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Structure−Property Relationship and Chemical Aspects of Oxide− Metal Hybrid Nanostructures Svetlozar Surnev,† Alessandro Fortunelli,‡ and Falko P. Netzer*,† †

Surface and Interface Physics, Institute of Physics, Karl-Franzens University, Graz A-8010 Graz, Austria CNR-IPCF, Istituto per i Processi Chimico-Fisici del CNR, I-56124 Pisa, Italy

‡

4.4.1. Reactions on “Inverse Model Catalyst” Surfaces 4.4.2. Reactions on Oxide Nanolayer Surfaces 4.4.3. Morphological Aspects during Reactions on Oxide Nanolayers 5. Synopsis Author Information Corresponding Author Notes Biographies Acknowledgments References

CONTENTS 1. Introduction 2. Experimental Aspects and Theoretical Methods 2.1. Experimental Aspects: Fabrication Procedures 2.2. Theoretical Methods 2.2.1. Density Functional Theory (DFT) Approaches 2.2.2. Local and Semilocal xc-Functionals 2.2.3. Hybrid xc-Functionals 2.2.4. Density of States (DOS) 2.2.5. Post-DFT and DFT+U Methods 2.2.6. Requirements for Theoretical Methods 2.2.7. Structure Prediction 3. Atomic Structure Concepts 3.1. Interface Geometry 3.1.1. TM−O (100)-Derived Structures 3.1.2. TM−O (111)-Derived Bilayer Structures 3.1.3. Hexagonal O−TM−O (111) Trilayer Structures 3.1.4. Single-Layer Structures with Unusual Building Elements 3.1.5. Early Transition Metal Oxides 3.2. Chemical Interactions at the Interface 3.3. External Variables and Surface Oxide Phase Diagrams 3.4. Effects of Low Dimensionality 3.5. Defects as a Structure-Stabilizing Concept 4. Structure−Property Relationships 4.1. Electronic Structure of Oxide Ultrathin Films 4.1.1. Dependence of the Work Function on the Oxide Film 4.1.2. Band Structure and Metallization 4.1.3. Excited States 4.2. Vibrational Properties 4.2.1. Thickness-Dependent Vibrational Properties 4.2.2. Vibrational Properties of 2D Oxide Nanolayers: Local Building Block Concepts 4.3. Magnetic Properties 4.4. Chemical Properties © XXXX American Chemical Society

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1. INTRODUCTION The concept of coupling a nanometer or subnanometer scale oxide phase to a metal surface to form a metal−oxide hybrid material with novel properties is an attractive approach in the fundamental search for emergent phenomena in low-dimensional materials, and it may open new avenues of materials design with a view toward nanotechnology applications.1 In general, these systems are prepared under extreme (ultrahigh vacuum, UHV) conditions and do not occur in nature, and from this perspective they may be regarded as artificial materials, with the potential of on-purpose tailoring their properties.2 Although the systems of study in this Review have mainly a reductionist model character, elements of them are inherent components in already existing applications, in particular in systems where the oxide−metal interface is of relevance. Examples of the latter comprise the fields of advanced heterogeneous catalysis, modern electronic device technology including magneto-resistive and spintronic concepts, gas sensor devices, (multi)functional coatings or corrosion inhibition, and environmental chemistry related products. With increasing miniaturization of system components in nanotechnology applications, the interfaces become more and more prominent and eventually system dominant. Thin films of oxides epitaxially grown on metal surfaces are popular systems for model studies of fundamental surface properties of oxide materials and have been so for the last two decades or so, but they have also entered the market of advanced technologies in many diverse areas. From a practical point of view, metal-supported oxide thin films have technical advantages over bulk oxide single crystal surfaces in

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Special Issue: 2013 Surface Chemistry of Oxides

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Received: July 31, 2012

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dx.doi.org/10.1021/cr300307n | Chem. Rev. XXXX, XXX, XXX−XXX

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interactions between the charge-separated oxide phase with the induced image charges in the support, etc. These can be helpful in drawing expected trends of a given property (such as charge transfer or chemical bonding) or to produce general correlations between physical observables.10 In many cases, however, the complex character of these multicomponent (and thus multifunctional) materials with several different physical interactions competing in determining the observed phenomena calls for quantitatively accurate and reliable computational predictions, able to resolve experimental ambiguities or to more and more often anticipate experiments. Under this point of view, the latest years have witnessed an impressive progress. The advances in hardware and software (in particular in bandstructure methods) have enabled usually reliable a priori calculations of most equilibrium properties of oxide-on-metal systems, in terms of both energetics (thermodynamics of phase transformations, lattice constants, phonon spectra) and electronic structure (charge distributions and redistribution, local density of states, and the associated STM simulated images), so that computational approaches are now an indispensable tool in structure characterization and assignment. Finally, what is equally fundamental, the accurate prediction of the topology of local minima and saddle points in the energy hyper-surface allows researchers the possibility, which is unique to theoretical modeling, of performing computational “gedanken experiments” and thus to disentangle via in-depth analyses the origin of the various basic interactions in these complex systems.11 Despite some limitations of current approaches, theoretical simulations therefore play an increasingly important role in characterizing and predicting novel structures and properties of oxide nanostructures, and a comparable role is expected in the near future in the investigation of electronic excited states and kinetic processes. The organization of this Review is as follows. In section 2, experimental aspects of the preparation of metal-supported oxide nanostructures are presented, and the theoretical methods for the description of oxide nanostructures are introduced and critically evaluated. The atomic structure of oxide nanophases, the central property around which this Review is composed, is discussed in section 3; here, it is attempted to introduce some systematics into the complex structural variety of oxide nanophases and to identify major driving forces. In section 4, structure−property relationships are discussed: electronic, vibrational, magnetic, and chemical property aspects are treated in separate subsections. Finally, a brief synopsis concludes this Review.

fundamental scientific oxide surface studies due to their ease of preparation and because the presence of the conducting metal substrate avoids charging problems in charged-particle probe techniques. For recent reviews on oxide thin films, we refer to those of Nilius,3 Freund and Pacchioni,4 and Giordano and Pacchioni5 for experimental and theoretical aspects, respectively. In this Review, we will address ultrathin films and metalsupported structures of oxides in the extreme nanoscale regime. Nanostructures of a given oxide electronically and elastically coupled to a different metal surface in the two-dimensional (2D) thin film limit of 1−3 (perhaps up to 5) monolayers (ML) as well as one-dimensional (1D) oxide line structures and (quasi)zero-dimensional (0D) oxide clusters or nanodots will be examined in terms of their novel emergent properties, caused by the hybrid character and the low dimensionality of the systems. From a different viewpoint, the oxide nanolayers coupled to a metal surface may be regarded as the realization of an isolated oxide−metal interface in the 2D limit, which lacks on the one side the oxide bulk component but becomes accessible to examination by surface science tools. The focus of this Review will be limited to hetero-oxide systems on metal single crystal supports, studied by the socalled surface science approach, as these represent the best characterized systems from both experimental and theoretical points of view. The surface science approach involves in situ ultrahigh vacuum studies, with atomic scale preparation control and atomic resolution in the characterization. We will concentrate on hetero-oxide hybrid systems and will neglect intrinsic “surface oxide” phases in this Review, referring the reader instead to the review of Lundgen et al. on this topic.6 The metal support surfaces used for hetero-oxide nanostructure fabrication are mainly of the group VIII and group Ib noble metals, because these are the most appropriate single crystal surfaces that allow the fabrication of well-defined interfaces to the oxide nanostructures due to their oxidation resistance during the oxidizing nanostructure preparation conditions. The oxides treated here will be predominantly of the binary oxide type, because ternary or more complex oxide systems have only scarcely been investigated in the extreme nanoscale regime; as a notable exception, we mention here the work of the Widdra group on BaTiO3 on Pt supports.7,8 As a result of the nanoscale nature of the oxide−metal hybrid structures, scanning tunneling microscopy (STM) is the major experimental characterization technique, which is responsible for much of the progress in oxide nanostructure research during the past decade. The real space view of local atomic structures, phase transformations, and interface properties is invaluable despite some difficulties in the interpretation of STM images of oxide systems. Some aspects of STM imaging of oxide nanostructures have been discussed in a recent review by some of us.9 Nevertheless, the value of the variety of areaaveraging spectroscopy techniques, if applicable to a homogeneous set of nanostructures over larger surface areas, should not be underestimated, because only a combination of several complementary techniques can provide the detailed insights necessary to understand the complex electronic, magnetic, and chemical properties of these hybrid systems. This leads us to the importance of theoretical modeling in oxide nanostructure research. Theoretical modeling plays a manifold role in this field. First, theoretical analysis can set up a framework of basic principles governing the emergent phenomena in oxide nanostructures, such as the continuity of the electric potential, the polarization

2. EXPERIMENTAL ASPECTS AND THEORETICAL METHODS 2.1. Experimental Aspects: Fabrication Procedures

There is a wealth of preparation techniques, which have been applied for the growth of ultrathin oxide films, ranging from physical vapor deposition (PVD) under highly controlled ultrahigh vacuum (UHV) conditions, to chemical methods, involving chemical vapor deposition (CVD) and liquidprecursor-based techniques.12 Here, we will briefly summarize preparation routes within a surface science perspective, focusing mainly on the PVD of metal atoms onto a dissimilar singlecrystal metal substrate. In the so-called reactive evaporation (RE) method, the metal deposition is performed in a background pressure of oxygen. Here, low deposition rates and partial pressures of the oxidizing gas are typically employed, B


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In this case, the smooth Cu3Au(100)-O surface not only prevents intermixing between evaporated metal atoms and substrate, but importantly it acts as a reservoir of subsurface oxygen atoms stored near the surface. The latter can be released in a controlled fashion by thermal treatment, promoting homogeneous oxidation of the metal deposit. This preparation recipe has been also successfully applied for growing atomically flat NbOx and MoOx overlayers.27 A completely different route for the fabrication of oxide nanostructures with a low dimensionality is also worth mentioning. It involves the direct sublimation of tungsten trioxide to form monodispersed (WO3)3 clusters in the gas phase, as reported by Dohnalek et al.28,29 The (WO3)3 clusters can then be deposited on a metal surface, where they condense and self-assemble under suitable annealing conditions, leading to the formation of well-ordered 2D W oxide layers.30

with the oxidation process between the gas species and the metal atoms taking place at the substrate’s surface. Conversely, in the postoxidation (PO) method, the metal atoms are first deposited in UHV onto the metal substrate and subsequently oxidized. The choice of the preferred method depends on the specific system under investigation: factors such as the rate of oxygen dissociation on the metal template, oxidation reactions at the substrate surface, and diffusion processes into the bulk need to be taken into account. In the case of low-dimensional oxide systems exhibiting a complex phase diagram, where a coexistence of different phases is possible, the selection of one rather than the other approach may be crucial in the effort to obtain atomically smooth surfaces and to drive the assembly of the overlayer into a single phase.13 In particular, the PO method has proven beneficial for the fabrication of onedimensional oxide nanostructures on stepped metal surfaces: here, the steps are first decorated by metal atoms, which then become selectively oxidized.14 Note that the RE approach is not applicable here, because massive restructuring (such as faceting) of the bare stepped substrate may occur during the oxygen exposure. Molecular oxygen is mostly employed as the oxidizing agent in both the RE and the PO methods, and the oxygen content in the oxide system is controlled through the oxygen pressure and the substrate temperature applied during the growth. Typically, the reactivity of metals such as Ni, Al, or Mg is sufficient to ensure formation of fully oxidized films by use of molecular oxygen in a high-vacuum environment (10−8 mbar ≤ p ≤ 10−6 mbar). However, this is not always the case, and particularly for transition metals that may adopt more than one oxidation state, alternative approaches are more often followed, in which activated oxygen is supplied during the growth. This can be accomplished by using an electron cyclotron resonance (ECR) oxygen plasma source,15−18 gaseous NO2,15,19 or ozone O3.17,19 A commercial thermal cracker can be also used as a source of atomic oxygen.20,21 Because the O2 dissociation step is often rate limiting, the atomic oxygen assisted deposition may facilitate the reaction between individual metal and O atoms before metal cluster formation occurs. In this way, a higher degree of order and a better control on the stoichiometry and morphology of the oxide−metal interface can be achieved, as recently demonstrated for the monolayer of NiO on Ag(001).21 A further alternative is provided by the use of a dedicated highpressure cell, as that employed by Guimond et al.22 to create flat, (001)-oriented, films of vanadium pentoxide V2O5, by using O2 pressures in the mbar range. A modification of the traditional PO approach has been reported by Matsumoto et al.23 for the growth on a Pt(100) surface of a titanium oxide (3 × 5) structure, identified as a Ti2O3 monolayer. Here, an ordered Pt3Ti surface alloy with c(2 × 2) structure was formed prior to the oxidation step with ozone. A similar procedure, but using molecular oxygen, had been previously adopted by Knight et al. to grow epitaxial manganese oxide by oxidation of the Cu(100)c(2 × 2)-Mn substitutional surface alloy.24 A different approach has been developed by Niehus et al.25,26 They noted that the evaporation of V metal onto the bare metal surfaces (Au, Cu, or Cu3Au) gives rise to massive surface alloying, and that both the PO and the RE approaches thus lead to rough vanadia films with poor order. In contrast, they observed that on a Cu3Au(100) alloy surface substrate very well-ordered two-dimensional oxide overlayers can be prepared provided that the surface is preexposed to oxygen prior to V dosing and sequential annealing.

2.2. Theoretical Methods

Theoretical and computational tools have become an essential component in the investigation of supported oxide nanostructures, as in many other fields of science and technology. This is due to the fact that these phases often present unusual or even unprecedented structural motifs and to the difficulty in clarifying the basic physics of these systems.31 We will here provide a brief overview of these methods, their usefulness, and possible limitations. The first distinction that is needed is that between Hartree− Fock (HF) and post-Hartree−Fock (post-HF) approaches, on the one hand,32,33 and density-functional theory (DFT)34 and post-DFT approaches, on the other hand. In post-HF approaches, variational or perturbative methods are applied to expand the electronic wave function in terms of excitations with respect to the Hartree−Fock reference.35 PostHF methods are considered to be less suited for describing metallic systems. One difficulty common to nonmetallic systems lies in the inefficient description of short-range correlation or the Coulomb hole: in post-HF methods, the two-body electronic wave function is expanded in terms of products of single-electron basis functions, and this expansion converges very slowly for small values of the distance between the coordinates of the two electrons. Furthermore, for systems such as those composed of transition metal atoms, in which electronic states of different atomic character are close in energy (say nd10, nd9(n + 1)s1, and nd8(n + 1)s2 in an end-ofrow transition metal), it is difficult at the ab initio level to find the correct balance in the description of differential correlation effects. The major issue comes from long-range correlation effects, such as screening. Metals in fact present a vanishing band or HOMO−LUMO gap, that is, the energy difference between their highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO). This makes them quite different from insulating systems in terms of localization properties of the electronic wave function36 and convergence of post-Hartree−Fock methods, as it entails a wealth of low-energy electronic excitations and the need of a delicate balance in their description to achieve chemical accuracy. It may be recalled that the Hartree−Fock approximation is known to describe very poorly the electron gas or “jellium” model, which is the simplest model of metallic bonding (an idealized system corresponding to a homogeneous electron gas moving in a constant external potential), by overestimating the width of the conduction band and presenting a diverging density of states (DOS) at the Fermi C


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level.37 This is cured at the post-HF level, but at the price of a heavy computational effort. In short, quasi-degeneracies and vanishing gaps cause issues to post-HF approaches, so that they are not much used in the field of supported oxide nanostructures and will not be discussed further. 2.2.1. Density Functional Theory (DFT) Approaches. The problems associated with short-range and some long-range correlation effects can be overcome by density-functional theory (DFT). DFT in its Kohn−Sham (KS) version38 is a single-particle (or one-electron) method.39 The Kohn−Sham equations are therefore formally equivalent to the Hartree− Fock ones and read: ⎛ ℏ2 2 ∇ + V (r) + ⎜− ⎝ 2m = λiφi(r)

τσ(r) =

occup

∑

(1/2)|Δφiσ (r)|2

i

(2)

where σ is the spin (the Laplacian of the total electron density can be also used). One nonempirical example of meta-GGA approaches is given in ref 47. The third rung of DFT somewhat improves upon the second one, providing usually better values of atomization energies and also of structural parameters.48 2.2.3. Hybrid xc-Functionals. Other very common xcfunctionals, corresponding to the fourth rung of the Jacob’s ladder, are named hybrid xc-functionals, and consist of a mixture of GGA and Hartree−Fock exchange.49−51 The related formula for the exchange and correlation energy reads:

⎞

∫ |rρ−(r′r)′| dr′ + Vxc[(r)](r)⎟⎠φi(r)

Exc hybrid = aEx HF + (1 − a)Ex GGA + EcGGA

(3)

where Exchybrid is the xc-functional for the hybrid approach, ExHF is the Hartree−Fock exchange, ExGGA is a GGA exchange functional, EcGGA is a GGA correlation functional, and a is a parameter ranging between 0 and 1 (an analogous equation holds for the xc-potential). It should be noted that the admixture of a fraction of HF exchange is not an ad hoc trick, but is formally justified by adiabatic connection arguments, that is, a rigorous formal connection between the noninteracting Kohn−Sham system and the physical one.52 It may indeed be recalled from a historical point of view that the first DFT approach, which went beyond LDA, was an exact exchange plus an orbital-dependent correlation functional whose basic ingredient was the kinetic energy density for the occupied Kohn−Sham orbitals reported in eq 2; it was fairly accurate for systems with an appreciable band gap and was influential in the successive developments of DFT.45,53 It should also be noted that the Hartree−Fock exchange is a functional not of the oneelectron density but of the one-electron density matrix.54 This is advantageous in some respects. In particular, local or semilocal xc-functionals inherently suffer from the selfinteraction correction (SIC) error, that is, an erroneous attribution of correlation energy to single-electron systems.55 The use of the density matrix instead of the density allows hybrid xc-functionals to correct for a great part of this error, which is negligible for spin-compensated metals, but can be substantial for systems with unpaired/localized electrons. From a practical viewpoint, however, the evaluation of the Hartree− Fock exchange is numerically harder than that of a functional of the electron density. For example, while HF and later hybrid xcfunctionals have been implemented in periodic codes using localized basis functions since a long time,56 their implementation in periodic codes using extended or delocalized (plane waves) basis sets is recent and computationally demanding.57 Moreover, auxiliary basis sets for expanding the electron density can be utilized to achieve linear scaling,58 but the extension to the density matrix is more complicated.59,60 Two developments of hybrid xc-functionals have been recently proposed. The first one takes advantage of so-called range-separation or screening.61 In these approaches, the 1/r12 term in the Hartree−Fock exchange is partitioned into a shortrange and a long-range component, and the latter (long-range) component in the exchange is screened; that is, only its shortrange part is considered, thus allowing a much more efficient numerical evaluation even for metallic systems.56,62 The corresponding exchange and correlation energy thus becomes:61

(1)

where (−((ℏ2)/(2m))▽2) is the electron kinetic energy, V(r) is the external electrostatic potential due to the nuclei, ∫ (∑Nj = 1[φj*(r′)φj(r′)])/(|r − r′|) dr′ = ∫ (ρ(r′)/(|r − r′|)) dr′ is the Coulomb operator, describing the average Coulomb field generated by the electrons, Vxc[ρ(r)] is the exchangecorrelation (or xc-) operator, which is a functional of the total electron density, ρ(r), thus an xc-functional. The eigenvalues {λi} of the KS operator are the one-electron or orbital energies, while its eigenvectors {φi(r)} are called one-electron wave functions or orbitals. {λi} and {φi(r)} play a fundamental role in any theoretical treatment, as they allow one to formally associate an energy and a wave function to each electron. It has been proven38 that there exists an exact Vxc[ρ(r)] xc-functional such that, although intrinsically single-particle, DFT can produce the exact ground-state energy of the complete manybody Schrödinger equation. Because of their simplicity and effectiveness, DFT methods are ubiquitous in the theoretical study of supported oxide nanostructures, and we will discuss them in some detail. The basic problem of DFT is that the exact xc-functional is unknown. Even worse, in the few exactly soluble model cases in which it has been possible to derive the correct expression for the xc-functional, this has proved to be extremely complicated. It is thus generally thought that knowledge of the exact xcfunctional is tantamount to the exact solution of the Schrödinger equation and is thus practically out of question. The goal of the present research is to derive increasingly accurate approximations to the exact xc-functional. This goal is often framed within the so-called Jacob’s ladder,40 which allows one to rank the various approximations to the exact xcfunctional according to a well-defined hierarchy. 2.2.2. Local and Semilocal xc-Functionals. The first two rungs of the ladder are the Local Density Approximation (LDA), in which the xc-functional is composed of the Slater exchange (simply proportional to the third root of the electron density)39,41 plus a correlation functional parametrized on accurate results derived for the homogeneous electron gas, and the Generalized Gradient Approximation (GGA), in which LDA is corrected via terms depending on the gradient of the electron density.42−46 LDA and GGA are called local and semilocal xc-functionals, respectively. The third rung in the Jacob’s ladder corresponds to the so-called meta-GGA xcfunctionals, which make use not only of the local value of the electron density and its gradients, but also of the kinetic energy density for the occupied Kohn−Sham orbitals: D


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Figure 1. The typical appearance of a projected density of states (PDOS) plot is shown in the upper panel in the specific case of a TiOx/Pt(111) phase, where a Gaussian broadening of 0.1 eV of the energy levels has been used. The one-particle levels, visualized as delta peaks, are shown in the lower panel. The occupied orbitals are those below the Fermi level, and the virtual ones are those above it. Semicore levels corresponding to the 3s states of Ti atoms (a) are shown at much lower binding energies to highlight the difference between 3-coordinated and 4-coordinated species. The region of the PDOS around the Fermi level is shown in (b), where the contributions coming from the different atoms are depicted in different colors.

of the exchange operator allows one to include exact exchange over a distance of a few chemical bonds, but reduces the issues connected with Hartree−Fock exchange in metallic systems. A source of inaccuracy in the DFT approach up to the fourth rung included is represented by the lack of long-rangecorrelation dispersion effects. The origin of these effects lies in the dynamical polarization of the electron cloud induced by the presence of electrons belonging to nearby species. As the corresponding contributions to the total energy decay analytically as inverse powers of atom−atom distances, they are awkward to describe in standard DFT, because the electron density decays exponentially as a function of the distance from the atoms,36 so that its local or semilocal functionals will also decay exponentially rather than inverse-power. Despite the fact that these terms are quantitatively smaller than typical covalent or electrostatic bonding contributions, there are contentions that they can sometimes play a role in the energetics of oxideon-metal systems.70 The most common and rather effective way to account for these effects is to add semiempirical terms to the total energy in the form of C6[n]/r6 tail corrections,71−73 where r is an interatomic distance and C6[n] is a numerical coefficient, and the open debate is on how to choose the C6[n] coefficients in the physically most consistent way; see, for example, ref 74. In any event, DFT in one of its variants is nowadays the method of choice for band structure calculations in the field of supported oxide nanostructures. Being a single-particle (or oneelectron) method, DFT can scale linearly with the size of the system; that is, it presents the lowest computational cost among first-principles approaches. Moreover, it is thought to efficiently describe short-range correlation effects and long-range effects associated with electron-gas bonding, which are both difficult to deal with by using post-HF first-principles approaches. Dispersion effects are easily included via empirical corrections.

Exc ωhybrid = aEx HF,SR (ω) + (1 − a)Ex GGA,SR (ω) + Ex GGA,LR (ω) + EcGGA

(4)

where ω is the screening parameter, ExHF,SR(ω) is the shortrange part of the Hartree−Fock exchange, ExGGA,SR(ω) is the short-range part of the GGA exchange, ExGGA,LR(ω) is the longrange part of the GGA exchange, and EcGGA is the GGA correlation functional. Another very promising approach is represented by the socalled local hybrid xc-functionals, in which the Hartree−Fock component is modulated by a space-dependent factor, see, for example, refs 63, 64 and references therein, in such a way that different parts of the system are described by a different combination of Hartree−Fock and GGA exchanges, as pioneered long ago in ref 65. The advantage of using hybrid xc-functionals is recognized when treating insulators or in general systems with a substantial HOMO−LUMO gap, such as oxides, but also organic molecules, for whose energetics the introduction of a Hartree−Fock term is often beneficial.49−51 Because of the reduced SIC error, also the description of the magnetic properties of transition metal complexes is often improved by the use of hybrid DFT approaches,66 as well as systems with unpaired50,67 or localized electrons. As mentioned above, these advantages can be undermined in the case of metal (or in general gapless) systems by an erroneous treatment of the electron-gas bonding: this reverberates in instabilities and pathologies even in finite metallic systems,68,69 so that standard hybrid xc-functionals should be used with care when studying metals. Range-separated or local hybrid xc-functionals could be helpful also in this respect (apart from being computationally cheaper), as screening or damping the long-range component E


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2.2.4. Density of States (DOS). The fact that DFT is basically a single-particle approach greatly simplifies not only the associated computational effort, but especially the theoretical analysis of the chemical interactions and bonding. One of the most useful tools in this connection is provided by the definition of the density of states (DOS) and projected or local density of states (PDOS or LDOS) of the system. A pictorial definition of the DOS is given in Figure 1. One starts from the spectrum of one-electron energy levels. These are then broadened usually with a Gaussian smearing technique and plotted with the energy as the x-axis and the density of levels in arbitrary units as the y-axis. The partition line between occupied and unoccupied levels is called “Fermi energy” according to a solid-state nomenclature (see Figure 1). The occupied orbitals determine the ground-state total electron density and first-order response properties, while the unoccupied (or “virtual”) orbitals are connected with excited states and higher-order response properties of the system (polarizability, etc.). A very useful further step is taken by partitioning the one-electron orbitals into atomic components: φi(r) =

atoms

∑

φi A (r)

A

approaches, even though they can affect the experimental images; see ref 77 for a detailed discussion of this topic. 2.2.5. Post-DFT and DFT+U Methods. Apart from socalled ab initio or post-Hartree−Fock approaches, firstprinciples methods that go beyond DFT also exist. The fifth and final rung of Jacob’s ladder in fact utilizes all of the Kohn− Sham orbitals, unoccupied as well as occupied, leading to generalizations of the random phase approximation (RPA).80−82 RPA provides an approximation to the response function of the real (interacting) system starting from the DFT (noninteracting) one, leading to corrections to the total energy83,84 and at the same time allowing one to access excited states. It can be thus framed within the time-dependent density functional theory (TDDFT).85−87 A further step is taken by using the response function provided by RPA or TDDFT to calculate the dielectric function, which can then be used to screen the Coulomb potential in generalizations of the Kohn− Sham equations, leading to the so-called GW approaches.88 This class of approaches is computationally very demanding but can provide accurate corrections to quasi-particle energies, and thus band gaps as well as excited states. A different path to the solution of the Schrödinger equation passes through methods of statistical integration such as the Quantum Monte Carlo one.89 Note that all of these methods can describe long-range correlation energy terms such as those involved in dispersion interactions from first-principles. Despite the intrinsically high computational demands of these methods, orders of magnitude higher than standard DFT, their use is constantly increasing and is becoming more widespread.90 A different possible extension of DFT is the so-called DFT +U method.91 This approach was developed to correct for the deficiencies of LDA and GGA xc-functionals in describing localized (or strongly correlated) electron systems. In these systems, electron−electron scattering is not weak as in ideal metals where electrons spend most of their time in the internuclear regions: rather, electrons are here confined in localized inner-shell orbitals, as it happens with incomplete dor f-shells. Therefore, electron−electron repulsion (giving rise to correlated motion) is here non-negligible with respect to the kinetic energy of electrons, and it can be so strong as to destroy the metallic character of the system and produce insulators (the Mott−Hubbard metal−insulator transition92). For obvious reasons, such systems are awkward to describe within a oneelectron formalism, but their physics is conveniently modeled by an effective Hamiltonian such as the Hubbard one:

(5)

φiA(r)

where is the component (projection) of orbital φi(r) on atom A, usually obtained via a projection of the wave function onto basis functions localized on the atoms (e.g., via the Lowdin procedure). The DOS can then be expressed as a sum of atomic PDOS, as illustrated in Figure 1. This is very useful as one can formally distinguish the behavior of the electronic structure of a single element within a complex multicomponent system, and PDOS plots are ubiquitous in the theoretical studies in this field. The LDOS is analogously defined as the DOS resolved in real space. The LDOS has a nice bearing with experiment, as in a first approximation it can be linked to STM images. As recalled in the Introduction, scanning tunneling microscopy (STM) is the major experimental characterization technique in the oxide nanostructure field. In an STM experiment, one measures the current flowing between the sample and a conducting tip scanning over the sample surface, thus roughly speaking obtaining images of the local sample conductivity. At the theoretical level, STM images are usually simulated by applying the Tersoff−Hamann approach,75,76 which links the experimentally observed contrast to the modulation of the LDOS at the bias energy with respect to the Fermi level at a given distance from the sample: I(r) ≈

2

∑ |φi(r)| i

i : E Fermi − eVbias ≤ λi ≤ E Fermi

Hubbard Ĥ =

nn ‐ sites

∑ ij

tij +

sites

∑

Unjni (7)

ij

̂ Hubbard93

(6)

where the Hubbard Hamiltonian H is composed of a kinetic energy term tij, that is, an electron hopping between nearest-neighbor sites, which models the band structure of the system, plus an electronic repulsion term, usually limited to the on-site component Unini. The idea of the DFT+U method is to combine the versatility of the DFT approach with a mean-field Hubbard-like correction due to electron−electron interactions, that is, an energy penalty for the double occupation of specific orbitals. In detail, the DFT+U energy expression reads:

where I(r) is the tunnelling current, EFermi is the Fermi energy, and Vbias is the bias. The Tersoff−Hamann approach usually provides a qualitatively reliable picture of the surface topography, but it should also be recalled that in general it does not fully capture the quantitative details of corrugation.77 Moreover, even the qualitative picture can break down if a strong chemical interaction exists between the STM tip and the sample. A fully quantitative treatment of STM imaging going beyond perturbation theory requires explicit conductance calculations, which can in principle be realized via a nonequilibrium Green’s-function approach; see, for example, refs 78, 79. It can be finally noted that finite-temperature, finitecurrent, and finite-tip-size effects are not included in all of these

E DFT + U = E DFT + E Hubbard[{ni}] − Edc[{ni}]

(7)

where EDFT is the usual DFT total energy, EHubbard[{ni}] is the on-site Hubbard correlation functional, depending on the set of {ni}, the generalized atomic occupations for the “Hubbard” F


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the splitting) of bonding and antibonding Shockley states at the two metal surfaces on the opposite sides of the slab.107 Embedding techniques, see, for example, ref 108, whose application in this field has been rare so far, may then become competitive in terms of CPU effort. Moreover, in studying polar phases (i.e., phases that present a net dipole moment perpendicular to the surface) and periodic 3D modeling, it is advisable to include a dipole correction into the Hamiltonian109 as it effectively cancels spurious interactions due to the periodic boundary conditions in planar polar slabs.110 The size of the unit cell is often such that only DFT approaches are feasible for routine calculations with the present computational resources. Given the issues of post-HF approaches in the description of these systems, this is not a serious limitation. In recent years, the systematic application of DFT to oxide systems such as ultrathin films has indeed allowed successful explorations of the geometric and band structure (i.e., the electronic ground state) of such systems, including spectroscopic properties such as phonons111,112 and STM imaging,113 even though cases in which the accuracy of the approach is still not sufficient to predict consistently the energy ordering between different phases are known.70,114 It should be noted that so far electronic ground-state properties of supported oxide nanostructures have been mostly investigated, whereas post-DFT methods will be needed to investigate response properties involving electronic excited states. (ii) One main issue is connected with the composite or multifunctional character of supported oxide nanostructures. The match of energy scales of the different components is crucial in this context, in the sense that any reliable theoretical approach should predict with confidence (or at least with a homogeneous error) the electronic structure of both the metal and the oxide components, plus their interaction, and this is more difficult to achieve than simple relative accuracy. The issue raises from the fact that chemical bonding in metals is qualitatively different from that in the (possibly insulating) oxide component. Within DFT this translates into the fact that local or semilocal approaches (with further improvements coming from the inclusion of kinetic energy density corrections at the third rung of Jacob’s ladder and possibly semiempirical dispersion corrections) can be very accurate for describing the metal support, but present issues of too low band gaps and SIC for oxides with localized electrons. Vice versa, standard hybrid xc-functionals can be very accurate for bulk oxides, but can have serious errors in the metal component as discussed above. Still limiting to properties connected with the electronic ground state, the experience so far accumulated to solve this issue has recurred to screened hybrid xc-functionals or the DFT +U approach,98,115 finding usually a fair agreement with experiment. Unfortunately, local hybrid xc-functionals63−65 in which the space-dependent modulation of the HF exchange seems rather appropriate for multicomponent systems have not yet been tested. 2.2.7. Structure Prediction. To conclude this section, we will briefly discuss the problem of structural prediction via theoretical approaches, which is the starting point of any in-

atoms, that is, the atoms with strongly correlated electrons, and Edc[{ni}] is the mean-field approximation to the on-site correlation functional, which has to be subtracted to avoid double counting. The generalized atomic occupations {ni} are evaluated, for example, by Lowdin projecting the Kohn−Sham orbitals onto atomic-like basis functions, or via similar procedures. As these can be cheaply evaluated in terms of one-electron operators, the cost of a DFT+U calculation is only slightly larger than a standard DFT one (ignoring possible issues of convergence of the iterative self-consistent procedure). However, one has the bonus that the spectrum of one-electron eigenvalues is modified by a term U(1/2 − ni), which opens a gap between occupied (ni ≈ 1) and unoccupied (ni ≈ 0) orbitals, and the DFT band gap problem is often effectively cured. This addition thus improves upon the description of localized electronic states within DFT and is widely used, for example, to treat magnetic insulators.94 The U value can in principle be evaluated a priori,95−97 but it is common practice to derive it empirically by comparison with experiment; see, for example, ref 98. One of the advantages of the DFT+U approach is that it is possible to add the U term selectively on certain atoms. In this way, one can describe composite systems such as metal/oxide ones by simultaneously retaining a satisfactory treatment of the metal component (described via a GGA DFT method) together with a fair description of the oxide component (via the DFT+U approach). In this connection, it can be added that the U value for the oxide layer at the interface with a metal support should be reduced with respect to the value appropriate for the bulk because of screening effects.98−100 In general, polarization effects at the oxide/ metal interface are called image-charge screening. These effects are active at both the static and the dynamic levels and correspond to a stabilization of charge-separated species interacting with a highly polarizable environment (such as a metal surface) due to the formation of image or mirror charges in the metal.101 The DFT+U approach is thus especially advantageous when dealing with ultrathin oxides supported on metal surfaces. In fact, the DFT+U method can be framed within the dynamical mean field theory (DMFT),102−104 an approach in which the calculation on a physical system is mapped onto an impurity problem, which is then solved via methods of statistical physics. Although very promising, such approaches are still in a pioneering stage and have never been applied in this field, so they will not be further discussed. 2.2.6. Requirements for Theoretical Methods. Coming now to the field of supported oxide nanostructures, the requirements that theoretical approaches need to satisfy to make accurate and reliable prediction on such complicated multifunctional systems are not trivial. The two major ones are: (i) From a computational point of view, both the presence of a semi-infinite metal support and the size of the typical unit cells (which can be considerable even for commensurate systems105 and much larger for incommensurate systems such as Moiré phases) pose serious constraints on the choice of the theoretical approach. The semi-infinite metal support is usually modeled with periodic slab models, with usually a very limited (4−8) number of slab layers. This can be accurate for adhesion energies,106 but more subtle properties can require a higher number of layers. For example, more than 20 layers can be needed to damp the interaction (and thus G


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thermodynamics,123 which, starting from first-principles energy evaluations, calculates the Gibbs free energy of the system (G) as a function of temperature (T) and oxygen pressure (PO2), according to a standard expression:

depth investigation of structure−property relationships. The procedure commonly employed so far for structural prediction421 can be summarized in the following scheme: (i) collect all of the experimental data concerning the system of interest: STM images at high (possibly atomistic) resolution, spectroscopic data, chemical/physical measurements, etc.; (ii) on the basis of guesses “inspired” from the experimental data, build up trial structures; (iii) perform DFT local relaxations on the trial structures to verify that they correspond to stationary points on the energy hypersurface; (iv) compare the energies of the trial structures to single out the lowest-energy structure and the optimal system stoichiometry in the given thermodynamic conditions; (v) once the putative global minimum (or a set of candidate global minima) has been found, compare simulated and experimental values of system properties; (vi) perform a dynamic evolution of the system at moderate temperature (usually via Car−Parrinello simulations116) to verify whether the proposed structure is robust or evolves toward a lower-energy configuration. It should be noted that inspired guesses can be based on the analysis of experimental data, but can be equally inferred by using guidelines derived by previous studies on similar systems: as an example, the building principles derived for ultrathin TiOx/Pt(111) films have been used to predict a priori the structure of another ultrathin oxide phase, which has then been confirmed at the experimental level.117 An alternative to this common scheme consists of performing an unbiased search (“from-scratch”), by applying, for example, a density-functional global optimization (DF-GO) procedure. This procedure has a general validity and does not rely on any external input information (even the system stoichiometry can be treated as a variable and used as a search parameter). It has successfully been applied to the investigation of gas-phase silicon,118 metal,119 metal oxide,120 and oxidesupported metal121 nanoclusters. It can be also applied to the structural investigation of supported oxide nanostructures such as ultrathin films: its only limitation is the rapidly growing computational effort with increasing system size; see, for example, ref 122. Regarding the matching of theoretical and experimental data, it should be noted that reproducing STM images represents a first goal in the procedure of structural characterization, with STM images usually simulated by applying the Tersoff− Hamann approach75 as discussed above. However, obtaining an STM image in agreement with the experiment is not a guarantee of the correctness of the developed structural model, as different structures can give rise to the same STM pattern. The validity of the proposed structural model has thus to be checked by comparing other simulated properties (like photoelectron spectra or vibrational data, which have a computational cost comparable to that of STM simulations) to the experimentally derived quantities. Once the structure of the system has been derived and carefully checked against all of the available experimental data, a further step forward consists of considering the dependence of the structure on the thermodynamic conditions, as it is known that the variation of temperature and oxygen pressure determines a change in the equilibrium stoichiometry of the oxide. This is the aim of first-principles (or ab initio)

G(T , pO2 ) = Etot + PV + Fvib

(8)

where Etot is the total energy of the various competing phases and is directly obtained from DFT total energy calculations, while the second term is the classical PV contribution. Fvib accounts for the vibrational contribution to the Helmholtz free energy: this term can also be calculated within the DFT approach, or, alternatively, it can be derived from experimental data of phonon spectra. By using this expression, in conjunction with a reasonable estimate of the chemical potentials of the chemical species involved in the process, the phase diagram of the system can be routinely calculated. Finally, the growth kinetics should in principle be simulated, as it is known that some experimentally observed phases are only kinetically stable.124 Simulating a growth process is an extremely complicated task, and to the best of our knowledge has not been conducted so far at a fully realistic level. Some kinetic information can be derived on the basis of the thermodynamic data.125 As an example, NiO ultrathin films on Ag(100) grow initially in a precursor (2 × 1) phase, which is higher in energy with respect to a pseudomorphic (100)1 × 1 phase. This has been rationalized by showing that, when growing on a step edge of the Ag(100) surface, the (2 × 1) phase results as kinetically favored,126 but it disappears after annealing in favor of the thermodynamically more stable (100)1 × 1 phase.

3. ATOMIC STRUCTURE CONCEPTS The spatial arrangement of the atomic constituents, that is, the geometric structure, is the most evident physical property of a material. The atomic structure is determined by the intrinsic parameters of the respective electron system and by external degrees of freedom, such as the boundary conditions and the thermodynamic environment. In this section, we will adopt a holistic approach of structure description of oxide nanosystems, in that we treat the geometry as the central property, and discuss structure concepts with their major determining factors, thereby emphasizing novel concepts with regard to the known oxide bulk structure geometries. To facilitate a systematic discussion, in a reductionist analysis we investigate the main factors that determine the geometry of oxide−metal nanostructures separately and sequentially, but keeping in mind the importance of their interactions and synergistic effects. The latter concern, for example, the delicate interplay between structural, electronic, and magnetic degrees of freedom, is significant in highly correlated oxide systems and in some cases decisive in stabilizing a particular nanostructure geometry. In the following, the parameters considered as structure-relevant in the present context of oxide−metal hybrid systems will be first identified: they will be listed and defined and subsequently illustrated with the help of prototypical oxide-on-metal systems. In view of the preponderance of surface and interface effects in determining the stability and properties of nanomaterials in general, the geometry (section 3.1) and chemistry (section 3.2) at the metal−oxide interface plays a major role for the structure of the metal-supported oxide. The interface parameters may be separated into elastic effects involving the symmetry and the lattice parameters of the metal substrate in relation to those of H


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and “broken symmetries” develop, which provide a gradual matching of lattices and thus promote the epitaxial order in the growing film;149,150 this will be further discussed below. Another interesting observation is that the symmetry can be conserved across the oxide−metal interface for oxide nanolayers, which form high-order commensurate or incommensurate superlattices, thus for moderately or weakly interacting systems; this is encountered for ceria overlayers on Pt(111), Rh(111), Ru(0001), and Ni(111) surfaces149,151−153 and for Fe oxide on Pt(111).154 Let us consider first the case of the ceria nanolayers. Ceria grows on Pt(111)151 and Rh(111)152 in the form of ordered CeO2 islands at elevated substrate temperatures, with (111) faces parallel and rotationally aligned to the main azimuth directions of the substrate, as evident from the LEED pattern in Figure 2a, which has been designated as a (1.4 × 1.4) pattern. This pattern consists of the Rh(111)1 × 1 reflections (the outer hexagon) and the CeO2 (111)1 × 1 reflections (inner hexagon), which are superimposed incoherently.152 The STM image of Figure 2b shows ceria island structures with hexagonal shapes and an average height of 5−6 Å, suggesting that they

the oxide overlayer, and chemical effects involving the affinity of the substrate toward oxygen and the oxide’s metal element and the possible oxidation state(s) of the metal constituent of the oxide. The former give rise to the lattice (mis)match at the interface and related strain effects, whereas the latter determine the chemical bonding at the interface, charge transfer between metal and oxide phases, and possible intermixing or doping of interface-near regions. The thermodynamic and kinetic variables during the oxide fabrication process comprising substrate temperature, oxygen pressure, metal coverage, and deposition rate as well as the detailed preparation procedures have a significant influence on the geometries of oxide nanostructures, their morphologies, and on the corresponding phase stabilities. This complex of parameters will be discussed in connection with the surface phase diagrams of oxide−metal hybrid systems (section 3.3). Moreover, the phenomenon of kinetic stabilization of particular oxide phases and global versus local minimum structures in the potential energy hypersurfaces will be examined. The dimensionality of oxide nanophases is another structure-determining aspect in oxide nanostructure systems (section 3.4): the structural flexibility in connection with confinement and finite size effects gives rise to novel metal−oxygen building blocks and arrangements, which are particular to low-dimensional systems. We will investigate these effects in two-dimensional (2D) oxide nanolayers, onedimensional (1D) oxide nanowires, and (quasi)-zero-dimensional (0D) oxide nanoparticles (clusters). The formation of defects as a structure concept in oxide nanophases as a means for strain relief will be introduced (section 3.5), and the condensation of defects into ordered arrays will be discussed in relation to the formation of nanopatterned oxide surfaces as substrates for the directed growth of hierarchical nanoparticle superlattices. 3.1. Interface Geometry

It is a well-known fact that the symmetry and crystallographic orientation as well as the lattice constant of the substrate are decisive parameters for epitaxial growth of thin films.127 Accordingly, it is expected that the symmetry of the substrate is imprinted onto the overlayer; for example, on (100) surfaces of fcc support surfaces the growth of square oxide overlayer lattices is expected, and on fcc (111) or hcp (0001) surfaces hexagonal overlayer orientations should be formed. Indeed, this expectation is fulfilled in a number of examples. Epitaxial rock salt structure MgO(100), NiO(100), CoO(100), or MnO(100) oriented thin films have been reported to grow on Ag(100) surfaces,128−134 MgO(100) on Mo(001),135 and the formation of MgO(111) on Ag(111),136,137 NiO(111) on Au(111),138,139 wurtzite ZnO(0001) on Ag(111),140 and corundum structure V2O3(0001) surfaces on Pd(111),141 Rh(111),142 Re(0001),143 and Au(111)144,145 substrate surfaces has been communicated. However, for systems with larger lattice mismatch, symmetry conserving epitaxially ordered single crystalline surfaces are only observed for thicker oxide films,141,142 whereas coherent growth structures down to the oxide−metal interface are only realized if the lattice mismatch between substrate and overlayer is small, of the order of a few percent or less.127 The surprising observation of epitaxial growth for large lattice-mismatched systems, such as NiO(100), CoO(100), and MnO(100) on Pd(100) (lattice mismatch 7%, 9%, and 14%, respectively),146−148 thus deserves further examination. It has been noticed that in these poorly lattice-matched systems, intermediate layers at the interface with particular geometries

Figure 2. (a) (1.4 × 1.4) LEED pattern of ∼1 ML of CeO2 on Rh(111) (electron energy E = 84 eV). (b) STM image of 0.9 ML CeO2/Rh(111) after annealing at 585 °C (sample bias VS = +1.98 V; tunneling current IT = 1.03 nA). (c) STM image of ∼0.5 ML CeO2/ Rh(111) after annealing at 585 °C (VS = +0.89 V; IT = 0.83 nA). Reprinted with permission from refers 149 and 152. Copyright 2002 and 2004 Elsevier. I


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Figure 3. (a) STM image (55 × 55 Å2 ; VS = +0.9 V; IT = 0.3 nA) of a FeO(111) monolayer on Pt(111). An atomic periodicity of 3.1 Å is modulated by a large 25 Å periodicity creating a Moiré superstructure.157 (b) Model of the FeO(111) monolayer on Pt(111), as first proposed by Galloway et al.419 The overlayer has a lattice constant of 3.11 Å and is rotated by 1.3° against the [110] direction, forming a |8 2/−1 10| or (√84 × √84)R10.9° coincidence structure with coincidence site 1 and the large unit cell indicated. Sites 2−4 indicate sites of other coincidence structures in the second and third FeO bilayers. Reprinted with permission from ref 157. Copyright 1998 American Physical Society.

FeO−Pt interface, but this doesn’t seem to be energetically costly because the interfacial bonding is only weakly directional (thus the almost incommensurate interface). Instead, the lateral lattice expansion and the concomitant reduction of the vertical interlayer spacing, as measured by Kim et al.159 using X-ray photoelectron diffraction, lead to a reduction of the polarity of the bilayer and thus to a reduction of the electrostatic energy of the system. This polarity compensation, in conjunction with the low surface energy of the oxygen terminated surface, stabilizes the hexagonal FeO bilayer system. Incidentally, an epitaxial FeO(111) monolayer has also been observed on a Pt(100)-hex substrate;160 the latter is a reconstructed surface with a quasihexagonal layer of Pt surface atoms resting on top of the square (100)1 × 1 bulk lattice. Two FeO(111)-type bilayer superstructures have been detected by STM and LEED, which have been described as c(10 × 2) and (9 × 2) coincidence lattices on the dereconstructed Pt(100)1 × 1 surface, where the hexagonal reconstruction has been lifted during the FeO overlayer growth.160 The FeO(111)-type coincidence structures found on Pt(100) have a correspondence in similar structures of Co oxide on Pd(100)161 and Ir(100),162 which will be discussed later in this section. For a large lattice mismatch between the oxide overlayer and the substrate, the strain term for a symmetry-conserving interface structure may increase the total energy of the system such that a “broken-symmetry” situation may win the competition for the lowest energy structure of the system. This is observed in many cases, in particular in the oxide nanolayer regime, and will be discussed in the following. A pertinent question concerns the values of the lattice constants, which have to be taken to derive the lattice mismatch and the interfacial strain: traditionally, the bulk lattice constants have been considered.127 However, in the limiting case of one or two monolayers, the hypothetically isolated overlayer (i.e., without the substrate) would not display the bulk lattice constant, but would exhibit a lattice shrinkage. Intuitively, this may be understood by a bond conservation picture involving the number of coordinated atoms: the metal−oxygen bond length in a diatomic molecule is shorter than in a two-dimensional layer, where it again is shorter than in a three-dimensional solid. This problem has been addressed recently in a quantitative way using DFT+U calculations.98,163 for NiO on Ag(100) and

consist of two hexagonal O−Ce−O triple layers in a CeO2(111) stacking sequence. The double (111)-type triplelayer slab with oxygen termination on both interfaces and its rotational alignement to the substrate has been rationalized by DFT/GGA calculations and found to be energetically favorable.149 The higher resolution STM image of Figure 2c reveals the hexagonal lattice and a mesh of bright lines forming a superlattice with a periodicity of ∼19 Å, which has been interpreted in terms of a Moiré superstructure, created by the interference between the Rh(111) and the CeO2 overlayer. Both Moiré and LEED patterns are consistent with a coincidence structure with 7aRh = 5aCeO2 yielding a periodicity of 18.83 Å, in good agreement with experiment (if the CeO2 lattice is allowed to contract by 1.56%). This rotational symmetry alignment and the formation of a coincidence lattice involving moderate overlayer strain at the metal−ceria interface apparently make the growth of ordered ceria films possible on such substrates, which are formally lattice mismatched by as much as 30−40%.151−153 It has to be added, however, that the ceria films remain granular with increasing thickness on the poorly lattice-matched metal substrates and that smooth and continuous ceria films are difficult to obtain, although some progress toward the latter has been made by using a kinetically limited growth recipe.155 In contrast, the growth of CeO2(111) on Cu(111) has been reported to yield good long-range order and a smooth and atomically flat oxide film morphology.156 Fe oxide grows in epitaxial FeO(111) bilayer form on Pt(111),154,157 despite the ∼10% lattice mismatch between substrate and overlayer. Beyond monolayer film thickness, the FeO(111) growth is terminated, and three-dimensional Fe3O4(111) islands nucleate. The FeO(111) layer consists of a hexagonal close-packed Fe−O bilayer that is polar, that is, with a permanent dipole moment perpendicular to the surface, and laterally expanded as compared to bulk FeO and slightly rotated against the Pt substrate; the Fe atom layer is at the Pt interface, and the outer surface is oxygen terminated. Depending on the preparation conditions, four coincidence structures with slightly different lattice constants and rotation misfit angles with respect to the Pt(111) substrate have been detected by STM measurements,157,158 for example, a (√84 × √84)R10.9° superstructure as shown in Figure 3.157 The lattice expansion of the FeO layer increases the lattice misfit at the J


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Pd(100). Without the metal substrate, the free-standing NiO monolayer typically exhibits a contraction as compared to the bulk, with a Ni−Ni distance at the energy minimum close to 2.8 Å, which corresponds to a 5−6% contraction from the experimental value and the optimized bulk value obtained in a hybrid DFT study. If the monolayer is deposited on a metal support, the oxide−metal interaction will increase the coordination number again, thus counteracting this contraction. It is predicted that NiO on Ag(100) adopts a 2% compression and on Pd(100) a 3.6% compression. The resulting interfacial stress is thus considerably smaller than the value of approximately 7% estimated on the basis of the NiO−Pd lattice mismatch, if one uses the respective bulk lattice constants. An important point which should be mentioned in this connection is that, when evaluating interfacial stress, the “stickiness” of the supported oxide nanophase, i.e., the second derivative of the effective oxide plus metal/oxide-interaction energy, should be considered, as discussed in ref 98. The geometries of metal-supported oxide nanostructures exhibit a fascinatingly large, even confusing variety of novel structures for oxide−metal systems across the Periodic Table, but with some intuition and if viewed cum grano salis some common structure concepts may be extracted. For the oxides of the second half of the 3d transition metal (TM) series (Mn− Ni), which include the stable rock-salt structure TMO monoxides, planar essentially nonpolar (100)-derived TM−O overlayer structures, polar (111)-type (quasi-)hexagonal TM− O bilayers, and (111)-derived O−TM−O hexagonal trilayer structures may be identified. The latter are nonpolar with a formal TMO2 stoichiometry and with a formal oxidation state of the TM atom higher than 2+; their phase stability requires higher oxidizing conditions than the TM−O structures and appears to increase in going from right to left in the Periodic Table. There is also another category of TMO structure concepts, which consists of single layers but with mixed trigonal and/or quadratic TM−O building blocks and which includes a periodic buckling that results in uniaxially ordered structures. The early 3d TM oxides (e.g., Ti, V) are characterized by a high flexibility of oxidation state and a rich variety of different oxide nanophases, including some of the above-mentioned structure units but also particular novel and nanophase-specific building blocks. Mn oxide is a borderline case between early and late TM oxides, and a very complex 2D phase diagram has been reported (for example, on the Pd(100) substrate13). Because the focus of this subsection is on the interface geometry, we will emphasize in the following mainly elastic aspects and the reduction of interfacial strain as the most important structure determining driving force. 3.1.1. TM−O (100)-Derived Structures. As pointed out, a NiO(100) monolayer has an effective lattice mismatch to Ag(100) of only 2%; accordingly, a NiO(100)1 × 1 monolayer structure has been generally observed.21,126,164−167 For low coverages, the NiO(100)1 × 1 phase grows in the form of islands, although the detailed morphology depends on the preparation conditions. Figure 4 shows STM images of NiO(100)1 × 1 structures on Ag(100) obtained after deposition of 0.65 ML NiO at room temperature followed by annealing at 600 K in oxygen.126 Well-ordered rectangular islands with dark contrast in the STM are visible (lower panel), due to NiO(100) islands embedded into the Ag substrate. The high-resolution image (upper panel) displays the atomically resolved (1 × 1) structure (the Ni atoms are imaged bright), and a mosaic contrast modulation with darker patches is also

Figure 4. STM images of 0.65 ML NiO/Ag(100), deposited at room temperature and postannealed at 600 K in 1 × 10−7 mbar O2. (a) (200 × 200 Å2; VS = −0.85 V; IT = 0.1 nA; Fourier filtered image). (b) (1500 × 750 Å2; VS = −1 V; IT = 0.1 nA). Reprinted with permission from ref 126. Copyright 2011 American Physical Society.

apparent. DFT calculations126 have revealed that the origin of the dark-contrast defects is due to the slight strain in the NiO(1 × 1) layer, which may be released by the insertion of occasional patches of a NiO bilayer underneath the monolayer, as this increases the optimal lattice parameter of the resulting composite system and thus lowers its total energy. On Pd(100), the effective lattice mismatch of a NiO(100) monolayer is ∼3.6% (see above), and a NiO(100)1 × 1 interfacial layer has not been detected. Instead, a wetting monolayer with a well-ordered c(4 × 2) superstructure with respect to the Pd(100) substrate has been observed.168,169 This c(4 × 2) structure has been rationalized first on the basis of LEED I(V)170 and STM measurements169 by a model involving a slightly distorted NiO(100) layer, into which cation vacancies are introduced periodically, thus defining the c(4 × 2) superstructure; this leads to a formal stoichiometry of Ni3O4 for this layer. The NiO(100)-derived c(4 × 2) Ni3O4 vacancy model has been confirmed subsequently by DFT hybrid calculations.171 Interestingly, the c(4 × 2) oxide monolayer structure appears to be of a more general phenomenon, because it is not unique to the NiO/Pd(100) system, but has also been observed for MnO/Pd(100)112 and CoO/Pd(100);172 recently, a c(4 × 2) Co3O4 monolayer has also been reported to grow on a monolayer of metallic Co, as an interlayer, on a Ir(100) substrate.173 In the present dicussion, we will use the c(4 × 2) Co3O4 monolayer on Pd(100) as a prototypical example to investigate the geometry and electronic properties of this oxide monolayer structure. Figure 5 displays STM images of the c(4 × 2) Co oxide on Pd(100) and a corresponding LEED picture, obtained after deposition of 0.8 ML Co in 1 × 10−6 mbar oxygen at room temperature followed by annealing at 570 K in oxygen.172 The large-scale STM image of Figure 5a shows large flat terraces covered by the c(4 × 2) wetting layer; some isolated second K


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Figure 6. Schematic drawing of the top (left) and side (right) views of the structural optimized model of the Co3O4 monolayer on Pd(100), corresponding to the most stable RH1 lateral registry (O atoms (small spheres) on top of Pd surface atoms). The optimized positions refer to the most favorable FM ordering, but a schematic explanation of the tested antiferromagnetic (AFmagnetic 1 and AFmagnetic 2) is also given: “+” and “−” signs represent Co spins up and down, respectively. Co1 and Co2 denote two distinct types of Co atoms, and the arrows (left panel) show the planar displacements of Co2 and O atoms adjacent to the Co vacancy. The vertical bucklings b (right panel) quantify the out-of-plane deviation fom the perfect planar monolayer, and z the distance between overlayer and substrate. Reprinted with permission from ref 172. Copyright 2010 Elsevier.

Figure 5. STM images of the c(4 × 2) Co oxide structure on Pd(100): (a) (2000 × 2000 Å2 ; VS = +2.0 V; IT = 0.1 nA). (b) (430 × 430 Å2 ; VS = +2.0 V; IT = 0.1 nA). (c) (150 × 150 Å2 ; VS = +1.0 V; IT = 0.1 nA). In panel (c) the rectangular (A) and associated rhombic (B) c(4 × 2) unit cells are drawn. (d) LEED pattern of the c(4 × 2) surface (electron energy E = 104 eV). Reprinted with permission from ref 172. Copyright 2010 Elsevier.

layer islands, apparently with a different structure, are also visible. The higher resolution images (b) and (c) display atomic corrugation and indicate the well-ordered c(4 × 2) periodicity, which is also confirmed by the LEED pattern in panel (d). The rectangular and rhombic c(4 × 2) unit cells are indicated in (c) as A and B, respectively; a step edge is visible in the lower part of (b). The c(4 × 2) Co oxide structure on Pd(100) has been modeled by GGA and screened hybrid DFT calculations using the original pseudomorphic (100)-type cation vacancy model of Agnoli et al.170 as a starting point, and the structurally optimized model for the Co3O4 monolayer is shown in Figure 6.172 Several overlayer-substrate registries and different magnetic orderings, two distinct antiferromagnetic (AF) as well as a ferromagnetic configuration (FM), have been tested. Figure 6 contains the most stable, so-called RH1 lateral registry with O atoms located on top of the Pd surface atoms and Co sitting in the 4-fold hollow sites. Independent of the functional approach used in the calculations (PBE or hybrid HSE), the RH1 solution with FM arrangement turns out as the most stable result. The creation of the Co vacancy induces a general rearrangement of the atomic positions, ultimately resulting in the formation of zigzag Co2 chains and single Co1 atoms sandwitched between two nearest neighbor vacancies. The structural relaxation destroys the ideal coplanarity of the CoO monolayer by inducing a vertical buckling as sketched in the right part of Figure 6. The electronic and magnetic properties of the Co3O4 phase are substantially modified as compared to an ideal CoO(100) monolayer, containing two magnetically and electronically inequivalent Co atoms. The latter are reflected in the STM data, as revealed in Figure 7,172 where experimental and simulated STM images are compared. The comparison of the analysis of the densities of states (DOS) and the experimental STM image indicates that the bright protrusions correspond to the Co atoms. Both types of Co atoms are imaged by STM at low bias, with the Co2 atoms

Figure 7. Comparison between simulated and experimental (inset) STM images of Co3O4 on Pd(100) at sample bias VS = +0.08 V. Co atoms appear as bright maxima, whereas dark depressions correspond to the Co vacancies. The rhombic c(4 × 2) unit cell and the Co2 zigzag chains are indicated. Reprinted with permission from ref 172. Copyright 2010 Elsevier.

forming well-defined Co2−Co2 zigzag chains and the Co1 atoms sitting in the center of the rhombic unit cell defined by the Co vacancies. The latter are imaged as dark depressions in Figure 7. The O atoms are not detected in STM because their projected DOS is much lower than the corresponding density of Co1 and Co2 states. The introduction of metal cation vacancies and the deviation from the 1:1 stoichiometry of CoO to Co3O4 thus appear to provide a favorable mechanism to stabilize a pseudomorphic (100)-type monolayer on Pd(100), ensuring a partial compensation of the compressive strain associated with the large lattice mismatch of the oxide/metal epitaxy. In the case of Ni3O4 and Co3O4, the resulting c(4 × 2) structures provide wetting monolayers with excellent long-range order, indicating that this strain compensation mechanism is very effective. For the c(4 × 2) Mn3O4 monolayer on Pd(100), a formally ∼14% lattice mismatched system (using the bulk lattice constants for this estimate in the absence of values for a free-standing MnO monolayer), the discussed strain compensation by Mn vacancy L


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oxygen at room temperature on to Pd(100) displays the c(4 × 2) structure as wetting layer and a square second-layer island of the NiO(100) type. For the formation of Co oxide and Mn oxide thin films on Pd(100), the growth sequence beyond the monolayer stage is more complicated, and a variety of structures have been observed in subsequent layers, until the growth converges to bulk type oxide phases.147,174 This more complex behavior in the latter cases is, in the end, the result of the larger lattice mismatch at the interface. 3.1.2. TM−O (111)-Derived Bilayer Structures. Hexagonal close-packed metal−oxygen bilayer structures, which are related to the (111)-type plane stacking of the rock-salt lattice, have been detected for oxide nanolayers on both fcc (100) and (111) substrate surfaces. A paradigmatic example is the hexagonal bilayer of FeO(111) on Pt(111),157 which has been discussed above. Whereas the lattice parameters of this FeO(111) bilayer are modified with respect to the corresponding bulk values to compensate for layer polarity, the oxygen and Fe layers remain strictly hexagonal and planar. The interaction of the different lattices of overlayer and substrate gives rise to a Moiré pattern, which is visible in STM (see Figure 3). Moiré patterns have also been invoked to explain the complex STM contrast of the so-called “wagon-wheel” or “pinwheel” structures observed for reduced V-oxides on Pd(111) and Rh(111) surfaces,201,175 for TiOx/Pd/TiO2(110),176 and for TiOx ultrathin films on Pt(111).177 Figure 10a shows an STM image of the “wagon-wheel” V-oxide on Rh(111), where the contrast lines giving rise to the “wagon-wheel” motif are indicated.175 The sharp LEED pattern of Figure 10b indicates good overlayer ordering, and the autocorrelation diagram of Figure 10c confirms the Moiré and primitive unit cells of the system. This “wagon-wheel” V-oxide phase has been interpreted in terms of a (7 × 7)R21.8° superstructure of a VO(111) bilayer, with V atoms at the Pd interface and O atoms terminating the outer vacuum surface. The VO(111) bilayer model was tested by DFT/GGA calculations, and the simulated STM images showed reasonable agreement with the experimental ones (perfect agreement, though, has not been achieved due to computational limitations associated with the flexibility in registry with the substrate of this high order commensurate overlayer).175 A different interpretation of the “wagon-wheel-like” phase of TiOx on Pt(111) has been given recently by Barcaro et al.124 This structure with a formal stoichiometry of TiO1.2 is composed of a close-packed arrangement of Ti atoms with an oxygen layer on top, but there are inequivalent 4-fold coordinated and 3-fold coordinated Ti atoms in the layer with different (electronic) contrast in the STM. The dark features in the STM images, marking the hubs of the “wagon wheel”, are in this model real holes, in which the substrate is exposed. Using this model, a successful agreement between simulated and experimental STM images has been obtained.124 On (100) substrate surfaces, the oxide (111) bilayer structures display significant distortions from the ideal hexagonal bilayer symmetry. Oxide bilayer structures of the (111)-type have been observed for FeO on Pt(100)1 × 1,160 CoO on Pd(100),161 CoO on Ir(100)1 × 1,162 and EuO on Ni(111) surfaces.178 For FeO(111) on Pt(100), two superstuctures, c(2 × 10) and (2 × 9), have been communicated, for CoO(111)/Pd(100) a (9 × 2) superstructure, and for CoO(111)/Ir(100) a c(10 × 2) superstructure have been reported. This apparent similarity in the superperiodicities suggests a common building principle. We illustrate the latter

formation, however, seems to be incomplete. This is suggested by the STM measurements presented in Figure 8.13 In Figure

Figure 8. STM images of the c(4 × 2) Mn3O4 monolayer on Pd(100). (a) (500 × 500 Å2 ; VS = +0.5 V; IT = 0.2 nA). (b) (200 × 200 Å2 ; VS = +0.5 V; IT = 0.2 nA). The c(4 × 2) structure occurs in domains parallel and perpendicular to Pd step edges and is separated by disordered domain boundary regions. Two step edges are apparent at the upper left and lower right-hand side corners of panel (a). Reprinted with permission from ref 13. Copyright 2009 IOP Science.

8a, the ordered c(4 × 2) Mn3O4 structure is seen to grow in elongated narrow domains with two orthogonal orientations, parallel and perpendicular to the Pd substrate step edges (in the upper left and lower right-hand side corners of the figure). The c(4 × 2) domains are separated by disordered stripes of domain boundaries (see Figure 8b), which may serve as a means of stress relief. There is a tendency of the Mn3O4 c(4 × 2) structure to nucleate preferentially at the step edges of Pd(100), which also provide a natural route for strain relaxation. Recent experiments of the growth of the Mn3O4 nanostripes on vicinal stepped Pd surfaces, discussed later in this Review, support this conjecture. The c(4 × 2) TM3O4 structures are stable phases only in the first monolayer, which emphasizes the additional importance, apart from the elastic strain relief, of electronic interactions across the interface. For Ni oxide on Pd(100), the c(4 × 2) Ni3O4 monolayer provides a suitable template for the growth of NiO(100) oriented films.146 Figure 9 shows an STM image, which illustrates this point: 1.2 ML Ni evaporated reactively in

Figure 9. STM image of 1.2 ML Ni evaporated onto Pd(100) reactively in oxygen at room temperature (165 × 165 Å2 ; VS = +1 V; IT = 0.1 nA). The square bright contrast is a second layer island of the NiO(100) type. Reprinted with permission from ref 146. Copyright 2006 Elsevier. M


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of oxygen. Occasionally, the two structures have been observed in coexistance at the surface, which is in agreement with the very similar total energy found in the DFT+U calculations. Figure 11 presents STM images of the (9 × 2) CoO

Figure 11. STM images of the (9 × 2) CoOx monolayer on Pd(100). (a) (400 × 400 Å2 ; VS = +0.1 V; IT = 0.6 nA). The black arrows indicate an antiphase domain boundary; at the upper right corner an area with a c(4 × 2) structure is visible. (b) (40 × 40 Å2 ; VS = +0.01 V; IT = 1.0 nA). B designates the (9 × 2) unit cell of the coincidence lattice, C the quasihexagonal unit cell of the overlayer. Reprinted with permission from ref 161. Copyright 2011 American Institute of Physics.

structure,161 taken under different tunneling conditions. The large scale image (a) shows rows of rod-like maxima separated by darker stripes along the [011] direction, which resolve into distinct atomic protrusions along the orthogonal [011] direction at low sample bias (image b); the latter protrusions are modulated dark and bright, and four brighter maxima are separated by four darker maxima. The unit cell analysis (Figure 11b) gives the (9 × 2) unit cell B with respect to the underlying Pd(100) substrate, corresponding to a (8 × 9) coincidence lattice, while the primitive unit cell C, joining nearest-neighbor protusions irrespective of their intensity, defines a quasihexagonal overlayer lattice. On the basis of the suggestions from the experimental STM images, the (9 × 2) Co oxide monolayer has been modeled by a hexagonal close-packed Co layer in contact with the metal surface, with the O atoms in the layer above occupying one-half of the hollow sites of the Co lattice; see Figure 12.161 Because of the (8 × 9) coincidence lattice, the in-plane positions of the Co(O) atoms with respect to the Pd surface atoms underneath vary: the Co atoms at the boundary of the unit cell are in a bridge position between two Pd atoms, while the Co atoms in Figure 10. (a) STM image of the “wagon-wheel” V-oxide phase on Rh(111) (80 × 80 Å2 ; VS = +2 V; IT = 0.1 nA). The “wagon-wheel” and two unit cells A, B are drawn. (b) LEED pattern of the surface of (a) (electron energy E = 60 eV). (c) Autocorrelation diagram of the STM image (a). Moiré unit cell A and primitive unit cell B are indicated. Reprinted with permission from ref 175. Copyright 2005 American Physical Society.

with the help of Co oxide on Pd(100),161 where high-resolution STM data together with extensive DFT+U calculations are available. The case of Co oxide on Pd(100) is particular, because two well-ordered wetting monolayers have been found: the c(4 × 2) Co3O4 phase, which has been discussed above, and the (9 × 2) phase. The latter is generated experimentally at a slightly higher coverage of Co (1.0 ML as compared to 0.75 ML for the c(4 × 2)) and a somewhat lower chemical potential

Figure 12. (a) Schematic model of the (9 × 2) CoO unit cell on Pd(100) in side view, showing the Pd−Co−O stacking. (b) Top view of four (9 × 2) unit cells. Pd atoms are depicted in white; O atoms in yellow; Co atoms in blue and red, reflecting the AF-3 magnetic ordering of the CoO monolayer. Reprinted with permission from ref 161. Copyright 2011 American Institute of Physics. N


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complicated by the incommensurate nature of the overlayers and the resulting Moiré structures in STM. This is exemplified by the so-called HEX-I phase of Mn oxide on Pd(100), as presented in Figure 13.13 Figure 13a shows a wide scan STM

the middle of the cell occupy hollow sites of the (100) surface. In the relaxed (9 × 2) structure, the Co atoms are positioned at about 2.13 Å above the Pd surface, while the O atoms are only about 0.5 Å above the Co atoms. The interplane Co−O distance and consequently the polar character of the layer are thus appreciably reduced with respect to the (111) stacking in the bulk, 1.2 Å. Both the Co and the O layers develop a wavy modulation of the height with an amplitude of 0.12 and 0.18 Å, respectively; the bilayer geometry becomes significantly perturbed, giving rise to a nonuniform and locally modulated polarity in the layer. The wavy modulation in height and polarity is at the basis of the brightness modulation in the STM images, where the intensity of the protrusions correlates with the height of the O atoms in the cell. As indicated by the comparison of simulated and experimental STM images, it is the O atoms, not the Co atoms, that are mostly visible as maxima in the STM at low positive bias voltages.161 As revealed by the analysis of the projected DOS, this is due to the height difference between Co and O species and to the fact that the tunneling states on the O atoms are much more diffuse as they carry a negative charge. It is worth pointing out that this is in contrast to the c(4 × 2) phase, where the Co atoms are imaged bright in STM at similar bias conditions, because the Co and O ions lie at roughly the same height and Co is in a higher oxidation state. As Co is a magnetic atom, different spin arrangements in the (9 × 2) have been tested: the antiferromagnetic order AF-3,179 in which ions with parallel spin orientation are arranged in a zigzag fashion running parallel to the short side of the unit cell, see Figure 12b, was found to be the lowest energy ground state. 3.1.3. Hexagonal O−TM−O (111) Trilayer Structures. Extending the rock-salt (111) stacking of bilayers by another oxygen layer yields O−TM−O trilayer systems, where the TM layer is sandwitched between two hexagonal oxygen layers. Such trilayer systems have been shown to be stable as the intrinsic surface oxide phases on various noble metal surfaces.180−184 The formal stoichiometry of the oxide trilayer is TMO2, but because the oxygen layer at the interface is shared with the metal support, the effective TM oxidation state is lower. It has been suggested that the hexagonal Co oxide phases observed on Pd(100) in the 1−2 ML regime172 and the quasi-hexagonal c(8 × 2) CoOx layer on Ir(100) at around 1 ML oxide coverage185 may be of the O−Co−O trilayer type, but no rigorous experimental or theoretical analyses have been performed. The O-rich phase that has been observed upon oxidation of the FeO(111) bilayer on Pt(111) at elevated O2 pressures has been interpreted as a O−Fe−O trilayer structure.158,186 DFT/GGA calculations indicated that the oxidation reaction occurs site specific within the large FeO(111) Moiré unit cell, yielding close-packed FeO2−x islands with a (√3 × √3)R30° superstructure, as a result of relaxation effects within the trilayer.158 The high catalytic activity of this FeO2−x trilayer phase in the low temperature CO oxidation reaction will be discussed in section 4.4.2. The hexagonal phases of V oxide on Pd(111)141,187 and of Mn oxides on Pd(100)13,112 observed at higher oxygen chemical potentials are well-documented examples of monolayer phases, which have been interpreted within the TMO2 trilayer concept. It appears that a certain propensity toward higher oxidation states of the metal cations is beneficial for the formation of O−TM−O trilayers in oxide heteroepitaxy, as it is encountered toward the left of the transition metal series in the Periodic Table. The analysis of the TMO2 trilayers is

Figure 13. (a−c) HEX-I distorted Mn oxide phase on Pd(100). (a) Large-scale STM image (2000 × 2000 Å2 ; VS = −1 V; IT = 0.15 nA). (b) LEED pattern (electron energy E = 96 eV). (c) High-resolution STM image (100 × 100 Å2 ; VS = +0.5 V; IT = 0.13 nA). Inset: (20 × 17 Å2 ; VS = +0.6 V; IT = 0.15 nA). (d) LEED pattern of the HEX-I undistorted Mn oxide (E = 90 eV). Reprinted with permission from ref 13. Copyright 2009 IOP Science.

image of the Pd(100) surface after evaporation of Mn with an oxygen pressure in the range (1−5) × 10−6 mbar, 670 K substrate temperature. Two domains consisting of stripe structures are recognized, which are rotated by ∼90° with respect to each other. The higher resolution image in Figure 13c reveals that the stripes consist of a complicated pattern of maxima, which are separated by darker troughs. The periodicity of the troughs as measured from the STM image is 22−23 Å. In the high-resolution image of the inset of Figure 13c, a quasihexagonal arrangement of maxima is recognized with a lattice constant of ∼3.1 Å. The LEED pattern from this surface (Figure 13b) displays characteristic elongated reflections due to unresolved double spots, which may be interpreted in terms of a quasi-hexagonal overlayer lattice on a square substrate. The analysis of this LEED pattern yields a distorted hexagonal overlayer with b1 = 2.94 Å and b2 = 3.14 Å, which is incommensurate in the b1 direction, but with a row matching condition along the b2 direction. This Mn oxide phase has been designated as HEX-I distorted. The STM image of Figure 13c may be understood in terms of a Moiré pattern as a result of the interference of the Mn oxide overlayer lattice with the square Pd(100) surface lattice. A geometrical simulation, taking b1/aPd ≈ 16/15, displays modulations in the form of broad lines in the correct direction with an average periodicity of ∼22 Å, as seen in the actual experiment.13 Under preparation conditions similar to those discussed above, another Mn oxide phase has been occasionally observed on Pd(100), which appears to be related to the distorted HEXO


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I, but with a perfect hexagonal pattern in LEED (see Figure 13d), corresponding to b1 = b2 = 3.14 Å. This undistorted HEX-I phase has been observed less frequently than the distorted HEX-I and thus seems to be less stable, requiring particular conditions for kinetic stabilization. It has been conjectured that the larger lattice mismatch in the b1 direction (formally 14% against 7%) and the related interfacial strain are responsible for the lower stability of the undistorted HEX-I as compared to the distorted one.13 The HEX-I Mn oxide phases on Pd(100) have been analyzed by screened hybrid DFT by Franchini et al., using the O−Mn− O trilayer model.112 Because of the incommensurate nature of the overlayer, the computational modeling had to be done with an unsupported trilayer setup, considering only freestanding thin layers. A rigid sphere model of the MnO2 trilayers is given in Figure 14a,b.112 As a direct consequence of the free-standing nature, the optimized planar lattice constant, a = 2.89 Å, is considerably smaller than the corresponding calculated bulk constant of 3.09 Å; the interlayer distance is also reduced by 25% for the same reason. To simulate the experimentally observed distorted HEX-I structure, the energy required to distort the perfect hexagonal trilayer by shrinking the lattice constant b by 7% has been computed: the energy cost amounts only to 70 meV, indicating that the distorted layer is only marginally less stable than the undistorted one (see the phase diagram of Figure 14c). Considering that the model neglects the effect of the Pd(100) substrate and that the distortion allows for a better lattice matching with the substrate, it is concluded that the distorted trilayer model of the HEX-I is supported by the DFT calculations.112 The soundness of the MnO2 trilayer model has also been tested by comparing the experimental phonon data, measured by HREELS, with the theory predictions. DFT locates the distinct dipole active mode between 72 and 75 meV, which compares very well with the measured value of 70.5 eV (Figure 14d). The displacements corresponding to this mode are sketched in the inset, showing that this particular vibration is linked to the alternatively stacked hexagonal layers of Mn and O atoms perpendicular to the surface. Such a mode is not compatible with a MnO(100)like monolayer. The trilayer model is also consistent with O 1s core level photoemission data, which give evidence of two O 1s components separated by ∼0.4 eV: the stronger spectral component at 529.1 eV is assigned to the O surface layer, whereas the weaker component at higher binding energy is related to the O layer at the interface to the Pd(100).13 In principle, the CeO2(111) nanolayers observed on various noble metal surfaces, as discussed above, may be placed also in this trilayer category, because they consist of hexagonal layers in a O−Ce−O stacking sequence. However, in the fluorite-type lattice of CeO2, the O−Ce−O hexagonal stacking is the most stable bulk configuration, with the (111) surface as the most stable bulk termination. Interface or thin layer confinement effects therefore need not to be invoked to justify the occurrence of CeO2 trilayers at metal surfaces. 3.1.4. Single-Layer Structures with Unusual Building Elements. The oxide nanostructures discussed up to this point may be thought of as being constructed by either rock-salt structure (100)-derived or (111)-derived building units. In this subsection, we examine structures containing a mixture of such building blocks or other atypical structure elements, arranged within a single layer. We begin by considering the structure of a Ni oxide monolayer on Rh(111).188 Here, the lattice symmetry of the substrate favors the polar (111) orientation, whereas

Figure 14. (a,b) Top view of geometrical models of unsupported MnO(111) HEX-I undistorted (a) and HEX-I distorted (b) phases (red sphere, O atoms; light gray, Mn atoms). Dashed lines indicate the 2D unit cells. (c) Thermodynamic DFT phase stability diagram of explored HEX-I models in equilibrium with an O particle reservior controlling the chemical potential μO; the top scale is converted into oxygen pressure. Note the close energetic correspondence of the two trilayer models, whereas a (111)-type bilayer model is significantly less stable. (d) Comparison between the measured HREELS phonon value (vertical bar) and PBE (dashed line) and HSE (full line) predicted dipole active modes for the HEX-I phase. The inset schematically depicts the atomic displacements of the calculated phonon peaks. Mn and O atoms are sketched as “●” and “○”, respectively. Reprinted with permission from ref 112. Copyright 2009 American Physical Society.

polarity considerations would predict a lower surface energy for a nonpolar (100)-type arrangement. Experimentally, the P


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oxidation of a Ni layer of up to one monolayer in 5 × 10−8 mbar oxygen on Rh(111) causes the formation of a 2D oxide phase with a uniaxially ordered (6 × 1) structure as shown in Figure 15a: several equivalent domains are recognized in the image. The high-resolution STM image Figure 15b exhibits zigzag ridges, which are separated by 6 Rh lattice constants while being pseudomorphically aligned along the [110] substrate direction, yielding the (6 × 1) periodicity, which is easily recognized in the LEED image, Figure 15c. Between the

ridges, a hexagonal arrangement of protrusions is visible in the STM image. Besides (6 × 1), other ridge periodicities have also been encountered as minority structures, such as (4 × 1) and (2 × 1), but the (6 × 1) is the dominant order parameter. Uniaxially ordered 2D Ni oxide monolayer structures with (6 × 1) and (2 × 1) periodicities have also been observed on the terraces of a vicinal Rh(111) surface,189 and a (7 × 1) reconstruction of Ni oxide on Pt(111), with a very similar STM appearance as shown in Figure 15, has been reported by Hagendorf et al.190 This suggests that these uniaxial structures derive from a more general concept of polarity cancellation, strain relief, and/or low surface energy effects. DFT/GGA calculations have predicted a model of the Ni oxide (6 × 1) structure on Rh(111) consisting formally of a single NiO monolayer with a Ni5O5 stoichiometry, as presented in Figure 16.188 The structure displays a significant corrugation of ridges

Figure 16. Top view (a) and side view (b) of the structural model of the (6 × 1) Ni5O5 layer on Rh(111) (Ni(Rh) atoms, orange(white); lower O atoms, dark red; higher O atoms, bright red). Simulated STM images for the occupied (c) and unoccupied (d) states. Reprinted with permission from ref 188. Copyright 2012 American Chemical Society.

of 1.2 Å along the [110] direction, and the simulated STM images in Figure 16c,d show bright stripes, separated by darker regions as experimentally observed. The model contains two different structure motifs: a triangular 3-fold coordinated oxygen atom and a rectangular 4-fold coordinated oxygen atom, related to (111) and (100) type building blocks of rocksalt structure NiO, respectively. According to the model of Figure 16a,b, the local building blocks of the oxide layer are correlated to specific sites on the Rh(111) surface, with the three lower-lying O atoms positioned on top of Rh surface atoms thereby anchoring the NiO layer to the substrate. The two higher O atoms connect the network with structural flexibility, and they appear in both triangular and rectangular arrangements. The compression of the Ni−Ni distances, with 2.50 Å in the direction perpendicular to the Rh rows and 2.70 Å along the rows as compared to 2.96 Å in bulk NiO, leads to a significant stress in the rectangular units and to the local deformation of the layer with a uniaxial buckling (see black arrows in Figure 16a,b) and the formation of pronounced troughs in the [110] substrate directions.188 The latter structural deformation is thus based on the large lattice mismatch between NiO and Rh(111) and on the directional nature of the interfacial bonding. Somewhat related to the uniaxial structures of NiO on the (111)-type substrates is the NiO (2 × 1) structure, which has

Figure 15. Large-scale (a) and high-resolution (b) STM images of the Ni-oxide (6 × 1) structure on Rh(111). (a) (1000 × 1000 Å2; VS = +1 V; IT = 0.2 nA). (b) (50 × 50 Å2; VS = +0.008 V; IT = 1.5 nA). (c) LEED image (E = 150 eV). Reprinted with permission from ref 188. Copyright 2012 American Chemical Society. Q


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been observed to grow in the first monolayer on Ag(100), under particular kinetic conditions.126,164,165 Figure 17a displays

surrounded by three O ions (tricoordinated Ti) and the stripes are separated by dislocation lines with Ti species surrounded by four O ions (tetra-coordinated Ti).117 Through the periodic formation of the dislocation lines with tetra-coordinated structure units, the stoichiometry required by the preparation parameters can be established, while, at the same time, the stress generated by the lattice mismatch can be released. In such a way, the oxide overlayer is flexible to react to modifications in stoichiometry due to the thermodynamic preparation conditions, and ordered phases with more reduced stoichiometries, such as TiO1.14, have indeed been reported.117 Another type of exotic planar oxide layer geometry has been realized for ZnO nanolayers on Ag(111) and Pd(111) substrate surfaces. ZnO in its stable bulk wurtzite structure can be visualized schematically as a stack of alternating hexagonal planes of O and Zn ions along the c axis, and two polar surfaces with O and Zn termination, the (0001)-O and the (0001)-Zn surface, have been reported on bulk single crystals with unreconstructed (1 × 1) structure.191 This simple bulk termination was a surprising result because of the polar nature of the two surfaces, and several stabilization mechanisms that are compatible with the observed (1 × 1) structures have been proposed.192−194 On Ag(111) supported ZnO, however, a new nonpolar structure has been detected for films up to 2−3 ML thickness, which has been rationalized by a structure of the hexagonal boron-nitride type, consisting essentially of planar Zn−O hexagons and thereby avoiding any dipolar effects perpendicular to the surface (see Figure 18).140 The formation

Figure 17. (a) STM image of 0.66 ML NiO on Ag(100), after deposition at room temperature and annealing to 450 K in 1 × 10−7 mbar oxygen (200 × 100 Å2; VS = −0.5 V; IT = 0.4 nA). The image shows two orthogonal domains of the (2 × 1) structure and islands of bare Ag(100) areas (A) and a single island of the NiO(100) 1 × 1 structure (B); the latter is a minority structure at the given preparation conditions. (b) Structure model of the NiO (2 × 1) phase (top and side views; Ag atoms, white; Ni atoms, blue, yellow; oxygen atoms, red). (c) Simulated unoccupied state STM image of the (2 × 1) phase. Reprinted with permission from ref 126. Copyright 2011 American Physical Society.

Figure 18. Bulk ZnO wurtzite structure (a) and planar boron-nitride type structure model for ZnO monolayers on Ag(111) (b) (Zn atoms, yellow; O atoms, red). Reprinted with permission from ref 140. Copyright 2007 American Physical Society.

of planar hexagonal structures as seen for ZnO on Ag(111) appears to be a more general mechanism of polarity compensation for polar oxide nanostructures of the MO stoichiometry. Goniakowski et al.137 have investigated by theory the stability of a Ag(111) supported MgO(111) monolayer, which had been found before experimentally by Kiguchi et al.,136 and have predicted a novel stabilization mechanism of polar oxide orientation based on a hexagonal graphite-like structure, an analogue to the boron−nitride structure, where the O ions and the Mg ions are coplanar and the polarity is removed. The graphite-like planar structure concept seems to be realized also for ZnO nanolayers on Pd(111),195 where, as a function of coverage and depending on oxygen pressure, two well-ordered Zn oxide phases with (4 × 4) and (6 × 6) coincidence structures have been observed. Figure 19 presents STM images and corresponding LEED patterns of the (4 × 4) phase (a,b) and the (6 × 6) phase (c,d) of 0.6 and 0.9 ML, respectively, Zn oxide on Pd(111).195 The (4 × 4) phase

a high-resolution STM image of the NiO (2 × 1) phase on Ag(100).126 The structure model of the (2 × 1) phase, first suggested by Caffio et al.165 and validated recently by DFT+U calculations,126 consists of trigonal 3-fold coordinated Ni(O) atoms in a distorted quasi-hexagonal structure, reminiscent of the NiO(111) stacking, in which two types of O atoms are present, belonging to ridges and valleys running along the ⟨110⟩ substrate directions; the latter provide the uniaxial ordering parameter and the line-type (2 × 1) corrugation in STM (Figure 17b,c). It was found that the NiO (2 × 1) phase represents only a local minimum with respect to the NiO(100)1 × 1 ground-state configuration, but that during growth the (2 × 1) phase is kinetically stabilized.126 The zigzag like Ti6O8 monolayer observed on Pt(111)106 is, strictly speaking, a bilayer structure, which however also contains two different types of Ti atom coordination, related to (111)-type and (100)-type structure motifs. The structure consists of quasi-hexagonal (111)-type stripes with Ti species R


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objects of study for nanostructure research. The class of vanadium oxides is used as a benchmark system here to illustrate the behavior of these materials. V oxide bulk systems occur as single valency oxides with oxidation states ranging from V2+ to V5+, in the form of VO, V2O3, VO2, and V2O5, as well as in substoichiometric situations with mixed valencies containing V5+ and V4+ or V4+ and V3+.196 The properties of bulk terminating V oxide surfaces have been reviewed by Hermann and Witko197 and by Surnev et al.196 The growth properties of V oxide nanostructures have been investigated in detail on Pd and Rh single crystal surfaces:198 while the complex diagramatic phase behavior will be treated in a subsequent subsection in relation to the thermodynamic growth variables, we will focus here on singular structure aspects. Under highly oxidizing growth environment, different V oxide structure concepts are realized on Pd and Rh surfaces in the 2D monolayer limit. On Pd(111), a rather open network structure with a (4 × 4) superlattice is formed (Figure 20a),141 which according to DFT analysis is made up of tetrahedral V− O building blocks, interacting with Pd surface atoms alternatingly with one or two oxygen atoms (Figure 20b);199 at the corners of the hexagon depicted in Figure 20b, the apex

Figure 19. STM images of the (4 × 4) (a) and (6 × 6) phase (c) of Zn oxide on Pd(100), and the corresponding LEED patterns (b,d): (a, 65 × 65 Å2; VS = +0.01 V; IT = 3 nA), (c, 50 × 50 Å2; VS = +0.1 V; IT = 2 nA), (b, electron energy E = 70 eV), (d, electron energy E = 72 eV). Reprinted with permission from ref 195. Copyright 2010 American Chemical Society.

displays a network structure with a honeycomb lattice consisting of 12 protrusions in STM (Figure 19a), whereas the (6 × 6) phase exhibits a hexagonal structure with a lattice constant b = 3.2 Å, which is identical to the in-plane lattice parameter of the ZnO(0001) surface, and a Moiré pattern with a larger periodicity specifying the (6 × 6) coincidence lattice (Figure 19c). The theoretical analysis of the (6 × 6) phase based on a graphite-like hexagonal Zn−O structure model gives a stable configuration in DFT+U, and the corresponding STM image simulations show excellent agreement with the experimental images.195 The graphite-like ZnO layers remain planar in the simulations up to 4 monolayers, but develop a rapidly increasing corrugation in the average heights between Zn and O atoms with increasing layer numbers, signaling the transformation of the graphite-like configuration into a wurtzite-like structure. This transformation is easily conceivable by inspecting Figure 18. The transition from planar hexagonal into the bulk-type wurtzite structure has also been observed experimentally to occur in the 3−4 ML range for ZnO on Ag(111).140 The open network (4 × 4) structure occurring at submonolayer coverage on Pd(111) was found to be challenging for the theoretical analysis.195 A variety of different model structures have been tested, which, although providing local energy minima in DFT, did not concur with the experimental STM profiles. Eventually, the depolarization mechanism proposed to explain the (1 × 1) structure of the O-terminated ZnO (0001) bulk surface, involving a termination with hydrogen atoms,193 has been successful in yielding an Hterminated Zn6(OH)5 structure, which is both an energetically stable configuration and gives satisfying agreement with the experiment.195 3.1.5. Early Transition Metal Oxides. The high structural flexibility together with several stable oxidation states provide the ingredients for a rich variety of structural and electronic degrees of freedom that render early TM oxides fascinating

Figure 20. (a) STM image of the (4 × 4) V5O14 phase on Pd(111) (110 × 110 Å2 ; VS = +0.5 V; IT = 0.1 nA). (b) DFT model of (4 × 4) V5O14 /Pd(111) (Pd atoms, gray; V atoms, green; O atoms, red). Reprinted with permission from ref 199. Copyright 2003 American Chemical Society. S


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feature of the DFT analysis, the (√13 × √13)R13.8° V oxide phase can be created by the same pyramidal O4VO building blocks in a slightly different corner-sharing arrangement, yielding a V6O18 stoichiometry, which breaks down to the same VO3 as for the (√7 × √7)R19.1° phase. These two 2D V oxide phases on Rh(111) have a very similar total energy and are thus observed frequently in coexistence at the surface.142,198 The formal overall stoichiometries of the 2D V oxide monolayers at high oxygen chemical potentials μO on Pd(111) and Rh(111) surfaces of V5O14 = VO2.8 and V3O9 = V6O18 = VO3, respectively, indicate that charge transfer from the metal substrate across the interface to the oxide layer is necessary to stabilize the structures. Given the maximum 5+ oxidation state of vanadium, such stoichiometries can only be realized by electron transfer from the metal to the oxide via the anchoring O atoms. This gives evidence of the importance of the electronic metal−oxid interface coupling in determining the oxide nanolayer geometry, as emphasized in the next subsection. At intermediate μO during V oxide deposition or by reducing the (4 × 4) V5O14 layer by hydrogen, a (2 × 2) V oxide monolayer structure has been reported on Pd(111).141 Figure 22a shows an STM image of a 2D island of the (2 × 2)

O atom of the tetrahedron is at the interface (not seen), whereas the tetrahedra at the sides of the hexagon are bonded with two oxygen atoms to surface Pd atoms. The overall stoichiometry of the (4 × 4) V oxide structure derives to V5O14, and the structure parameters of the model have been confirmed subsequently by quantitative LEED I(V) analysis.200 Conversely, under similar oxidizing conditions, two closely related, ordered V oxide phases are formed on Rh(111) with (√7 × √7)R19.1° and (√13 × √13)R13.8° superlattices.142 Figure 21 displays the structural details of the (√7 × √7)R19.1° V

Figure 21. The (√7 × √7)R19.1° V oxide phase on Rh(111). (a) Large-scale STM image of 2D oxide islands (1000 × 1000 Å2 ; VS = +2 V; IT = 0.1 nA). (b) High-resolution STM image (50 × 50 Å2; VS = +0.75 V; IT = 0.2 nA). The (√7 × √7)R19.1° unit cell and the Rh(111) substrate direction are indicated. The inset shows a DFT simulated STM image. (c) DFT model of the (√7 × √7)R19.1° phase, top view. The unit cell and pyramidal units are drawn (Rh atoms, gray; V atoms, green; O atoms, red). (d) Side view of the (√7 × √7)R19.1° model. The inset shows a detailed view of the pyramidal O4VO unit. Reprinted with permission from ref 198. Copyright 2006 IOP Science. Figure 22. (a) STM image of 0.25 ML V oxide on Pd(111), showing a flat monolayer island of the (2 × 2) structure (200 × 200 Å2; VS = +0.5 V; IT = 1.0 nA). (b) High-resolution STM image of the honeycomb (2 × 2) V oxide (78 × 78 Å2; VS = +0.04 V; IT = 1.0 nA). The inset shows a simulated STM image with maxima corresponding to the V atoms. (c) Structure model (side and top views) of the 2D surface-V2O3 phase (Pd atoms, white; V atoms, light grey; O atoms, dark grey). Reprinted with permission from ref 141. Copyright 2001 Elsevier.

oxide phase. The large-scale STM image of Figure 20a shows the 2D island morphology at submonolayer oxide coverage, whereas the high-resolution image of Figure 21b reveals a hexagonal lattice of atomic protrusions in a honeycomb arrangement. Extensive DFT calculations, based on experimental STM, core level photoemission, and high-energy electron energy loss spectoscopy (HREELS) phonon data, have established the structure model of the (√7 × √7)R19.1° phase as shown in Figure 21c,d.142 It corresponds to a V3O9 stoichiometry with tetragonal pyramidal O4VO building blocks (see Figure 21d), with the V atom at the center, four basal oxygen atoms, which form the contact with the substrate, and a double bonded vandyl-O (VO) at the apex. The pyramids are linked together via the basal O atoms in a corner sharing way, thus creating the (√7 × √7)R19.1° unit cell (Figure 21c). A simulated STM image according to this model is contained in the inset of Figure 21b, showing excellent agreement with the experiment. Interestingly, and a convincing

structure at a submonolayer coverage (0.25 ML), and Figure 22b reveals the well-ordered hexagonal honeycomb lattice at higher resolution. The atomic structure of this (2 × 2) phase has been rationalized in form of a V2O3 bilayer, which consists of two V atoms per unit cell located in the 3-fold hollow Pd(111) sites and three O atoms above V−V bridge sites (see Figure 22c). The simulated STM pattern (inset of Figure 22b) according to this model agrees well with the experimental one and reveals that the V atoms are imaged as bright protrusions in T


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the empty state STM image.201 This (2 × 2) V oxide monolayer phase on Pd is a paradigmatic example in oxide nanostructure research, because it established for the first time that metal-supported oxide nanostructures exhibit very different structural properties as compared to those of their respective oxide bulk materials.202 To emphasize this difference from the bulk V2O3, the (2 × 2) V oxide on Pd(111) has been designated as surface-V2O3. It is worth noting in this context that the DFT derived surface-V2O3 structure parameters have been confirmed impressively by X-ray photoelectron diffraction measurements,203 which established an important test case for the successful applicability of standard DFT methods to modeling the structure of oxide nanophases. The XPS core level spectra of surface-V2O3 provided another interesting observation: the V 2p3/2 binding energy was found at 514.3 eV,204 which is about 1 eV smaller than the value of the corresponding V2O3 bulk compound (in conventional XPS wisdom, this would suggest a VO stoichiometry). The low value of the V 2p XPS binding energy is of course due to a proximity effect, with the metal surface coupled to the oxide layer with associated interfacial bonding and final state screening effects.205 It indicates that XPS core level binding energies taken from bulk reference compounds have to be treated with caution for estimating the stoichiometry or oxidation state of metal-supported oxide nanostructures. The reduction of the highly oxidized √7 and √13 “VO3“type nanostructures on Rh(111) leads to a complex phase diagram with a series of structures, but for the purpose of the present discussion the (9 × 9) structure, obtained at intermediate to low μO, is our focus. Figure 23a gives an

Figure 23b displays perfect agreement with the experiment, with every contrast detail being truthfully reproduced. A careful examination of the (9 × 9) model shows that locally ordered (2 × 2) arrangements of hexagons within the (9 × 9) unit cell are present (dashed in the figure), which suggests a certain similarity to the (2 × 2) surface-V2O3 phase on Pd(111). A pertinent question arises then: Why does such a complicated network structure evolve on Rh(111), to realize the simple V2O3 stoichiometry? We have argued recently that the absence of a long-range ordered (2 × 2) surface-V2O3 phase on Rh(111) is due to interfacial strain.206 On Pd(111), the favorable V−O bond distances project exactly onto a hexagonal (2 × 2) honeycomb lattice, while on Rh(111) the corresponding (2 × 2) lattice is mismatched by 2.3%. It has thus been proposed that only small local (2 × 2) regions can be formed, which have to be connected by eight-fold and five-fold rings to release the interfacial strain and to yield the (9 × 9) periodicity. Some support for this proposition has been received from growth experiments of V oxide on a stepped Rh(15 15 13) surface, which is a vicinal to Rh(111).420 Under suitable thermodynamic conditions, the (9 × 9) structure has been observed on the wider Rh(111)-type terraces, whereas a (2 × 2) V2O3 structure has been found on the narrow terraces, where the step edges provide additional strain relaxation. The so-called k-Ti2O3 nanolayer (k stands for kagomé) reported on Pt(111)124 has a structure that is related to the surface-V2O3 phase on Pd(111), with a honeycomb lattice of Ti atoms at the Pt interface and bridging O atoms forming the outer surface. However, whereas the surface-V2O3 honeycombs fit exactly onto a (2 × 2) lattice on Pd(111), the k-Ti2O3 honeycomb phase is incommensurate on Pt(111), with an approximate (2.15 × 2.15) superstructure derived from the LEED pattern. From large-scale STM images,124 it appears that the k-Ti2O3 phase is a domain structure, where relatively small domains of order are separated by defected domain boundaries, which presumably provide the necessary strain relief. 3.2. Chemical Interactions at the Interface

Chemical interactions at the metal−oxide interface play certainly an important role in many of the systems investigated in the previous subsection. This was mentioned explicitly for the case of V oxides on Pd and Rh substrates, where electron charge transfer across the interface is necessary to rationalize the favored structure concepts. Here, we will examine, in accord with our reductionist concept, how the interface chemistry is selective in stabilizing particular oxide structures at the nanolayer scale, and we will stress examples where the interfacial bonding is a dominant structure-determining element. In this context, the arguments of Campbell207 are worth mentioning, who demonstrates using general thermodynamic reasoning that the adhesion energy at the metal−oxide interface provides an extra stabilization term to the total free energy of transition metal oxide systems wetting their own metal. This term can play a decisive role in stabilizing ultrathin oxide films of thickness of the order 1−2 nm in relation to their bulk phases. Although the thermodynamic treatment of Campbell applies strictly only for thin slices of bulk-type oxide phases covering their own metal, the importance of the metal−oxide interface bonding (essentially the adhesion energy) for oxide nanolayer stabilization is clearly a general principle. The interface chemistry, as the term is understood and employed here, is determined by the chemical affinity of the

Figure 23. (9 × 9) V oxide phase on Rh(111): (a) large-scale STM image (2000 × 2000 Å2; VS = +2 V; IT = 0.05 nA). The inset shows a high-resolution image (60 × 60 Å2; VS = +2 V; IT = 0.1 nA). (b) DFT model of the (9 × 9) V36O54 phase. (9 × 9) (solid) and (2 × 2) (dashed) unit cells are indicated; the circle encloses a hexagonal V−O structure element (Rh atoms, gray; V atoms, green; O atoms, red). The inset is a simulated STM image. Reprinted with permission from ref 175. Copyright 2005 American Physical Society.

overall STM impression of the (9 × 9) phase on Rh(111), showing compact monolayer islands with rounded boundary lines, whereas the high-resolution image of the inset displays an intriguing STM contrast, with excellent order and hexagonally arranged larger and smaller depressions, the larger ones defining the (9 × 9) supercell.175 The DFT/GGA structure model in Figure 23b involves planar V−O hexagons as structure motifs (see circle in the figure), which are connected in a very complex way, with eight-fold and five-fold rings between, to establish the observed (9 × 9) periodicity. The unit cell stoichiometry according to the model is V36O54, which reduces again to V2O3. The DFT simulated STM image of the inset of U


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bridging site (as shown in Figure 24A) was calculated to be the thermodynamically most stable situation210 (see also section 4.2.2). The role of the metal−oxygen bond strength at the interface in determining the atomic structure of silica overlayers on various metal substrates has also been emphasized recently by Yu et al.,211 with strong metal−oxygen bonds favoring crystalline structures, whereas weaker metal−oxygen bonds promote vitreous silica films.212 The decisive role of the interface chemistry for structure selection has been demonstrated recently for a Co oxide nanolayer grown on a chemically modified Ir(100) surface: the Ir(100)1 × 1 surface has been covered by a pseudomorphic monolayer of Co, onto which the Co oxide has been deposited.173 On the bare Ir(100)(1 × 1) surface, Co oxide grows in a c(10 × 2) monolayer phase,162 which is structurally closely related to the CoO (9 × 2) (111)-type bilayer phase on Pd(100),161 as discussed in section 3.1. In contrast, on the metallic Co interlayer on Ir(100), the Co oxide forms a (100)derived c(4 × 2) Co3O4 cation vacancy structure, as on Pd(100).172 Gubo et al. have investigated the thermodynamic stability of the c(4 × 2) Co3O4 layer in relation to the CoO c(10 × 2) and a c(8 × 2) O−Co−O trilayer phase: Figure 25

substrate and the overlayer metals to each other and to oxygen. Electronegativities or the propensity to alloy formation of the elemental constituents of the interface may be taken as a first guideline to estimate the chemical interactions. Directional covalent bonding at the interface can be an essential ingredient for the evolution of a particular stable structure. In this context, the formation of an ordered silica nanolayer on a Mo(112) substrate may be cited. The preparation of ordered thin films of silica has been a desired but elusive endeavor for a long time, but has finally been achieved on a Mo substrate by Schroeder et al. 208 and Chen et al.209 In the structure model of Weissenrieder et al.,210 based on experimental data of STM, infrared reflection absorption spectroscopy, X-ray photoelectron spectroscopy, and DFT/GGA calculations, the silica consists of a single hexagonal layer, forming a 2D network of corner-sharing SiO4 tetrahedra (Figure 24). One of the O atoms anchors the structure to the protruding Mo atoms of Mo(112) surface, whereas the other three form Si−O−Si bonds with neighboring tetrahedra and constitute the hexagonal network. The strongly directional Si−O−Mo bond is clearly a stabilizing factor for this 2D silica layer, and the registry of the overlayer with the interface O atom in a Mo

Figure 25. Phase diagram of single-layer Co oxides on (a) bare Ir(100) and (b) 1 ML Co/Ir(100). G is the surface free energy, μO the chemical potential of oxygen, and pO the corresponding oxygen pressure at the temperature of the sample preparation, 320 and 670 K. Reprinted with permission from ref 173. Copyright 2012 American Physical Society.

displays a phase diagram of single-layer Co oxides on the bare Ir(100) surface (a) and on 1 ML Co/Ir(100) (b).173 On bare Ir(100), the c(4 × 2) structure is never energetically favorable (Figure 25a), but on Co/Ir(100) the c(4 × 2) Co3O4 phase is energetically most stable in the experimentally relevant parameter space of pressure and temperature (Figure 25b). Gubo et al. have proposed that the higher stability of the c(4 × 2) structure on the Co interlayer is due to the stronger oxygen−Co substrate bond, which is particularly important for the Co3O4 phase as compared to the competing c(10 × 2) and c(8 × 2) phases, where a wider variety of bonding configurations exist at the interface. Gubo et al. thus concluded that the interface chemistry is the relevant parameter for the crystallographic orientation of the Co oxide films (i.e., (100)-

Figure 24. Three models of the single layer SiO2 film on Mo(112). Model A is the thermodynamically most stable configuration. The Si4O10 surface unit cell is indicated. Reprinted with permission from ref 210. Copyright 2005 American Physical Society. V


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enables the growth of a rather defect-free oxide overlayer. Figure 26 presents a structure model of a MoO terminated

type on Co/Ir versus (111)-type on bare Ir) and not the surface strain and stress.173 The highly oxidized V oxide monolayer phases on Pd(111) ((4 × 4) V5O14) and Rh(111) (√7 and √13 VO3) have been introduced before, and the necessity for electron charge transfer donation from the metal to the oxides to stabilize their structures has been pointed out. An open question is, however, why do such different structures and (slightly) different stoichiometries with different metal−oxygen coordination spheres occur on the two substrates? In view of the small difference in lattice constants of ∼2.2% between Pd and Rh, elastic strain effects are unlikely, leaving the interface chemistry for further consideration. The different formal stoichiometries of VO2.8 versus VO3 indicate a somewhat lower oxidation state and oxygen content of the (4 × 4) phase on Pd than the √7 (√13) phase on Rh, but the closer inspection of the structure models (see Figures 20b and 21c,d) reveals that the (4 × 4) phase has less oxygen-substrate bonding than the √7 (√13) phase. In the (4 × 4) network, the tetrahedral V−O units interact with the Pd(111) surface alternatingly via one or two O atoms, whereas the tetragonal pyramids of the √7 (√13) structure are bonded by the four basal oxygen atoms to the Rh surface atoms. Because Rh is known to have a higher affinity to oxygen than Pd,213 it has been conjectured that this higher oxygen affinity of Rh is responsible for the structural differences of the highly oxidized V oxides on the two substrates.198,206 The so-called zigzag phases of TiOx on Pt(111)106 and VOx on Pd(111)198 reveal another aspect of oxide surface chemistry: while displaying incommensurate structures with an intriguing similarity in the STM images, suggesting similar structural arrangements, the chemistry of the metal cations leads to different stoichiometries and layer terminations. The V6O14 zigzag layer on Pd(111) has been interpreted in terms of a O− V−O/Pd stacking, with a layer of O atoms at the Pd interface and vanadyl-type VO groups at the surface forming the zigzag structure motif:199 this reflects the oxidation chemistry of the V atoms with their possibility of high oxidation state. The zigzag Ti6O8 phase on Pt(111) can be represented by the same metal−oxygen backbone as the V6O14 zigzag, but the corresponding DFT/GGA derived model indicates a Ti termination at the Pt interface with a O−Ti/Pt stacking and the absence of titanyl TiO species at the surface;106 the latter is in accord with the established chemical behavior of Ti. The different oxidation states of Ti versus V and the different oxygen affinity of Pd versus Pt (Pd−O being greater than Pt− O) thus leads, despite apparent similarities, to this interfacechemistry-based structural diversity. Strong chemical interactions at the oxide−metal interface may cause interdiffusion of the constituents and the formation of a mixed or so-called “reacted” interface. Shao et al. have investigated recently the growth of CaO films on Mo(001) and have reported considerable interdiffusion of Mo ions from the support into the ad-layer for low film thickness, giving rise to the formation of a (2 × 2) superstructure.214 The Ca/Mo mixed oxide grows pseudomorphically with a low defect concentration on the Mo(001) surface due to a well-matched lattice parameter to the substrate; the latter is in contrast to the large lattice mismatch of 8% of pure CaO(100) with respect to Mo(001). DFT/GGA calculations215 have provided mechanistic insights into the phase transition initiated by the diffusion of Mo into the oxide layer. The replacement of 25% of Ca ions in CaO by Mo results in a rock-salt type Ca3MoO4 phase, which has a very good lattice match with the Mo(001) and

Figure 26. Structure model of a MoO terminated 4-layer Ca3MoO4 mixed film with an oxidized Mo plane at the interface to the Mo(001) support. The right model shows the relaxation of O ions around the Mo impurity as indicated by the arrows. Reprinted with permission from ref 215. Copyright 2011 American Physical Society.

Ca3MoO4 film with an oxidized Mo plane at the interface to the Mo(001) support.215 Thermodynamic stability considerations confirm that a mixed Ca3MoO4 four-layer system grown on an oxidized Mo surface with MoO surface termination is indeed the most stable configuration for the range of used experimental parameter conditions. One of the driving forces for the formation of the mixed oxide is the oxidation of the Mo atoms and the higher oxygen content of the mixed Ca3MoO4 phase. Another one is the reduction of the lattice constant of the mixed phase caused by the alternation of Ca and Mo cations, whereby the Mo ions occupy less space than the Ca cations: this leads to a reduction of interfacial strain. The interdiffusion of Mo into the CaO matrix may therefore be considered as a chemical means of strain relaxation. The chemical interactions at the oxide−metal interface can be employed to trigger more exotic growth mechanisms for the fabrication of oxide nanostructures. This has been illustrated by the experiments of Dohnalek et al.,30 where the reactivity of surface Pt atoms has been used to chemically activate (WO3)3 clusters and to drive them via a self-assembled condensation reaction into a 2D W oxide nanolayer on Pt(111). The (WO3)3 clusters have been generated in the gas phase by direct thermal sublimation of WO3 powder and deposited via a molecular beam on to the Pt surface at 700 K. The surface reaction leads to a ring-opening of the (WO3)3 clusters and to their condensation into a 2D oxide layer with a zigzag chain structure and with a c(4 × 2) periodicity, whereby part of the W atoms get reduced to the 5+ oxidation state. The structure of this 2D W oxide layer has been elucidated with the help of DFT calculations. The c(4 × 2) structure is limited to the first monolayer at the Pt interface; in the second layer the (WO3)3 remain intact and form an ordered (3 × 3) array of molecularly bound (WO3)3 clusters.30 3.3. External Variables and Surface Oxide Phase Diagrams

The growth of thin films is intrinsically a kinetic process, but the thermodynamic variables and experimental parameters during growth exercise a profound influence on the geometry and stoichiometry of oxide overlayers and nanostructures. W


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interface-stabilized and cannot grow beyond monolayer coverage, and that the dissociation of O2 from the gas phase is kinetically hindered on these √7 (√ 13) V oxide monolayer surfaces. The continuing deposition of V atoms leads to a reduction of the V oxide at the surface until the most reduced V oxide phase has been reached (wagon-wheel VO), from where on metallic surface V deposits develop, which promote O2 dissociation again leading to further oxidation. Eventually, after a few monolayers of film thickness, the oxide converges to a bulk-type V2O3 phase, which is the stable V oxide under the respective experimental conditions.216,217 The gradual reduction of the oxidation state of V oxide monolayers on Rh(111) from 5+ to 2+ as illustrated in Figure 27b is the result of the flexibility of both V oxidation state and V−O coordination spheres, which are able to generate different building blocks that can be combined at the surface to create structures with variable stoichiometries.175 The highly oxidized tetragonal O4VO pyramids (Figure 21) and the planar hexagonal V−O rings (Figure 23) have already been introduced in a previous subsection. The structures in this reduction/ oxidation sequence are connected by the evolution of the local arrangement of common V−O structure motifs, with the reduction (oxidation) proceeding mainly by a progressive loss (gain) of vanadyl VO groups. The reactions can be steered into both directions by decreasing (increasing) μO via H2 (O2) gas-phase exposures and/or annealing at higher substrate temperatures.218 The phase diagram of Ti oxide nanolayers grown on Pt(111) as observed by Granozzi et al.177 is characterized by a number of understoichiometric TiOx (1.2 < x < 1.5) phases, which are shown with their STM fingerprints in Figure 28. The structures

Besides the temperature of the substrate, the oxygen pressure during oxidation, and the deposition rate of the parent metal of the oxide, the coverage or surface concentration of the metal is also an important parameter, because it reflects back on the stoichiometry of the growing oxide phase. The possibility of different oxidation states of the metal cations and the generally observed flexibility of structure elements in nanolayers add additional parameters for phase diversity. Nanoscale TM oxides may develop therefore a complex conglomerate or sequence of different phases on a given substrate during nanolayer growth or as a function of the thermodynamic variables, such as the chemical potential of oxygen μO. To describe the experimentally observed oxide phases on a given metal substrate, we will use so-called surface oxide phase diagrams, which in the present understanding represent the projection of the stability ranges of oxide phases onto the relevant parameter space. V oxides on Rh(111) display a rich phase behavior with many different structures developing as a function of oxide coverage and chemical potential of oxygen. The oxide coverage is determined by the amount of the evaporated metal (in, e.g., a reactive PVD process), but it reflects back also on the stoichiometry of the oxide, in particular during nanolayer growth, that is, in the first few monolayers. Figure 27a lists V

Figure 27. (a) Sequence of V oxide structures forming during growth on Rh(111). The coverage is expressed in terms of monolayer equivalents (MLE) of evaporated metal atoms; (b) V oxide monolayer structures with their unit cell stoichiometries forming on Rh(111) as a function of the chemical potential of oxygen μO. Reprinted with permission from refs 142 and 175. Copyright 2004 and 2005 American Physical Society. Figure 28. Surface phase diagram of TiOx on Pt(111) as a function of Ti coverage and oxygen pressure during annealing. The various phases are categorized according to their visual appearance in the STM images (k = kagomé; z = zigzag; w = wagon-wheel). Reprinted with permission from ref 124. Copyright 2009 American Chemical Society.

oxide structures on Rh(111) as a function of the coverage of the deposited V metal, and Figure 27b contains V oxide phases as a function of μO. The two diagrams thus represent the sequence of V oxide phases developing during growth on Rh(111) and the evolution of phases during oxidation/ reduction of the first monolayer, respectively. The phase behavior during growth (Figure 27a) reveals an interesting kinetic phenomenon:142 with increasing V coverage during a constant μO (high enough to support the highly oxidized √7 or √13 VO3-type phases in the first monolayer), a phase transformation into structures with a lower oxidation state (“oblique”, hex-VO2, wagon-wheel) takes place after the completion of the first oxide monolayer (corresponding to a metal coverage ∼0.6 MLE). Without going into the details of the respective oxide structures, this behavior has been explained by the fact that the √7 (√13) monolayer structures are

are either commensurate or incommensurate with hexagonal or rectangular unit cells, but as a common building principle the formation of bilayers with Ti at the Pt interface and oxygen forming the topmost surface can be identified, with defects of Ti vacancies or irregular Ti−O distributions within dislocation lines providing the required stoichiometry control imposed by the external parameters. The energetics of the various TiOx−Pt systems has been derived from DFT/GGA calculations,124 and the detailed energy analysis allows one to estimate the different components of the interfacial interactions. The energetics has been separated into different contributions: the total energy of X


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Figure 29. Schematic phase diagram of 2D Mn oxides on Pd(100), presented in the form of STM images as a function of oxygen pressure and of the chemical potential of oxygen μO. The nominal coverage of Mn is 0.75 ML. Note that nine different structures are observed, because the “chevron” phases consist of two related structures. Reprinted with permission from ref 13. Copyright 2009 IOP Science.

the combined oxide−metal system, the adhesion energy, that is, the energy required to separate the oxide and the Pt slabs, and the pure oxide energy of formation. Whereas the total energy of the combined oxide−metal system is to a good approximation directly proportional to the stoichiometry, with k-TiO1.5 being the most stable structure, the variation of the adhesion energy is more subtle and depends on the local epitaxial relationship, with directional bonding between Ti−O structural units and Pt surface atoms playing an important role.124 Mn oxide is a borderline case between early and late TM oxides, and the bulk phase diagram contains the divalent rocksalt-type MnO as well as higher oxidation state oxides as stable compounds.219 On a Pd(100) substrate, an amazing succession of Mn oxide structures has been observed as a function of the chemical potential of oxygen μO, which is a good demonstration of the complex phase behavior of such 2D nano-oxide systems. Nine different MnOx structures have been detected in the monolayer phase on Pd(100) as a function of μO, as illustrated by their respective STM fingerprints in Figure 29.13 In the “oxygen-rich” μO regime, the HEX-I and HEX-II structures have been interpreted on the basis of extensive experimental data (obtained by LEED, STM, XPS, HREELS) and screened hybrid DFT calculations in terms of hexagonal O−Mn−O trilayer structures with formal MnO2 stoichiometries (see section 3.1).112 In the “intermediate” μO regime, the c(4 × 2) Mn3O4 cation vacancy structure is the central phase, which is similar to the respective c(4 × 2) Ni3O4 and Co3O4 structures on Pd(100); the latter has also been discussed in detail above. The creation of cation vacancies as a structure concept has an interesting variant in the so-called “chevron” MnOx structures, which occur in the same μO regime as the c(4 × 2) (the name “chevron” is derived from the characteristic antiphase domain boundaries forming “chevron-like” structure motifs13). The “chevron” phases may be thought of as being generated from the c(4 × 2) structure by a simple vacancy propagation mechanism: this is illustrated in Figure 30. The top (a) and bottom panels (c) show the experimental LEED pictures and the corresponding simulated reciprocal lattice patterns, respectively, of the c(4 × 2), “chevron I” and “chevron II” phases (from left to right). The center panels (b) are the corresponding real lattice models, with the grid of lines representing the Pd(100) lattice and the black dots the Mn vacancy lattice, defining the oxide unit cells. In the c(4 × 2) structure, b1 and b2 define the rhombic primitive unit cell. If the c(4 × 2) vacancy at b2 is shifted to a neighboring antiphase position of the Pd lattice as shown in the middle panel of Figure 30b, the “chevron I” structure is obtained. Repeating this vacancy shift in the “chevron I” structure as depicted in the right panel of Figure 30b gives the “chevron II” structure. This

Figure 30. LEED patterns (a), real lattice models (b), and reciprocal lattice patterns (c) of the c(4 × 2) (left panel), chevron-I (middle panel), and chevron-II (right panel) structures. The respective LEED energies in (a) are 116, 108, and 60 eV. Reprinted with permission from ref 13. Copyright 2009 IOP Science.

propagation of Mn vacancies in the MnO(100)-type overlayer corresponds to a reduction of the number of vacancies and thus formally to a decrease of the oxidation state of the oxide layer. In fact, the stoichiometry reduces from Mn3O4 in the c(4 × 2) to Mn6O7 in the “chevron I” and to Mn10O11 in the “chevron II” structure, thus approaching the MnO stoichiometry. The “chevron I” Mn6O7 vacancy model has been tested by DFT calculations and was found to be a stable structure.112 The MnOx phases in the “oxygen-poor” μO regime, designated with fanciful names such as “waves”-like or “labyrinth”-like structures reflecting their appearance in the STM images or for the lack of better knowledge, have been less well characterized structurally, mainly because they were difficult to obtain as a single phase and occurred often in coexistence at the surface, thus precluding spectroscopic specification. The HEX-III structure observed at the low μO end of the phase diagram corresponds to a (√3 × √3)R30° superstructure of a MnO(111) surface and is possibly based on a (111)-type Mn−O bilayer.13 In conclusion, the low-dimensional Mn oxides on Pd(100) illustrate impressively how elastic, electronic, and chemical degrees of freedom conspire to produce a very complex phase behavior, with several different phases coexisting and creating what may be regarded somewhat loosely as a structurally degenerated ground state; moreover, kinetic stabilization may Y


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add further to this structural multiplicity. It is noted parenthetically that a much simpler phase behavior has been reported for the late TM oxides such as Ni oxide or Co oxide,146,147 where only one or two wetting layer structures are observed in the 2D nanolayer limit. 3.4. Effects of Low Dimensionality

Most of the oxide−metal systems, which have been discussed in the previous subsections, are monolayer type systems, in the sense that they measure only one or two unit cells in the direction perpendicular to the surface and thus are effectively two-dimensional. They derive their stability from this 2D confinement as a result of the enhanced possibilities for polarity compensation in the case of polar layers,220 their structural flexibility to accommodate interfacial strain, and the chemical interactions at the interface. Consequently, many of these 2D oxide phases are only stable in the 2D limit and are self-limited in the third dimension: they cannot grow into thicker 3D layers. However, they often provide the basis for the growth of a sequence of interlayers with modified structures that provide a graded interface for the subsequent epitaxial growth of bulk phases. Such a sequence of interlayer structures has been found, for example, in the epitaxial growth of corundum-type V2O3(0001) films on Pd(111).141,187,202 Another example is Fe oxide on Pt(111), where a 1−2 ML FeO(111) phase mediates the growth of epitaxial Fe3O4(111) films.154 For the case of titania nanosheets assembled on Pt(110), the reduced dimensionality rather than the interactions with the substrate has been invoked to explain the stabilization of the 2 ML titania sheets in a lepidocrocite structure:221 the DFT calculations have shown that an isolated lepidocrocite nanosheet is the most stable 2D titania phase (the lepidocrocite structure can be generated from an anatase (001) bilayer by a shift of the upper layer by half a unit cell). Specific interface interactions are a weaker effect and are responsible only for a structure modulation due to a (14 × 4) coincidence structure on the Pt(110).221 In the following, we will concentrate on 1D nanowires and (quasi-)0D nanodots (including molecular-type clusters) of oxide−metal hybrid systems. Vicinal metal surfaces are suitable templates for the self-organized growth of 1D nanostructures. Such an approach has been employed by Schoiswohl et al.14,225 to fabricate 1D Ni oxide nanowires on vicinal Rh(111) surfaces, by decorating first the surface steps with pseudomorphically coupled monatomic rows of Ni adatoms, followed by oxidation of the Ni nanowires. The exclusive oxidation of the Ni rows is made possible by the enhanced chemical reactivity toward oxygen of the coupled Rh−Ni nanowires at the steps, which form 1D NiO2 lines after oxidation as evidenced by STM, XPS, and screened hybrid DFT analysis.14 Mn oxide nanowires have also been successfully fabricated on vicinal surfaces of Pd(100).222,223 Figure 31a displays an STM image of monatomic rows of Mn oxide nanowires decorating the step edges of a Pd(1 1 17) surface; they have been prepared by deposition of 0.1 ML Mn at 300 K followed by oxidation in 1 × 10−8 mbar oxygen at 470 K. The MnOx nanowires are readily recognized at the step edges by their brighter contrast in STM, while at the (100) terraces a p(2 × 2) structure of chemisorbed oxygen is visible. The monatomic MnOx rows are attached to the Pd steps in a (×1) periodicity, thus they are pseudomorphically coupled, and their formal stoichiometry is MnO2, with every Mn atom coordinated to four O atoms, two at the lower and two at the upper step edge, as indicated by

Figure 31. (a) STM image of MnO2 nanowires decorating the step edges of a Pd(1 1 17) surface (80 × 80 Å2; VS = +0.01 V; IT = 0.1 nA) . The nanowires at the steps have an in-row periodicity of 2.75 Å (×1), whereas the (100) terraces are covered by a p(2 × 2) structure of chemisorbed oxygen. (b) Mn 2p3/2 XPS core level spectra of the metallic (lower curve) and oxidized (upper curve) Mn nanowires on Pd(1 1 17). Reprinted with permission from ref 222. Copyright 2010 Elsevier.

DFT calculations.224 The oxidation of the Mn nanowires is confirmed by the Mn 2p3/2 XPS core level spectra in Figure 31b: the lower spectrum from the as-deposited metallic Mn nanowires, which displays a complex spectral structure due to final state effects in the photoemission process, is shifted by 0.8 eV to higher binding energy after oxygen exposure (upper curve), which is a clear spectral fingerprint of the oxidation of the Mn atoms.222 The 1D oxide−metal hybrid structures exhibit novel chemical and possibly also novel magnetic properties, which are still largely unexplored. They are interesting model systems for advanced catalysis studies, and this will be discussed further below in section 4.4. An obvious question is whether there are restrictions to this fabrication approach of oxide nanowires by decorating stepped metal surfaces in terms of the materials involved, that is, what influence has the chemical nature of the oxide and/or the metal support. An answer to this question has to consider both energetic and kinetic effects, with the latter particularly difficult to gauge. While refraining from a general answer, we would like to report a couple of unsuccessful fabrication attempts from our own laboratory to give an idea of the problems that may be encountered. The preparation of oxide nanowires on stepped Rh(111) surfaces, which was successfully achieved for Ni oxide, was unsuccessful for V oxide, because the Rh step edges could Z


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not be decorated by V adatoms (which is the first step of the preparation procedure). The physical origin of this failure has been investigated using DFT/GGA.225 It was found that the energy barriers for surface adatom diffusion of V across the Rh(111) terraces are much higher than those of Ni adatoms, and that the energy benefit for the attachment of adatoms to the step edges is smaller for V than for Ni adatoms. In addition, the propensity for subsurface alloy formation of V adatoms on Rh(111) at elevated temperature226 is detrimental to the attempt to overcome the kinetic barriers by increasing the surface temperature. Thus, V oxide nanowires could not be prepared on stepped Rh surfaces under the kinetic conditions screened in our experiments. Another attempt that has been unsuccessful was to create 1D Ce oxide nanowires on a vicinal Au(788) surface,227 which was undertaken to prepare an actual 1D model catalyst system. Flat islands of ceria could be grown on the terraces of the Au(788) surface (see Figure 32), which display a certain order suggesting preferential nucleation of ceria on the herringbone reconstruction of the Au(111)-type terraces. The ceria islands have a fairly uniform size of ∼5 nm but irregular shapes, and the desired decoration of the Au step edges with 1D structures could not be obtained, although

different recipes and a range of kinetic parameters have been tried. The exact reasons for this failure are unclear, but possibly the easy formation of a Ce−Au surface alloy precludes the diffusion of Ce adatoms across the Au terraces to the step edges. The lowering of the surface temperature to below room temperature to avoid surface alloy formation might be a way out and a route to achieve step decoration, but this was not possible in our experimental setup. Finite size effects exert a profound influence on the physical and chemical properties of low-dimensional oxide systems. A striking example has been found recently during the growth of a c(4 × 2) Mn3O4 overlayer on vicinal Pd(11N) surfaces.224 We recall that the Mn3O4 structure is a rocksalt-type (100) monolayer with 25% of Mn cation vacancies, which define the c(4 × 2) superstructure.228 The formation of cation vacancies and the associated lattice relaxation is a means to release the compressive strain, which builds up due to the mismatch between rocksalt TMO(100) and Pd(100) lattices. While this mechanism appears to work well for the c(4 × 2) structures of Ni3O4 and Co3O4, for the case of Mn3O4 the strain relaxation is incomplete leading to a domain surface with small regions of ordered c(4 × 2) domains separated by disordered domain boundaries, as shown in Figure 8. If, however, the c(4 × 2) Mn3O4 layer is deposited on a Pd(1 1 17) surface, a vicinal of Pd(100) with 23.4 Å (∼8.5 atoms) wide (100) oriented terraces separated by monatomic steps, a perfectly ordered nanopatterend surface is obtained as shown in Figures 33a and b, where STM images of the bare Pd(1 1 17) and the c(4 × 2) Mn3O4 covered surfaces are compared. The clean Pd(1 1 17) surface (Figure 33a) exhibits only moderate step periodicity order (see the terrace width distribution histogram in Figure 31c), as a result of the weak step−step interactions across the wide terraces.229,230 This situation is changed dramatically on

Figure 33. STM images of the clean Pd(1 1 17) surface (a) and the c(4 × 2) Mn3O4 covered Pd surface (b): (a) (200 × 200 Å2; VS = +0.0005 V; IT = 0.1 nA); (b) (200 × 200 Å2; VS = +0.0003 V; IT = 0.1 nA). (c) Terrace width distribution histogram. (d) Line scans across the autocorrelation plots perpendicular to the step edges of the STM images of the clean Pd(1 1 17) surface (a) and the c(4 × 2) covered Pd surface (b). Reprinted with permission from ref 224. Copyright 2012 IOP Science.

Figure 32. STM images of ceria nanoislands on Au(788), prepared by deposition of 0.14 ML of Ce at room temperature followed by oxidation at 200 °C, 1 × 10−7 mbar oxygen ((a) 1000 × 1000 Å2, (b) 500 × 500 Å2, VS = +2.2 V; IT = 0.02 nA). Note the regular arrangement of islands in lines perpendicular to the step edges, which may be caused by preferential nucleation on the herringbone reconstruction of the Au(111)-type terraces. AA


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the c(4 × 2) Mn oxide covered surface (Figure 33b), where a defect-free, highly odered surface with perfect long-range step periodicity is observed, as testified by the very narrow terrace width distribution histogram in panel (c). A closer inspection of the terrace width periodicity in panel (d) using the line scans across the autocorrelation plots of the STM images of Figure 31a and b reveals that the periodicity of the clean and the oxide covered surface is different, and that the average terrace width of the latter is expanded as compared to the pristine Pd(1 1 17) surface. The analysis of the long-range order parameters by high-resolution spot-profile analysis (SPA-) LEED gives a quantitative measure of the terrace expansion with a mean terrace width of 28.5−29 Å, corresponding to an addition of 2 atomic rows per terrace and a rearrangement of the original (1 1 17) Pd surface to a (1 1 21) surface. The deposition of Mn3O4 nanostripes thus leads to a major reconstruction of the Pd(1 1 17) substrate and to the stabilization of a very regular defect-free Pd(1 1 21) surface. The physical origin of this phase transformation has been traced down in DFT by investigating the stability of freestanding and Pd(11N) supported c(4 × 2) Mn3O4 nanostripes as a function of the corresponding widths.224 It turns out that particular stable configurations are obtained for so-called “magic” widths, where the oxide overlayer is terminated laterally by full MnO2 lines. The most stable adsorbed nanostripe has a Mn14O20 unit cell across the terrace, which fits both onto Pd(1 1 19) and onto (1 1 21) terraces, but not onto Pd(1 1 17). The preference for the experimentally observed reconstruction into Pd(1 1 21) has been attributed to kinetic effects at the high oxidation temperature. The fundamental agreement between theory and experiment is demonstrated in Figure 34, which shows the experimental and

form, or they may be generated at the surface via some sort of self-assembly process. Schoiswohl et al.238 have observed the formation of small V oxide clusters by spontaneous selfassembly on a Rh(111) surface. The clusters have been formed by the spontaneous condensation of adatoms after deposition of V atoms onto an O-covered Rh(111) surface followed by gentle annealing: as shown by STM, they consist of star-like hexagonal structures (Figure 35a), which have a well-defined

Figure 35. V6O12 cluster molecules on Rh(111), formed at the surface by spontaneous aggregation of vanadium and oxygen adatoms. (a) STM image (63 × 63 Å2; VS = +0.5 V; IT = 0.1 nA). The inset shows a DFT-simulated image. (b) Relaxed DFT model geometry of the V6O12 clusters in top and side views. Reprinted with permission from ref 198. Copyright 2006 IOP Science.

orientation with respect to the Rh(111) substrate. Highresolution STM images and DFT/GGA calculations have identified the clusters as planar V6O12 units (Figure 35b), where the six V atoms are located in the 3-fold hollow sites of the Rh(111) surface and the oxygen atoms occupy Rh top sites.238 The planar geometry of the V6O12 is the result of the stabilizing effect of the metal surface, because planar VxOy oxide clusters for x > 3 are unstable in the gas phase.237 The V6O12 clusters become mobile at elevated temperatures and condense into extended 2D V oxide structures.218 A different approach to create surface supported oxide clusters as model catalysts has been adopted by Dohnalek et al.:28,29 W3O9 clusters have been generated in the gas phase by direct thermal sublimation from WO3 powder and have been deposited onto a substrate surface. The cyclic (WO3) 3 (=W3O9) clusters are the dominant subliming species in the gas phase,239 and they can be deposited onto TiO2 in essentially monodisperse intact form.29 Deposited on a thin FeO(111)/ Pt(111) layer, however, the clusters interact strongly with the Fe ions, leading to cluster dissocation and to a significant restructuring of the FeO(111) layer.240 This strong cluster/ support interaction has a significant influence on the catalytic activity of the (WO3)3 clusters: whereas (WO3)3/TiO2 is catalytically very active for the dehydration of alcohols,241 (WO3)3/FeO(111) does not efficiently catalyze alcohol dehydration.240 Taking on the cluster fabrication approach of Dohnalek et al.,28 Wagner et al. have studied recently the structure and bonding of (WO3)3 clusters on nanostructured Cu−O surfaces.242 The (WO3)3 clusters have been deposited at low temperature (

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