. An early investigation of Pt reactionwith SiO2 conducted by Lamber et al. [99] also confirmedthe formation of a cubic platinum-rich Pt3Si phase, which isintermediate to the monoclinic Pt3Si phase. The hydrogen-induced metal silicide formation has been reported in manyother metal/SiO2 systems. For example, a Ni3Si compound wasobserved after prolonged heating of a Ni/SiO2 model systemin a hydrogen atmosphere [609]. A Pd2Si phase was revealedby TEM, TED, and Convergent Beam Electron Diffraction(CBED) after heating a Pd/SiO2 system in H2. They found that

metal–support interaction is influenced by the pre-treatmentof SiO2 support, indicating that the presence of OH groupson the silica facilitates chemical metal–support interaction andformation of a metal silicide [561]. Other Pd silicides, such asPd4Si and Pd3Si, have been observed after reduction of Pd/SiO2catalysts in H2 up to 600 ◦C [610,611].

6. Summary

In the present review, two fundamental questions atmetal/oxide interfaces were addressed. (1) What is theelectronic interaction during metal/oxide interface formation?(2) What kind of chemical reaction occurs during metal/oxideinterface formation?

First, contact between a metal and an oxide results in chargeredistribution at the interface at the local range and/or the longrange. The electron redistribution is driven by principles ofenergy minimization of the system and continuity of electricpotential in the solid.

Local charge redistribution occurs within a few atomiclayers close to the interface, which includes polarizationof metal electron orbitals, formation of image chargesin the metal, MIGS at interfaces, and interface bonding.The local electronic interaction is particularly importantat metal/insulating oxide systems, e.g., metal/Al2O3(0001),metal/MgO(100), and metal/SiO2. Two aspects are critical inthe process. One of the important factors is the electronegativityof the metal; the oxide’s surface property including surfacedefects, surface stoichiometry, surface termination, and surfacehydroxylation is the other one.

The long-range electronic interaction at metal/oxideinterfaces is analogous to that at metal–semiconductorjunctions. The equilibration between EF of the metal andoxide contacting phases induces charge transfer between themetal and the space charge region in the oxide. This spacecharge transfer relies on the electronic structures of the twocontacting phases, i.e. the surface work function (EF ) of thesolids. In cases of metals supported on mixed conductingoxides, such as TiO2(110) and SrTiO3(100), the spacecharge is very important. In oxide film-based model systems,e.g., metal/Al2O3-film/support, metal/MgO-film/support, andmetal/SiO2-film/support, electrons can tunnel through the thininsulating oxide layers, which allow long-range charge transferto occur between the lower conductive support and the metaloverlayers. In this case, the charge transfer depends on the EFof the support and the metal overlayers.

Secondly, when putting a metal onto an oxide surface, themetal and substrate atoms may be involved in various masstransport processes as shown in Fig. 17. One of the basicprocesses is the surface diffusion of metal adatoms on oxides,which enables the atoms to find the energetically preferablesurface adsorption sites. Metal adsorption on an oxide surface isclosely correlated with the local electronic interaction betweenthe metal and oxide, which depends on metal electronegativityand oxide surface properties.

Metal and/or substrate atoms can diffuse across the interface,which results in interfacial reactions. The reactions are driven


by thermodynamic forces but may be kinetically limited atrelatively low temperature (<1000 ◦C). For mixed conductingoxides, such as TiO2 and SrTiO3, the interface reactionsinvolve ionic defect diffusion processes in the substrates, whichare closely related to defect chemistry of the oxides. Forexample, the production and diffusion of oxygen vacanciesare critical in metal oxidation on TiO2 and SrTiO3 surfaces.Diffusion of Ti interstitials Tin+i (n ≤ 4) dominates in metalencapsulation reactions, oxygen-induced restructuring, andbulk-assisted reoxidation reactions on TiO2 surfaces. Transportof Sr vacancies and oxygen vacancies between the surfaceand bulk of SrTiO3 determines the surface reconstruction andthe formation of new surface phases at SrTiO3 surfaces. Thediffusion of ionic defects in oxides can be promoted or retardedby an interfacial electric field. At the metal/oxide interfaces,mass transport is coupled with the electronic interaction.As has been shown in many cases, reactions at metal/TiO2and metal/SrTiO3 interfaces show a strong dependence onbulk electronic structure (free electron density or EF ) ofboth contacting phases. The metal–oxide interface reactionmechanism is consistent with the observed oxidation reactionsof metal surfaces and etching reactions of semiconductorsurfaces. This behavior can be explained in the frameworkof the generalized Cabrera–Mott theory, i.e. mass transportoccurring in a bi-phase (solid–solid or solid–gas) interfacereaction may be dependent on the bulk electronic structureof the two contact phases. Thus, interface reactions can besuccessfully tuned by controlling EF of the metal and/or oxidephases, such as doping of oxides, alloying of metals, applicationof external electric field, etc.

Metal/oxide interface is a fast-growing research field andmany systems have been investigated. In this paper metals ona few typical oxides (TiO2, SrTiO3, Al2O3, MgO, and SiO2)have been reviewed, demonstrating the critical factors in de-termining the metal–oxide interactions and the possible routesto desirably control the formation of metal/oxide interfaces.The discussed metal/oxide interfaces can be classified intothree systems: (1) metals on mixed conducting oxides includ-ing metal/TiO2(110) and metal/SrTiO3(100); (2) metals on in-sulating oxides, such as metal/Al2O3(0001), metal/MgO(100),and metal/SiO2(0001); (3) metals on supported thin oxide films,e.g., metal/Al2O3-film/support, metal/MgO-film/support, andmetal/SiO2-film/support. Most of the previous research effortsperformed in these systems have focused on the adsorption andbonding of metals on the oxide surfaces, which show strongdependence on the oxide surface properties. Some experimentshave shown that the nucleation, growth, and epitaxy of met-als on oxides can be deliberately controlled through the sur-face modification of the oxide supports, e.g. the surface recon-struction, change in surface termination, and variation in sur-face point defects. Surface hydroxyl groups are often presenton many oxide surfaces and the surface hydroxylation is an-other important parameter to regulate the metal–oxide interac-tion [319,459,460]. It is clear that more such studies should bedone in future.

Recent results indicate that the formation of metal/oxide in-terfaces is not only correlated with the oxide surface proper-

ties but also strongly influenced by the oxide bulk properties,such as the bulk ionic defect and bulk electronic structure [13,28,85,89]. This point is particularly important for systems ofmetals on mixed conducting oxides, where space charge trans-fer and mass transport at the interfaces are feasible. We showthat the interactions and, in particular, chemical reactions atthe interfaces are determined by the defect chemistry and EFin the oxide bulk and the experiments can be explained by theconcept of the generalized Cabrera–Mott theory (see Figs. 13–15). Although these results were observed at metal/TiO2 andmetal/SrTiO3 interfaces it seems to be very likely that similarresults occur on the surfaces of other mixed conducting oxides,such as CeO2, ZrO2, ZnO, SnO, etc.

For metals on supported thin oxide films, the film supportsare of great importance. Some example showed that electronicinteraction between the conductive support underneath the filmand the metal overlayer on top of the film occurs via electrontunneling through the thin oxide films (see Fig. 41). Chemicalreactions, mainly alloy formation, between the metal overlayersand the support may promote the interdiffusion through thethin oxide films. Apparently, the charge transfer and masstransport processes are sensitively dependent on the thicknessof the thin oxide films. Therefore, the choices of the conductivesupport and the thickness of the overgrown oxide film arecritical parameters controlling the metal–oxide interactions.Nucleation, growth, electronic state, and surface chemistry ofthe metal deposit on the thin oxide films have to be relatedto properties and variations of the underlying support and filmthickness [159].

It is highly expected that the macroscopic properties of themetal/oxide interfaces in both model and real systems, suchas catalytic performance, electrical properties, and mechanicalstability, can be tuned based on the above concepts developedfrom the studies on the model systems.

Acknowledgements

The authors would like to thank Prof. Manfred Ruhle for hissupport when this project was performed. We owe warm thanksto Prof. Xinhe Bao, Prof. Robert Schlogl, Dr. Dangsheng Sufor their support during the writing of this manuscript; to Dr.Jason White for critical reading of the manuscript; and to manycolleagues for providing us electronic versions of their figures.We would also like to thank Prof. Ulrike Diebold for havinginvited us to write this review and critically reading the article.

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Glossary

AEM: analytic electron microscopeAES: Auger electron spectroscopyAFM: atomic force microscopyALD: atomic layer depositionALE: atomic layer epitaxyBE: binding energyCB: conduction bandCBM: conduction band minimumCNL: charge neutrality levelCSD: chemical solution depositionCTEM: conventional transmission electron microscopyCVD: chemical vapor depositionDFT: density functional theoryDOS: density of statesEELS: electron energy-loss spectroscopyEF : Fermi energyELNES: electron energy-loss near-edge structureEPR: electron paramagnetic resonanceEXAFS: extended X-ray absorption fine structureFTIR: Fourier transform infrared spectroscopyFWHM: full width at half maximumGIXS: grazing incidence X-ray scatteringHAS: helium atom scatteringHREELS: high resolution electron energy-loss spectroscopyHRTEM: high resolution transmission electron microscopyHTR: high temperature reductionIR: infrared spectroscopyIRAS: infrared reflection-adsorption spectroscopy

ISS: ion scattering spectroscopyLDA: local density approximationLEED: low energy electron diffractionLEIS: low energy ion scatteringMBE: molecular beam epitaxyMD: molecular dynamicsMEIS: medium energy ion scatteringMIGS: metal-induced gap statesML: monolayerMOS: metal–oxide–semiconductorMOVPE: metal-organic vapor phase epitaxyMS: metal–semiconductorPES: photoemission electron spectroscopyOR: orientationPVD: physical vapor depositionRHEED: reflection high energy electron diffractionSBH: Schottky barrier heightSEXAFS: surface extended X-ray absorption fine structureSFM: scanning force microscopySMSI: strong metal–support interactionSTM: scanning tunneling microscopySTS: scanning tunneling spectroscopyTEM: transmission electron microscopyTED: transmission electron diffractionUHV: ultrahigh vacuumUPS: ultraviolet photoelectron spectroscopyVB: valence bandVBM: valence band maximumXRD: X-ray diffractionXPS: X-ray photoelectron spectroscopy

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