Oxide Surfaces (The Chemical Physics of Solid Surfaces Vol. 9)

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THE CHEMICAL PHYSICS OF SOLID SURFACES


Volume 1 CLEAN SOLID SURFACES Volume 2 ADSORPTION AT SOLID SURFACES Volume 3 CHEMISORPTION SYSTEMS Volume 4 FUNDAMENTAL STUDIES OF HETEROGENEOUS CATALYSIS Volume 5 SURFACE PROPERTIES OF ELECTRONIC MATERIALS Volume 6 COADSORPTION, PROMOTERS AND POISONS Volume 7 PHASE TRANSITIONS AND ADSORBATE RESTRUCTURING AT METAL SURFACES Volume 8 GROWTH AND PROPERTIES OF ULTRATHIN EPITAXIAL LAYERS Volume 9 OXIDE SURFACES


EDITED BY

D.P.WOODRUFF B.Sc. (Bristol), Ph.D., D.Sc. (Warwick) Professor of Physics, University of Warwick

VOLUME 9

OXIDE SURFACES

2001 ELSEVIER AMSTERDAM - LONDON - NEW YORK - OXFORD - PARIS - SHANNON - TOKYO

ELSEVIER SCIENCE B.V. Sara Burgerhartstraat 25 RO. Box 211, 1000 AE Amsterdam, The Netherlands © 2001 Elsevier Science B.V. All rights reserved This work is protected under copyright by Elsevier Science, and the following terms and conditions apply to its use: Photocopying Single photocopies of single chapters may be made for personal use as allowed by national copyright laws. Permission of the Publisher and payment of a fee is required for all other photocopying, including multiple or systematic copying, copying for advertising or promotional purposes, resale, and all forms of document delivery. Special rates are available for educational institutions that wish to make photocopies for non-profit educational classroom use. Permissions may be sought directly from Elsevier Science Global Rights Department, PO Box 800, Oxford 0X5 1DX, UK; phone: (+44) 1865 843830, fax: (+44) 1865 853333, e-mail: [email protected]. You may also contact Global Rights directly through Elsevier's home page (http://www.elsevier.nl), by selecting 'Obtaining Permissions. In the USA, users may clear permissions and make payments through the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, USA; phone: (+1) (978) 7508400, fax: (+1) (978) 7504744, and in the UK through the Copyright Licensing Agency Rapid Clearance Service (CLARCS), 90 Tottenham Court Road, London W1P OLR UK; phone: (+44) 207 631 5555; fax: (+44) 207 631 5500. Other countries may have a local reprographic rights agency for payments. Derivative Works Tables of contents may be reproduced for internal circulation, but permission of Elsevier Science is required for external resale or distribution of such material. Permission of the Publisher is required for all other derivative works, including compilations and translations. Electronic Storage or Usage Permission of the Publisher is required to store or use electronically any material contained in this work, including any chapter or part of a chapter. Except as outlined above, no part of this work may be reproduced, stored in a retrieval system or transmitted in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, without prior written permission of the Publisher. Address permissions requests to: Elsevier Science Global Rights Department, at the mail, fax and e-mail addresses noted above. Notice No responsibility is assumed by the Publisher for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions or ideas contained in the material herein. Because of rapid advances in the medical sciences, in particular, independent verification of diagnoses and drug dosages should be made.

ISBN 0-444-50745-0 (Vol. 9) ISBN 0-444-41971-3 (Series) First edition 2001 Library of Congress Cataloging in Publication Data A catalog record from the Library of Congress has been applied for. © T h e paper used in this publication meets the requirements of ANSI/NISO Z39.48-1992 (Permanence of Paper). Printed in The Netherlands.

Contributors to Volume 9 A. BARBIER

CEA-Grenoble, Departement de Recherche Fondamentale sur la Matiere Condensee, SP2M, 17, rue des Martyrs, 38054 Grenoble Cedex, France

M.A. BARTEAU

Center for Catal3l;ic Science and Technology, Department of Chemical Engineering, University of Delaware, Newark, DE 19716, USA

S.A. CHAMBERS

Environmental Molecular Sciences Laboratory, Pacific Northwest National Laboratory, PO Box 999, MS K8-93 Richland, WA 99352, USA

C.C. CHUSUEI

Department of Chemistry, P.O. Box 30012, Texas A&M University, College Station, TX 77842-3012, USA

B.G. DANIELS

Surface Science Research Centre and Chemistry Department, Manchester University, Manchester M l 3 9PL, United Kingdom

U. DIEBOLD

Department of Physics, Tulane University, New Orleans, LA 70118, USA

R.G. EGDELL

Inorganic Chemistry Laboratory, University of Oxford South Parks Road, Oxford 0X1 3QR, United Kingdom

H.-J. FREUND

Fritz-Haber-Institut der Max-Planck-Gesellschaft, Department of Chemical Physics, Faradayweg 4-6, 14195 Beriin, Germany

D.W. GOODMAN


B.E. HAYDEN

Department of Chemistry, University of Southampton, Highfield, Southampton SO 17 IB J, United Kingdom

V.E. HENRICH

Department of Applied Physics, Yale University, P.O.Box 208284, New Haven, CT 06520, USA

K. HERMANN

Fritz-Haber-Institut der Max-Planck-Gesellschaft, Faradayweg 4-6, 14195 Berlin, Germany

H. KUHLENBECK

Fritz-Haber-Institut der Max-Planck-Gesellschaft, Department of Chemical Physics, Faradayweg 4-6, 14195 Berlin, Germany

R. LINDSAY

Surface Science Research Centre and Chemistry Department, Manchester University, Manchester Ml3 9PL, United Kingdom

D.C. MEIER


C. NOGUERA

Laboratoire de Physique des Solides, UMR CNRS 8502, Universite Paris-Sud, 91405 Orsay, France

G. PACCHIONI

Dipartimento di Scienza dei Materiali, Universita di Milano-Bicocca, Istituto Nazionale per la Fisica della Materia, via Cozzi, 53, 20125 Milano, Italy

G. RENAUD

CEA-Grenoble, Departement de Recherche Fondamentale sur la Matiere Condensee, SP2M, 17, rue des Martyrs, 38054 Grenoble Cedex, France

T. RISSE

Fritz-Haber-Institut der Max-Planck-Gesellschaft, Department of Chemical Physics, Faradayweg 4-6, 14195 Berlin, Germany

G.S. ROHRER

Department of Materials Science and Engineering, Carnegie Mellon University, Pittsburgh, PA 15213-3890, USA

J.L. DE SEGOVIA

Institute de Ciencia de Materiales de Madrid, CSIC E-28049 Cantoblanco Madrid, Spain

A.B. SHERRILL

Center for Catalytic Science and Technology, Department of Chemical Engineering, University of Delaware, Newark, DE 19716, USA

G. THORNTON

Surface Science Research Centre and Chemistry Department, Manchester University, Manchester Ml3 9PL, United Kingdom

E.M. WILLIAMS

Surface Science Research Centre, The University of Liverpool, Liverpool L69 3BX, United Kingdom

M.WITKO

Institute of Catalysis and Surface Chemistry, Polish Academy of Sciences, ul. Niezapominajek, 30239 Cracaw, Poland

Preface During the late 1960s and 1970s the commercial availability of ultra-high vacuum (UHV) systems allowed the development of a plethora of new techniques which were devised to probe materials in a surface-specific fashion, and this in turn led to the creation of modem surface science; the study of the structural, electronic and chemical properties of extremely well-characterised surfaces on an atomic scale. When David King and I first conceived this series of volumes in the later 1970s our objective was to recognise the growing maturity of this new scientific discipline which was already starting to apply these techniques in a combined fashion to understand surface processes. In the 20 years since the first volume was published, this perception has certainly proved to be well-founded, and while new techniques have continued to appear, they have rapidly been assimilated into the general armoury of methods (the increasing pervasiveness of scanning probe microscopies is very evident in the current volume), and it is the combination of methods which has proved most effective. The topic of the present volume, Oxide Surfaces, has been in many ways a 'Cinderella' of surface science; while the importance of oxide surfaces has been long recognised, the perceived difficulty of preparing and probing oxide surfaces (especially insulating oxide surfaces) meant that there was far more attention devoted in the mainstream of the subject to metal and semiconductor surfaces. There were notable exceptions, and some of the authors contributing to this volume have made major contributions to the topic throughout the whole period of development of modem surface science. The study of oxide surfaces has been a major growth area in the last few years, partly encouraged by new methods including the use of epitaxial oxide films grown on metallic substrates. These films are more amenable to study by electron spectroscopies, but raise new issues of surface characterisation and associated properties. This volume provides a view of some of the main areas of development and of recent progress in the study of well-characterised oxide surfaces. The first chapter by Henrich, one of the pioneers of modem surface studies of oxides, and co-author of the first text on the subject, provides an overview of the subject and relates the remaining chapters to this overview. Chapters 2 to 4, by Noguera, by Pacchioni and by Hermann and Witko, are concemed with the theory of oxides surfaces; they cover a range of materialsfi*omsimple rocksalt stmctures such as MgO through to the complexity of transition metal oxides, and also present some complementary methods of modelling and calculation. These theoretical studies also address the key issue of surface defects, and cover some aspects of adsorption at oxide surfaces. In some ways oxide surfaces is a topic in which theory was, for some years, ahead of experiment, and hence unchallenged. This was especially tme in the predictions and

measurements of surface structure, but substantial experimental advances have been made in the last few years in quantitative structure determination of oxide surfaces and this progress is surveyed in Chapters 5 and 6; Renaud and Barbier focus on the achievements of surface X-ray scattering methods (an approach which circumvents one of the main problems of surface charging when using electron probes), while Lindsay, Daniels and Thornton concentrate on other methods, and especially on adsorbate structure determination. A closely related issue is the structural characterisation of thin epitaxial oxide films, and this is described in Chapter 7 by Chambers. Crucial complementary structural information on oxide surfaces structures, especially of Ti02, has been provided in the last few years by scanning probe microscopies, and this is the focus of Diebold's Chapter 11. Because of the important role played by surface stoichiometry of oxides in both structure and reactivity, this work also relates quite closely to the content of Chapters 8 (Freund, Kuhlenbeck and Risse), 9 (Chusuei, Meier and Goodman), 10 (Sherrill and Barteau) and 12 (Rohrer) which all address different aspects of the chemistry and reactivity (including photochemistry) of oxide surfaces (including those of thin films), an area which constitutes one of the primary motivations for the study of oxide surfaces. Chapters 13 (Hayden) and 14 (Egdell) discuss two general aspects of oxide surfaces, their vibrational and electronic properties; while the titles suggest an emphasis on physical properties, vibrational spectroscopies, in particular, are also an important probe of adsorbate interactions and are especially relevant to chemical properties. The final chapter by de Segovia and Williams describes another topic of especial relevance to oxide surfaces and closely related to their photochemistry, that of DIET (desorption induced by electronic transitions).

January 2001

D.P.Woodruff

Contents Preface

ix

Chapter 1 (V.E. Henrich) Metal oxide surfaces and interfaces: concepts and issues 1. Introduction 2. Interesting bulk properties 2.1 Electronic properties 2.2 Crystal structure 3. From bulk to surface 4. Electrons on oxide surfaces 5. Oxide-gas, -liquid and -solid interfaces 6. Adsorption on metal oxides 7. Endnote Acknowledgements References

1 1 2 2 8 9 16 21 25 31 33 33

Chapter 2 (C. Noguera) Clean oxide surfaces: a theoretical review 1. Introduction 2. Numerical methods 2.1 Classical approaches 2.2 Semi-empirical quantum methods 2.3 The ab initio HF method and methods beyond 2.4 The ab initio Density Functional Theory and methods beyond 2.4.1 The Self-Interaction Correction (SIC) method 2.4.2 The GW method 2.4.3 The LDA+U method 2.5 General features for simulating surfaces 3. Review of literature 3.1 Fluorite MO2 and anti-fluorite M2O structures 3.2 Rocksalt structure 3.3 Corundum structure 3.4 Rutile structure 3.5 Wurtzite structure 3.6 Perovskite structure 3.7 Other oxides 4. Discussion of non-polar stoichiometric surfaces 4.1 Atomic structure 4.2 Energetics 4.3 Electron distribution 4.4 One particle excitations 4.5 Two particle excitations 5. Non-stoichiometric surfaces 5.1 Atomic relaxation 5.2 Charge and spin distribution

35 35 36 36 37 38 39 40 41 42 43 44 45 45 49 51 53 54 57 59 59 61 62 65 68 69 70 70

5.3 Spectroscopic signature 5.4 Energetics 6. Polar surfaces 6.1 Criterion for surface polarity 6.2 Models of electronic structure in iono-covalent materials: application to polar surfaces 6.3 Surface processes relevant for polarity healing 6.4 Summary 7. Conclusion References

72 73 75 76 78 80 83 85 86

Chapter 3 (G. Pacchioni) Theory of point defects at the MgO surface 1. Introduction 2. Quantum chemical description of ionic crystals 2.1 Cluster models 2.2 Embedding schemes 2.3 Solutions of the Schrodinger equation 3. Role and nature of defects at the MgO surface 3.1 Low-coordinated cations 3.2 Low-coordinated anions 3.3 Hydroxyl groups 3.4 Anion vacancies 3.4.1 Charge distribution 3.4.2 Hyperfme interactions 3.4.3 Stability and formation energies 3.4.4 Energy levels and ionisation potentials 3.4.5 Barriers to vacancies diffusion 3.4.6 Optical spectra 3.4.7 Chemical reactivity 3.5 Cation vacancies 3.6 Divacancies 3.7 Impurity atoms 3.80" radical anions 3.9(lll)microfacets 4. Conclusions Acknowledgements References

94 95 96 98 99 101 103 107 109 110 111 113 115 116 117 118 120 121 122 124 126 128 129 129 130

Chapter 4 (K. Hermann and M. Witko) Theory of physical and chemical behaviour of transition metal oxides: vanadium and molybdenum oxides 1. Introduction 2. Vanadium oxides 2.1 Vanadium oxide bulk systems 2.1.1 Geometric structure 2.1.2 Electronic properties

136 139 140 140 142

2.2 Vanadium oxide surfaces 2.2.1 Geometric structure 2.2.2 Physical and chemical properties (a) Properties of clean surfaces (b) Surface oxygen vacancies (c) Adsorption 3. Molybdenum oxides 3.1 Molybdenum oxide bulk systems 3.1.1 Geometric structure 3.1.2 Electronic structure 3.2 Molybdenum oxide surfaces 3.2.1 Geometric structure 3.2.2 Physical and chemical properties (a) Properties of clean surfaces (b) Surface oxygen vacancies (c) Adsorption Synopsis Acknowledgement References

148 148 152 152 157 162 168 170 170 172 172 172 175 175 182 186 190 191 191

Chapter 5 (R. Lindsay, B.G. Daniels and G. Thornton) Geometry of adsorbates on metal oxide surfaces 1. Introduction 2. Magnesium oxide - MgO 2.1MgO(100)-H2O 2.2MgO(100)-C2H2 2.3MgO(100)-CO 2.4MgO(100)-CO2 2.5MgO(100)-Ca 2.6MgO(100)-Ag 2.7MgO(100)-Pd 2.8MgO(100)-Ni 2.9MgO(100)-Fe 3. Nickel oxide-NiO 3.1NiO(100)-NO 3.2NiO(100)-CO 3.3NiO(100)-H2S 3.4NiO(lll)-NO 3.5NiO(lll)-CO 4. Titanium dioxide - Ti02 4.1TiO2(110)-HCOOH 4.2 Ti02(l 10) - H3CCOOH 4.3 Ti02(l 10) - H5C2COOH 4.4 Ti02(l 10) - H5C6COOH 4.5 Ti02(l 10) - bi-isonicotinic acid 4.6TiO2(110)-Na 4.7 Ti02(l 10) - Na/C02 4.8 Ti02(l 10) - (NH4)6Mo7024.4H20

199 202 202 206 207 208 209 210 212 212 213 213 214 216 217 218 219 219 221 223 224 224 225 226 227 228

4.9 Ti02(l 10) - {Rh(CO)2Cl}2 4.10TiO2(110)-SO2 4.11TiO2(110)-K 5. Aluminium oxide - AI2O3 5.1a-Al2O3(0001)-Cu 5.2 y-like AbOsCOOO 1) - di-tert-butyl nitroxide 6. Chromium oxide - Cr203 6.1Cr2O3(0001)-CO2 6.2Cr2O3(0001)-NO 7. Iron oxide - FexOy 7.1 FeO(lll)-H5C6CHCH2 7.2 Fe304(l 11) - H5C6CHCH2

230 231 232 233 233 234 235 235 236 236 237 237

8. Zinc oxide-ZnO

238

8.1 ZnO(loiO) - HCOOH

238

8.2ZnO(10i0)-CO2

^^^

8.3ZnO(10i0)-C6H6

^^^

8.4ZnO(10i0)-C5H5N

240

8.5 ZnO (lOiO) - C6N3H5,C7N2H6, C7N2H6, C7N3H7

241

8.6ZnO(000i)-CO

242

8.7ZnO(000i)-CO2

243

8.8 ZnO (OOOi) - HCOOH 8.9ZnO(000i)-C5H5N 8.10 ZnO (OOOi)-K 8.11ZnO(0001)-CO 8.12ZnO(0001)-CO2 8.13 ZnO(OOOl) - HCOOH 8.14 ZnO(OOOl) - CH3OH 8.15ZnO(0001)-C6H6 8.16 ZnO(OOOl) - CeHsOH 8.17 ZnO(OOOl) - C5H5N 9. Copper (I) oxide - CU2O 9.1 C u 2 0 ( l l l ) - C H 3 0 H Acknowledgements References

243 244 244 246 246 247 247 247 248 248 249 249 249 250

Chapter 6 (G. Renaud and A. Barbier) Atomic structure of oxide surfaces by surface X-ray scattering 1. Introduction 2. X-ray scattering by surfaces 2.1 Grazing incidence X-rays 2.2 Basic X-ray scattering 2.3 Diffraction by a surface 2.4 Instrumental considerations 3. Surfaces of oxide single crystals

256 258 258 259 260 262 262

3.1 Specific sample preparation and requirements 3.2 MgO(OOl) surface 3.3 a-Al2O3(0001)(lxl) and reconstructed surfaces 3.4 TiO2(110Hlxl) and (001Hlx3) surfaces 3.5 NiO(l 1 l)-p(2x2) reconstructions 3.5.1 NiO( 111 )-p(2x2) structure after air annealing 3.5.2 NiO(l 1 l)-p(2x2) structure after controlled reduction 3.5.3 NiO(l 1 l)-p(2x2) comparison and conclusion 3.6CoO(lll)andMnO(lll) 3.7 ZnO(OOOl) and ZnO (1010) 3.8 SrTiOsCOOl) 3.9Cr2O3(0001Hlxl) 4. Structure of thin oxides films 4.1 NiO(l 1 l)-p(2x2) thin film structure 4.2 AI2O3 layer on NiAl 4.3 Fe203 and Fe304 layers 4.4 Cr203 layers 5. Conclusions and future outlook Acknowledgements References

262 263 266 273 276 277 279 281 281 284 286 288 289 289 292 293 293 294 295 296

Chapter 7 (S.A. Chambers) Structure of thin epitaxial oxide films and their surfaces 1. Introduction 2. Oxidation chemistry and its effect on surface and film structure 2.1 Fe304(l 11) and a-Fe3O4(0001) on Pt(l 11) and a-AÔsCOOOl) 2.2 Cr02 on Ti02(l 10) & (100), and a-Cr203 on a-Al203 3. Lattice mismatch and residual strain fields 3.1 (a-Fe203/a-Cr203)n on a-Al203 4. Crystal symmetry mismatch 4.1 Fe3O4onMgO(001) 5. Conclusions Acknowledgements References

301 303 303 309 313 314 316 317 320 323 323

Chapter 8 (H.-J. Freund, H. Kuhlenbeck and T. Risse) Molecules on well-structured oxide surfaces 1. Introduction 2. Surfaces of oxides with rocksalt structure 3. Surfaces of oxides with corundum structure 4. Surfaces of oxides with rutile structure 5. Surfaces of oxides with layered structures 6. Synopsis References

326 327 341 359 363 366 367

Chapter 9 (C.C. Chusuei, D.C. Meier and D.W. Goodman) Atomic-scale chemical and electronic structure studies of well-defined metal oxide surfaces 1. Introduction 2. Magnesium oxide 2.1 Adsorption of benzene on MgO/Mo(l 00) studied by MIES/UPS 2.2 MIES and TPD analysis of CH3OH and D2O adsorption on MgO/Mo(l00) 2.3 Surface defect characterisation of MgO by MIES 3.Titania 3.1 Structural and electronic properties of Au on Ti02( 110) 3.2 Adsorption of CO on Ai/Ti02 4. Thin-film alumina and mixed-metal oxides 4.1 STM imaging of AI2O3 thin films 4.2 NO adsorption on MgO-NiO and CaO-NiO 5. High pressure STM studies of supported gold clusters 5.1 Imaging Au on Ti02(l 10) at elevated pressures with variable temperature and pressure STM 5.2 Imaging solution-deposited Au6(PPh3)6[BF4]2 on Ti02(l 10) 6. Summary Acknowledgements Reference

373 374 374 378 381 385 385 388 392 393 395 399 400 402 404 405 405

Chapter 10 (A.B. Sherrill and M.A. Barteau) Principles of reactivity from studies of organic reactions on model oxide surfaces 1. Introduction and scope 2. Titanium dioxide as a model for transition metal oxides 3. Case study I: formic acid as a probe of surface properties 3.1 Formic acid decomposition on metal oxides 3.2 The participation of oxygen vacancies in formic acid decomposition 3.3 Tracking formic acid decomposition with scanning probe microscopy 3.4 Catalytic reactions of formic acid on titanium dioxide (110) 3.5 Extension to higher carboxylic acids 4. Case study II: thermal and photoreactions of methanol on Ti02 4.1 Thermal reactions of methanol on Ti02 single crystals 4.2 Thermal reactions of methanol on Ti02 powders 4.3 Thermal reactions of methanol on Sn02 - another rutile oxide 4.4 CH3OH decomposition on fluorite metal oxides 4.5 Decomposition of higher alcohols on titanium dioxide single crystal surfaces 4.6 Higher alcohols: extension to poly crystalline Ti02 and other materials 4.7 Photoreactions of alcohols on Ti02 5. Summary References

409 409 412 412 415 418 420 422 424 425 428 429 430 432 435 437 439 440

Chapter 11 (U. Diebold) The structure of Ti02 surfaces 1. Introduction 2. Bulk structures 2.1 Bulk defects 3. The structure of the rutile Ti02(l 10) surface 3.1 The rutile TiOiCl 10)(lxl) surface 3.1.1 Bulk truncation 3.1.2 Relaxations 3.1.3 Appearance in STM and AFM 3.1.4 Surface defects Step edges Oxygen vacancies created by annealing Oxygen vacancies created by other means Line defects Impurities Crystallographic shear planes 3.2 Reconstructions 3.2.1 Reconstruction under reducing conditions: the structure(s) of the (1x2) phase 3.2.2 Restructuring under oxidising conditions 3.3 Recommendations for surface preparation 4. The structure of the rutile (100) surface 4.1 The TiO2(100)(lxl) surface 4.2 Reconstructions 4.2.1 The microfacet model of the rutile TiO2(100)(lx3) surface 4.2.2 Is the simple microfacet model valid? 5. Rutile (001) 6. Vicinal surfaces 7. Anatase surfaces 7.1 Anatase(lOl) 7.2 Anatase (001) 8. Conclusion References

443 444 445 446 447 447 449 451 455 455 457 459 459 460 461 463 463 465 469 470 470 472 472 473 474 475 476 476 478 479 480

Chapter 12 (G.S. Rohrer) The anisotropy of metal oxide surface properties 1. Introduction 2. Surface character-property relationships 2.1 Particulate oxidation catalysts 2.2 Single crystal surfaces 2.3 Thin films 3. Factors determining surface morphology 4. Factors influencing surface stoichiometry 4.1 The solid-vapour equilibrium 5. Reaction anisotropics on M0O3 surfaces 6. The orientation dependence of the photochemical reactivity of Ti02

485 486 486 488 489 490 496 497 501 506

7. Conclusion Acknowledgement References

510 510 511

Chapter 13 (B.E. Hayden) Vibrational spectroscopy of oxide surfaces 1. Introduction 2. Theoretical and experimental considerations 2.1 High resolution electron energy loss spectroscopy (HREELS) 2.2 Reflection absorption infrared spectroscopy (RAIRS) 3. Surface vibrational modes of the substrate 4. Vibrational spectroscopy of adsorbates 5. Vibrational spectroscopy of adsorbates on oxide supported metals References

514 515 516 521 530 532 539 546

Chapter 14 (R.G. Egdell) Electronic structure of oxide surfaces 1. Introduction 2. Electronic structure of oxides: general considerations 2.1 Models of oxide electronic structure 2.2 Covalency and band dispersion in oxides 2.3 Surface states in oxides 3. Experimental techniques for probing electronic structure of oxides 3.1 Photoemission and inverse photoemission spectroscopies 3.2 Electron energy loss spectroscopy 3.3 Scanning tunnelling spectroscopy 4. Selected case histories 4.1 The rocksah MgO(lOO) surface 4.2 Transition metal rocksalt monoxide (100) surfaces 4.3 Sn02 4.4 Ti02 4.4.1 Comparison of Ti02 and Sn02 4.4.2 Ti02(l 10) 4.4.3 Ti02(l00) 4.5 WOsandNaxWOs 4.5.1 WO3(001) 4.5.2 NaxWOsCOOl) 5. Concluding remarks References

550 551 551 552 553 554 554 556 557 558 558 565 572 577 577 578 585 588 588 593 598 598

Chapter 15 (J.L. de Segovia and E.M. Williams) Desorption induced by electronic transitions, DIET, at oxide surfaces 1. Historical introduction to ESD and PSD 2. Ejection mechanisms: the Feibelman-Knotek model 3. DIET at oxide surfaces 3.1 Oxidised surfaces

608 614 617 618

3.2 Stoichiometric oxides 3.2.1 The Ti02 model surface 3.2.2 The TiO2(100)(lx3) and (110)(lx2) reconstructed surfaces 3.2.3 Some further oxide surfaces studied by DIET 4. Adsorption and reactivity at oxides by DIET 4.1 Small molecules 4.2 Large molecules 5. Summary of relevant kinetics parameters 6. Conclusions Acknowledgements

618 618 622 623 625 625 631 63 8 638 638

Index

645

This Page Intentionally Left Blank

Oxide Surfaces D.P. Woodruff, editor © 2001 Elsevier Science B. V. All rights reserved.

Chapter 1

Metal oxide surfaces and interfaces: concepts and issues Victor E. Henrich Department of Applied Physics, Yale University, P.O. Box 208284, New Haven, CT 06520, USA 1. INTRODUCTION The physics and chemistry of metal oxide surfaces, from the perspective of surface-science experiments and theoretical calculations on well characterized single-crystal samples, was reviewed up to 1992 in Ref 1. Since that time, however, several new areas have emerged in the study of oxide surfaces, and powerful new techniques are being applied to oxides. In 1992, only a few groups had used scanning probe microscopies (SPM) to study metal oxides. A tremendous amount has now been learned from such studies. One of the most important discoveries is that metal oxide surfaces can be far more complex than had been anticipated [2]. Another area that had barely begun in 1992 was the growth of epitaxial oxide films on either single-crystal metal or other oxide substrates. That field has now yielded detailed studies of the surface and chemisorption properties of insulating oxides that cannot be studied by many techniques on bulk crystals because of surface charging. Interesting surface structures have also been produced that are not stable on bulk crystals. Still another new area of oxide surface science is in situ studies of oxide surfaces and chemical reactions in aqueous solution; such oxide-aqueous solution interfacial properties are extremely important in geochemistry [3]. The present volume summarizes the current state of the art in many aspects of the chemistry and physics of metal oxide surfaces. While it cannot be exhaustive, it should prove to be a valuable resource for those interested in the present — and future — of metal^ oxide surface research. The technological and commercial interest in metal oxide surfaces and interfaces has also continued to grow. For decades, oxide surfaces have played a key role in corrosion protection, catalysis, sensors, fuel cells, ceramics, etc. Over the last few years, totally new devices and technologies that rely on the properties of oxide surfaces and interfaces have emerged. Non-volatile ferroelectric field-

effect transistors have been fabricated that will be extremely useful in memory appHcations [4]. Novel metal-oxide-metal spin valve structures are being explored for high-density magnetic storage readout [5]. The properties of oxideoxide interfaces have been invoked in using the ferroelectric field effect to drive a quantum phase transition between normal and superconducting modes in thin oxide superconductor films in order to produce a new type of solid-state switch (an idea that was first proposed in 1956, but has only recently been experimentally realized) [6,7]. The list goes on and on, but the point is that metal oxide surfaces and interfaces will play an ever more important role in a wide raage of future technologies. Before delving into the detailed material contained in the other chapters of this volume, there are some important concepts and issues — some of them unique to oxides — that should be discussed. This chapter highlights a few of those and considers ways in which one should think about the properties of metal oxides. For more information, the reader should consult Refs. 1 and 8 - 1 1 .

2. INTERESTING BULK PROPERTIES Before discussing the surface properties of metal oxides, it is instructive to consider some of the important bulk electronic and geometric properties that characterize them and determine their behavior. 2.1. Electronic properties Perhaps the most important property of metal oxides that determines their electronic structure is their ionicitv. The driving force for this is the desire of oxygen ions in crystal structures to have a closed shell 2s^p^ electronic configuration, thus making the oxygen ions formally O^" (i.e., the "oxide" ion). (Formal ionic charge is an extremely useful chemical concept; it is commonly used in describing ionic materials, and we shall use it here. However, the actual distribution of electron density in a crystal structure is too complicated to fit neatly into such a simple concept [1].) The most highly ionic metal oxides are nontransition-metal oxides such as MgO, whose electron charge density in the crystal closely resembles that of an assembly of Mg^^ and O^" ions. It should be noted that the "free" O^" ion does not exist, however; it is unstable by 7.5 eV compared to the free O atom unless it is stabilized by surrounding positively charged ligands [12]. Covalent bonding is much more important in the later transition-metal oxides such as NiO and La2Cu04 [9]. However, even these relatively covalent metal oxides have a significant ionic component to their bonding and thus also posses the properties that derive from ionicity. tonicity in metal oxides is also discussed in Chaps. 2, 4 and 14. Since they are composed of charged ions, an extremely important property of metal oxides, as for all ionic compounds, is their stoichiometry. Because

Coulomb forces are both large and long range, there is a strong driving force for any ionic material to maintain local charge neutrality. This neutrality must be maintained, both in the bulk and at surfaces and interfaces, on a spatial scale of only a few atoms. As will be discussed throughout much of this book, charge neutrality plays a dominant role in both surface geometric and electronic properties of metal oxides. However, it is equally important in the bulk. For example, a single missing O ion (an "F center") in MgO can exist as a neutral center (i.e., two electrons reside at the defect to compensate for the charge on the missing O^" ion), or it can have only one electron, or no electrons, associated with it. In the last two cases, the defect has a net charge of 1+ or 2+, respectively. Such charged F centers exist, but they must be accompanied by defects of opposite sign nearby in order to maintain local charge neutrality. Charge neutrality is also responsible for the increased catalytic activity of some metal oxides. This is the case when MgO is substitutionally doped with lithium. Since the only possible ionic state for lithium is Li^, an adjacent oxygen ion must be O" (formally). Since it does not have a closed shell electronic configuration, O" is much more chemically reactive than O^", and it is believed that the unique catalytic properties of Li-doped MgO result from the presence of a LiÔ" site [10]. This and other defects in MgO are discussed further in Chap. 3. The ionic nature of metal oxides manifests itself in strong spatial localization of electronic charge in the crystal structure. Since both the cations and anions prefer to exist in a particular valence state in a given crystal structure, it costs energy to add electrons to or remove them from an ion. Thus the energy state of a particular electron is strongly dependent on the number and state of other electrons around it; this is referred to as electron correlation [1,9]. Correlation is the antithesis of simple free-electron band theory, in which electronic states are calculated for a particular crystal structure, regardless of electron occupation, and electrons are then placed in those states until the proper total electronic charge is obtained. As an example, consider electrical conduction in a rocksalt transition-metal oxide such as NiO. This is shown schematically in Fig. 1. The ground state electron configuration of the nickel ions is Ni^^ [the oxygen ions are (formally) taken to be O^"], or 3d^ (d" in Fig. 1). Since only the cations have incomplete shells, electrical conduction corresponds to removing a d-electron from one cation and placing in on a nearby cation: d" d"" -^ d^'^ d"^\ As shown in the lower half of Fig. 1, there is an energy cost for this process of U, which is referred to as the Coulomb repulsion energy, or "Hubbard U". For NiO, U is about 5 - 7 eV. Since this is an enormous energy compared to kT at room temperature (0.025 eV), NiO an insulator even though its 3d shell is only partially occupied. The Hubbard U has no analogy in free-electron band theory.

Charge transfer energies U and A d" d" - d"-' cf ^'

d"-d"^^L

•J

L = hole on O " ligand

n-1

2w

ja+1

d"-' d^ W + W

A

d'd^^^L

u n n

d d

Fig. 1. Schematic representation of the Coulomb repulsion energy, U, and the charge transfer energy, A, in an ionic compound. The lower figure shows the energy levels of the system for those excitations (see text).

Also shown in Fig. 1 is the excitation of an electron from an O^" ion to an empty cation orbital; this is referred to as an interatomic charge transfer excitation (usually called simply a "charge transfer excitation"). In this process, an electron is removed from an O^" ion (L denotes a hole on an oxide ion) and is placed on a cation, making its configuration d^^^. The energy required for this process is the charge transfer energy. A; for NiO, A - 4 eV. This excitation does have an analog in band theory; A is the bandgap energy. However, it cannot be calculated from free-electron band theory since correlation is still present in the change of valence state of both the anion and the cation. Free-electron band theory is able to adequately describe the ground states of some metal oxides (e.g., non-transition-metal oxides and d^ transition-metal oxides), but, because of correlation, it cannot be used for excited states [1,9]; see Chap. 2. However, the terminology and general concepts of band theory are still useful, and it is common to use them in discussing oxides; we will do so in this chapter. For simple d^ metal oxides such as MgO and Ti02, the (predominately) O 2p orbitals correspond to the valence band, and the lowest energy empty cation orbitals correspond to the conduction band [9]. The correspondence becomes less clear in d"" transition-metal oxides, for example, since the highest energy occupied orbitals may be on the cations, and the oxide may be an insulator although the d-orbitals are only partially occupied (e.g., NiO). However, in cases where the use of band terminology would not cause confusion, it is often used. A useful distinction exists between oxides of the transition metals and those of the pre- or post-transition metals (i.e., non-transition-metal oxides) [1]. Nontransition-metal oxides are ones in which the cation bonding orbitals are of s or p character. In the alkali metal oxides, the cations can only exist in a 1+ valence state; for alkaline earth oxides, the cations are always 2+. The important oxides of Group IIIB metals, AI2O3 and In203, can only have 3+ cations. One important exception to this is tin, which can have either a 2+ or 4+ valence. (This may not be surprising, however, since Sn is in Group IVB along with diamond and the highly covalent semiconductors silicon and germanium.) Thus both SnO and Sn02 are stable bulk compounds, and some of the defect properties of Sn02 more nearly resemble those of transition-metal oxides than of other non-transition-metal oxides. Transition-metal cations, on the other hand, in which bonding occurs via delectron orbitals, can generally exist in more than one stable oxidation state; the partially filled d-orbitals also give the oxide a component of directional covalent bonding. For example, molybdenum can exist as Mo^^, Mo^^, Mo"^^, Mo^^ and Mo^^; manganese has stable valences states of Mn^^, Mn^^, Mn"^^, Mn^^ and Mn^^. This opens up a whole new dimension of complexity in both bulk and surface electronic structure. Bulk phases of most of the oxidation states are sta-

ble, each one having a different geometric and electronic structure than the others. New types of bulk defects also become allowed. For example, if an O atom is removed from M0O3, the two electrons that must be left behind in the vicinity of the O-vacancy in order to maintain local charge neutrality can occupy dorbitals on two neighboring Mo ions, changing them from Mo^^ to Mo^^; or they could both reside on the same Mo ion, making it Mo"^^. The polyvalent oxides of molybdenum and vanadium are discussed in Chaps. 2, 4 and 12. Oxides of the lanthanide rare earth elements share some of the properties of transition-metal oxides, at least for cations that can have two stable valence states. (None of the lanthanide rare earth cations have more than two ionic valence states.) Oxides of those elements that can only have a single ionic valence are subject to the limitations imposed on similar non-transition-metal oxides. One actinide rare-earth oxide, UO2, has understandably received quite a bit of attention from surface scientists [1]. Since U can exist in four non-zero valence states, UO2 behaves more like the transition-metal oxides. The electronic properties of rare-earth oxides differ from those of transition-metal oxides, however, because of the presence of partially filled f-electron shells, where the f-electrons are spatially more highly localized than are d-electrons. In some transition-metal oxide systems, including those of Ti, V and Mo, a far larger number of stable bulk phases can exist than there are cation valence states (see Chaps. 4, 11 and 12). This occurs through the formation of structures, called Magneli phases, in which single-crystal regions of one cation valence structure are separated by regularly spaced shear planes on which lower valence cations exist [9]. An example is the class of compounds Tin02n-i, in which regions of the rutile Ti02 structure are separated by shear planes occupied by Ti^^ cations, having a local structure similar to that of corundum Ti203. These planes have been observed intersecting the surface of heavily reduced Ti02 crystals by STM; see the discussion in Chaps. 11 and 14. The rutile-based Magneh phases Tin02-n that have been prepared include: TigOis, TiyOo, Ti60ii, Ti509,Ti407andTi305[13].

Another result of the existence of multiple cation valence states in transition-metal oxides is that the population of those states can easily be changed by optical excitation; that is shown as the charge transfer excitation in Fig. 1. The change in population of cation valence states changes the local electronic structure, which can change the chemical activity of surfaces. For example, the corundum oxide hematite, a-Fe203, contains all Fe^^ cations. When light of energy greater than A is incident on the material, electrons can be excited from O^' anions to Fe cations, changing their valence state to Fe^' [14]. This is the other stable valence state for iron, and the excited state may exist long enough to produce measurable changes in chemical activity. This particular system is of in-

terest in solar photochemistry, since A lies in the visible region of the solar spectrum; this will be discussed in more detail in § 6 below. Some transition-metal oxides exhibit metal-insulator transitions as a function of temperature, pressure, composition, etc. that may produce large changes in bulk electronic structure (see Chap. 4). In corundum V2O3, for example, there is a metal-insulator transition at about 150 K in which the sample's conductivity changes by seven orders-of-magnitude within a few degrees K [9]. The electronic behavior becomes more complex if V2O3 is doped with up to 2 at. % Cr (which goes in substitutionally). Cr-doped V2O3 exhibits the same large metalinsulator transition near 150 K, where it goes from a low-temperature antiferromagnetic insulating phase to a high-temperature paramagnetic metallic phase [15]. However, it displays an additional transition from the paramagnetic metallic phase to a higher-temperature paramagnetic insulating phase at a temperature (around room temperature) that depends on the Cr content. The latter transition produces about a two order-of-magnitude change in conductivity. In this particular system, there is a structural phase change (from monoclinic to corundum) at the low-temperature transition, but no symmetry change across the high-temperature one. This high-temperature transition has recently been investigated for use in positive-thermal-coefficient thermistors [16]. The strong electron correlation effects in metal oxides complicate the interpretation of excited-state spectra in a wide range of experimental techniques [1]. In materials that are well described by itinerant electron band models, such as many metals and elemental semiconductors, electrons can be excited, or added to or removed from the material, without significantly altering the energy level structure of the remaining electrons. In correlated materials, however, changing the state of an electron changes the energies of other electrons in the system, and the concept of a single-electron density-of-states becomes meaningless [9]. Consider two common, powerful experimental methods that are used to determine the spectrum of filled and empty electronic levels in solids: photoemission, and inverse photoemission (or Brehmstrahlung isochromat spectroscopy) [17]. In photoemission, an incident photon excites an electron from its ground state to an allowed, normally empty state above the vacuum level, from which the photoexcited electron is emitted into vacuum, where its kinetic energy is measured with an electron spectrometer. Without correlation, the initial state of the electron can be determined by subtracting its kinetic energy in vacuum from the incident photon energy, properly taking the work function of the surface into account. However, if the initial state of the system (i.e., the atom or ion that will be excited) contained n electrons, then the final state contains only {n-l) electrons. When the photoexcited electron is removed, the other electrons in the vicinity (on both the excited atom and its neighbors) change their energies. Thus the state that the photoemitted electron senses will not be the initial ground state

of the system. In late transition-metal oxides such as NiO, for example, photoemission spectra of both valence and core-level electrons exhibit multiple features (sometimes called the "main line" and "satellites", or "well screened" and "poorly screened" final states) that arise from different electron configurations seen by different exiting electrons [1,18]. Thus the interpretation of experimental photoemission spectra is not straightforward; it is usually necessary to compare the experimental spectra with the results of a theoretical model that includes both the n- and (/7-l)-electron states [18]. Additional aspects of photoemission from oxides are considered in Chap. 14. In inverse photoemission, an electron is added to the system in some allowed state above the vacuum level, and the energy of the photon emitted as it falls into some lower energy allowed (normally empty) state is measured. Without correlation, it is straightforward to determine the spectrum of empty electronic states in such an experiment in a manner analogous to that for photoemission. However, the ground, or initial, state of the system has n electrons, while the final state has (n+\). In correlated materials, the addition of the extra electron perturbs the energies of other electrons, so that the empty-state levels involved in the process are not those of the ^-electron ground state. Again comparison with calculations for both the n- and (/i+l)-electron states is required. 2.2. Crystal structure Metal oxide crystal structures consist of metal cations coordinated in various ways by oxygen ligands [1,9,19]. Most common oxides have their cations octahedrally coordinated with six O ions, although in different crystal structures the octahedron will be distorted in different ways. Such structures include rocksalt, rutile, corundum, anatase and molybdenum trioxide. A few metal oxides have tetrahedrally coordinated cations, such as the wurtzite structure of ZnO. The spinel and inverse-spinel structures have a mixture of octahedrally and tetrahedrally coordinated cations. In the perovskite structure, ABO3, the B cation is octahedrally coordinated, while the A ion has twelve nearest-neighbor O ligands. Other ligand coordinations also occur, such as the planar configuration in some high-Tc oxide superconductors. However, the predominance of octahedral and tetrahedral ligand coordination leads to one of the common ways of graphically representing metal oxide crystal structures. Instead of showing the individual ions, the structure is drawn as a collection of octahedra or tetrahedra, where adjacent polyhedra share common corners, edges or faces [9]; this type of representation is shown in Fig. 2 for three oxide crystal structures whose cations are octahedrally coordinated. In the rhenium trioxide structure, all octahedra share only comers. In rutile, both cornerand edge-sharing are present, and corundum contains a mixture of edge- and face-sharing. This is a widely used representation (see Figs. 2 and 10 in Chap. 12), and it is extremely useful for understanding bulk structures such as shear

IV Vlv V y

(a)

(b)

\ /

M' \

Z-

(c)

Fig. 2. Connected polyhedra representation of the bulk crystal structures of (a), rhenium trioxide; (b) rutile; and (c) corundum. [From Ref 9, by permission of Oxford University Press.]

planes in Magneli phases, etc. However, it is not a particularly good starting point for picturing atomic arrangements on surfaces. A better approach for that purpose is a ball-and-stick model, where all atoms or ions, and their relative positions, are represented [19]. (The one misleading feature of such models for ionic crystals is that the sticks which hold the balls together may be mistaken for directional chemical bonds.) Another good way to schematically show crystal structures is with juxtaposed spheres having the proper relative ionic radii [1,19], such as those in Figs. 3 - 5 below; this method is used extensively in other chapters in this volume. The bulk crystal structures of many oxides are discussed at length in Refs. 9 and 19. Our interest here is in determining the geometric structure of various surfaces that can exist on bulk oxide crystals. 3. FROM BULK TO SURFACE When one considers how a semiconductor crystal will separate in order to create a surface along a particular crystallographic plane, one of the considerations is to break the smallest number of directional covalent bonds. Ionic crystals must be considered somewhat differently [1]. Li a highly ionic oxide such as MgO, the ions are nearly spherical distributions of electronic charge; there are no directional bonds. Since cations in ionic crystals are stabilized by their neighboring anions, and visa versa, one should think in terms of retaining as many ligands around each surface ion as possible (i.e., minimizing coordinative un-

10

saturation). Symmetry dictates that, where possible, the two opposing surfaces formed by cleaving along a plane should be identical. A third, and very important, consideration is not to induce any net electric dipole moment normal to the surface. That will be the case if each plane parallel to the surface has zero net charge; such a plane is said to be "charge neutral". (In the terminology of semiconductor surfaces, that is referred to as "autocompensation" [20]; see also Chaps. 7 and 11 in this volume.) Surfaces that do have a net dipole moment are called "polar" surfaces, and the dipole electric field leads to a divergence in the energy of the crystal [21], making the surfaces unstable; see Chap. 2. (There are indications from work on iron oxides that the situation may be more complex if there is a large amount of interlayer relaxation at the surface [22].) A good example of these principles is the rocksalt structure; we will consider MgO to be specific. In the bulk crystal structure, both cations and anions have six nearest-neighbor ligands in an octahedral configuration; this is the ideal configuration for stabilizing both the O^' and Mg^^ ions. Consider three lowindex surfaces of this crystal: (100), (110) and (111). The only way to form a (100) surface is to cleave the crystal between two bulk (100) planes. The resulting surface consists of an alternating square array of Mg and O ions, with each having five nearest-neighbor ligands (see Fig. 3 below). Since each atomic plane parallel to the (100) surface is atomically flat and contains an equal number of Mg and O ions, it is charge neutral, and there is no net dipole moment in any direction. Both of the surfaces created by the cleave also have exactly the same structure. For the (110) surface, the only possible cleavage plane produces two identical surfaces which have alternating lines of Mg and O ions, with each ion having four nearest-neighbor ligands. Planes parallel to the surface are charge neutral, so there is no net dipole moment. The (111) surface is fundamentally different than either the (100) or (110). There is also only one plane along which one can fracture the bulk crystal structure, but it lies between a plane of Mg^' ions and a plane of O^" ions. The two surfaces produced are fundamentally different in that one consists of an entire plane of Mg ions, while the other contains only O ions. Since planes parallel to the surface have a net charge, and the sign of that charge alternates from plane to plane, there is also a large net dipole moment normal to the surface. To make matters worse, each ion in the surface plane has only three nearest-neighbor ligands, compared to the six that it would prefer to have for stability. Needless to say, the rocksalt (111) surface is not stable; it must either reconstruct or be stabilized by the presence of charged species adsorbed on the surface. If an ionic crystal has two opposite polar faces, then, in order for the total crystal to be charge neutral, the two opposing faces must be of opposite types. This effect can readily be observed macroscopically; a well known case is ZnO. ZnO has a tetrahedrally coordinated wurtzite crystal structure; in the [0001] direction, it consists of alternating planes of Zn^^ and O^' ions [1]. When a sample

11

is prepared with opposite {0001} faces, one of them must consist of exposed O ions [the (OOOT) face], and the other will have exposed Zn ions [the (0001) face]. [Both surfaces will relax so as to try to increase the coordination of the exposed surface ions; this relaxation is greater on the (0001) face since the smaller Zn^^ ions can more easily move down between the larger O^" ions.] The O- and Znfaces etch differently; aa HCl etch produces a macroscopically smooth (0001) surface, but a rough, matte (OOOT) surface [23]. This is, in fact, a simple way to determine the orientation of a sample. The above considerations are borne out experimentally on most rocksalt ionic compounds. For example, when magnesium metal is burned, the tiny MgO smoke particles that are formed are almost perfect cubes (see Fig. 2.4 in Ref. 1). The need to form a non-polar surface and to maximize the ligand coordination of surface ions makes the (100) surface energy much lower than that of other possible surfaces in the rocksalt structure. This is also manifest in the cubic shape of grains of table salt, NaCl. The (110) surface of MgO, whose ions are only four-fold coordinated, is also much less stable than the (100) surface [24]. The various types of atomic reconstructions and rearrangements that exist on oxide surfaces, and their implications for surface electronic structure, are discussed in Chaps. 2, 6 - 8, 11, 12 and 14 in this book and will not be discussed here. However, one interesting point should be mentioned in connection with the MgO surfaces described above. It is possible that the reduced ligand coordination of surface ions could lead to charge transfer between them, with a consequent change in the valence state of surface ions. For MgO, one might expect surface ions to be less ionic than those in the bulk, say Mg^ and O". This was studied for MgO (100) by electron energy loss spectroscopy (EELS) [25]. The experimental results suggested that the valence state of the surface ions was very nearly Mg^^ and O^'; i.e., that little such charge transfer actually occurs. This concept will be considered again below in connection with steps on oxide surfaces. It is not always straightforward to determine what the surface structure of a crystal will be along a particular crystallographic plane. A good example of this is the (0001) surface of the very important corundum structure. The corundum crystal structure is trigonal and can be considered as planes of nearly closepacked O ions perpendicular to the three-fold [0001] axis, separated by planes of metal cations. A first glance at the structure would thus suggest that the crystal would have to be separated between a plane of O anions and one of metal cations, which would result in polar surfaces similar to rocksalt (111). A closer look at the structure, however, reveals that the plane of cations in not atomically flat; one-half of the cations lie closer to one of the 0-ion planes, and the other half lie closer to the other 0-ion plane. If one separates the crystal between the two planes of cations, keeping each metal cation with the O plane to which it is

12

closest (i.e., more tightly bound), the resulting surfaces are charge neutral, have no net dipole moment, and both faces of the cleave have identical structures. Experimental evidence on the corundum metal oxides that have been studied to date suggests that this is indeed the case [1]. In the above discussion, we considered how a crystal should cleave along a particular plane that we chose. However, the crystal may have other ideas; in fact, it is not in general possible to predict from the simple arguments above how easily a crystal will separate along a particular plane. Rocksalt is a simple case, and many ionic rocksalt compounds cleave well — and only — along (100). Some metal oxides having the corundum structure, including Ti203 and Cr203, cleave well along the (10T2) plane (which is the plane that contains the empty octahedral cation sites in this structure) [1], while corundum itself, aAI2O3, does not exhibit any easy cleavage planes. An even more striking case is the rutile structure of Ti02. Its most stable surface plane is (110), which is shown in Fig. 5 below and will be discussed in connection with surface defects. In the bulk rutile structure, the Ti cations are octahedrally coordinated with O ions. On the (110) surface, one-half of the cations still have the bulk six-fold ligand coordination, with the other half having five ligands. This surface has been thoroughly studied by a variety of techniques, as is describe in later chapters, and the predicted structure is indeed observed. However, it is not a cleavage plane; there are no easy cleavage planes in rutile. In spite of this, the considerations of surface stability discussed above are important even for surfaces that must be prepared by methods other than cleaving. Surface morphology is considered more generally in Chap. 12. A few metal oxides have layered bulk structures; two of these — V2O5 and M0O3 — are discussed at length in Chaps. 4, 8 and 12 below. In these compounds, there is strong ionic-covalent bonding within a layer consisting of several planes of atoms, but between these layers there is only weak van der Waals bonding (the layers are charge neutral and do not contain any net dipole moment). The oxides thus peel apart hke mica or graphite. The resulting surfaces expose almost entirely O ions, have a structure almost exactly the same as in the bulk crystal, and are chemically quite inert. The criteria given above for stabiHty of crystal surfaces can also be applied to steps on surfaces. Again consider the rocksalt MgO structure for example. Figure 3 shows an idealized ball model of this (100) surface, including missing Mg and O ions and a monatomic height step. The ions on flat terraces each have five nearest-neighbor ligands. In order to form a step from one terrace to another, the ligand coordination of the atoms along the top edge of the step must be reduced. However, the most stable steps should be those in which the ligand coordination is reduced as little as possible. For rocksalt, that is the step in the [010] direction shown. Each ion along the top edge of the step has only four

13

Rocksalt

[010] [001]

Fig. 3. Model of the rocksalt (100) surface. Large circles are O anions, small circles are metal cations. A [010] step to another (100) terrace is shown, as are both missing anion and missing cation point defects.

ligands; however, the stoichiometry of the crystal is maintained along the step, and the step-edge ions can thus retain the same valence states that they have on the terraces (nearly Mg^^ and O^" in this case). This is indeed the step direction that is observed on MgO (100) surfaces. Another example of steps on oxide surfaces is shown in Fig. 4 for the corundum structure. (One anion point defect is also shown in the figure.) As mentioned above, some corundum metal oxides cleave well along the (10T2) plane, along which the empty octahedral cations sites in the crystal are arranged; cleaving along this plane minimizes coordinative unsaturation of the surface ions. From applying the above principles, it was concluded that the most stable monatomic height step on this surface would be that in the [0221] direction shown [26]. As for the rocksalt case, the step-edge cation coordination is reduced from five on the terraces to four along the top of the step. However, the correct stoichiometry of the crystal is maintained, so that no change in ion va-

14

Corundum

(1012)

[0221] Fig 4.

Model of the comndum (1012) surface, including a [0221] step to another (1012) terrace, and an 0-vacancy point defect.

lence state is required. Subsequent AFM measurements on cleaved Cr203 (lOT 2) showed that the predominant step direction was indeed [0221] [27]. The most common type of nonstoichiometry observed on metal oxide surfaces is the formation of O-vacancy point defects associated with surface reduction [1]. The electronic properties of the metal cations in a metal oxide play an important role in the geometric structure of reduced surfaces. Non-transitionmetal oxides, whose cations have only one stable valence state, have a limited amount of freedom in the defect structures that can be formed. Good examples of that are IMgO and AI2O3. In either oxide, the removal of a single O ion would require that two electrons occupy the defect site, as discussed for bulk defects above (unless oppositely charged defects are present nearby to maintain local charge neutrality). Since the lowest-lying empty state on the cation is several eV above the energy of the O 2p orbitals, they are not accessible. This problem can be circumvented, however, if both anion and cation vacancies are created in

15

Stoichiometric proportion. This appears to occur for MgO; when MgO (100) surfaces are bombarded by inert gas ion ions (which reduces virtually all transition-metal oxide surfaces by preferential removal of O atoms), Auger and XPS spectra indicate very httle deviation of the surface stoichiometry from that of the bulk [1]. Also, no significant changes are observed in the surface electronic states in this process [28]. The complexity of surface defect structures on MgO will be treated thoroughly in Chap. 3. For AI2O3, slightly reduced surfaces can be produced that exhibit complex reconstructions; one of these yields (Vsi x V3I) R ± tan"\V3/ll) LEED patterns, for example (see Chap. 6). Such reconstructions are believed to be ordered structures of point defects [1]; however, the surface stoichiometry is again close to that of the bulk. (The electronic properties of these surfaces have not been studied in detail.) The situation is very different for transition-metal oxides, where it is energetically easy to change the population of surface cation d-orbitals. When the surface of a transition-metal oxide is reduced, the electrons associated with the O vacancies can often populate normally empty d-orbitals on adjacent cations; an example of this is given in the discussion of surface electronic structure in § 4 below. A system whose reduced surface structures have been studied in some detail consists of the oxides of iron. Single crystals of both a-Fe203 (hematite) and Fe304 (magnetite) have been reduced and the resulting surface structural changes monitored by LEED or SPM. The structure of the surface layer, which is generally only 1 - 2 monolayers thick, depends on the preparation conditions (i.e., temperature, oxygen partial pressure, time). On a-Fe203 (0001) surfaces, overlayers of Fe304 (111) and Fei.^Ô (111) have been observed [29], On Fe304 (111) surfaces, an Fci.Ô (111) overlayer structure has also been identified [30]. Unlike the case of non-transition-metal oxides, where the surface cannot really change composition appreciably due to charge balance constraints, thin layers of entirely different metal oxides can be formed on transition-metal oxides. Thus transition-metal oxides offer a much richer array of surface defect properties than do non-transition-metal oxides. Various types of defects on oxide surfaces will be treated in depth in many chapters in this book. An interesting, but troublesome, surface structural effect occurs on one of the important Cu-oxide high-Tc superconductors, YBa2Cu307.;,. The oxygen atoms are quite mobile in that crystal structure, and bulk samples can be easily oxidized or reduced by annealing in an appropriate ambient. However, when single crystals are cleaved in vacuum, some of the O atoms evaporate from the surface, leaving the surface reduced relative to the bulk. In order to study the properties of stoichiometric YBa2Cu307-:v surfaces in vacuum, it was found necessary to cleave the samples at temperatures near 20 K in order to retain the bulk oxygen stoichiometry at the surface [31].

16

4. ELECTRONS ON OXIDE SURFACES Simply from symmetry considerations, the electronic structure of any surface, where the atoms or ions are necessarily coordinatively unsaturated, should be different from that of the bulk. The magnitude and type of the differences between surface and bulk electronic structure depend on the particular oxide, of course; this is addressed in Chaps. 2, 4 and 14 in this volume. However, a few general observations can be made. Since most "surface-sensitive" techniques sample at least a few atomic planes into the sample, it is difficult to experimentally separate the electronic structure of the outermost plane of atoms from that of the planes below. Theoretical calculations are able to clearly separate surface from bulk electronic structure, of course; it is common to calculate a separate electronic density-ofstates for each plane in the crystal structure ("layer density-of-states"). Significant changes from the bulk electronic structure are sometimes found for the surface planes in calculations. However, it is difficult to confirm those results experimentally [1]. In some oxides, the bandgap at the surface has been observed to narrow compared to that of the bulk. The measured core-level binding energies of partially coordinated surface atoms are often shifted, by as much as an eV, from their bulk values [32]; these are referred to as "surface core-level shifts". However, the experimental separation of surface from bulk electronic structure is at present far from satisfactory. The one case in which surface electronic structure can be clearly observed experimentally is for non-stoichiometric surface defects. As in the case of the bulk structure, the surface stoichiometry of an ionic compound plays a crucial role in its surface electronic properties because of the necessity of maintaining local charge neutrality. As a specific example, consider the (110) surface of rutile Ti02; this surface is shown in Fig. 5, and is also discussed in Chaps. 10, 11 and 14. As discussed above, half of the cations on the stoichiometric surface retain their bulk six-fold ligand coordination, while the other half are five-fold coordinated. Rows of "bridging" O ions lie above the six-fold cations. Two Ovacancy point defects are shown schematically in Fig. 5, one a missing bridging O ion, and the other an in-plane O ion. The formal valence of all of the Ti ions on the stoichiometric surface is 4+, and the O ions are taken to be 2-. The stoichiometric surface is charge neutral and has no net dipole moment (to see this, one must consider the plane of five- and six-fold Ti ions, as well as the O ions in that plane and the bridging O ions on either side of it). When one bridging O ion is removed, creating an 0-vacancy point defect, there must be (formally) two electrons remaining at the defect site in order to maintain charge neutrality. They cannot occupy orbitals on neighboring O^" ions because of their closed-shell configuration; the first available empty O orbital would be 3s, which is several eV higher in energy. However, the 3d orbitals on

17

Rutile

[110]

Fig. 5. Model of the rutile (110) surface. Two types of 0-vacancy point defect are shown.

the two (now five-fold) Ti ions immediately adjacent to the vacancy are empty, and it does not cost much energy to place one (or more) of the electrons in them. Thus the schematic picture of such a point defect consists of two Ti^^ ions, each having a 3d^ electron configuration, adjacent to the O vacancy [1]. Needless to say, this population of some of the Ti 3d orbitals gives a very different electronic structure than that of the stoichiometric surface (or the bulk crystal). Stoichiometric Ti02 is an insulator with a charge-transfer bandgap of about 3.1 eV at room temperature [9]. Figure 6 presents a series of angle-integrated photoemission spectra, taken with the photon energies indicated, for an ntype bulk reduced Ti02 (HO) sample whose surface is very nearly stoichiometric [33]; its Fermi level, Ep, is pinned at the bottom of the Ti 3d band. (The location of EF was determined from spectra taken from a gold foil in electrical contact with the Ti02 sample.) There is no emission from the bulk bandgap region. The emission from about 3 to 9 eV is from the O 2p orbitals (i.e., the valence band); emission below that band is inelastic background. The creation of surface O-vacancy defects drastically changes the electronic structure of the surface. Figure 7 presents similar spectra for the sample after its surface has been

18

N(E)

10 5 lUnding Energy (cV)

0

Fig. 6. Series of angle-integrated UPS spectra form stoichiometric Ti02 (110) taken for photon energies between 30 and 76 eV.

slightly reduced by bombardment with inert Ar* ions. In addition to the O 2p emission, electrons are now emitted from a band of states lying slightly below EF in the bulk bandgap. Those are the electrons that are localized at the Ovacancy defects. By following the intensity of that "Ti 3d" emission as a function of photon energy across the Ti 3p -> 3d optical absorption threshold (i.e.,

19

N(E)

15

10

5

Binding Energy (eV) Fig. 7. Series of angle-integrated UPS spectra form reduced Ti02 (110) taken for photon energies between 30 and 76 eV.

by using "resonant photoemission" [1]), the states can be shown to be predominantly of Ti 3d character, as predicted above. This type of O-vacancy defect, with accompanying changes in the population of empty orbitals on neighboring metal cations, is typical of virtually all transition-metal oxides and will be discussed in depth for specific oxides in the following chapters. It is the most important type of defect on metal oxide surfaces.

20

As the photon energy is varied in both Figs. 6 and 7, the emission from the 0 2p band is seen to change amplitude and shape. Since there are no oxygen intra-atomic optical absorption transitions in the photon energy range used here, that emission should not change (i.e., "resonate") if the states were of purely oxygen character. However, in any solid there must be overlap between the electronic orbitals of neighboring atoms or ions; that is chemical bonding. In Ti02, both the O 2p and the Ti 3d orbitals overlap in the region of space between the ion cores; this is referred to as hybridization, or covalent bonding. Thus the "valence band" contains electrons with primarily O 2p character, but also some Ti 3d (and, it turns out, Ti 4s,p) character [33]; the resonant effects seen in this photon range are due to the presence of the Ti 3d wavefunctions in that band. The amount of Ti 3d hybridization in the O 2p band is believed to be of the order of 10 - 20 %. Similarly there is some admixture of O 2p orbitals in the normally empty Ti 3d band (although that hybridization is not seen in a filledstate experiment such as photoemission). This way of looking at hybridization shows why the idea of a formal valence on cations and anions cannot be an exact one. In discussing the electronic structure of 0-vacancy point defects above, we said that the O 2p orbitals were completely occupied. That is not quite true, because hybridization puts some of the O 2p wavefunction into empty states above Ep. Thus there are not quite eight 2s,p electrons on each O ion, and thus their valence state is slightly less negative than 2-. Similarly there is some population (in the valence band) of Ti 3d orbitals, so the Ti ions actually have a valence slightly less positive than 4+. That is one way of saying that the bonding in Ti02 is not purely ionic, but that it also contains some covalent character. There has been recent interest in a different type of defect on transitionmetal oxide surfaces: photoexcited reduced cations. The interest in this comes from atmospheric chemistry and will be discussed further in § 6 below. The idea is to use photons whose energy is greater than the bandgap of the oxide to create charge-transfer excitations of the type shown in Fig. 1, as mentioned in § 1 above. The systems that that have been investigated to date on single-crystal samples include Ti02 [34] and a-Fe203 [14]. The latter was chosen because it is a significant component of tropospheric aerosol particles and thus has the potential to do solar photochemistry. In a-Fe203, high signal-to-noise UPS spectra are able to resolve the changes in valence-band emission when the sample is illuminated with greater-than-bandgap light, compared to the sample in the dark. In the ground state of a-Fe203, all cations are Fe^^. Increased emission at the upper edge of the valence band is seen upon illumination, at the same energy that Fe^^ states are found on a-Fe203 surfaces that are reduced by removal of surface O atoms. Population of the states is completely reversible. These are transient, short-lived excitations, but it appears that enough of them can be ere-

21

ated to observe photochemical effects in the adsorption of SO2 on the surface; this is described in § 6 below. 5. OXIDE-GAS, -LIQUID AND -SOLID INTERFACES Several of the chapters in this book are concerned with the electronic and geometric properties of clean metal oxide surfaces; i.e., the oxide-vacuum interface. Theoretical and experimental determination of clean surface properties is crucial to understanding the physics of metal oxides. This physics can be rather complicated, due to the ionic nature of the system and the strong effect of electron correlation. Some of this work is at the cutting edge of current solid-state physics. Many other chapters are concerned with how metal oxide surfaces react chemically with atoms and molecules. Chaps. 6 - 9 deal with oxide films grown epitaxially on bulk metals or other oxides. The interfacial properties of oxides are where much of their technological, commercial and environmental utility lies, and these fields will be expanding tremendously over the next few years. It is important to consider the connection between the two types of studies. One often refers to the "pressure gap" that separates vacuum studies of chemisorption and catalysis from commercial catalytic reactions, which generally run above —often well above — atmospheric pressure. There is simply no way to properly simulate high pressure conditions in a surface analysis system. Reactions can be run in an attached reaction chamber, which is then pumped out and the sample transferred, under vacuum, into an analysis system equipped for electron, ion and photon spectroscopies. However, except for some optical and x-ray methods that can be performed in situ, the surface analytical tools are not measuring the system under reaction conditions. This gap is well recognized, and both the low- and high-pressure communities keep it in mind when comparing their results. A similar gap exists between clean surface experiments and any adsorption or interface study. For example, an oxide surface may exhibit relaxation or reconstruction due to the reduction in symmetry and ligand coordination at the surface. In the case of ZnO (0001), for example, the small Zn ions relax into the close-packed plane of O ions in order to be more completely surrounded by oxygen, and to reduce the surface dipole moment. But if an adsorbate — almost any adsorbate — is placed on such a surface, it will become an additional ligand for some of the surface ions. Adsorbates therefore, although they are not exactly the atoms or ions that would be present in the bulk structure, do increase the ligand coordination of surface ions. In general, this tends to reduce the amount of relaxation or reconstruction that was present on the clean surface. Thus, one cannot assume that the surface structure that was determined for the oxide-vacuum interface remains when a gas, liquid or other solid is placed on the surface.

22

In general, the surface will tend more toward an ideal termination of the bulk lattice. Similar effects can also occur in surface electronic structure when a moiety is weakly physisorbed onto the surface. The surface core-level shifts measured at the vacuum interface are reduced when atoms or molecules are physisorbed onto the surface. Changes may also occur in the valence electronic structure upon physisorption, such as the disappearance of intrinsic surface states on metals and semiconductors. In chemisorption, adsorbate-substrate interactions are sufficiently strong that one can no longer assume that the surface structure — electronic or geometric — that existed in vacuum remains unchanged. Relaxation and reconstruction that existed in vacuum will be altered not only because of a change in ligand coordination, but because of changes in electronic structure, whether substrate-adsorbate orbital overlap in an acid/base interaction or actual charge transfer in a redox reaction. Vacuum surface properties certainly serve as a guide for what will happen at an interface. However, one should not consider a chemisorbed adsorbate as a perturbation on the clean surface; the adsorbate-substrate complex should be considered as a new system. Many chapters in this book consider chemisorption and catalysis on metal oxide surfaces. Some of the basic ideas of adsorption on oxides will be discussed in § 6 below. An obvious example of the strong effect that chemisorbed molecules can have on surface properties can be seen in Figs. 6 and 7 above for 0-vacancy point defects on Ti02. In Fig. 7, the band of electron emission just below Ep originates from electrons trapped in d-orbitals on Ti cations adjacent to the defect. When the reduced surface is exposed to O2 at room temperature, the defect bandgap emission band almost completely vanishes; the photoemission spectra look very similar to those from the stoichiometric surface in Fig. 6. In this simple case, it is believed that O2 dissociates at 0-vacancy defect sites, removing electrons from the adjacent cations in order to become O^' ions. While there have been thousands of studies of the interaction of gaseous atoms and molecules with oxide surfaces, much less work has been done on the oxide-liquid interface. As mentioned above, this is largely because of the limited number and type of experimental techniques that can penetrate the liquid and access the interface. The greatest interest in oxide-liquid interfaces is from the geochemistry community, where oxide-water interfaces are crucially important in both terrestrial and atmospheric geochemistry. The current state of metal oxide-aqueous solution interfaces was recently reviewed in the report of the Department of Energy, Chemical Sciences Division, Workshop on Metal Oxide Surfaces and Their Interactions with Aqueous Solutions and Microbial Organisms [3]. That report is much broader than the scope of this book, and it addresses a wide range of problems and possibilities in that field. Suffice it to

23

say that oxide-water interfaces in nature are extremely complex, and it will take many years of research before we fully understand them. One of the most active current research areas in metal oxide surface science is oxide-solid interfaces. Work to date falls into three broad categories. The earliest work was the growth of metal films on single-crystal oxide substrates. One of the motivations there was to produce model systems to study the basic properties of oxide-supported metal catalysts. Quite a few groups have been involved in that work, and several papers have reviewed parts of the field [35-37]. An example of this approach is presented in Chap. 9 here. The second area of interest is growth of epitaxial metal oxide films on single crystals of metals. That work was initiated largely in order to expand the types of surface analytical techniques that could be used to study metal oxide surface properties. In the discussion of photoemission spectra on Ti02 in § 4 above, it was mentioned that the normally insulating Ti02 samples were slightly bulk reduced to make them A?-type; this is so that their surfaces would not charge up electrically during photoemission measurements, or other measurements that involved the use of charged particles. While that technique works well for many transition-metal oxides, it does not work for most non-transition-metal oxides. For example, two extremely important oxides are MgO and AI2O3; the former has a bandgap of 7.8 eV, the latter 9.5 eV. They are thus very good insulators, and they cannot be doped sufficiently by reduction or the addition of impurities to make them reasonably good electrical conductors. An alternative approach is to grow thin single-crystal films of those oxides on conducting (usually metal) substrates. If the oxide film thickness is kept below 2 - 3 nm, quantum mechanical tunneling through the film will remove any excess charge that might tend to build up on the surface. Not all oxides or crystal faces can be grown in this way, but there has been a great deal of success for some metal oxides. The quality of films obtainable is not always as good as that for single crystals, due to lattice mismatch and strain between the film and the substrate, the consequent formation of domains in the oxide film, etc. However, interesting surface structures, including polar surfaces, that would not be stable on the bulk oxide can be produced in thin films. Some publications have reviewed that work as well [38,39]. Chapters 6 - 9 here treat various aspects of the growth of metal oxides on metal substrates. Chapters 8 and 9 consider adsorption and reaction on some of those films. The last area is oxide-oxide interfaces. While metal oxide films have been deposited commercially on oxide substrates for many years, and many such systems are of technological importance, the surface science of oxide-oxide interfaces is in its infancy [39,40]. Some oxide-on-oxide growth is discussed in Chaps. 6, 7 and 9. However, its future importance warrants some additional comments on the field. The constraints on the structure of oxide-oxide interfaces are more stringent than for oxide-metal interfaces since both components

24

of the interface are (at least partially) ionic compounds. Not only must there be geometric alignment between substrate and overlayer atoms at the interface, but ligand charges must be matched appropriately. As an example, consider the simple case of the formation of an edge dislocation in a rocksalt compound such as MgO. Such dislocations have a very high energy of formation and rarely occur. They can be pictured by removing one-half of an atomic plane of ions from the bulk structure and then healing the cut by moving the adjacent planes together to fill the gap [41]. Consider removing part of a (100) plane from MgO. When the planes adjacent to the cut are brought together, the entire half-plane would consist of Mg^^ (or O^') ions on one plane directly opposite Mg^^ (or O^) ions on the other plane. There would thus be a large net Coulomb repulsion between the half-planes, which would tend to blow the crystal apart. While that is an extreme example, similar effects can occur at the interface between two different ionic materials, whether they have different geometric structures or simply different lattice constants. A stable interface can only develop if both geometrical and electronic constraints are met. The electronic properties of oxide-oxide interfaces affect a wide range of processes and physical phenomena; the basic physics and chemistry of oxideoxide interface formation are also extremely interesting, as they push the limits of solid-state theory. Several recent technological developments illustrate the importance of oxide-oxide interfaces. Field-effect transistors (FETs) have been successfully fabricated by using metal oxides for both the active conduction medium and the ferroelectric gate [4,42-45]. Such ferroelectric field-effect transistors are useful in memory applications since they are non-volatile; memory retention of the order of hours has been demonstrated in Lai.;cCa;cMn03 LaAlOs devices, and for more than ten days at room temperature in (Pb,La)(Zr,Ti)03 - Lai.;,Sr^Cu04 devices. In such devices, the electronic properties of the oxide-oxide interface are crucial [46], since "any imperfections at the interface, such as the formation of undesirable phases or electronic trapping states, will seriously degrade the performance of the device" [4]. A different type of interfacial effect occurs in superconductor-normal-superconductor (SNS) junctions. In SNS junctions fabricated with high-Tc cuprate superconductors and normal oxide barriers (e.g., CaRuOs, La;^Sri.j,Co03, etc.), the junction resistance is found to be far higher than expected theoretically, and its dependence on device fabrication parameters indicates that it is primarily of interfacial origin [47]. Oxide-oxide interfaces are also extremely important in optical guided wave device structures [48]. Recent work on the growth of perovskites on MgO showed that the quality of the interface depended strongly on which perovskite plane [BaO or Ti02 for the (100) plane of BaTi03] was grown first on the MgO (100) substrate [49]. The reason for the drastic difference in growth mode arises from interface electrostatics, i.e., ion-ion near-neighbor interactions. Very recently, novel ferromagnetic spin order has been observed in

25

LaFeOs - LaCrOs superlattices; the individual oxide layers were between 2.3 and 16.1 A thick [50]. With layers that thin, the quality and properties of the oxide-oxide interface are crucial. Other applications were mentioned in §1 above. Although there have only been a few surface science studies of the atomicscale properties of oxide-oxide interfaces to date, some of which are discussed in Chap. 7, extremely high quality interfaces have been fabricated for many years. A good example is superlattices of Fe304 and NiO, which have been grown with individual oxide layer thickness a small as one Fe304 unit cell and two NiO unit cells [51]. Figure 7.18 in Ref. 1 shows the high quality of the resulting superlattices from their low-angle x-ray diffraction spectra. 6. ADSORPTION ON METAL OXIDES Since the surfaces of metal oxides contain both negatively charged anions and positively charged cations — often with more than one cation valence state present —many avenues of adsorption are available. The various chemisorption and physisorption possibilities on metal oxides are discussed at length in Ref. 1; they will only be briefly summarized here. Using the scheme of Ref. 1, a major distinction can be made between: Non-dissociative or molecular adsorption and Dissociative adsorption. While this distinction is generally a straightforward one, it says nothing about the chemical nature of the bonding or about charge transfer between the surface and the adsorbate. There are four main types of interaction that can occur: weak electrostatic or dispersion acid/base or donor/acceptor oxidation/reduction with electron transfer and oxidation/reduction with oxygen transfer.

26

The first of these implies physisorption only. It dominates for many atoms and molecules at low temperatures, but it is of little interest catalytically. Therefore it has not been as thoroughly studied as the other types. In acid/base, or donor/acceptor, reactions, bonding results form the overlap of filled orbitals on the "donor" and empty orbitals on the "acceptor". Surface cations are generally Lewis acids and act as electron acceptors, while surface O ions are Lewis bases and can donate electrons to acceptor adsorbates. In lower oxides of the transition metals (i.e., in which the cations are in an oxidation state lower than their maximal valency), cations may also be able to donate electrons in an acid/base reaction. Although one talks of donating and accepting electrons in acid/base reactions, the electrons are in no sense free, and there is no actual electron transfer involved. This type of bonding can be either molecular or dissociative. Actual electron transfer does occur in oxidation/reduction, or "redox", reactions. In this type of reaction, there is a change in the oxidation state of the adsorbate. A simple example is the chemisorption of an alkali atom, in which it becomes a 1+ ion, transferring its outer electron to empty electron orbitals of the substrate. It is the large electric dipole moment created by this charge transfer process that lowers the work function of surfaces on which alkali atoms are adsorbed (e.g., "cesiation") by up to several eV. This type of bonding is generally strong, and it can also be either molecular or dissociative. A unique property of metal oxides is their ability to also exchange O ions with adsorbates; this property makes them very useful as partial oxidation catalysts. Such ion transport is a type of redox reaction in that charge (here in the form of an O^" ion) is really transferred from the substrate to the adsorbate. A simple example of a redox reaction with O ion transfer is the oxidation of CO: CO + O'lattice ^ CO2 + 2 e". The oxide surface becomes reduced in this process; if this were a step in a catalytic reaction, some other O-containing species would have to donate an O atom back to the substrate in order for the reaction to continue. Like the other types of reaction, O-transfer adsorption can be either molecular or dissociative. An interesting example of this type of redox reaction — and the role that surface defects can play — is the interaction of CO with the surface of V2O3. This has been recently studied on both pure and Cr-doped V2O3 (0001) singlecrystal surfaces by ultraviolet photoemission, among other techniques [52]. V2O3 doped with 1.5 at.% Cr is an insulator at room temperature, with the V 3d electrons (each cation has a 3(f configuration) forming a band that lies just below Ep. This is shown as the heavy curve in the UPS spectra in Fig. 8(a). The other curves in Fig. 8(a) show how the spectrum changes when the stoichiomet-

27

Clean surface 100 L CO 103 L CO

N(E)

AN(E)

Binding Energy [eV]

Fig. 8. (a) UPS spectra of stoichiometric, insulating Cr(l .5%): V2O3 (0001) at 273K, both clean and exposed to CO (hv = 40.8 eV). (b) Differences created by subtracting the clean surface spectrum from the CO-exposed spectra.

ric, low-defect-density surface is exposed to the amounts of CO indicated at room temperature. Up to 10^ L, there is virtually no effect on the spectra; this can be seen most clearly in the difference spectra (where the clean-surface spectrum is subtracted from each of the other spectra) in Fig. 8(b). However, above a certain threshold, the UPS spectrum changes dramatically. Emission from the O 2p band decreases, while there is increased emission from the V 3d band and at several eV below the O 2p band; the emission near 11 eV is at the location expected for the CO 4a molecular orbital. XPS spectra of the V 2p core levels

28

show that some of the V^^ cations have been reduced to lower valence states in the process; carbon is also present on the surface. That this process involves irreversible removal of O ions from the V2O3 surface can be seen by heating the surface to remove the adsorbed moieties. This is shown in Fig. 9, where the surface was first exposed to 10"^ L CO at room temperature, and then heated sequentially to the temperatures shown; a clean-surface spectrum is included for reference. After heating to 473 K, no carbon remained on the surface, and heating to higher temperatures produced no additional changes in the UPS spectra. However, instead of simply desorbing the adsorbed moiety and returning to something close to the clean surface, more complex changes occur in the spectra. While emission in the CO 4a and O 2p band regions does move in the direction of the clean surface spectrum, the intensity of emission form the V 3d band increases, indicating further reduction of the surface. Coupled with the absence of carbon above 473 K, this suggests that the CO desorbed as CO2, taking surface lattice O ions with it. The threshold behavior seen in Fig. 8 suggests that CO does not react readily with the nearly defect-free stoichiometric surface, but that some O ions must be removed, forming 0-vacancy point defects, before a significant reaction can

10-* L CO at 273 K 323 K 373 K 423 K 473 K Clean surface

N(E)

14

12

10

8

6

4

2

Binding Energy [eV]

Fig. 9. UPS spectra of CO-exposed Cr(1.5%):V203 (0001) heated to 473 K; the clean surface is included for comparison, (hv = 40.8 eV)

29

take place. This can be confirmed by exposing a CO-reduced and then annealed surface (i.e., the final state in Fig. 9) to additional CO; this is shown in Fig. 10. While the spectral changes are smaller here, the difference spectra in Fig. 10(b) show that CO interacts for all exposures used; there is no reaction threshold as there was for the stoichiometric surface. When the surface in Fig. 10, after exposure to 10^^ L CO, is heated to remove the adsorbed species, the changes are similar to those in Fig. 9, with the intensity of the V 3d band increasing as CO2 is desorbed from the surface, removing still more lattice oxygen. When the above cycle is repeated several times, the reduction of the surface continues to increase, as shown in Fig. 11. The dashed spec-

CO-reduced surface 100 L CO lO^LCO

N(E)

0-

•«••••««•«

^'•""^vj^:.

.,jîr^

' • ^ i . ^ '

AN(E)

(b) Differences 1—^ 14

12

10

1—'—r 4 2

Binding Energy [eV] Fig. 10. (a) UPS spectra of slightly reduced, insulating Cr(1.5%):V203 (0001) exposed to CO at 273 K (hv = 40.8 eV). Prior to the CO exposures shown in this figure, the (stoichiometric) surface was reduced by one cycle of CO exposure and desorption. (b) Difference created by subtracting the first spectrum from successive CO-exposed spectra.

30

• Clean surface • One CO/desorb cycle "Two cycles -Three cycles - Four cycles / ^^

N(E)

10

8

6

4

2

Binding Energy [eV] Fig. 11. UPS spectra showing the effects of repeated cycles of CO exposure (10"* L at 273 K) followed by desorption at 473 K (hv = 40.8 eV). The first spectrum is of clean, stoichiometric Cr:V203; successive spectra were taken following desorption of the CO dose.

tra in Fig. 11 are taken sequentially after each cycle of exposure to 10"^ L CO followed by desorption to 473 K. The four cycles shown have reduced the sur face sufficiently that it should be thought of as a surface layer of a lower oxide of vanadium rather than just reduction of V2O3. (Similar experiments on pure V2O3 showed that the effect was not due to the presence of substitutional Cr.) There has been recent interest in a somewhat different aspect of adsorption and reaction on metal oxides: photocatalysis. The interest stems partially from that role that some transition-metal oxides can play in photochemical reactions in the atmosphere. Atmospheric aerosol particles can act as substrates to catalyze heterogeneous photochemical reactions in the troposphere. Most tropospheric aerosols are silicates, aluminosilicates and salts whose bandgaps are larger than the cutoff of solar radiation in the troposphere (about 4.3 eV); they are thus unable to participate directly in photoexcited reactions. However, transition-metal oxides that have much smaller bandgaps also occur as aerosols — the most prevalent ones are the oxides of iron and manganese — and these materials may thus undergo charge-transfer excitations (discussed above) in the pres-

31

ence of sunlight through electron-hole pair creation, which enable photochemical reactions for which the material would otherwise be inert. Some of the earliest work on such heterogeneous photocatalysis using transition-metal oxides was performed by Yates' group using Ti02; Ref. 34 is an excellent review of the photochemistry of that material. Chapters 10 and 12 in this volume also discusses photochemistry. Recent experiments have focused on the potential role that hematite, aFe203, might play in the oxidation of SO2 in the formation of acid rain [14,53]. Single-crystal hematite was used as a model for atmospheric aerosol particles, and solar radiation was simulated by use of a focused 200 W Hg(Xe) arc lamp (the integrated intensity reaching the sample for photon energies greater than the charge-transfer bandgap of hematite, 2.2 eV, was roughly 70 times the solar flux in that wavelength range). Figure 12(a) shows photoemission spectra for stoichiometric a-Fe2O3(0001) at 300 K, clean and following exposure to 10^ L SO2 while being irradiated by the Hg(Xe) lamp. The dominant changes produced after SO2 exposure are increases in intensity at 4.5 and 10 eV. These can be seen more clearly by subtracting the clean-surface spectrum from the spectra after SO2 exposure; this is done in Fig. 12(b) both for exposure during UV irradiation [i.e., the data in Fig. 12(a)] and for exposure in the dark. The shape of the two differences is the same, indicating that the same moiety adsorbs both in the dark and under illumination; that moiety is believed to be either SOs^' or S04^". However, the amount of SO2 adsorbed is much greater under illumination. The increased S02-hematite interaction is believed to result from the formation of transient Fe^^ ions on the a-Fe203 surface by the greater-than-bandgap illumination (see § 4 above); these Fe^^ cations are much more reactive chemically than are the maximal valency Fe^^ cations that are present in a-Fe203 in the dark. The adsorption that does occur in the dark is believed to do so at a small density of Fe^^ sites associated with surface O-vacancy sites that are always present on the surface. The range of adsorption processes that can occur on metal oxide surfaces is very broad; these will be discussed in many of the other chapters in this book. Chapter 5 considers in detail the atomic positions of adsorbed moieties on several different oxide surfaces. The use of vibrational spectroscopies as a complement to electronic techniques is discussed in Chap. 13. Chapter 15 considers desorption from oxide surfaces induced by incident electrons or photons. 7. ENDNOTE All of the knowledge that has been gained about the properties of metal oxide surfaces and interfaces, particularly over the past 8 - 1 0 years, would fill many volumes this size. There are also several dozen research groups currently

32

a-Fe2O3(0001) exposed to SO2at300K

(a)

UPS spectra

N(E)

0-

- | — I — I — \ — I — I — I — I — I — I — I — I — I — I — r

16

14

12

10 8 6 4 Binding Energy [eV]

(b)

2

differences (xl2)

UV-irradiated during dose NoUV irradiation AN(E)

\r

1

16

\

14

I

I

12

I

I

I

i~

~i—I—\—I—I—\—I—r

6 4 10 8 Binding Energy [eV]

2

0

Fig. 12. (a) He ii UPS spectra for stoichiometric a-Fe203 (0001) at 300 K, clean and following exposure to 10^ L SO2 [while irradiated by the Hg(Xe) arc lamp]; (b) Differences between the S02-exposed spectrum and the clean-surface spectrum (x 12), displaying features arising from adsorbate molecular orbitals.

33

involved in expanding that knowledge base even further. However, the authors who have contributed to the current volume, and the range of topics covered here, comprise a representative cross-section of the field. This volume should thus serve both as a useful portal for researchers wanting to enter the field, and a resource for those of us already in it. ACKNOWLEDGMENTS Some of the research reported here was partially supported by NSF Grant CTS-9610140, and by The Petroleum Research Fund Grant 28797-ACS. REFERENCES [I] [2] [3]

[4] [5] [6] [7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18]

V.E. Henrich and P. A. Cox, The Surface Science of Metal Oxides, Cambridge University Press, Cambridge, 1994. D.A. Bonnell, Prog. Surf. Sci., 57 (1998) 187. G.E. Brown, Jr., V.E. Henrich, W.H. Casey, D.L. Clark, C. Egleston, A. Felmy, D.W. Goodman, M. Gratzel, G. Maciel, M.I. McCarthy, K.H. Nealson, D A . Sverjensky, M.F. Toney and J.M. Zachara, Chem. Rev., 99 (1999) 77. S. Mathews, R. Ramesh, T. Venkatesan and J. Benedetto, Science, 276 (1997) 238. S. Araki, M. Sano, S. Li, Y. Tsuchiya, O. Redon, T. Sasaki, N. Ito, K. Terunuma, H. Morita and M. Matsuzaki, J. Appl. Phys., 89 (2000) 5377. CH. Ahn, J.-M. Triscone, N. Archibald, M. Decroux, R.H. Hammond, T.H. Geballe, 0. Fischer and M.R. Beasley, Science, 269 (1995) 373. C.H. Ahn, S. Gariglio, P. Paruch, T. Tybell, L. Antognazza and J.-M. Triscone, Science, 284(1999)1152. C. Noguera, Physics and Chemistry at Oxide Surfaces, Cambridge University Press, Cambridge, 1996. P.A. Cox, Transition Metal Oxides, Clarendon Press, Oxford, 1992. H.H. Kung, Transition Metal Oxides. Surface Chemistry and Catalysis, Elsevier, Amsterdam, 1989. H.-J. Freund, H. Kuhlenbeck and V. Staemmler, Rep. Prog. Phys., 59 (1996) 283. C. Kittel, Introduction to Solid State Physics, Fifth ed., John Wiley & Sons, New York, 1976. T.B. Reed, Free Energy of Formation of Binary Compounds, MIT Press, Cambridge, 1971. D.S. Toledano, E.R. Dufresne and V.E. Henrich, J. Vac. Sci. Technol., 16 (1998) 1050. H. Kuwamoto, J.M Honig and J. Appel, Phys. Rev. B, 22 (1980) 2626. H.R. Kokabi, M. Rapeaux, J.A. Aymami and G. Desgardin, Mater. Sci. Eng. B, 38 (1996) 80. D.P. Woodruff and T.A. Delchar, Modern Techniques of Surface Science, Cambridge University Press, Cambridge, 1986. GA. Sawatzky, in; J. Kanamori and A. Kotani (Eds.), Core Level Spectroscopy in Condensed Systems, Springer Verlag Series in Solid State Sciences, Vol. 81, SpringerVerlag, Berlin, 1988.

34

[19] R.J. Lad, in: W.N. Unertl (Ed.), Handbook of Surface Science, Vol. 1, Elsevier, Amsterdam, 1996, p. 185. [20] C.B. Duke, Chem. Rev., 96 (1996) 1237. [21] There is, however, some recent theoretical work on ways to deal with polar surfaces. See, for example, J.H. Harding, Surf. Sci., 422 (1999) 87; and, M.H. Finnis, phys. stat. sol. A, 166(1998)398. [22] S.A. Chambers and S.I. Yi, Surf. Sci., 439 (1999) L785. [23] R.R. Gay, M.H. Nodine, V.E. Henrich, H.J. Zeiger and E.I. Solomon, J. Amer. Chem. Soc, 102(1980)6752. [24] V.E. Henrich, Surf Sci., 57 (1976) 385. [25] V.E. Henrich, G. Dresselhaus and H.J. Zeiger, Phys. Rev. B, 22 (1980) 4764. [26] V.E. Henrich, personal communication. [27] R.J. Lad and M.D. Antonik, Ceramic Trans., 24 (1991) 359. [28] L.H. Tjeng, A.R. Vos and G A Sawatzky, Surf. Sci., 235 (1990) 269. [29] R.J. Lad and V.E. Henrich, Surf Sci., 193 (1988) 81. [30] N.G. Condon, P.M. Liebsle, T. Parker, A.R. Lennie, D.J. Vaughan and G. Thornton,, Phys. Rev. B, 55 (1997) 15885. [31] J.G. Tobin, C.G. Olson, C. Gu, J.Z. Liu, F.R. Sohal, M.J. Fluss, R.H. Howell, J.C. O'Brien, H.B. Radousky and P A . Sterne, Phys. Rev. B, 45 (1992) 5563. [32] PH. Citrin and G.K. Wertheim, Phys. Rev. B, 27 (1983) 3176. [33] Z. Zhang, S.-P. Jeng and V.E. Henrich, Phys. Rev. B, 43 (1991) 12004. [34] A.L. Linsebigler, G. Lu and J.T. Yates Jr., Chem. Rev., 95 (1995) 735. [35] M. Baumer and H.-J. Freund, Prog. Surf Sci., 61 (1999) 127. [36] T.E. Madey, U. Diebold and J.M. Pan, in: H.J. Freund and E. Umbach (Eds), Adsorption on Ordered Surfaces of Ionic Solids and Thin Films, Springer, Berlin, 1993, p. 147. [37] C.T. Campbell, Surf. Sci. Reports, 27 (1997) 1. [38] S.C. Street, C. Xu and D.W. Goodman, Annu. Rev. Phys. Chem., 48 (1997) 43. [39] S.A. Chambers, Surf Sci. Reports (in press). [40] R.J. Lad, Surf Rev. Lett., 2 (1995) 109. [41] A.L. Ruoff, Introduction to Materials Science, Prentice-Hall, Englewood Cliffs, NJ, 1972. [42] Y. Watanabe, Phys. Rev. B, 57 (1998) R5563. [43] S.A. Campbell, D.C. Gilmer, X.-C. Wang, M.-T. Hsieh, H.-S. Kim, W.L. Gladfelter and J. Yan, IEEE Trans. Electr. Dev., 44 (1997) 104. [44] Y. Watanabe, M. Tanamura and Y. Matsumoto, Jpn. J. Appl. Phys., 35 (1996) 1564. [45] Y. Watanabe, Appl. Phys. Lett., 66 (1995) 1770. [46] T.S. Kim, C.H. Kim and M.H. Oh, J. Appl. Phys., 76 (1994) 4316. [47] K. Char, L. Antognazza and T.H. Geballe, Appl. Phys. Lett., 63 (1993) 2420. [48] A.M. Glass, Mater. Res. Soc. Bull., 13 (1988) 16. [49] R. A. McKee, F.J. Walker, E D . Sprecht, G.E. Jellison, Jr., LA. Boatner and J.H. Harding, Phys. Rev. Lett., 72 (1994) 2741. [50] K. Ueda, H. Tabata and T. Kawai, Science, 280 (1998) 1064. [51] D M . Lind, S.D. Berry, G Chem, H. Mathias and L.R. Testardi, Phys. Rev. B, 45 (1992) 1838. [52] D.S. Toledano, V.E. Henrich and P. Metcalf, J. Vac. Sci. Technol. A, 18 (2000) 1906. [53] D.S. Toledano and V.E. Henrich, J. Phys. Chem. (in press).


35

Chapter 2

Clean Oxide Surfaces: a theoretical review Claudine Noguera Laboratoire de Physique des Solides, UMR CNRS 8502, Universite Paris-Sud, 91405 Orsay, France 1. I N T R O D U C T I O N The last five years have witnessed a tremendous effort to better produce, characterise and study insulating oxide surfaces. Several reasons stem for the rapid development of this field. They are related to experimental considerations — a better control of the fabrication of surfaces, a more thorough use of advanced spectroscopic or structural tools — but also to a more accurate recognition of the technological importance of high quality oxide surfaces in catalysis, magnetic recording, as sensors or as constituents of artificial nano-materials. A simultaneous effort to simulate these surfaces has been performed. While most theoretical works quoted in the Henrich's and Noguera's books [1,2] five years ago, referred to semi-empirical methods, first principles approaches are now routinely used to calculate the ground state properties of surfaces. The delay, with respect to the simulation of semi-conductor surfaces, may be assigned to the larger computational resources necessary to quantum-mechanically treat oxygen and transition metal atoms. However, ab initio methods find their limitations as soon as defects break the symmetry of the surface or when large cell reconstructions take place. In addition, most of them are based on a variational principle which limits their use to the prediction of ground state properties. Accounting quantitatively for quasi-particle spectra, optical excitations or d -^ d excitations on oxide surfaces remains a challenge. The same is true as regards analytical theories describing the basic microscopic processes at work on these surfaces. These considerations will be examplified here on three aspects of oxide

36

surface physics, namely (i) the non-polar stoichiometric surfaces, (ii) the oxygen deficient non-polar surfaces and finally (iii) the polar surfaces. The specific properties of the first family are induced by the low coordination of the surface atoms and are shared by other low-dimensional oxide systems, such as non-supported clusters or ultra-thin films [3]. In the second family, the presence of oxygen vacancies is responsible for strong electron redistributions in the surface layers which can result in a higher surface reactivity and/or surface reconstructions. A similar situation is met in the third family due to the strong electrostatic polar instability [4]. After a presentation of the available numerical methods in Section 2 and a review of simulation works in Section 3, we will consider successively the three families of surfaces in Sections 4, 5 and 6, from the point of view of their generic properties and the theoretical arguments to understand them. Finally, we will conclude on open questions for future investigations. 2. NUMERICAL METHODS The first historical step in the simulation of insulating oxides relied on the idea that most of these materials are highly ionic. It has led to the development of classical models of cohesion in which the leading forces result from Coulomb interactions between ionic charges. These approaches usually give a quite reasonable description of atomic positions and vibrations, but cannot yield information on electronic degrees of freedom, which require quantum treatments. As a function of time, the latters have evolved from semi-empirical — tight-binding or semi-empirical Hartree-Fock (HF) — to first principles methods, in which there are no adjustable parameters. In parallel, quantum-chemistry all-electron approaches allow a treatment of electron correlations beyond the HF approximation, but only for small finite size systems. We will stress the strengths and limitations of these methods when applied to oxide surfaces, and we will also place special emphasis on some methods which can describe excited state properties, or ground state properties of highly correlated oxides. 2.1. Classical approaches In classical approaches, the lattice energy E, which is the diff'erence between the total energy of the system and that of its component ions at infinity, is expanded into a sum of pair, triplet, etc terms:

E = lEEij + l Z Eijk + ...

(1)

Keeping only the Eij terms yields the so-called pair-potential approximation. The pair interactions include Coulomb (l/r^j dependence), van der

37

Waals {1/rfj dependence) and short range terms. The latters result from PauU's principle, which impedes the overlap of closed electronic shells. It may be taken under a Lennard-Jones ( l/rjj variation), or a under a Born Mayer form {Bij exp{—rij/p)). The parameters entering Eij are fitted, so that the properties of the perfect lattice are reproduced. In the applications to surfaces, they are kept equal to their bulk values, in particular the ionic charges Qi. Many body effects are included in the triplet and higher order terms in Eq. (1). The triplet terms, for example, give energy variations with bond angles. In addition, low-coordinated atoms, which are submitted to non-vanishing electric fields, may be strongly polarized. The shell model [5] was designed to describe these effects, in an approximate way. Rather than being written under empirical forms, potentials can also be derived from the functional expression of the free ion energies as a function of the electron density [6]. Corrections may also be included to account for the instantaneous environment in which an ion finds itself, through modifications of the ionic radii, shapes and charge distributions [7,8]. A good review of the state-of-the-art interatomic potentials in solid state chemistry can be found in Ref. 9. In order to predict the ground state geometry, energy minimization with respect to the ionic positions can be performed by various techniques: static ones — steepest descent, conjugate gradient methods — or dynamical ones — damped dynamics, or molecular dynamics [10]. The latters allow an exploration of the configurational space, through the technique of simulated annealing. Classical simulations have long been performed to predict surface phenomena. Due to the increasing availability of first principles calculations for medium size systems, they are nowadays essentially restricted to large size systems, like zeolites, grain boundaries, complex interfaces, etc. 2.2. Semi-empirical quantumi methods At the lowest level of sophistication of quantum treatments, the tightbinding method and the semi-empirical HF method reduce the complexity of the interacting electron system to the diagonalization of an effective one-electron Hamiltonian matrix, whose elements contain empirical parameters. The electronic wave functions are expanded on a minimal basis set of atomic or Slater orbitals centered on the atoms and usually restricted to valence orbitals. The matrix elements are self-consistently determined or not, depending upon the method. In the tight-binding method, the elements of the Hamiltonian matrix are treated as adjustable parameters to be fitted to experimental or first-

38

principles calculation results. To estimate the total energy, an additional short-range repulsion term is added to the electronic contribution. The dependence of the parameters upon interatomic distances are fitted in order to account for the bulk equilibrium structure and elastic constants. It is assumed that the parameters are transferable from bulk to surfaces [11]. This method has been widely used in the field of semi-conductors, and has yielded a good qualitative description of their surface properties. However, in more ionic systems, it completely disregards the variations of the electrostatic potential with the local environment of the atoms (see below), which is the reason for some wrong predictions made in the past. The semi-empirical HF approaches can be viewed as self-consistent tightbinding methods, since they incorporate the self-consistent relationship between charges and electrostatic potentials in the expression of the Hamiltonian matrix elements. The empirical parameters either fix the spatial dependence of the basis set wave functions, or are incorporated directly into some Hamiltonian matrix elements [12-14]. In the application to oxide surfaces, these methods can account for the shifts of the atomic orbital effective levels, i.e. for the surface modification of electronegativity, for example (Section 4.4). This is a very important point, both for obtaining reliable eigen-energy spectra, and to properly describe the reactivity of certain types of atoms as a function of their environment. It can also account for the surface charge redistribution, and for its influence on the structural degrees of freedom, contrary to classical approaches. These methods can be combined with geometry optimization as well as with molecular dynamics algorithms, with forces obtained from the gradients of the total quantum energy [10]. This equally applies to all quantum methods, quoted in the following. 2.3. T h e ab initio H F method and methods beyond The HF method treats electron-electron interactions at a mean field level, with the Hartree and exchange interactions exactly written. The method can be implemented either in its spin restricted form (RHF), for closed shell systems, or in the unrestricted form (UHF) for open-shell or strongly correlated systems. In the first case, the one-electron orbitals are identical for electrons of both spin directions, while UHF can account for a non-uniform spin density. The one electron orbitals, which are determined in the course of the self-consistent resolution of the HF equations, are expanded on an over-complete basis set of optimized variational functions. The HF method has been implemented on periodic systems [15], including bulk and surface crystalline materials. It has proved very useful in the description of magnetic insulators, but it has also successfully been used for describing surface properties of a large number of simple oxide surfaces.

39

One of the very serious problems with the HF approach is the unscreened nature of the Coulomb and exchange interactions. The bare value of a typical intra-atomic Coulomb integral U is in the range 15-20 eV, while screening in solids weakens it by more than a factor of 3. A large overestimation of the HOMO-LUMO (HOMO= Highest Occupied Molecular Orbital; LUMO=Lowest Unoccupied Molecular Orbital) gap in insulators results. However, there have been quite accurate predictions of gap values or d ^ d excitation energies, based on HF total energy differences (Sections 4.4 and 4.5). Quantum chemical approaches have long tried to solve N-electron Hamiltonians of finite systems, beyond HF, at levels of increasing sophistication. The essential ingredient is the expansion of the ground state and excited state N-electron wave functions on a basis set of Slater determinants, built from one-electron HF eigen-functions. This so-called configuration interaction method (CI) includes explicitely electron correlation effects. Depending upon the basis set size and the number and nature of configurations kept, ground state and some types of excited state properties may be calculated [16]. By essence, these methods can only treat finite size systems containing a limited number of electrons, although some hybrid DFT-CI schemes (DFT== Density Functional Theory, see Section 2.4) try to deal with extended systems [17]. Bulks or surfaces are represented by a "quantum part" containing a small number of atoms, which is embedded in an array of point charges. Varying the cluster size gives hints on the validity of the representation. However, due to the heavy demand in computational resources, these methods, usually, can only treat few atoms. 2.4. The ab initio Density Functional Theory and methods beyond The DFT provides another route to include correlation effects. It gives a prescription for calculating the ground state energy of an assembly of atoms at fixed positions, as a function of the electron density n{r) [18]. Assuming that there exists a system of non-interacting electrons with the same n{r)j the ground state energy of the interacting system may be expressed as a function of the eigen-functions and eigen-values of the non-interacting system. The latters are solutions of one-electron self-consistent equations, named Kohn Sham equations, which include a Hartree potential, an exchange term, due to the Pauli's principle and a correlation term. No exact expression exists for the correlation energy functional. The Local Density Approximation (LDA) replaces the exchange and correlation potential Vxc, which is by essence non-local, by a local potential built from the properties of the homogeneous electron gas (HEG). At each point f of the inhomogeneous system, the exchange correlation energy is

40

equal to that of the HEG: £'^^^(n), with the uniform electron density n functionally replaced by the actual density n{r). In open-shell or magnetic systems, it is possible to build an exchange and correlation functional which depends on both the electron density and the spin density. This is the Local Spin Density Approximation (LSDA). Excellent reviews of DFTLDA or LSDA may be found in Ref. 19-21 One step beyond the LDA, the GGA (generalized gradient approximation) represents a semi-local approximation, in which not only the density but its gradient are locally taken into account [22-24]. The GGA corrects a large part of the systematic overestimation of the LDA cohesive energies. In particular, at surfaces, the GGA yields a much better description of adsorption phenomena than the LDA. In order to solve the Kohn Sham equations, an expansion of the oneelectron wave functions on a basis set is performed. Both localized basis sets and plane wave ones are currently used. Localized basis sets have the advantage of their small size. However they are attached to the atomic positions, which yields non-zero Pulay forces in geometry optimization and molecular dynamics. Plane waves, on the other hand, provide a uniform sampling of space, whatever the specific conformation of the system; they are independent of the atomic positions, but they require the use of pseudopotentials to mimick core electrons and a very large number of vectors is necessary in standard surface calculations. The density functional theory is a theory for the ground state. It cannot predict excitation properties. In particular, there is no theoretical justification to identify the Kohn-Sham eigen-values with quasi-particle energies, although this is currently done. In addition, DFT-LDA fails in localizing electrons in highly correlated systems. Several schemes have been proposed, in order to overcome these drawbacks. However, their computational cost is usually very high and, at present, they cannot be used routinely in surface physics. Despite this fact, we will now present them, because there is no doubt that an improved understanding of oxide surface properties will be gained thanks to them, in the coming years. 2.4-1' The Self-Interaction Correction (SIC) method In the mean-field approximation for electron-electron interactions, the Hartree U[n] and exchange Ex energies read: rrr ^ f i'^ A ,nir)n{r') U\n] = d^rd^r'^-^—^ ^^ J 2\r-r'\ Ex = - Y: /dVdV^-^^^-^^^^'^^^^^^^^^^^"!j^^^^^^^^^

(2)

41

Ex involves a sum over orbital (a, a') and spin (cr) quantum numbers, of a product of one-particle orbitals "ipao- with occupation numbers faa obeying Fermi statistics. U[n] includes a spurious Coulomb self-interaction, as can be realized if Eq. (2) is applied to a single electron system. This Coulomb self-interaction is exactly cancelled out by the a = a^ terms in Ex (selfexchange term). In the DFT-LSDA approach, the local approximation to the exchange functional achieves only partial cancellation of the self-Coulomb term. Only for orbitals which are delocalized over the whole system does this self-interaction vanish. In the general case, for finite systems and localized states in extended systems, it leads to systematic errors, which have been summarized in Ref. 25. This work proposes a method for SIC, in which the LSDA exchange-correlation energy functional £'^^^^[n>|^,nj, which depends upon the electron densities for t and I spins, is replaced by: Ell^ = E ^ f A ^ , „ j _ Y: U[n,,,] + El!^^[n,,,,

0]

(3)

a,cr

Modified Kohn-Sham equations result which contain an orbital-dependent local potential. The correction usually leads to a greater localization of the electronic wave functions. It was shown that SIC heals most of the LSDA drawbacks, including a large part of the gap problem in insulators and it corrects the bad shape of the exchange hole around electrons. In addition, in atoms, the negative of the SIC eigen-energies closely approximates the relaxed excitation energies for electron removal. 24.2. The GW method The calculation of a quasi-particle (QP) spectrum in a crystal requires the resolution of an equation of the type: - ^ + 14xt(r) + VH(r-) (4) in which E^j^ and (t>^%{'f^ are the quasi-particle energy and wave function in a band n at a given k point in the Brillouin zone. Kxt(r) i^ ^^^ ionic potential, V]i{r) is the Hartree potential and E is the self-energy operator, which is in general non local, non-Hermitian and energy-dependent. It contains the exchange and correlation effects and is state [nk) dependent. By comparison, in the Kohn Sham equations, E is replaced by T4c(r)^(^"" f^), and the resulting eigen-values fail in reproducing quasiparticle energies in solids, as has been stressed by many authors [26-28]. The estimation of E(r, r';£^^^) is a very difficult task. A possible approximation is the GW one, [29,30] in which a perturbation expansion for

42

the self-energy is constructed and stopped at the first order: E(f, P; Lj) = i /_"^J ê+i^^'

G{r, ?• u + J) W{r, P; a;')

(5)

In Eq. (5), G is the one-particle Green's function, W is the screened Coulomb interaction and (5 = 0"^. The real part of the self-energy contains a screened exchange contribution, which requires an explicit calculation of the dielectric matrix of the system, and a Coulomb-hole term which takes into account the actual presence of the quasi-particle (excess electron or hole) in the system and its screening by the surrounding electrons. Although many improvements have been proposed for the calculation of the dielectric response function from first principles [31,32], still this stage is both computer time and memory very demanding when a large basis set is used for the description of the electronic structure. Efficient GW methods have thus been developed, in which a model dielectric function is used to mimick the screening properties of the system under study [28,33,34]. It should be noted that the resolution of Eq. (4) may be performed self-consistently or in a perturbative way with respect to E(r, r';£J^^) — 14c(^)- In the second case, HF as well as DFT-LDA eigen-states and eigenenergies may be used as starting points for the implementation of the GW approximation. 24.3. The LDA-hU method The LDA+U method assumes an approximate non-local and frequency independent form for T,{r^r^; ^nk)^ which makes it a HF-type theory with a screened exchange. The idea consists in describing delocalized s — p electrons by an LDA-type approach, and in adding a Hubbard term ^C/ E [ui the occupancy operator of the di orbital) to the d- or /-electron Hamiltonian [35-37]. This procedure stabilizes the occupied LDA bands, and destabilizes the unoccupied ones. This jump in energy between occupied and unoccupied states, which exists in an exact DFT theory, as well as the correlated jump of potential when the number of electrons goes through an integer value [38], are absent in the LDA, which thus misses an important contribution to the gap value. Taking into account exchange interactions (J) and the non-sphericity of electron-electron interactions (dependence of U and J upon the orbitals i), the orbital-dependent potential reads (C/eff = U — J/2):

43

Fig. 1: Domains of the projected atomic structure of the Q;-Al203-(\/3l x \/3l)R±9° reconstruction, where the unit cells as well as domain walls are drawn. The two constituting Al planes are shown separately, with evidence of one being much better ordered than the other. Numerical relaxation has shown that the ordered layer could be associated to the outermost layer and the more disordered one to the layer adjacent to the substrate (from Ref. 43).

Via{r)

= VLDA{r) +

E{UiJ-Ues)n^J-a 3

+ YliUij

- Jij - Ues)nja

+ C/effCo ~ ^*>) ~~ 7*^

(^)

The LDA+U theory may be regarded as an approximate GW method [37]. The screened Coulomb and exchange parameters U and J are usually estimated in a supercell approximation [39]. However, there is some arbitrariness in the choice of the localized orbitals when performing the partitioning of the Hamiltonian. A further step in the improvement of LDA+U consists in adding dynamical effects — frequency dependence in S(r, r';a;). This may be performed using a DMFT-type approach (DMFT= Dynamical Mean Field Theory) [40] as part of the so-called LDA++ approaches [41]. 2.5. General features for simulating surfaces Whatever the choice of the quantum mechanical method, two representations are currently used to simulate semi-infinite surfaces: slabs or clusters. Only the CI method is restricted to cluster geometries. In slab calculations, a finite number of layers mimicks the semi-infinite system, with a two-dimensional (2D) translational periodicity. A minimal thickness dmin is required, so that the layers in the slab centre display bulk characteristics. Practically speaking, dmin should be at least equal to twice the damping length of surface relaxation effects, which depend upon the surface orientation. In plane wave codes, the slab is periodically repeated

44

in the direction perpendicular to the surface, in order to make finite the required number of plane waves. The procedure generates spurious electric fields and energy contributions: in order to minimize them, one has to introduce thick enough vacuum regions between the repeated slabs and has to construct symmetric slabs with a zero total dipole moment. This question is particularly crucial for polar orientations. Cluster calculations seek to determine surface properties by simulating only the local active sites [42]. They are usually implemented when it is believed that local bonding prevails in the electronic structure, when the property under study is local by essence, or when one wishes to model excited states through CI methods. It has been realized that a correct modelling requires the effect of the electrostatic potential created by the ions outside the cluster: the clusters are thus embedded in an array of point charges and/or in effective core potentials to describe atoms at the fringe of the clusters. A good embedding is more and more necessary as the cluster size decreases. Treating non-stoichiometric surfaces requires special care in several respects. First, the long range elastic deformations (surface stress) induced by the defects are not easy to account for in a realistic way, neither in periodic calculations, nor in cluster models. When isolated vacancies are simulated in slab calculations, the size of the 2D unit cell, which is, most of the time, fixed by computational rather than physical constraints, corresponds to vacancy densities by far larger than realistic concentrations. At present, large size reconstructions, such as the (\/3T x \/31)R9^ on the a-Al2O3(0001) surface (Fig. 1) are still beyond the possibilities of quantum simulations. The presence of oxygen vacancies is usually associated to a strong redistribution of electrons, some of them being trapped at the vacancy site. When using localized basis sets to expand the electronic wave functions, it is mandatory either to place virtual orbitals at the location of the missing oxygen or to use strongly polarized orbitals on the neighbouring cations, in order to span the whole space for the electron density (Section 5.2). Finally, in order to account for strong electron redistributions, selfconsistent approaches which properly solve the Poisson's equation have to be used. The same is true for polar surfaces, whose stability relies on a well-defined electrostatic condition. 3. REVIEW OF LITTERATURE This section summarizes the present stage of our knowledge on clean oxide surfaces, obtained from numerical works. It is presented according to the oxide bulk crystal structures, with special emphasis on those oxides which have been more thoroughly studied.

45

3.1. Fluorite MO2 and anti-fluorite M2O structures The anti-fluorite structure consists in an fee oxygen lattice, in which every 0 - 0 pair is bridged by two metal ions, of charge 4-1. It is found in highly ionic oxides [44], like alkali-oxides Li20, Na20 and CS2O. The fluorite structure is similar, with a mere exchange of oxygens and cations, as in Zr02, Ce02, UO2, etc. The non-polar surfaces of lowest Miller indices are the (111) and (110) surfaces. The former is more dense than the latter, with a single broken bond per formula unit, to be compared to 2 on the (110). The (100) surface is polar. Various works have shown that the (111) surface has always a lower energy than the (110), whether Li20 [45-47], Ce02 [48,49], UO2 [50], or Zr02 [45] are concerned. Relaxation eflFects depend upon the oxide: on 1102(111), inward atomic displacements take place, while on the (110) surface, there are small vertical atomic displacements associated to large inplane displacements [50]; on tetragonal ZrO2(001), oxygens move inwards and cations outwards [51], while the (110) and (101) surfaces display some rumpling, with no preferred outward displacements of the surface oxygens [52]. The relaxation strength is an increasing function of the unrelaxed surface energy. The modifications of electronic structure at the Li20 surfaces include a shift of the oxygen Is core levels, and a narrowing of the oxygen-derived valence band (VB) width, especially on the (110) surface. No surface states are found in the bulk band gap [46,47]. In an effort to interpret STM images (STM= Scanning Tunneling Microscopy) of U 0 2 ( l l l ) , electronic charge density maps were obtained, using a DFT-LDA-i-U approach, and confirm the high degree of ionicity of the oxide [53]. Noticeable charge redistributions take place at the tetragonal (001) Zr02 surface [51]. 3.2. Rocksalt structure The rocksalt structure consists in two interpenetrating fee lattices of anions and cations, in which all atoms are in an octahedral environment. It is met in alkaline-earth oxides (MgO, CaO, SrO, BaO) and in some transition metal oxides like TiO, VO, MnO, FeO, CoO, NiO, etc, with cations in a +2 oxidation state. The non-polar surfaces of lowest Miller indices are the (100) and (110) surfaces: they have neutral layers, with a^ many cations as oxygen ions, and their outermost atoms are 5- and 4-fold coordinated, respectively. Actually, planar surfaces can only be produced along the (100) orientation. The polar direction of lowest indices is (111): it has an hexagonal 2D unit cell, three-fold coordinated surface atoms and equidistant layers of either metal or oxygen composition. As regards the (100) surface, all theoretical works predict very small atomic displacements, on MgO [45,54-69], as well as on CaO, SrO, BaO

46 0.10 0.08 0.06 0.04 0.(»

1«llay«rofMnpty| •bovt 0 sltet ml

îiiwHwibw [t

4.0

a«^

I' turfac* layer

"^2.0

1 0.0 I 4.0

1 St •ubturtac* lay«r

» 2.0 UJ H 0.0

"I

•"I

»W|W"I

|

»»l»i

a 2.0

CM

O 0.0

N I ¥* d excitation energies in NiO (001) monolayer {Z = 4), NiO (001)/MgO (001) bilayers {Z = 5) and MgO (001)/NiO (001)/MgO (001) trilayers (Z = 6), as a function of the Ni coordination. Lines are drawn to guide the eye (From Ref. 84).

are never simultaneously present in the system. In order to interpret optical absorption experiments or EELS measurements, excitonic effects, have to be explicitely taken into account. The gap value obtained in this way will be denoted G*. Embedded-cluster CI techniques may be used to simulate charge transfer excitations. This was done in a study of planar [253] or defective [254,255] MgO surfaces. The results support the argument of decreasing G* as Z decreases. They are consistent with the arguments of Madelung potentials developed above. Total energy difference techniques also yields values of G*. This procedure was followed in order to determine charge transfer gaps in various geometries of the NiO(100) monolayer, thanks to symmetry breaking tricks [84]. An increase of G* with the coordination number of the Ni ions was found in a series including the unsupported NiO monolayer, a NiO/MgO bilayer and a MgO/NiO/MgO trilayer, which again may be rationalized with Madelung potential arguments. On the other hand, a variation of G* with an opposite slope is found between the monolayer and bulk NiO {G* =5.1 eV and 4 eV, respectively), whose origin remains at present unclear. 5. NON-STOICHIOMETRIC SURFACES Non-stoichiometry in oxides is not a rare phenomenon, and especially at surfaces, where the process of annealing in ultra high vacuum after surface cleaning, induces a desorption of oxygen. Depending upon the annealing temperature and time, surfaces with various oxygen contents may be produced, whose exact stoichiometry is not easily quantified. However, some bench marks can be found when the vacancies order and give rise

70

to reconstructed structures, which allows to prepare surfaces in a roughly reproducible manner. There is a peculiarity of oxide surfaces, which is not yet understood, that no intrinsic reconstruction has ever been observed, at variance with metals or semi-conductors. In this section, we will discuss energetic, structural and electronic signatures of isolated or ordered oxygen vacancies at surfaces. We will only consider the case of neutral oxygen vacancies, in which neutral atoms are removed from the surface and leave two electrons behind them. 5.1. Atomic relaxation Structural relaxations are induced around vacancies by the breaking of oxygen-cation bonds. The importance and sign of the atomic displacements depends upon the crystal structure, the surface orientation, and more generally upon the coordination number of the missing oxygen. The general trend is a displacement of the neighbouring cations away from the vacancy, leading to a contraction of the remaining cation-oxygen bonds. This effect is qualitatively similar to relaxations on stoichiometric surfaces, as discussed in Section 4.1. It relies on the property of interatomic distances to decrease when the coordination number Z of the partner atoms is reduced. When a vacancy is created on an MgO(lOO) surface, for example, the neighbouring cations, which are five-fold coordinated on the perfect surface, become four-fold coordinated, and their displacement away from the vacancy shortens some surface Mg-0 bonds [69,91,95]. This effect is also found on SrTiO3(100) [210] and on TiO2(110) [143,249]. Its strength increases close to surface defect sites — step edges, kinks —, whose Z is smaller. However, there are exceptions to this rule, when a cation-cation bond across the vacancy site may be formed. This is the case in bulk Si02, where the creation of an oxygen vacancy breaks two 0-Si bonds. The atomic displacements of the two neighbouring silicon atoms are directed inwards in such a way that the Si-Si distance across the vacancy site is reduced by about 0.5 A with respect to its value in the perfect oxide [256]. 5.2. Charge and spin distribution The spatial distribution of the two electrons associated to a neutral oxygen vacancy depends upon the degree of covalency of the oxygen-cation bond, and on topological factors: the oxide crystal structure and the coordination of the vacancy site. In some cases, as in MgO, the electrons are trapped in the vacancy site by the strong electrostatic potential exerted by the neighbouring cations, thus forming a surface F centre (denoted Fg centre). In more covalent oxides, Ti02 and SrTiOs for example [89,210], the electron redistribution is more diffuse and reaches the neighbouring

71 d2

43

Fig. 13: Characteristics of the saddle point configuration in the diflFusion path of an oxygen vacancy on MgO(lOO) (top view), (a) atomic configuration; (b) iso-density surface for the vacancy gap state, plotted at about 15% of the maximum state density. The magnesiums and oxygens are shown as large and small circles, respectively (from Ref. 69).

cations and sometimes the second neighbour oxygens. In all cases, the electron density is asymmetric with a larger contribution above the surface plane than below. Similar distortions are also found when the vacancies are located at surface defect sites. However, the precise localization of the excess electrons is not straightforwardly found in numerical simulations, and contradictory predictions have been made as a result of different choices in the internal parameters of the computation. One of them concerns the basis set on which the wave functions are expanded. According to whether it consists of plane waves or localized functions, and, in this last case, according to whether or not virtual orbitals are introduced at the site of the vacancy, space is thoroughly spanned or not and electrons can or cannot redistribute freely in order to minimize energy. The absence of orbitals at the vacancy site, for example, prevents a correct description of an Fg centre et constrains the excess electrons to redistribute on the neighbouring cations [92]. To examplify the difficulty, we show on Fig. 13 the excess electron distribution during the migration of an oxygen vacancy in the MgO(lOO) surface plane: when the oxygen atom reaches the saddle point between two lattice positions, the excess electrons are equally shared between these two sites [69]. It is quite clear that describing this distribution with a localized basis set requires an extreme care. In addition, when localized basis sets are used, it is important to introduce polarization orbitals on the neighbouring cations, to account for their response to the strong electrostatic fields created by their highly asymmetric environment. Part of the differences found between non-stoichiometric Sn02 and Ti02 surfaces was for example

72

assigned to the larger polarisability of the former [166]. The MuUiken charges are also extremely sensitive to the choice of the basis set. This is due to the fact that bond charges, which are equally shared between the two partner atoms, result from the overlap between the neighbouring orbitals [98]. To my knowledge, no Bader analysis — which is basis set independent — ha.s been performed on non-stoichiometric surfaces. In any Ccuse, a consensus now exists that the direct view of charge density maps give more reliable — although not quantitative — information than a MuUiken charge analysis on these surfaces. The electron distribution also depends sensitively upon the treatment of exchange and correlations. For example, a comparison of DFT-LSDA and UHF simulations of non-stoichiometric bulk Ti02 shows that the electron density is larger in the vacancy site in UHF than in LSDA [161]. Similarly, on the reconstructed (2 x 1) TiO2(110) surface, a symmetry breaking induced by spin degrees of freedom — and consequently absent in DFTLDA calculations — takes place in which the excess charge is gathered on a single titanium, rather than equally shared between two titaniums. It was argued that the complete filling of spin-unpaired orbitals is more effective in reducing on-site repulsion — which is a dominant factor in narrow d band compounds — than the partial occupancy of spin-paired orbitals [161,249]. 5.3. Spectroscopic signature Spectroscopic experiments have revealed several specific signatures of the presence of oxygen vacancies, in the electronic spectra. The two electrons left by the missing oxygen fill one additional electronic level, usually located in the the gap of the oxide and which pins the Fermi level. The density of states is distorted, with a transfer of weight from non-bonding states at the top of the VB to bonding states at the bottom of the VB. Finally, core level shifts on the neighbouring cation(s) take place. The numerical studies, quoted in Section 3, support these observations. Due to the difficulties in reproducing correctly gap widths by standard ab initio methods, the position of the vacancy level e^ in the band gap is only given reliably through its distance in energy to the top of the VB. One should note that several factors affect the value of Cy. First, a good prediction requires a correct account of self-consistency effects, because the electron redistribution induces non-negligible electrostatic potentials. All the effects mentioned above — basis set limitations, choice of the exchangecorrelation functional, etc —, which may change the electron localization, can also shift Cy, since energies and wave functions are simultaneously determined in the self-consistent procedure. In addition, by modifying the Madelung potential at the vacancy site, relaxation effects around a vacancy

73

5r

0.0

Fig. 14: Band structure of a fully oxygen defective ( 1 x 1 ) MgO(lOO) surface along the three symmetry lines J-F-M of the 2D Brillouin Zone, as obtained through the FP-LMTO calculation (Full Potential- Linear Muffin-Tin Orbital method). The dashed horizontal line represents the Fermi level, black dots (star) indicate the energy positions of the filled (empty) Bloch states at F calculated in a (2v^ x 2\/2) supercell. The dashed line in the gap of the projected bulk bandstructure gives the dispersion of the F^ centre band. The dashed-dotted line is used for the surface conduction band of lowest energy (from Ref. 69).

can move Cy, This is especially important when the vacancy is located at step edges or kinks, where large atomic displacements take place. It was shown that optical transitions associated to F centres on sites of low coordination of MgO are red shifted with respect to their bulk value, due to a reduced Pauh repulsion in the excited state [96]. The nature of the gap states may change as a function of the vacancy concentration c. On MgO (100), at low values of c, the Fg levels broaden into a band, which is completely filled. The band width was interpreted within a simple tight-binding model, involving effective hopping processes between F^ orbitals on neighbouring vacancies [69]. Above a critical concentration, this band overlaps the surface conduction band, and the surface becomes metallic (Fig. 14). Two types of quantum states, available for the excess electrons, thus coexist on the surface: the Fg orbitals close to the vacancy sites, and surface CB states localized on the surface magnesiums. On rutile Ti02 (HO) surfaces with large vacancy concentrations, it was shown that the metaUic state is less stable than an insulating spin polarized state [249]. 5.4. Energetics The energetics of non-stoichiometric surfaces can be characterized in two different ways. One consists in defining the vacancy formation energy Eyf. It is the energy required to extract one neutral oxygen atom from the

74

surface and to send it at infinite distance into vacuum, either as an isolated atom or as part of an oxygen molecule. In the first case, for example, Eyf is equal to: ^vf =

Est - Enan-st "

-j^

NE{0)

(13)

with Est^ Enon-st ctud £"(0) the energies of the stoichiometric surface, the non-stoichiometric surface with N vacancies, and the neutral oxygen atom, respectively. Most of the simulations quoted in Section 3 give values of Eyf^ which allows to discuss its variations as a function of the oxide, the surface orientation and the coordination Z of the vacancy site at the surface. For a given oxide, E^f systematically decreases as Z gets smaller, i.e. as the number of broken oxygen-cation bonds gets smaller. This is checked, for example, for various vacancy sites on the MgO(lOO) surface [69,91,95] , and on inequivalent oxygen sites on the ¥205(010) surface [229]. Another approach consists in comparing the energetics of surfaces with different oxygen contents, to understand which configuration is the most stable under given experimental conditions. The surface is then assumed to be in contact with an external reservoir. The relevant thermodynamical potential is no longer the internal energy E of the surface, but rather its grand potential Vt\ n = E + PV-TS-Y:fîNi

(14)

i

For typical pressures P and temperatures T, the PV and —TS terms can usually be neglected, f] is a function of the chemical potentials //j of the different species i, which are not independent variables. For a binary system XnOm, for example, the surface is in equilibrium both with the bulk oxide and with the outer oxygen atmosphere. The first condition yields the relationship niix + mjio = I^XnOm between the cation and oxygen chemical potentials and the bulk energy per formula unit /J^XnOm • Knowing the internal energy for a given surface termination, Q can be calculated, according to Eq. (14), for the whole range of accessible values of /JLOStability of non-stoichiometric Al2O3(0001) [127,128], Fe2O3(0001) [122] and SrTi03(lll) and (110) [220] surfaces has been discussed in this context, and a typical Q — f{fio) graph is shown in Fig. 4. The nature and strength of vacancy-vacancy interactions have scarcely been investigated, although they drive the thermodynamics of vacancy ordering on defective surfaces, and are thus key quantities to understand the numerous reconstructions observed on surfaces annealed in vacuum [1,2]. However, the ordering process may be inhibited by kinetic effects, relying on parameters, such as the activation energy for vacancy diffusion and the temperature. On MgO(lOO), vacancies have been found to weakly repell

75

each other, in a non-monotonic way as a function of their distance [69]. The overall repulsion is typical of an hybridization process — between two Fg orbitals — in which both bonding and anti-bonding states are filled. The non-monotonic behaviour is a consequence of the surface topology, the second and third neighbour interactions being differently mediated, by magnesiums and oxygens, respectively. On this surface, no strong tendency towards order could be found, in agreement with experimental observation. On SrTiO3(100), the same problem was addressed [94], but the interactions are of a different nature due to the more diffuse character of the electron redistribution. An interesting point concerns the comparison between the Ti02 and SrO terminations. On SrO, the excess electrons are shared between the vacancy site in the outermost layer and the titanium located in the sub-surface layer, thus producing a vertical electric dipole. On the Ti02 termination, the excess electrons are spread over the vacancy site and its two titanium neighbours in the surface layer [199], resulting in an horizontal quadrupolar distribution. This difference manifests itself in the sign of vacancy-vacancy interactions on the two terminations, which are found to be repulsive on SrO, and either repulsive or attractive on Ti02, depending upon the relative orientation of the two quadrupoles under consideration [94]. Finally, activation barriers have been determined for the diffusion of a vacancy in the MgO(lOO) surface, between sub-surface and surface layers, and close to surface defects [69,95]. As expected, the barrier is nearly twice lower in the first case than in the second one. Moreover, it is of the order of .20 to .25 times the formation energy, a quite reasonable ratio, in view of similar results on metal surfaces [257]. 6. POLAR SURFACES Polarity represents a peculiarity of some surface orientations in compounds involving simultaneously atoms of different electronegativity. It has been thoroughly studied in semi-conductor compounds [236,258]. A polar orientation is such that each repeat unit in the direction perpendicular to the surface bears a non-zero dipole moment. An electrostatic instability results from the presence of this macroscopic dipole, which can only be cancelled by the presence of compensating charges in the outer planes. This can be achieved either by a deep modification of the surface electronic structure — total or partial filling of surface states, sometimes leading to surface metallization— or by strong changes in the surface stoichiometry — spontaneous desorption of atoms, faceting, large cell reconstructions due to the ordering of surface vacancies, etc. Polar oxide surfaces present a large diversity, compared to sp^ semi-conductors. They exhibit a vast

76

number of crystallographic structures — rocksalt, corundum, spinel, inverse spinel, wurtzite, perovskite, for the simplest ones — which reflect the subtle mixing of ionicity and covalency in the metal-oxygen bonding and the specificities of the d electrons in transition metal oxides. In addition, mixed valence compounds, such as magnetite Fe304, can form when metal atoms with several oxidation states are involved, and, playing with experimental parameters, such as temperature, partial oxygen pressure, etc., allows to stabilize oxides of different stoichiometrics. It is clear that these peculiarities demand a generalization, if not a total reconsideration, of some theoretical concepts, a work which is still in its first stages [4]. In this section, we will discuss successively classical electrostatic arguments, some models of electronic structure which are useful in this context and finally the various processes which can stabilise a polar surface, including non-stoichiometric reconstructions. 6.1. Criterion for surface polarity According to classical electrostatic criteria, the stability of a compound surface depends on the characteristics of the charge distribution in the structural unit which repeats itself in the direction perpendicular to the surface [259]. Type 1 or 2 surfaces — with neutral or charged layers, respectively — have a zero dipole moment fl in their repeat unit and are thus potentially stable. At variance, polar type 3 surfaces have a diverging electrostatic surface energy [181] due to the presence of a non-zero dipole moment not only on the outer layers (which would not distinguish them from non-polar rumpled or reconstructed surfaces), but on all the repeat units throughout the material. The total dipole moment is thus proportional to the number of repeat units and so is the electrostatic contribution to the surface energy per unit area (Fig. 15). This is the origin of the surface instability. The estimation of the dipole moment in the repeat unit is not always an easy matter. In binary compounds, the difference in electronegativity of the constituents readily points out the existence and sign of the charge transfer between the ions. In simple crystal structures and for orientations such that layers containing cations only and anions only alternate, the presence of a dipole moment in the repeat unit is thus unquestionable, whatever the charge values — provided they are non-zero. These surfaces are undoubtly polar surfaces. The same is also true for some surfaces of ternary compounds, such as the (110) and (111) faces of ABO3 perovskites. For example, the (110) and (111) repeat units contain alternating ABO and O2 layers in the first case, and alternating AO3 and B layers in the second one, which are undoubtly charged, because, in each case, one of the layers contains a single constituent.

77 -a 0 -a a -CG

4iio

(a)

l.4n(5R.

-GO-GO

-GG 4710 ^7 4jto_^2?j

(b) -47ia ^

J_4-

/?j+/?2

Fig. 15: Spatial variations of the electrostatic field E and potential F in a slab cut along a polar direction.

There are less obvious cases, among which is the (100) perovskite sur~ face. SrTiO3(100) , for example, presents alternating layers of SrO and Ti02 composition. If formal charges are assigned to the ions (Sr^"^, Ti'*"^ and 0^~), each layer is charge neutral, the repeat unit bears no dipole moment, and the orientation is considered as non-polar. This is the statement most often encountered in the literature. However, SrTiOs is not fully ionic. Its gap width, equal to 3 eV, places it on the border line between semiconductors and insulators and the Ti-0 bond presents a nonnegligible part of covalent character. The actual charges are thus likely not equal to the formal ones and Qsr + Qo and Qxi + 2Qo are not likely to vanish. SrTiO3(100) should thus be considered as a polar surface and this examplifies how careful one should be in the classification of surfaces. In addition, the surface orientation h is not always sufficient to fully characterize a semi-infinite system, especially when various terminations may be produced. In the rutile structure, for example, the bulk repeat unit in the (110) direction is made of three layers of O and (M0)2 composition, and there exist three chemically inequivalent terminations, which expose a single oxygen layer ( 0 / ( M 0 ) 2 / 0 sequence), two oxygen layers ( 0 / 0 / ( M 0 ) 2 sequence) or one mixed cation-oxygen layer ( ( M 0 ) 2 / 0 / 0 sequence). Only in the first case, the repeat unit bears no dipole mo-

78

merit. Similarly, on the basal (0001) face of the corundum structure, three chemically distinct terminations may be produced (Section 3.3). Only the M/O3/M termination is non-polar. According to classical electrostatics, polar surfaces are thus unstable. However, it can be shown [2,94] that the macroscopic dipole moment can be cancelled out by modifications of the charge densities on the outer layers. Assuming that m outer layers have a charge density CTJ different from the bulk {\aj\ 7«^ cr for 1 < j < m) and that the layer spacing is alternatively equal to Ri and i?27 the electrostatic condition for surface stability reads:

E ^j =

^m+l

— 5 -

(_!)"» _ -R2 - -Ri i?2 + -Rl

(15)

Several scenarios, in which either the charge or the stoichiometry of the layers is changed can thus lead to polarity healing: • one or several surface layers have their composition which differs from the bulk stoichiometry. Reconstruction or terracing will result, depending upon how the vacancies or adatoms order. If no order takes place, no information on the surface stoichiometry can be extracted from surface diffraction patterns, unless a quantitative analysis is performed. • foreign atoms or ions, coming from the residual atmosphere in the experimental set-up, provide the charge compensation. • on stoichiometric surfaces, charge compensation may be provided by the electron redistribution which takes place in response to the polar electrostatic field. This is well examplified in self-consistent electronic structure calculations. Which process actually takes place depends upon energetic as well as kinetic considerations. As will become clear in the following, if stoichiometric ideal polar surfaces are not observed, this is never because their surface energy diverges. There always exist enough electronic degrees of freedom in a material to reach charge compensation through the third mechanism. However, in most cases, the resulting surface energy is high and other processes may be more efficient. 6.2. Models of electronic structure in iono-covalent application to polar surfaces The characteristics of the charge distribution in the bulk and layers are thus key factors to understand the physics of polar zeroth-order description of the electronic structure of oxides

materials: in the outer surfaces. A is the fully

79

ionic model, in which formal charges are assigned to bulk as well as to surface atoms. This model successfully assesses whether, in a binary oxide, a bulk repeat unit bears a non-zero dipole moment, and can predict the stable surface terminations in some complex structures. It applies well to highly ionic materials, but cannot be justified in more covalent compounds or when surface charge redistributions take place. For example, according to it, the surface energy of stoichiometric SrTiOs and BaTi03(lll) and (110) surfaces should be infinite, as well as those of the oxygen or iron bilayer terminations of a-Fe203 (0001), in contradiction with both experiment or ab initio calculations. In addition, it may even give wrong predictions on the nature (polar or not) of a surface, in ternary compounds, as in the case ofSrTiO3(100) In tetrahedral semi-conductor compounds, on the other hand, the most widely used criterion for the stability of polar surfaces relies on specific fillings of dangling bonds — sps atomic orbitals which remain unpaired due to surface bond breaking. The criterion is referred to as the autocompensating model [260,261]. It states that a surface is stable if anionderived and cation-derived dangling bonds are occupied and empty, respectively, because such a filling induces a net lowering of the surface energy. According to electron counting rules, it has been realized that, at polar surfaces, the condition of auto-compensation implies that Eq. (15) is fulfilled, and the surfaces are then said to be "charge neutral" [262]. Later, the auto-compensation principle was generalized to describe mineral surfaces [263]. assuming that the electron number per bond is the same at the surface and in the bulk (but not necessarily equal to 2 as in tetraedral semiconductors). This principle has been applied for example to a-Al2O3(0001) [110], Fe3O4(100) [264-267] and a-Fe2O3(0001) [268,269]. This generalization may be useful in a number of cases. However, the vocabulary taken from the physics of semiconductors is highly unsuited to oxides, because, in most cases, dangling bonds have no physical significance. The conservation of the number of electrons per bond between bulk and surface is also not a well founded assumption. The description of the charge distribution in terms of electron transfer per bond [244] given in Section 4.3, allows to make a bridge between the two limits and can account for the metal-oxygen bonding in the whole range of ionicities. We will use it in the following to understand the mechanisms at work on polar surfaces. It has to be noted that, for this purpose, it will never be necessary to know quantitatively the values of the parameters AoiCj, neither in the bulk nor at surfaces. It is the principle of electron sharing per bond which is the key ingredient.

80

6.3. Surface processes relevant for polarity healing We will successively discuss surface relaxation effects, changes of covalency in the outer layers, partial fillings of surface states and stoichiometry changes. It is often thought that surface relaxation is one of the most efficient processes to stabilize polar surfaces, especially open surfaces, which experience large atomic displacements. Actually, two aspects of the problem should be distinguished. First, as far as the electrostatic criterion Eq. (15) is concerned, the macroscopic component of the dipole moment is entirely determined by the charges and layer spacings in the bulk repeat unit. Consequently, in the vicinity of the surface, a mere contraction or expansion of the interlayer spacings, not accompanied by a change in the charge densities, can never heal polarity However, once charge compensation is achieved — by whatever mechanism —, surface relaxation induces a lowering of surface energy, as it does on non-polar surfaces, and the strength of the effect increases as the coordination of the surface atoms gets lower [60] (c.f Section 4.1), as exemplified on the a-Al203 (0001) surface (Section 3.3), and on the reconstructed (2 x 1) NiO(lll) surface [270]. Since charge compensation requires a modification of the charge density, changes of covalency at the surface are often assumed to heal the polarity. With the help of the bond transfer model, one can show that this statement is incorrect, as far as semi-infinite polar surfaces are concerned. It is useful to make a distinction between weakly polar surfaces, in which the dipole moment in the repeat unit is entirely due to covalent effects, and truely polar surfaces whose dipole moment contains an integer contribution. As already said, in the fully ionic limit, the first ones are considered as nonpolar, while the second ones are recognized as polar. A prototype of the first family is SrTiO3(100). Applying the Bond Transfer Model [244], it is found that the charge densities per two-dimensional unit cell, on the Ti02 and SrO bulk layers, read (JB = IQxi + 2(5o| = QST + Qo — 2ATi-o — 8Asr-o, in terms of the electron transfers Axi-o and Asr-o per Ti-0 and SrO bonds. The dipole moment in the bulk repeat unit is thus non-zero. However, it has no integer contribution and depends only on covalent effects. On the (100) surface, due to the change in local environment, due to possible structural distortions and due to shifts of atomic levels, redistributions of charge take place (Fig. 16). On the Ti02 termination, for instance, assuming modified Afpi_o and Agj._o values yields atomic charges which differ from the bulk in two layers. However, it is found that a\ + a2 = Axi-o — 4Asr-o, which fulfills the condition ^1 + cr2 — —crB/2, Eq. (15), whatever the specific values of the A and A' parameters. On SrTiO3(100), the bond breaking mechanism, by itself.

81 Ti02

-fô

«-^-»

MgO(lll)

SrTiO3(001)

0 =oxygen

000000 0

• =titanium

# =strontium or magnesium

Fig. 16: Electron transfers per bonds introduced in the modelization of an SrTiO3(001) surface (left pannel) and of an MgO(lll) surface (right pannel). On SrTiO3(001), specific transfers Aî-o înd Asi._o (represented by thick arrows) are introduced inside the surface layer and in between the surface and sub-surface layers for T i - 0 bonds and Sr-0 bonds involving surface atoms. On M g O ( l l l ) , only the first inter-plane transfer is assumed to be modified, (from Ref. 4).

thus yields the charge redistribution which suppresses the divergent part of the electrostatic potential. The surface charges are different from the bulk, both because the bond covalency in the surface layers is different from the bulk (A' ^ A) and because the coordination of surface atoms is reduced with respect to the bulk. However, the model shows that only the second factor is effective for polarity healing, while the change of covalency plays no role in the charge compensation process. On the other hand, on truely polar surfaces, in which the dipole moment borne by the bulk repeat unit contains an integer contribution a change in covalency in the surface layers cannot either heal polarity, MgO(lll) is representative of this family. The charge density per ( 1 x 1 ) unit cell on bulk (111) layers is equal to GB — ±(2 — 6A). On a stoichiometric (111) magnesium termination, surface atoms are three-fold coordinated to atoms located in the underlying layers. Their charges QMgi are different from those in the bulk. But the same is also true for oxygen atoms in the sub-surface layer (n = 2) because A' 7^^ A (Fig. 16). (^Mgi and QQ^ are such that (Ji -h (J2 = 3A. This does not fulfill the electrostatic criterion, which, instead, requires that ai + a2 = 3A — 1. The surface layers have a deficit of one electron, which is independent of the values of A' and A. It should be noted that the same result is obtained in a fully ionic picture (A' = A = 0) or when neglecting the change of covalency (A' = A). Neither in this family, thus, can charge compensation be achieved by change of covalency at the surface. What we have demonstrated here, is very general: a mere change of

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(b)

I

E p

vty^ o.o

Energy ( e V )

-4.0

-

O.O Energy ( e V )

Fig. 17: Calculated DOS for five-layer symmetric MgO(lll) slabs, with oxygen (a) and magnesium (b) terminations. Plain and dashed lines represent the DOS in the slab and in the bulk, respectively. The top and lower panels refer to projections on majority and minority spin states, respectively. All DOS have been convoluted by a 0.1 eV wide Gaussian function (from Ref. 103).

covalency in the outer layer can never provide the compensating charges, because it concerns several layers whose contributions cancel out in the expression of the electrostatic criterion. On these stoichiometric polar surfaces, charge compensation can only be achieved by a partial/total filling of CB surface states or depletion of VB surface state. The Bond Transfer Model helps understanding that this filling / can be determined by ionic arguments, even in very covalent materials, and that it is independent of the specificities of hybridization in the outer layers [244]. In the case of MgO(lll), for example, folloviîng the charge analysis just given above, charge compensation is only achieved if / = 1/2. The analysis of the oxygen termination leads to the symmetric conclusion that a surface valence band has to be half-filled. It should be noted that the condition / = 1/2 is independent of the precise values of the electron transfers per bond. It would have been similar using a fully ionic picture. The result that we have obtianed is not restricted to ionic oxides such as MgO. On ZnO(OOOl) or SrTi03(lll) or (110), the occupancy of surface states is also dictated by the values of the formal charges and of the interlayer spacings Ri and i?2, despite the fact that these oxides are rather covalent. The ionic limit can thus be currently used to estimate the value of / required for charge compensation. Two cases occur. In the first case, / is non-integer. The surface states are only partly filled or empty and the surface layers have a "metallic" character ^ ^We use here the term "metallic" for the sake of simplicity, meaning that electronic excitations of zero energy can be

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(Fig. 17). This occurs on the rocksalt oxide (111) surfaces (/ = 1/2), the ZnO(OOOl) surface (/ = 1/4), the oxygen termination of corundum oxide (0001) surface (/ = 1.5), etc. These partial fillings were indeed found in the quantum calculations. For some other polar surfaces, / is integer, and the surface can remain insulating. This takes place for example, on the (111) or (110) polar surfaces of SrTiOa (/ = 1). Charge compensation can also be achieved by changing the atomic density in the outer layers. Again, according to the bond transfer model, the required modifications of stoichiometry are found to be entirely determined by ionic arguments, and not by the specificities of hybridization in the outer layers. For example, in the case of MgO(lll), in a two-dimensional supercell containing M surface Mg, from which one atom is removed, the electrostatic criterion requires that / = (M — 2)/(2M — 2). This relationship is independent of the degree of covalency of the Mg-0 bond Its simplest achievement, with no filled conduction band ( / = 0), is obtained for M = 2. This amounts to removing half of the surface atoms from the outermost layer. A similar reasoning, assuming modifications of stoichiometry in two surface layers instead of one, would lead to the octopolar surface configuration. It is thus possible to predict which stoichiometry, in the surface layers, compatible with an insulating band structure, yields charge compensation, only on the basis of ionic considerations, even in very covalent compounds. This does not mean that covalency effects are absent, or that they are the same as in the bulk, but rather that they cancel out in the condition for charge compensation. Charge compensated nonstoichiometric surfaces, such as the (2 x 2) octopolar reconstructed rocksalt (111) surfaces, have electronic properties which present no anomaly. However, due to the presence of surface atoms with low coordination numbers, interesting phenomena for applications, such as a reduction of the gap, an increase in basicity of surface oxygens, or an increase in acidity of surface metal atoms can be expected, as on open non-polar surfaces [58,59,271,272]. 6.4. Summary Polar oxide surfaces present a wide variety of electronic and atomic characteristics, which are dependent upon the crystal structure, the ionicity of the metal-oxygen bonding, the surface orientation and its stoichiometry. The nature of the microscopic processes responsible for the cancellation of polarity provides a means to introduce a classification among these surfaces. Weakly polar surfaces are met each time that the dipole moment in produced. However, when the Fermi level crosses a narrow band, as met on the oxygen termination of MgO(lll) for example, and when the surface metallization would survive in the limit of zero bond width (no electron delocalization, zero value of the resonance integrals), the system should rather be named an "open-shell" system

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the repeat unit is a function of the electron transfers per bond only and does not involve integer contributions, as in the case of SrTiO3(100). On stoichiometric surfaces, the charge compensation may be assigned to the effect of bond breaking. The surface LDOS differs from the bulk one because of the reduced local environment of the surface atoms (different effective atomic orbital energies, bond-length modifications, change in the coordination numbers). However, the terminations are insulating, with filled oxygen-derived surface states and empty metal-derived surface states. These surfaces have low surface energies and can be produced stoichiometric and planar. We have called them previously weakly polar surfaces^ because their polar instability is weak. For fully polar surfaces, the surface dipole in the repeat unit contains an integer contribution, independently of the values of the electron transfer per bond. For semi-infinite surfaces, it was shown that a mere change of covalency in the outer layers can never provide the compensating charges, because it modifies the charges of several types of atoms whose contributions cancel out in the expression of the electrostatic criterion. Similarly, and for the same reason, surface relaxation cannot provide the compensating charges. Only a modification in the filling of surface states or a change in the stoichiometry of the surface layers may yield charge compensation. It turns out that if stoichiometric polar surfaces are unstable, this is never because they present a diverging electrostatic surface energy. The compounds have enough degrees of freedom, and in particular enough flexibility of their electronic structure in response to the surface potential, to heal the polarity while remaining stoichiometric. If a partial filling of surface states is required — and, as shown above, this can be estimated assuming a fully ionic limit —, as in the rocksalt (111), wurtzite (0001) or (0001), etc surfaces, the terminations present a "metallic" character. This is usually not favourable, from an energetic point of view, as recognized in the expression of the auto-compensation principle. On other stoichiometric polar surfaces, an integer filling of surface states is required. This is the case on SrTiOa or BaTiOs (110) and (111) surfaces. The surfaces thus may remain insulating but a conduction state is located below the Fermi Energy Ep- Self-consistent calculations give a hint that this electronic structure is not associated with a high surface energy. They support preliminary experimental observations of non-reconstructed perovskite (111) and (110) surfaces. This suggests that the autocompensation principle (filled anion-derived surface states and empty cation-derived surface states) should be extended to include all insulating surface configurations, whatever the nature of the filled and empty states. However, in perovskite compounds, the absence of reconstruction is not necessarily

85

synonymous with stoichiometry, and further work should assess this point in the future, by determining quantitatively surface compositions. Polarity may also be healed by the removal of a certain percentage of atoms in the outer layers. When the vacancies order, most of the time, this leads to surface reconstructions. The surface concentrations compatible with a vanishing macroscopic dipole moment can be correctly estimated within a fully ionic picture. Low energy configurations are expected to be insulating, and indeed, on non-polar oxide surfaces, a correlation has been found between the stability of a surface orientation and the surface gap value [58,59]. It is clear that much work remains to be done to extend our understanding to polar surfaces of transition metal oxides in which the cations have partially filled d orbitals. An especially challenging issue is related to mixed valence metal oxides, such as Fe304, in which the cations exist under two oxidation states. In addition, considering the rapid development of ultra-thin film synthesis and characterization, a simultaneous effort should be performed on the theoretical side to settle the conditions of stability of polar films. More generally, on the experimental side, it seems that one of the present bottlenecks is in a quantitative determination of the surface stoichiometry, an information of prominent interest to interpret the presence or absence of reconstruction. 7. CONCLUSION We end this review on clean oxide surfaces by stressing some directions in which theoretical improvements can be expected in a near future One deals with the ah initio description of electronic excited states. These include the attachment or removal of electrons, the account of direct or inverse photo-emission spectra, and the electron-hole excitations of the d -> d or charge transfer type. Advanced methods are presently under development to account for them: the GW method, the SIC method, the LDA-hU method, etc. However, they imply an increased computation cost, which is not routinely accessible for complex systems, such as most oxide surfaces. These methods are also expected to open the field of strongly correlated materials, among which transition metal oxides, which have important technological applications: high-Tc superconductivity, giant magneto-resistance, etc. A second challenge is in the understanding of surface magnetism: elementary spin-spin interactions, atomic magnetic moments on surface sites, collective properties such as surface magnetic ordering, spin polarised transport across oxide ultra-thin barriers, etc. We have not developed this problematics here, because, until now, there have been very few related the-

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oretical studies. However, new beautiful experimental results have been recently published which should encourage theorists to join this field. Finally, understanding the reconstructions of oxide surfaces from the point of view of non-stoichiometry and/or polarity represents a necessary step for the production and use of high quality nano-structured surfaces. On the experimental side, it seems that one of the present bottlenecks is in a quantitative determination and control of the surface stoichiometry. On the theoretical side, there are strong limitations to apply ab initio methods to reconstructed surfaces because of their computational cost. In addition, it is not yet clear why intrinsic reconstructions do not exist and what are the driving forces for vacancy ordering. No doubt that these points will be seriously considered in a near future if one is to use reconstructed oxide surfaces as substrates to grow artificial structures. REFERENCES I] V.E. Henrich and P.A. Cox, The Surface Science of Metal Oxides, Cambridge University Press, Cambridge, 1994. 2] C. Noguera, Physics and Chemistry at Oxide Surfaces, Cambridge University Press, Cambridge, 1996. 3] C. Noguera, Surf. Rev. Letters , (2001) under press. 4] C. Noguera, J. Phys. Condensed Matter, 12 (2000) R367. 5] B. G. Dick and A. W. Overhauser, Phys. Rev., 112 (1958) 90. 6] W. C. Mackrodt and R. F. Stewart, J. Phys. C, 12 (1979) 431. 7] P. A. Madden and M. Wilson, Chem. Soc. Rev., 25 (1996) 339. 8] M. Wilson and P. A. Madden, Faraday Discussion, 106 (1997) 339. 9] A. M. Stoneham and J. H. Harding, Ann. Rev. Phys. Chem., 37 (1986) 53. 10] M. C. Payne, M. P. Teter, D. C. Allan, T. A. Arias and J. D. Joannopoulos, Rev. Mod. Phys., 64 (1992) 1045. II] J. P. LaFemina, Surf. Sci. Rep., 16 (1992) 133. 12] J. A. Pople, D. P. Sentry and G. A. Segal, J. Chem. Phys., 43 (1965) S129. 13] J. A. Pople, D. P. Sentry and G. A. Segal, J. Chem. Phys., 44 (1965) 3289. 14] W. A. Harrison, Electronic Structure and the Properties of Solids, Freeman and co, San Francisco, 1980. 15] C. Pisani, R. Dovesi and C. Roetti, Hartree-Fock ab initio treatment of crystalline systems. Lecture Notes in Chemistry, Vol. 48, Springer Verlag, Berlin, 1992. 16] A. Szabo et N.S. Ostlund, Modern Quantum Chemistry: introduction to advanced electronic structure theory, MacGraw Hill, 1989. 17] N. Govind, Y. A. Wang and E. A. Carter, J. Chem. Phys., 110 (1999) 7677. 18] P. Hohenberg and W. Kohn, Phys. Rev., 136 (1964) B864. 19] R. G. Parr and W. Yang, Density functional theory of atoms and molecules, Oxford University Press, Oxford, 1989. 20] R. Jones and O. Gunnarsson, Rev. Mod. Phys., 61 (1989) 689. 21] A. Nagy, Phys. Rep., 298 (1998) 1. 22] J. P. Perdew, J. A. Chevary, S. H. Vosko, K. A. Jackson, M. R. Pederson, D. J. Singh and C. Fiolhais, Phys. Rev. B, 46 (1992) 6671. [23] J. P. Perdew, K. Burke and Y. Wang, Phys. Rev. B, 54 (1996) 16533.

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Chapter 3

Theory of point defects at the MgO surface Gianfranco Pacchioni Dipartimento di Scienza dei Material!, Universita di Milano-Bicocca, Istituto Nazionale per la Fisica della Materia, via Cozzi, 53 - 20125 Milano, Italy

1. INTRODUCTION Oxides surfaces are finding continuous new applications in advanced technologies like in corrosion protection, coating for thermal applications, in catalysis as inert supports or directly as catalysts, in microelectronics for their dielectric properties; films of magnetic oxides are integral components in magnetic recording devices and many microporous materials are based on oxides. For all these reasons there is a considerable effort to better characterize the surface and the interface of oxide materials [1,2]. New experimental methods have been developed to growth well defined films of oxides on metal supports in UHV conditions. This approach allows in principle to overcome some difficulties connected to the use of electron spectroscopies for the study of insulating materials, as most of the oxides are. Recently, some very good review articles have been published on this subject [3-10]. It is no surprise that the increasing experimental activity has stimulated a parallel computational activity based on high-quality first principle calculations. One important information coming from calculations is the structure of the oxide surface. Oxide surfaces are often heavily reconstructed or simply relaxed compared to the truncated bulk, and the experimental determination of the surface structure is often not easy. To this end, reliable classical potentials have been developed, in particular for the study of ionic crystals and of covalent solids. Nowadays, first principle band structure calculations making use of large supercells can also be used. These methods, although quite expensive from the point of view of the size of the calculations, provide results which are in excellent agreement with the experimental determinations. Band structure calculations, usually based on plane waves basis sets and on the density functional (DFT) approach represent the most appropriate computational

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method to study solids. However, an alternative approach to band structure has been developed in the past two decades with the aim of describing in chemical terms interactions occurring at surfaces. This is known as the cluster approach [11] and it has been applied to a variety of problems in surface science. As it will be explicitly discussed below, it has advantages and disadvantages with respect to band structure supercell methods. One of the problems where the use of the cluster approach is more appealing is in the study of surface defects. Only recently it has been recognized by the surface science community that these centers are often the most interesting ones from the point of view of the physical and chemical properties of a material. Several chemical reactions taking place at an oxide surface are directly or indirectly connected to the presence of point and extended defects. Unfortunately, defect centers are elusive species because of their low concentration even in the bulk material, and their identification by spectroscopic methods can be rather difficult. This is even more dramatic when one is interested in surface defects because the distinction from bulk defects may be extremely subtle. For all these reasons the theoretical modeling of defect centers at the surface of oxides is attracting an increasing interest. In this review we will give a general overview of the most common defects occurring at the surface of ionic oxides and in particular of MgO, an ionic material which has been at the center of considerable attention in the last years as a model system, but also because of its potential and technological applications in catalytic processes. 2. QUANTUM CHEMICAL DESCRIPTION OF IONIC CRYSTALS The electronic structure of solids is usually described in terms of band structure. To this end, a unit cell containing a given number of atoms is periodically repeated in three-dimensions to account for the "infinite" nature of the crystalline solid and the Schrodinger equation is solved for the atoms in the unit cell subject to periodic boundary conditions [12]. This approach can also be extended to the study of point defects in the bulk or at the surface, although the ab initio simulation of defects in crystals is not simple. One should in fact include a "reasonable" description of the "infinite" host system while accounting for the rupture of the translational symmetry induced by the defect. One way to circumvent this problem is the supercell approach in which an artificial periodic structure is created where the defect is translationally reproduced in correspondence to a given super-lattice of the host. This procedure allows the use of efficient computer programs designed for the treatment of periodic systems and has indeed been followed by several authors to study defects using either DFT and plane waves approaches [13-15] or Hartree-Fock-based methods with localized atomic orbitals [16,17].

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The presence of the defect in the unit cell, however, results in a periodic repetition of the defect in the three directions of space, hence in an extremely high concentration. The only way to reduce the defect concentration is to increase the size of the unit cell, a solution which implies a very large computational cost. Nowadays, periodic calculations for supercells containing several tens of atoms are routinely done. Even for large supercells containing -100 atoms, however, the defect concentration is orders of magnitude larger than in real systems where 10^^-10^^ defects/cm^ are usually present [18,19]. At these concentrations one can reasonably assume that the defects are isolated and do not interact with each other. The supercell approach is therefore based on the assumption, largely verified, that neutral defects do not interact appreciably except when they are very close to each other [20], so that rapid convergence is achieved with increasing size of the supercell. With charged defects the supercell approach is feasible but less reliable because of the long range Coulomb interaction between the defects [21]. Methods to include correction terms to account for these spurious interactions have been proposed [22]. 2.1 Cluster models An alternative approach to the periodic band structure methods to study solids is the cluster approach [11]. Here one considers explicitly only a finite number of atoms to describe a section of the solid while the rest of the crystal (or of the amorphous material) is treated in a more or less simplified way (embedding). The main conceptual difference is that in the cluster approach one uses molecular orbitals, MO, instead of delocalized bands. The description of the electronic properties is thus done in terms of local orbitals, allowing one to treat problems in solids with the typical language of chemistry, the language of orbitals. This is particularly useful when dealing with surface problems and with the reactivity of a solid surface. In fact, the interaction of gas-phase molecules with a solid surface can be described in exactly the same way as the interaction of two molecules. Of course, also the cluster model is not free from limitations. The most serious one is that the effect of the surrounding is often taken into account in a more or less approximate way, thus leading to some uncertainties in the absolute values of the computed quantities. It is also possible that some properties are described differently depending on the size of the cluster used. It is therefore necessary to check the results versus cluster size and shape. Examples of this aspect will be discussed in more detail below. The advantages, beside a smaller computational cost, are that in describing point defects an infinite dilution is considered so that no mutual interaction of the defects is present in the model and that theoretical methods derived from quantum chemistry can be applied. This latter is an important advantage which should not be underestimated. In fact, in this way it is possible (a) to explicitly include correlation effects in the calculations through, for instance, a configuration

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interaction (CI) procedure (see below) and (b) to treat exactly the non-local exchange as in the Hartree-Fock formalism; in DFT in fact the exchange is taken into account in an approximate way through the exchange-correlation functional. We briefly mention here two examples where the use of quantumchemical techniques is extremely important for the correct description of the electronic structure of point defects. The first one is the calculation of excited state properties like for instance optical transitions and intensities, and therefore optical spectra. While some examples of excited-state calculations based on time-dependent DFT are now available [23], the application to solid state problems is still in its infancy. The use of cluster models and CI methods represents a valid approach to the study of excited states of point defects in sohd matrices, and excellent results have been obtained in recent years for the study of optical spectra in Si02 [24-27] and MgO [28,29]. The second problem is that of spin localization in paramagnetic defects. Recent DFT calculations based on supercells and plane waves have described the [A104]° center in silica, an Al impurity replacing a four-coordinated Si atom in the Si02 lattice, as a tetrahedrally coordinated Al atom with a hole delocalized over the four O neighbors [30-32]. This description contradicts the experimental evidence from ESR studies of a hole completely localized on a single O atom [33], a physical picture supported by HF cluster calculations [34,35]. The reason for the incorrect description of the center in DFT is the lack of self-interaction correction and the tendency to delocalize the hole. Recent results obtained at various levels of theory have clearly shown the important role of exact exchange for the proper description of this center [36]. Therefore, cluster calculations represent an alternative way of describing localized defects in ionic crystals. The problem is to introduce in a reasonable way the effect of the rest of the crystal. Completely different strategies can be adopted to "embed" clusters of largely covalent oxides, like Si02, or of very ionic oxides, like MgO. In Si02 and related materials the cluster dangling bonds are usually saturated by H atoms [37]. The positions of the cluster atoms can be initially those of one crystalline phase; then the H atoms are kept fixed while the position of all the other atoms of the cluster are reoptimized. The fixed H atoms provide a simple representation of the mechanical embedding of the solid matrix [24-26]. As an alternative, fully relaxed clusters have also been used [27]. It is believed that these latter systems are better models of the amorphous phase of the material. The saturation of the dangling bonds with H atoms is an important aspect of the embedding, but not the only one. In fact, in this way one neglects the crystalline Madelung field. While this term is of moderate importance in silica and other dielectrics, it may play a role in the description of charged defects and becomes essential in solids with more pronounced ionic character.

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2.2 Embedding schemes The very ionic nature of MgO implies that the Madelung potential is explicitly included. Indeed, several properties of MgO are incorrectly described if the long range Coulomb interactions are not taken into account [38]. A simple approach is to surround the cluster of Mg and O ions by a large array of point charges (PC) ±2 to reproduce the Madelung field of the host at the central region of the cluster [39]. However, the PC's polarize the oxide anions at the cluster border and cause an incorrect behavior of the electrostatic potential [40]. The problem can be substantially reduced by placing at the position of the +2 PC's around the cluster an effective core potential, ECP, representing the finite size of the Mg^"^ core [41]. No basis functions are associated to the ECP [42] which accounts for the Pauli or exchange repulsion of the O^" valence electrons with the surrounding. This is a simplified approach to the more rigorous ab initio model potential (AIMP) method [43,44] but is computationally simple and reliable. In the AIMP approach the grid of bare charges is replaced by a grid of AIMP's which account not only for the long-range Coulomb interaction but also for the quantum mechanical short-range requirements of exchange and orthogonality without introducing explicitly extra electrons in the model. The addition of the ECP's to the cluster results in a better representation of the electrostatic potential and in substantially different stabilities of the charge states of the defect. For instance, the ionization potentials of oxygen vacancies (the F centers described in the next section) are lowered by about 3 eV when the cluster is embedded in ECP+PC's. What is still missing from this simplified approach is the polarization of the host crystal by the defect. This effect can be particularly important for charged defects. The polarization, Epoi, induced by a charge on the surrounding lattice can be estimated by means of the classical Bom formula [45]: Epoi= - ( l - l / 8 ) q V 2 R

(1)

where 8 is the dielectric constant of MgO, q is the absolute value of the charge and R is the radius of the spherical cavity where the charge is distributed. Since a certain degree of ambiguity remains in the definition of R, this correction is only qualitative. A more refined approach which has been used for the study of the ground state of F centers in MgO [46] makes use of the ICECAP program [47]. In this approach instead of PC's the cluster is surrounded by polarizable ions described according to the shell model [48,49]; in this way the polarization response of the host is taken into account self-consistently up to infinite distance. In the shell model an ion is represented by a point core and a shell connected by a spring to simulate its dipole polarizability. A similar method has been applied recently to the study of energy states of defect sites at the MgO

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surface [50]. A quantum cluster has been embedded in a finite array of PC's placed at the lattice sites, Fig. 1. The part of the ions closest to the quantum cluster has been treated by the shell model in such a way that they interact among themselves and the quantum cluster via specific interatomic potentials. The novelty of the approach lies in the fact that the positions of the cores and shells of the ions are optimized in response to the changes in charge density distribution within the quantum cluster to minimize the total energy of the system [50]. Quantum Cluster

a),

Fixed ions (Region II)

Shell model (Region I)

Fig. 1. Clusters of surface and comer sites of the MgO surface with embedding in shell models and point charges. Reproduced from ref. [50]. Copyright 2000 Elsevier.

An alternative, more rigorous approach has been developed in recent years by Pisani and coworkers [51-53]. It is named perturbed cluster method and is based on the EMBED computer program [54]. With this approach the properties of neutral and charged defects at the surface of MgO have been studied at the HE and MP2 levels [55]. The method relies on the knowledge of the one-electron Green function G^ for the unperturbed host crystal which is obtained by means of the periodic program CRYSTAL [56,57]. A cluster (C) containing the defect is defined with respect to the rest of the host (H). The molecular solution for the cluster C in the field of H is corrected selfconsistently by exploiting the information contained in G^ in order to allow a proper coupling of the local wave function to that of the outer region. 2.3 Solutions of the Schrodinger equation Two general groups of methodologies are used to solve the Schrodinger equation in combination with cluster models, the Hartree-Fock (HE) approach and related methods to include correlation effects like M0ller-Plesset perturbation theory (MP2) or configuration interaction (CI) [58,59] and the Density Eunctional Theory (DET) approach [59,60].

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In HF theory the energy has the form EHF = ENUCL +

+ 1/2 - 1/2

(2)

where ENUCL is the nuclear repulsion energy, P is the density matrix, is the one electron energy including kinetic and potential energy terms, 1/2 is the classical coulomb repulsion term and -1/2 is the exchange energy resulting from the quantum nature of the electrons. Thus, the HF method does not include electron correlation which must be included explicitly through perturbation theory at the second (MP2) or higher order (MPS and MP4) or through configuration interaction (CI) techniques. In the Kohn-Sham formulation of DFT the exact HF exchange for a single determinant is replaced by a general expression, the exchange correlation functional, which can include both exchange and electron correlation energy terms: EKS

= ENUCL + + 1/2 + Ex[P] + Ec[P]

(3)

where Ex[P] is the exchange functional and Ec[P] is the correlation functional. HF can be seen as a special case of DFT where Ex[P] is given by the exchange integral -1/2 and Ec = 0. In this respect cluster models offer the opportunity to use the so-called hybrid approaches where the exchange interaction is described partially by the HF method. For example, in the popular B3LYP hybrid method the HF exchange is mixed in with the DF exchange while the correlation, Ec[P], is treated by the Lee, Yang and Parr, LYP, correlation functional [61]. In particular, the exchange treatment is based on the Becke 3-parameters approach [62] where a set of parameters is derived in order to accurately describe the thermochemistry of a large number of molecules. While this introduces some empiricism in a first principle theory, it has been shown that the B3LYP exchange-correlation functional describes in a much more accurate way the energetics of reactions at surfaces than most of the commonly used exchange-correlation functional. This has been proved in particular for the adsorption of metal atoms on MgO [63,64]. However, exceptions are also known and in general the choice of the exchange-correlation functional remains an open problem. The main difference between HF and DFT is that in the HF approach the objective is the determination of the wave function of the system while in DFT the quantity of interest is the charge density p; in fact, the Hohenberg-Kohn theorem [65] ensures that the energy E of a system is a functional of p which takes its minimum value EQ for the ground state density po. Despite this profound conceptual difference, the two methods, HF+correlation and DFT, provide in ultimate analysis very similar results and the choice of one approach or the other is largely matter of convenience.

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3. ROLE AND NATURE OF DEFECTS AT THE MgO SURFACE Often the most important properties of materials are directly or indirectly connected to the presence of defects and in particular of point defects [18]. These centers determine the optical, electronic and transport properties of the material and usually dominate the chemistry of its surface. A detailed understanding and a control at atomistic level of the nature (and concentration) of point defects in oxides is therefore of fundamental importance to synthesize new materials with well defined properties. This has lead in recent years to the birth of the new field of defect engineering. Of course, before to be created in controlled conditions point defects have to be known in all aspects of their physico-chemical properties. The accurate theoretical description of the electronic structure of point defects in oxides is essential for the understanding of their structure-properties relationship. MgO is a particularly well studied oxide; the structure of the (100) single crystal surface is extremely flat, clean, and stoichiometric. Recent grazing incident X-ray scattering experiments have shown that both relaxation, -0.56±0.4%, and rumpling, 1.07±0.5%, are extremely small [66]. However, no real crystal surface consists of only idealized terraces. A great effort has been undertaken in recent years to better characterize the MgO surface, in particular for polycrystalline or thin-film forms which in some cases exhibit an heterogeneous surface, due to the presence of various sites. All these sites can be considered as defects. The identification and classification of the defects is of fundamental importance. In fact, the presence of appreciable concentrations of defects can change completely the chemical behavior of the surface. A typical example is that of the reaction of CO on MgO (see § 3.1). Table 1 Summary of most important surface defects in MgO Defect Symbol Schematic description low-coordinated cation Mg'^nc (n=3,4) coordinatively unsaturated cation low-coordinated anion 0'-nc(n=3,4) coordinatively unsaturated anion hydroxyl group (OH) proton attached to O^' anion vacancy F"^-'nc(m=0,l,2;n=3,4,5) missing oxygen with trapped electrons cation vacancy V""-„c(m=0,l,2;n=3,4,5) missing cation with holes at O neighbors divacancy cation and anion vacancy VMgVo substitutional cation (M) or anion (X) impurity atoms M"VO"-; Mg^'-ZXôxygen radical hole trapped at O anion 0-nc (n=3,4,5) small ensemble of Mg3c or Osc ions (111) microfacets none

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While the single crystal (100) surface is completely unreactive, a defect-rich polycrystalline MgO surface exhibits a rich and complex chemistry when interacting with CO. In the following we provide a summary of the most important kinds of defects at the MgO surface, Table 1, and we discuss their electronic structure and chemical behavior, in view of the existing experimental and theoretical data. In a broad classification of the defect centers at the surface of MgO one can recognize at least 9 major kinds of irregularities: • Low-coordinated cations (§ 3.1). These are Mg^"^ ions with a number of neighbors lower than on the flat (100) terraces. To this category belong therefore four-coordinated ions located at step and edge sites, Mg^%c» and the three-coordinated ions located at comers, kinks, etc. Mg^'*"3c. • Low-coordinated anions (§ 3.2). The O^" sites exhibit a completely different chemistry when located in the bulk, at the five-coordinated terraces, O^'sc, or at irregular sites with lower coordination, 0^'4c and O^'sc, where the Madelung potential and the basicity of the site are different. • Hydroxyl groups (§ 3.3). H2O is almost ubiquitous and easily reacts with low-coordinated sites to form OH groups at the surface of MgO. These centers can act as nucleation centers in the growth of metal particles, induce asymmetries in the surface electric field, or exhibit a classical Broensted acid behavior. • Oxygen vacancies (§ 3.4). These are usually called color centers or F centers from the German word for color, Farbe. The vacancies can have different formal charges. The removal of a neutral O atom results in a neutral F center; the removal of a O' ion in a F"^ center (paramagnetic); the removal of a O^' in a F^"^ center. Since the surface O ions can be located at terrace, step and corner sites, also the corresponding vacancies can in principle form at five-, four-, and three-coordinated sites. Fig. 2. This makes the number of possible combinations very high and the identification of all species quite complex. In the following the subscript s distinguishes a surface F center, Fs"^, from the bulk counterpart, F"^. • Cation vacancies (§ 3.5). These are often classified as Vs centers, where the subscript s has again the meaning of distinguish between bulk or surface vacancies. Cation vacancies in MgO correspond to the removal of Mg, Mg"^, or Mg^"^ species and results in V, V and V^' defects. The V and V" centers are paramagnetic. • Divacancies (§ 3.6). Closely related to the O and Mg vacancies are the divacancies, or defect centers created by removing a neutral MgO unit. This process is thermodynamically more favorable and the existence of divacancies has been shown for MgO bulk. The formation of a divacancy on

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a (100) terrace results in a mini-step, a site where some of the surface atoms are three- or four-coordinated. Impurity atoms (§ 3.7). The presence of substitutional atoms can lead to modified chemical centers on the surface. The replacement of Mg ions by Ni ions, as in MgO-NiO solid solutions, introduces transition metal atoms in a MgO matrix and can alter the local properties of the material. Even more effective is the replacement of a divalent Mg^"*" cation by a monovalent dopant like Li^. In order to compensate the charge, some O^" anions at the surface become O", a paramagnetic species. Oxygen radical anions (§ 3.8). The O" center can be formed by various means on the surface, not only by doping the material with alkali metals ions. O' exhibits an high chemical reactivity and can be located at lowcoordinated sites as well as at regular terrace sites. ( I l l ) microfacets (§ 3.9). Small ensembles of three-coordinated atoms with a pyramidal structure resembling a reconstructed (111) polar surface can in principle form and have been indeed been proposed as the centers where oxygen exchange reactions occur.

Fig. 2. Schematic representation of oxygen vacancies at various sites of MgO. (a) terrace; (b) edge; (c) comer.

3.1 Low-coordinated cations As an example of a chemisorption reaction where the control on the surface defects is crucial we mention here the case of CO on MgO. This process has been studied on polycrystalline samples, on MgO thin films supported on a metal, and on MgO single crystals, three forms of the material where the surface morphology and the defect concentration are different. This experimental

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activity has been complemented during the last decade by an intense theoretical activity, largely based on cluster models (for a complete account see ref. 67). CO is known to adsorb on Mg cations with the molecular axis normal to the (100) MgO surface. Infra-red (IR) measurements at 77 K of the C-0 vibrational frequency on polycrystalline MgO have shown a small positive shift of about +10 cm"^ with respect to free gas-phase CO [68], at variance with metals where the shift is usually large and negative. The adsorption energy of CO on MgO powders determined experimentally from adsorption isotherms [69] or from integrated intensities of the IR spectra at various temperatures and pressures [70] is quite low, 0.15-0.17 eV, indicating a very weak adsorption. A quite different result was obtained in a study of the interaction of CO on ultra-thin MgO films grown in UHV on a Mo(llO) substrate [71]. Using IR reflection absorption spectroscopy (IRAS) it was found that CO has a frequency of 2178 cm"\ with a shift of +35 cm'\ and from temperature programmed desorption (TPD) spectra an heat of adsorption of 0.43-0.46 eV was estimated [71]. Thus, the results for the adsorption on a MgO(lOO) film look quite different from those of a polycrystalline sample; in particular, while a vibrational shift of ~+10 cm'^ has been detected on MgO powders and smokes, an about three times larger shift, +35 cm'^ [71], has been measured on the films. In a similar way, the binding energy of -0.15 eV deduced from measurements on MgO powders [69,70] is three times smaller than that obtained for the film, 0.43 eV [71]. This apparent contradiction has stimulated a considerable amount of theoretical work from the early 90's until very recently. A number of more and more refined calculations has been performed by several groups using various methods [72], from simple HF [73,74] to HF with extended inclusion of correlation [75,76], from LDA [77] to gradient corrected DFT [78,79]. However, the most sophisticated methods and calculations lead to the same conclusion: CO is bound to the terrace sites of the MgO surface through electrostatic forces, with a very weak binding energy of less then 0.1 eV, and a very small vibrational shift, of +10 cm'^ or less [67,72]. Eventually, the discrepancy between the ab initio calculations and the results on MgO thin films lead to the suggestion that surface defects could be responsible for the observed chemisorption properties of CO/MgO/Mo(l 10). This is connected to the electrostatic nature of the interaction [80] and to the role of low-coordinated sites. A direct relationship exists between the Madelung potential at a given site, Fig. 3, the strength of the surface electric field and the CO stretching frequency [80]. In fact, less coordinated cations generate a larger electric field which interacts with the CO dipole and multipole moments giving rise to an increase of cOb. This has permitted the rationalization of the presence of different peaks in the IR spectrum of CO adsorbed on MgO micro-crystals or MgO smokes [81] by means of simple electrostatic arguments. In fact, the adsorption of CO on low-coordinated, more exposed, Mg^"^ cations at

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edge and comer sites leads to larger binding energies and CO vibrational shifts [82]. For instance, CO frequency shifts of +31, +55, and +97 cm'^ were predicted at the HF level for CO adsorbed on Mg^'^sc (terrace), Mg^'^4c (edge) and Mg^"^3c (comer) sites [82], with a similar trend to that experimentally observed in IR for CO on various sites of MgO smokes, +13, +21, and +60 cm"^ [81,83]. The increasing binding energy and vibrational shift is a feature which is perfectly consistent with the electrostatic picture of the CO-MgO bonding. This idea has then been refined and lead to a reinterpretation of the TPD spectra of He et al. [71] based on accurate cluster model calculations on CO/MgO [75,84]. The results have shown that the binding energies for CO adsorbed on Mg^"^ cations at terraces, edge, and corners of MgO surfaces were 0.08, 0.18 and 0.48 eV; frequency shifts of +9, +27, and +56 cm"^ values were obtained for the same sites [75], Table 2. The computed adsorption energies have suggested that three peaks should be expected in a TPD spectmm of CO on MgO, around 40 K (terrace), 80 K (edge), and 160 K (comer). Only the last temperature is close to the regime where CO desorption was detected on MgO films [71]. On this basis it has been concluded that the oxide film contains an high density of defects. {Mgg }(step) \ { M g 3 } (comer)

{ M g , - } (step)

}(100)

Fig. 3. Schematic representation of a MgO surface; Mg^"^ ions are represented as small spheres, O^" anions as large spheres. The subscript indicates the coordination number of each cation site. The Madelung constants computed for these sites are: 1.681 Mgsc (terrace); 1.591 Mg4c (edge); 1.566 Mg4c (step); 1.344 Mgsc (comer), 0.873 Mgsc (step). The Madelung constant for a bulk Mg6c is 1.747.

The conclusive proof came only recently and was based on MgO single crystal experiments. A single crystal surface is expected to have a small number (almost negligible) of defects. On a UHV cleaved MgO single crystal [85,86], the CO thermal desorption spectra (TDS) showed a peak at T 57 K and a broad feature at 76 K. The adsorption energy, calculated via the Redhead equation

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[87], 0.14 eV, is very close to the earlier data of CO on MgO powder [69,70] and to the theoretical calculations [72]. The TDS peak at 57 K and the De value of 0.14 eV refer to CO molecules adsorbed on the five-coordinated Mg cations at the terrace sites; the broad feature around T = 76 K is most likely due to adsorption on surface defects, since it becomes much stronger after introducing additional defects through ion bombardment. The study of the vibrational modes of small adsorbed molecules provides therefore a very useful tool to detect and titrate low-coordinated cation sites in UHV cleaved single crystals as well as on high surface area polycrystalline materials. This does not apply only to CO but also to other molecules; for instance, excellent correlations between theory and experiment have been found for the v(CN) bands of CD3CN adsorbed on Mg4c^"^ and Mgsc^"*^ cations of MgO [88]. In the field of catalysts characterization the use of small unreactive "probe" molecules to identify coordinatively unsaturated sites is well established [89]. Not always, however, a direct correlation between the CO vibrational frequency, the strength of the interaction, and the surface electric field exists. Recent DFT cluster calculations [90] have shown that CO adsorbed on step sites gives rise to a relatively strong interaction but to a negligible CO vibrational shift; this is due to the inhomogeneity in the electric field above a MgO(lOO) step. This study [90] has permitted the complete attribution of the IR spectrum of CO adsorbed on MgO [81,83,91], Table 2. Table 2 Binding energies, De, and frequency shifts, Aco, for CO adsorbed on various sites of the MgO (100) surface from embedded cluster calculations De,eV Aco(C-0), cm'^ ^"^ CI [75] DFT [90] CI [75] DFT [90] exp. [81,91] terrace 0.08 0.01 -7 +9 +9/+13 edge 0.18 +27 0.17 +10 +16/+21 comer 0.48 0.34 +56 +49 +57/+60 — step 0.21 +1 +5 (a) Computed with respect to free CO (Oo: 2186 (CI), 2210 cm'^ (DFT); the experimental coo is 2170 cm'' but the experimental shifts are computed with respect to the CO cob 2143 cm'\

Another technique to identify the surface cations in MgO is connected to the use of the O2' superoxide anion and the EPR spectroscopy [92,93]. O2" can form at the surface of MgO by interaction of O2 with electron-rich centers see § 3.4.7). The electric field at the surface created by the Mg^"^ cation removes the degeneracy of the Tig orbitals giving rise to an EPR spectrum with three components of the g-tensor, gxx, gyy, and gzz- Depending on the cation where the superoxide anion is adsorbed, the splitting A between the TCgx and TCgy components determines the value of gzz, according to the following relation which neglects second order terms [1]:

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g,, = ge + 2X/A

(4)

Here X is the spin-orbit coupling constant for oxygen, 135 cm"\ and ge is the free spin value, 2.0023. From ab initio Hartree-Fock calculations the value of A, hence of gzz, has been determined theoretically for O2' adsorbed at various sites of the MgO surface; the comparison with the observed values has allowed the identification of various cations sites on the surface [94]. 3.2 Low-coordinated anions In the analysis of the results of CO adsorption on MgO we have seen that the low-coordinated Mg^^ cations, being more exposed, generate a larger electric field. Therefore, every molecule adsorbed on these cations will undergo a stronger polarization and a stronger interaction with the surface. This interaction however is largely electrostatic, and does not imply a significant charge transfer. In general the surface cations of MgO are not directly responsible for the chemical reactivity, unless the neighboring O anions are also involved. Things are different in transition metal (TM) oxides where the d electrons of the metal cation can be directly involved in the bonding. In general, the activation of an adsorbed molecule is connected to the transfer of electronic charge into the antibonding orbitals of the adsorbate, with consequent rupture or weakening of the bond, Fig. 4. In this respect, the O anions can play a very important role in the activation of adsorbed species. 3b^ ''''>••'"

-0.34 ay

^ /

+H--

« ^

^/

'••'11 '

/

0.02 au

4^^0

-0.36 au

i Mg-—O — M g Mg

Mg

Fig. 4. Schematic representation of the two-electron two-orbitals interaction between a donor sp orbital on a four coordinated 04c^' site of the MgO surface and the accepting orbitals of SO2. Reproduced from ref. [96]. Copyright 1994 Elsevier.

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As a simple example to illustrate the different reactivity of oxygen anions located at various sites of the MgO surface let us consider the interaction of CO2 and SO2. These two molecules can form surface carbonates, MgO + CO2 ^ MgCOs, and sulfites, MgO + SO2 -> MgSOs. In ab initio cluster model studies of these reactions [95,96] it has been found that the five-coordinated O^' ions at the (100) MgO terraces are very unreactive towards CO2 and SO2 and that only weakly bonded physisorbed species form. On the contrary, the 0^"4c ions at the step sites are very reactive and give exothermic reactions with formation of carbonates and sulfites. The O^sc ions at the (100) terraces of CaO exhibit a similar reactivity to that of the low-coordinated anions of MgO, and stable chemisorbed species form at the terraces of CaO. This result is consistent with the trend of surface basicity, MgO Mg—O—Mg

Mg + H2O

(5)

Anion vacancies on the MgO surface, the Fs^"*^ centers corresponding to the removal of O^', are difficult to observe, often invisible to many surface science techniques. However, their existence can be shown by doping the material with excess electrons to form the corresponding Fs"^ and Fs centers which can be detected either by EPR (Fs"") [122-125] or by optical (Fs"" and Fs) [126-128] spectroscopies. 3.4.1. Charge distribution The removal of an O atom or ion from the MgO surface results in a large relaxation of the lattice. For a neutral Fs center there is an outward displacement of the Mg^"^ ions and an inward movement of the second neighbor O^" ions by 12% with respect to the unrelaxed surface. The effect is much more pronounced for charged vacancies, and changes of - 5 % and -10% in the Mg-0 distances around the cavity are found for Fs"^ and Fs^"^ centers, respectively [39]. In the F"^ or F centers one or two electrons are associated with the defect. The extra electron(s) can either be delocalized over the 3 s levels of the Mg^^ ions around the cavity, with partial reduction of these ions and a change of their charged state from Mg^^ to Mg"^, or can be localized in the vacancy center. All theoretical studies agree with this second picture [13-17,39,99]. This is shown by the plots of the electron density maps and in particular by density difference maps which show very clearly the electron localization at the center of the vacancy [39,55]. This is true not only for the bulk but also for the surface of MgO, Fig. 5. The locaUzation of the electrons in the center of the vacancy is an indirect proof of the highly ionic nature of MgO. In fact, the electrons are trapped in the cavity by the crystalline Madelung potential. Calculations performed on cluster models have shown that in absence of the external field the electrons tend to distribute more over the 3 s levels of the Mg ions around the vacancy [38]. The localization of the electron in the center of the vacancy is

112

Fig. 5. Total density maps (above) and difference density maps (total minus ionic superposition below) for Fs, Fs^, and Fs^"^ defects at the MgO (100) surface (from left to right). In the plots below continuous lines correspond to charge accumulation, dashed lines to charge depletion with respect to the superposition of the ionic densities. Reproduced from ref. [55]. Copyright 1997 American Institute of Physics.

clearly shown also by spin density plots, Fig. 6, for the case of paramagnetic F"*" centers.

r? F^ > F^^. Neutral F centers are more stable because two electrons interact with the Madelung field of the crystal. F^"^ centers, on the other hand, are much less stable and are very electron deficient. The formation energies computed at the HF level for neutral F centers at bulk, surface, step and comer sites, 9.7, 8.1, 6.9, and 5.6 eV, respectively [139], are lower than those obtained

116

using the DFT approach [15] because of the absence of correlation effects. The trend in stability is directly connected to the energy required to remove an oxygen atom from the material. Thus, the formation energy of neutral vacancies follows the trend bulk (or sub-surface) > surface > step > comer; it follows that low-coordinated F centers are easier to form, hence thermodynamically more stable, than the bulk F centers. This is true also for charged vacancies but, in particular for the F^"*" centers, the difference in stability between low- and highcoordinated sites is not as high as for neutral F centers [139]. 3.4.4. Energy levels and ionization potentials An important question when one addresses the stabiUty of charged defects in insulators is that of the relative position of the neutral and charged states of a given center with respect to the vacuum level. This problem has been considered recently for the case of the MgO(lOO) surface [50]. Using cluster models embedded in polarizable ions and PC's, the ionization energies and electron affinities of oxygen vacancies located at the terraces and at the comers of the MgO surface have been determined. Fig. 7. A neutral Fsc center gives rise to an impurity level in the gap which is about 3 eV above the top of the valence band; a neutral Fsc center at a comer site lies more or less at the same energy. F"^ centers give rise to states closer to the valence band; for a terrace F"^5c the level is about 1 eV above the O 2p states, while for a F'^'sc defect the impurity level is nearly at the top of the valence band. The vertical ionization energy of a neutral Fsc defect is 3.4 eV; that of a F"^5c center 5.6 eV. Surface E, eV IP ^ I

~T

Corner

Exciton A (calc.)

IP

vacuum - level

(experimT

Exciton (experimicalc.)

-2-j -3

I

F centre

F centre

-4 -5 -6

F+ centre

Corner

Fig. 7. Energy levels of various defects at the MgO surface from embedded cluster DFT calculations. Reproduced from ref. [50]. Copyright 2000 Elsevier.

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It is interesting that similar IP's are found for F centers located at a comer site: the IP's of Fsc and F^3c centers are 3.4 and 6.6 eV, respectively [50,139], Table 3. Thus, not only electrons trapped at low-coordinated sites may exist, but they are even more strongly bound that at the corresponding terrace sites. This result is not obvious. In fact, one would expect the energy to ionize an F center to be proportional to the MP at that site. Assuming no geometrical relaxation, the MP is smaller at low-coordinated sites and increases for more coordinated sites, opposite to what found for the vacancy ionization energy. This behavior is due to the Pauli repulsion of the trapped electrons with the surrounding ions of the vacancy [139]. This is larger for high-coordinated F centers and decreases for low-coordinated defects where the trapped electrons can extend toward the vacuum region thus reducing both the Coulomb and the exchange repulsion. The Pauli repulsion destabilizes the F center impurity level and decreases the ionization energy for the high-coordinated sites. It follows that electrons associated to low-coordinated oxygen vacancies are more strongly bound than in the corresponding high-coordinated vacancy sites because of their reduced repulsion with the surrounding [139]. We will see below that this is also the physical origin for the different excitation energies of electrons trapped at F centers located at high- and low-coordinated sites. 3.4.5. Barriers to vacancies dijfusion The barrier for oxygen migration at MgO surfaces is another topic of interest for the understanding of the defect distribution. The F^"^ centers are expected to be involved in ionic migration; neutral or singly charged vacancies in fact should give rise to higher barriers because of the repulsion of the migrating O ions with the trapped electrons. For the migration of a sub-surface F^"^ vacancy to the surface a barrier of 2.9 eV has been computed with HF cluster models [139], slightly higher than experimental measurements or band structure calculations for the migration in the bulk [141-143]. Much lower barriers have been computed for vacancy migration from high-coordinated to low-coordinated surface sites, 1.6 eV for the migration of an O ion from a terrace to another terrace site, and 0.7 eV from a step site to a terrace vacancy [139]. The barrier of 1.6 eV for the O migration on the surface of MgO is close to that measured at the grain boundary of NiO, 1.8 eV [144]. On the surface, the diffusing oxygen atom from a filled to a vacant site on the (100) terrace follows a curved trajectory and at the midpoint (transition state) is 0.6 A above the top layer [139]. The same path has been found for the diffusion of a neutral oxygen vacancy on the MgO surface from band structure DFT-LDA calculations [20]. In this case a barrier of 2.6 eV has been estimated for the diffusion process, a value which drops to about 2 eV when gradient corrections are taken into account [20]. Since in this case one is dealing with a filled vacancy, it is not surprising that the barrier is higher than for migration involving Fs^"*" vacancies.

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All these results suggest that at high temperatures the O ions can migrate from low-coordinated sites to high-coordinated vacancies leaving an higher concentration of vacancies at the surface, step or comer sites [139]. 3.4.6. Optical spectra Both thermodynamic arguments as well as spectroscopic evidences suggest that oxygen vacancies at the MgO surface localize preferentially at lowcoordinated sites and that F centers at the terrace sites may not represent the dominant species [133]. Another possible indication of the location of F centers in MgO comes from the analysis of optical spectra. Bulk F and F"*" centers give rise to an intense absorption band around 5 eV with two components, one due to F"^ centers at 4.96 eV and one to neutral F centers at 5.03 eV [145]. Much less is known on the optical properties of surface F centers. Optical spectra of MgO surface F centers have been measured on polycrystalline materials [133,146], single crystals [126] and ultrathin MgO films [127,128,147]. Optical measurements on fine powder samples of MgO using diffuse reflectance technique showed a feature at 2.05 eV attributed to Fsc"^ centers [146]. A peak at 2.3 eV has been observed by Henrich et al. [126] on MgO single crystals using electron energy loss spectroscopy, EELS.

Fig. 8. Electron energy loss spectra of 15 ML thick MgO layers, (a) as grown; (b) after Ar"^ sputtering; (c) after additional deposition of 4 ML of Mg; (d) after deposition of 4 ML of Mg and post oxidation with O2 and consequent annealing. Reproduced from ref. [128]. Copyright 1999 Elsevier.

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The feature was tentatively assigned to surface F centers although according to Underhill and Gallon [148] the loss is due to a surface V center. High resolution EELS studies of the defects generated in an ultrathin film of MgO grown on Mo(lOO) show three features at 1.15, 3.58 and 5.33 eV [127]. The 5.33 eV band was assigned to bulk F centers, the band at 1.15 eV to surface F centers, and the band at 3.58 eV to F aggregates [127]. By reactive deposition of Mg in oxygen atmosphere Pfntir et al. [128,147] have studied the optical spectra of MgO films with large oxygen deficiency. They observed characteristic loss peaks at 2.1 and 3.3 eV which can be attributed to color centers due to surface oxygen vacancies, Fig. 8 [128]. It is interesting to note that optical absortion bands in the same regions, 380 and 520 nm, have been observed in the UV-vis spectra of polycrystalline MgO, Fig. 9 [133]. The complexity of the spectra and the difficult interpretation due to the low sensitivity to surface defects, makes a theoretical support highly desirable in order to provide a firm attribution of the observed bands. The calculation of electronic excitations in solids is very challenging, and only recently accurate configuration interaction (CI) calculations have been reported on this problem [28,29]. To this end, finite clusters embedded in ECP and point charges have been used. A calibration of the accuracy of the results, possible for bulk F and F^ transitions where assignments are unambiguous, have shown that CI calculations tend to overestimate the optical transitions of F centers in MgO bulk by about 15% because of limitations in the size of the basis set [28]. 1001

R%

800 1000 1200 1400 1600 1800 2000 wavelength (nm)

Fig. 9. UV-vis spectra of polycrystalline MgO. (a) As prepared; (b) UV-irradiated for 1 h at RT; (c) UV-irradiated for 1 h at RT in presence of D2; (d) UV-irradiated for 1 h at 77 K in presence of D2. Reproduced from ref. [133]. Copyright 1999 Elsevier.

Scaling the computed excitation energies for surface F centers in various sites by 15% the transitions energies and the corresponding assignments are [29]: bulk Fee and F6c"^ ~5 eV, terrace Fsc and Fsc^ ~3 eV, step F4C and F4c^ -2.2-2.5 eV,

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comer Fsc and F3/ -2.1-2.2 eV, Table 3. No evidence of transitions below 2 eV due to F centers has been found. Thus, the surface transitions require photon energies which are ~2-3 eV smaller than for the bulk [29]. The reason is that on the surface the excited state, more diffuse than the ground state, can extend toward the vacuum with less Pauli repulsion with the Mg^"^ ions than in the bulk, where the excited state is confined within the octahedral shell of Mg cations. The computed optical transitions imply that in the EELS experiment [128] not only five-coordinated F centers but also low-coordinated oxygen vacancies are present, in agreement with the results of other spectroscopic investigations [133]. In this discussion the possible role of F centers aggregates has not been considered, despite the fact that it has been suggested time ago that they are responsible for some of the optical bands observed in bulk MgO, a possibility which has been mentioned also for the surface transitions [127]. However, recent first principle band structure calculations have clearly shown that the interaction between neutral vacancies is repulsive [20]. Thus, the mechanisms which may lead to the formation of aggregates of oxygen vacancies are not clear. 3.4.7. Chemical reactivity It is only recently that the reactivity of F centers towards chemisorbed species has been investigated theoretically. Some of these studies address the mechanism and the barrier for molecular dissociation; among the systems investigated CH4 [17], H2 [149], and HCOOH [150]. The general conclusion emerging from these studies is that reactions involving bond breaking which do not occur on the regular surface, quite unreactive, may take place at defects and in particular at oxygen vacancies. Recently, an accurate study of the interaction of H2 with a Fs center located at the (100) terrace has been performed with embedded cluster HF calculations [149]. It has been found that the reaction leads to the formation of Fs(H) centers in two steps, a dissociative interaction followed by UV irradiation which removes a neutral H atom and leaves an electron trapped at the vacancy with an hydroxyl group nearby. However, the high energy barrier estimated for the process, about 1 eV, suggests that the isolated surface anion vacancy is not the most likely site for formation of Fs(H) centers and that other sites, like low-coordinated F centers or divacancies (§ 3.6), are probably more efficient in determining the reaction [149]. An interesting case of reactivity of color centers is that of the interaction with CO, O2 and N2 molecules. EPR studies have clearly shown that exposure of a MgO surface with an high concentration of F centers to these molecules results in the formation of stable or metastable radical anions [151-153]. The general reaction can be written as:

121

Fs^ + X2 -^ Fs'^ + X2

(8)

The formation of the X2' species (X2 = O2, N2, or CO) is unambiguously shown by the analysis of the EPR spectrum and by the hyperfine interactions. This process implies a net charge transfer from the vacancy to the adsorbed molecule, with formation of surface species like the CO" or N2' radical anions which do not exist as free molecules in gas-phase. The process has been investigated in detail in a series of theoretical studies [153-154]. The reaction proceeds via the transfer of one electron trapped in the surface cavity to the empty levels of the adsorbed molecule. The resulting surface complex, X2' / Fs^"^, is bound by electrostatic forces. The electron is completely transferred from the cavity to the molecule which becomes negatively charged and is attracted by the strong electrostatic potential towards the surface. Although the mechanism of the interaction is the same for the three molecules, the details of the energetics are different. O2 removes spontaneously the electrons trapped in the MgO oxygen vacancies to form the stable O2' superoxide anion [94,154]. On the contrary, CO" and N2" form only at finite temperatures and are metastable species [153, 154]. The different behavior can be rationalized in terms of the electron affinities of the three molecules, positive for O2 and negative for CO and N2. In recent years, a particular attention has been devoted to anion vacancies as possible nucleation centers for the growth of metal particles on oxides. MgO (100) is the best studied oxide surface from this point of view. It has been found that several metal atoms interact weakly with the MgO surface [155] giving rise to rapid diffusion and desorption after deposition. The nucleation and growth of metal particles is assumed to occur at defect sites [156], and the oxygen vacancies are among the best candidates. Some ab initio studies have shown indeed that the metal-oxide interaction is much stronger at the F centers than at the regular sites [157-160]. It is still debated, however, if these centers promote nucleation. For the case of Pt it has been proposed that better nucleation centers are the surface OH groups (§ 3.3) and the divacancies (§ 3.6) [112]. F centers have also been shown to activate the supported metal atoms or clusters increasing their catalytic activity in CO oxidation or acetylene trimerization reactions [161-163]. 3.5 Cation vacancies Among intrinsic defects the cation vacancy or V center is of particular interest because it favors the formation of trapped holes at neighboring O ions; at the surface this may result in particularly active O" species which can react and exhibit catalytic activity [164] (see also § 3.8). There are several studies dedicated to the V center in the bulk, both experimental [165,166] and theoretical [14,15,141,167,168], as reviewed by Chen and Abraham [169]. This center can exist in three charged states, V°, V", V^". The ground state of V°

122

corresponds to two holes located on oxygen ions at opposite sides of the vacancy, and the ground state is a triplet. The singly charged V defect is very stable and is the one which is better characterized, while mostly theoretical information exists on the diamagnetic V^' center. Compared to the F centers, much less work has been devoted to the Vs center. One reason is that these centers are supposed to be present in lower concentrations than the O vacancies [170,171] due to their high formation energy. According to various theoretical estimates the energy required to remove a Mg^^ cation from bulk MgO with formation of a V^" center is of the order of a 20-30 eV [14,21,141]. An ab initio HF cluster study of the electronic structure of V centers at the MgO (100) surface [39] has shown that the electrons associated to the V and V^' centers are occupying levels near the top O 2p valence band and that the associated holes are completely localized on the oxygen atoms around the vacancy [39], in analogy with the bulk electronic structure. Very recently, a combined periodic supercell and embedded cluster study has been performed at the HF level on the electronic structure of the cation vacancy at the MgO surface [172]. It has been found that the formation energy of the neutral V^ defect in the bulk, 15 eV, is higher than at the surface, 13.5 eV, which means that migration to the surface should be thermodynamically favored. The holes connected to the formal removal of a neutral Mg atom are localized on a p orbital pointing toward the cavity [172]. Except for a study of the interaction with isolated metal atoms [157], to the best of our knowledge there are no theoretical investigations of the reactivity of surface V centers in MgO. Clearly, this is a center which should be considered in more detail in future investigations of MgO surface defects. 3.6 Divacancies Early experiments on the EPR spectrum of paramagnetic centers at the surface of MgO by Lunsford and Jayne [130] have been interpreted in terms of an electron trapped at a surface divacancy because of the analogy with the spectrum of a Fc"^ center, attributed to a divacancy with a trapped electron in bulk MgO [173]. One of the characteristic features of the spectrum which lead to the conclusion that the unpaired electron occupies a Mg and O divacancy and not a single vacancy site is the anisotropy in the values of the hyperfine coupling constants with the ^^Mg ions. Recently, this asymmetry has been explained as due to the Fs'^(H) centers, an oxygen vacancy in proximity of an adsorbed proton [125], as described above. Nevertheless, one cannot rule out the possibility that divacancy centers exists at the surface of MgO. There are several arguments in favor of the divacancy [174]: (a) it is energetically much easier to form than a pair of isolated cation and anion vacancies; (b) due to the mobility of cation vacancies, recombination of isolated vacancies should occur at high temperatures, > 1000 K, the typical temperatures reached in the thermal

123

treatment of polycrystalline MgO; (c) the divacancy is intrinsically asymmetric, a fact that would easily explain the anisotropy of the EPR signal; (d) the divacancy is a neutral center and can be present in high concentrations without the need for compensating charges. These reasons have stimulated in the last few years a growing interest around this center [174] and its reactivity [112,175]. The divacancy is sometimes referred to as the mini-step because the removal of a MgO unit from the surface results in four-coordinated ions, Fig. 10. Two different types of divacancies can be formed, the pit, where a surface Mg and the O ion underneath are removed, and the tub, where both ions are removed from the first layer, Fig. 10. It has been found [174] that the divacancies are relatively stable defects which can form through the recombination of isolated vacancies with an energy gain of up to 12 eV. The order of stability is similar as for isolated vacancies, buUc < surface, so that migration to the surface is expected at high temperatures. Neutral divacancies have an appreciable electron affinity of almost 1 eV; this means that electrons can be trapped at these defects forming a paramagnetic center. The affinity for a second electron however is negative. This is different from single O vacancies; both Fs^"^ and Fs"^ vacancies have positive electron affinities and transform into the diamagnetic Fs^ center by electron trapping. Recently, however, it has been suggested based on spectroscopic measurements [164] that the paramagnetic divacancy, with a trapped electron, cannot be transformed into a diamagnetic center by further electron trapping, indicating that a divacancy with two trapped electrons is unstable in full agreement with the theoretical prediction.

Fig. 10. Schematic representation of a divacancy on the MgO surface. Top: pit conformation; bottom: tub conformation; x and x' indicate the positions of the removed atoms. Reproduced from ref. [174]. Copyright 1998 American Institute of Physics.

124

Of course, relaxation of neighboring ions makes an important contribution to the relative stability of neutral and charged divacancy centers. Thus, from a thermodynamic point of view the divacancy has been confirmed as an important defect center on the surface of MgO [174] and recent experimental data on the formation of surface O" ions by bleaching color centers with N2O point to the important role of these defects [164]. The study of the hyperfine coupling constants, however, resulted in a too small anisotropy of the signal, compared to experiment. In this respect the assignment of the observed EPR spectra [125] to surface divacancies with a trapped electron is not conclusive [174]. A detailed theoretical study of the reaction of a surface divacancy with H2, aimed at the identification of the Fs"^(H) color center observed in EPR, also come to the conclusion that the calculated magnetic parameters show discrepancies with the experiment [175]. Despite these open questions, there is no doubt that this defect center should be considered with great attention in future analysis of the reactivity of the MgO surface. 3.7 Impurity atoms Impurity atoms in MgO can be either cations replacing Mg^"^ or anions substituting O^'. In either case the surface reactivity is significantly modified by their presence. In the case of cation impurities, these can be of at least two kinds. If Mg^"^ is replaced by another cation with the same formal charge, e.g. Ca^"^, Ni^^ Co^\ etc., the main changes in the local electronic structure are connected to the different size of the ion, to small modifications in the chemical bonding with local increase or decrease of the covalent character, and to the presence of filled d orbitals. In fact, it has been shown by ab initio HF cluster calculations that, in bulk MgO, a cation like Ni^^ is stable against disproportionation into the +1 and +3 states [176]. If, on the other hand, a cation with a different formal charge e.g. Na"^ or Li"*^, substitutes Mg^"*^, then a deep change in the electronic structure occurs to compensate the missing charge. For monovalent cations this is the change in the anion formal charge from -2 to -1; the O' species, however, is a radical and its chemistry is fundamentally different from that of the corresponding fully reduced species, as it will be discussed in more detail in § 3.8. Let us consider first isovalent cations. MgO can dilute ions of similar size, as for instance Ni^"^ or Co^"^ forming NiO-MgO and CoO-MgO soUd solutions with an infinite range of composition. The effect of progressively replacing Mg^"^ by Ni^^ or similar cations (Co^"^, Cu^"^) on the surface properites has been investigated both experimentally [83,177,178] and theoretically [179,180]. The presence of Ni^^ cations diluted in the MgO matrix results in an efficient catalyst for nitrous oxide, N2O, decomposition; this has been attributed to the different bond strength of the Ni-0 and Mg-0 bonds at the surface [177]. Plane wave calculations on Ni-doped MgO have shown that the presence of Ni atoms on the

125

surface enhances the bonding of S-containing molecules like H2S through electronic states associated with the Ni 3d levels [181]. To check the role of the 3d orbitals of the TM cation in the chemistry of the surface the vibrational properties of adsorbed CO have been considered [83,178]. In fact, assuming that chemical forces do not play any role so that Ni^"*" and Co^"^ ions diluted in a MgO matrix experience an ionicity very similar to that of Mg^^, the frequency of CO adsorbed on Mg^"*" or Ni^"*" and Co^"^ should be practically indistinguishable. The experiments show a different situation. The peak of CO adsorbed on Ni^"*" falls at a distinctly lower frequency than that of CO on Mg^^ (2130 cm"^ versus

00

Fig. 11. Difference density maps for CO adsorbed on Mg "^ ions on a MgO (100) surface, (a) contribution from a orbitals; (b) contributions from n orbitals. Continuous lines correspond to charge accumulation, dashed lines to charge depletion with respect to the superposition of the non interaction CO and MgO densities. Reproduced from ref. [180]. Copyright 1993 Elsevier. b

( _,/' --W^^^^-

\

i

% ^--^—-' '''--'' '^—-' v,^ % %

y

Fig. 12. Difference density maps for CO adsorbed on Ni ^ ions on a MgO (100) surface, (a) contribution from a orbitals; (b) contributions from 7C orbitals. Continuous lines correspond to charge accumulation, dashed lines to charge depletion with respect to the superposition of the non interaction CO and MgO densities. Reproduced from ref [180]. Copyright 1993 Elsevier.

2144 cm'^ [83]); furthermore, even the intensity of the peak is not the same, and for Ni^'*"-CO is larger than expected on the basis of the stoichiometry. Similar results have been obtained for Co. A rationalization in terms of a small, but non

126

negligible metal dJi-CO 271* interaction has been proposed based on DFT cluster model calculations [180], see Fig. 11 and Fig. 12. If the role of the d-7i overlap is quite modest in the Ni^^—CO surface complex, this is no longer true when a stronger 7C-acceptor molecule like NO is used as a probe. In fact, on MgO at 77 K NO is so weakly bound that that only lateral interaction products form like the (N0)2 dimers; when Ni^"^ ions are diluted in the material, Ni^"^-NO complexes form and remain stable even at room temperature [178]. While the effect of cation impurities on the surface chemistry of MgO has been investigated in detail, very little is known about anion substitution. Defect formation and excitation energies for S^' and Se^-doped bulk MgO have been calculated [182,183] but there are no data for the surface. In the bulk it has been estimated that the presence of S^" or Se^" impurities result in a outward relaxation of the Mg^"*" neighbors of 6% and 8%, respectively [182]. A recent report of the 0^"-S^" exchange reaction on MgO has been reported [184]. The reaction involves adsorption of CS2 on MgO powders and the subsequent exchange reaction with formation of COS and of S^" ions probably located at the low coordinated sites. It has been found that the basicity of the MgO surface doped with sulfur ions is drastically modified with respect to that of pure MgO [184]. The inclusion of impurity atoms in MgO is much more interesting from a chemical point of view when alkali metals are used to replace Mg ions. In fact, this results in trapped-hole centers. The M'^/O' pairs have been extensively studied in the bulk of alkaline-earth oxides by optical studies, EPR and ENDOR measurements [185,186] as well as by embedded cluster calculations [187]. The LiVO" ions create an effective dipole which polarizes the surrounding lattice, with the two ions moving toward each other. The presence of an O' radical, however, is most interesting when one is dealing with surface properties. This center in fact is very reactive and is the subject of the next paragraph. 3.8 C radical anions Paramagnetic O' anions located at the surface of ionic crystals are capable of radical cleavage processes, such as hydrogen abstraction from methane and hydrocarbon molecules adsorbed on the surface. They can also facilitate UV induced homolytic splitting of hydrogen. This particular reactivity of the O" anions has been studied extensively by means of EPR and TPD [188]. In the previous discussion we have already encountered the O' ion in connection with other defects. In fact, a neutral cation vacancy in MgO, Vs, corresponds to two trapped holes on two surface O anions; in the charged Vs" center only one O' radical is present. In the previous section we have seen that doping an alkalineearth oxide with Li or Na results in M V O ' pairs. Indeed, some of the most efficient catalysts to promote the conversion of methane into higher hydrocarbon derivatives are based on alkali-doped metal oxides, including Na/CaO and Li/MgO. At high partial pressures of methane, gas-phase oxidation of methane

127

with O2 to give coupled products such as C2H4 or C2H6 proceeds with high conversion rate at temperatures above 500 °C [189]. At low methane partial pressures the catalyst is required to give acceptable conversion rate and selectivity. The reaction proceeds via C-H bond cleavage and the centers that facilitate C-H bond breaking are unlikely to be present on a perfect MgO surface. The surface activity has been attributed to the presence of O' radical ions formed upon doping with Li. It has been proposed that methane is adsorbed at the surface and hydrogen abstracted by the O" radicals, with formation of methyl radicals [190], although there are reports that the F centers do also play a role [191]. However, the mechanisms proposed to explain all the processes occurring in real catalytic conditions comprise more than 100 steps, many of which not yet completely clarified. The elementary steps of the reaction have attracted considerable theoretical interest in recent years [192-195]. The studies have been performed at various computational levels, HF, MP2, DFT [193-195], and explicitly correlated wavefunctions (MCSCF and CI) [192] and the surface has been represented by embedded clusters [192,194,195] or by periodic slabs [193]. All the calculations agree with the fact that the LiVO' complex is the active center although most recent studies suggest that other surface species may also play a role [195]. The activation energy for the C-H bond breaking is very dependent on the method used but also on the proper inclusion of surface relaxation along the reaction profile. The values range from 4-6 kcal/mol [192] to 18 kcal/mole [193] although the best estimates are probably consistent with a low barrier [194]. The O" ions can also be produced on the MgO surface by other methods, e.g. by excitation of low-coordinated surface anions, Oic, by UV light [164,196198] or by filling anion vacancies with paramagnetic O' species [164,199]. In the first case the process involves an ionization or a charge transfer process: Oic^- + hv -> Oic" + e-

(9)

where the electron is released to an electron acceptor; under O2 atmosphere the electron can be captured by the oxygen molecule with formation of the superoxide anion, O2'. In UHV conditions, short lived 0-Mg"^ pairs form in the excitation step. Recently, the energy required to ionize O^' anions from the MgO surface have been determined theoretically from DFT cluster calculations [50]. The computed IP of a five-coordinated anion is 6.7 eV, in excellent agreement with an experimental estimate of 6.7±0.4 eV [200]. On a O3C anion at a comer site the IP decreases to 5.6 eV [50]. Thus, it is conceivable that UV irradiation will preferentially produce O' radicals at low-coordinated sites. The second mechanism to generate O" radicals on the MgO surface implies first the creation of color centers and then their bleaching with N2O:

128

(10)

Fs"" + N2O - ^ Os" + N2

In this case the decomposition of the reactive N2O molecule results in the filling of a paramagnetic vacancy site with release of the unreactive N2 molecule. Clearly, the presence of O" species at the surface of MgO is not disconnected from the existence of other defects, low-coordinated anions or O vacancies. In fact, the complex interconversion of one center into another one is one of the reasons for the difficult identification of defect centers on oxide surfaces. 3.9 (111) microfacets In contrast to nonpolar (100) surfaces of rock-salt structured ionic materials, which are thermodynamically stable, (111) polar surfaces are unstable if they remain in the bulk terminated structure due to the diverging surface potential [2,201] and undergo substantial surface reconstruction. It has been predicted that the (111) surfaces of NaCl, NiO, MgO, etc. reconstruct from the p(lxl) bulk terminated into a p(2x2) structure. Fig. 13 [201]. This kind of reconstruction has been indeed observed on NiO(lll) grown on Au(lll) [202]. Polar surfaces can also be stabilized by the presence of oxygen vacancies, impurities, hydroxyl groups [203] or in general molecular adsorbates [204,205]; on NiO, by simple heat treatment, the OH groups are desorbed and the surface reconstructs exhibiting a diffuse p(2x2) structure [203].

PlirM;?!! -.-m^m^^m. -^^KT^B- ^^f^m^^m

4^B'^V

dHi^V

t^l^V'''

Fig. 13. Structure of a p(2x2) reconstructed MgO(l 11) surface.

Recently, the surface morphology and faceting of MgO(lll) surfaces has been studied with atomic force, scanning and electron microscopies [206,207]. It has been suggested that oxygen trimers reminiscent of cyclic ozone appear on the reconstructed surface [207]. A p(2x2) reconstruction leads to the formation of a series of micropyramids, and it is possible that local arrangements of ions at the surface of MgO assume a similar morphology. A complete account of the

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reconstruction of the polar surfaces of ionic crystals is outside the scope of this review, but here we simply want to mention the possible existence of microensembles of ions with (111) microface morphology. It has been suggested that a fundamental defect in MgO consists in a missing cation from a cube comer, leaving a (111) microface with three oxygen ions and one hole. The electronic structure of this center was then investigated theoretically in a pioneering study of Kunz and Guse [208]. The involvement of (111) microfacets was invoked also in a thermal desorption and isotope exchange reaction study of CO and CO2 on MgO where the condition needed for double exchange was identified by means of HF cluster calculations in small ensembles of three-coordinated O ions [209,210]. 4. CONCLUSIONS In this brief account we have listed a number of defects occurring on the surface of ionic oxides, in particular MgO. The experimental and theoretical activity aimed at the characterization of these defects has become quite intense in recent years, also thanks to the great advances in the preparation of oxide films in UHV under controlled conditions. Still, a detailed understanding of the relations existing between preparation methods and nature and concentration of surface defects is missing. Complex mechanisms control the interconversion of one defect into another and even the nature of a simple defect like the oxygen vacancy is still under discussion. A real control on the microscopic aspects of the chemistry of oxide surfaces will be possible only when the chemistry of defects, vacancies, irregularities will be completely identified. In this respect, ab initio calculations provide an important tool to interpret and complement the experimental information. In particular, local cluster models are particularly useful because of their computational simplicity and of the possibility to compute a great variety of observable properties, like IR, optical, photoemission NMR and EPR spectra. The design of "realistic" models of the surface defects and their validation through the calculation of observable quantities represents probably the best way to solve the problem. Acknowledgments I am indebted to A. M. Ferrari (Torino), T. Bredow (Hannover), C. Sousa and N. Lopez (Barcelona), L. Giordano, R. Soave and D. Ricci (Milano) for their substantial help in the study of defects in oxides. The cooperation with the groups of F. lUas (Barcelona), E. Giamello and C. Pisani (Turin), N. Rosch (Munich), and U. Heiz (Losanne) has been invaluable to elucidate many open questions and to link experimental with theoretical results. Financial support through the INFM Projects PAIS and PRA-ISADORA is gratefully acknowledged.

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Chapter 4

Theory of physical and chemical behavior of transition metal oxides: vanadium and molybdenum oxides K. Hermann" and M. Wîtko*' ^Fritz-Haber-Institut der Max-Planck-Gesellschaft, Faradayweg 4-6, 14 195 Berlin, Germany. Înstitute of Catalysis and Surface Chemistry, Polish Academy of Sciences, ul. Niezapominajek, 30239 Cracow, Poland.

1. INTRODUCTION Transition metal oxides are well known for their large variety of physical and chemical properties [1-3]. Many of these materials undergo phase transitions with interesting structural, electronic and magnetic behavior [3]. Some exhibit high temperature superconductivity and exciting optical properties or high catalytic activity [2]. Among these, vanadium and molybdenum oxides represent an important class of materials due to their large variety in crystal structures and physical/chemical properties. They cover a wide range of electronic properties from metals to semiconductors and insulators and are, therefore, used in many technological applications. Examples are vanadium oxide based electrical and optical switching devices, write-erase media, light detectors, critical temperature sensors, infrared spatial light modulators [4-9], and even vanadia constituents of surfaces of medical Ti-Al-V implants [10]. Further, molybdenum trioxide forms n- and p-type semiconductor phases [11] which are of technological interest. The rich and diverse chemistry of these oxides results mainly from two interrelated factors. First, the metal ions can adopt different formal oxidation states in the oxides, ranging from +2 to +5 in the case of vanadium and from +4 to +6 in the case of molybdenum. Second, the ions can exhibit quite different

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coordination geometries described by octahedra, pentagonal bipyramids, square pyramids, or tetrahedra, where the different local units share comers and/or edges and/or faces. Many of these oxides are described by a defined composition (single or mixed valency oxides) while others exhibit a wide range of compositions. Vanadium and molybdenum oxides are also extremely interesting from a chemical point of view because of their wide exploitation in many catalytic processes of industrial relevance. Among several factors responsible for the catalytic behavior we mention the mobility of surface/lattice oxygen atoms, the existence of Lewis acid-base sites as well as different reactivity of different crystal faces [2, 12-15]. The versatility of the properties of both oxides, when combined with other elements such as bismuth, cobalt, aluminum, or alkali metals, leads to their use as active and selective catalysts in many reactions belonging to redox or acid-base processes. Redox processes include reactions where oxygen is involved as well as those where only hydrogen participates. Examples of the former group are selective oxidation, oxidation of different functional groups, oxidative condensation or dehydrogenation or dehydrocondensation, oxyhydration, or epoxidation. Examples of the latter are hydrogenation, hydrogenolysis and reduction. Among acid-base reactions the most important are isomerization (both structural isomerization and ring contraction), disproportionation (metathesis), decomposition (dehydration, dehydrogenation, dehydrocyclization, dehydrocondensation) and etherification. Vanadia-based compounds are used as components of various catalysts in mild oxidation, ammoxidation, and dehydrogenation of hydrocarbons and other organic compounds [16]. They are also efficient catalysts in oxidation of S02to SO3, naphthalene or oxylene to phthalic anhydride and more recently n-butane to maleic anhydride [16]. Vanadia-based catalysts seem also promising for the oxidation of toluene to benzaldehyde, methanol to formaldehyde and to methyl formate, as well as for the removal of NOx by selective reduction with NH3 [4, 17]. Molybdenum oxides are used, either as pure materials or in combination with other elements (Bi, V, Co, Al) as highly active and selective catalysts for many reactions of very different type. Examples of commercial processes are isomerization, polymerization, and production of formaldehyde and acronitrile [2, 12-13]. In particular, molybdenum trioxide, M0O3, serves as a selective catalyst for the partial oxidation of hydrocarbons and alcohols [18-32]. Its catalytic behavior results from several interrelated factors, where the valence of the Mo ions together with their local environment and the type of exposed crystal plane are the most important. Each of the differently oriented M0O3 surfaces contains different active oxygen sites that participate in various elementary steps of catalytic reactions. The role of these surface sites has become the subject of lively discussion in recent years, see Ref. [12] and references therein. For

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example, in selective oxidation of propene to acroleine [19-27] two subsequent conversion steps proceed on different crystal planes of M0O3 [22]. Hydrogen abstraction takes place at the basal (001) or (100) crystal planes whereas nucleophilic addition of oxygen occurs at the basal (010) plane [25, 26]. In both steps, surface oxygen sites form Lewis centers with strong basic character [2, 33] leading to surface hydroxyl or water species and to the oxygenated precursor of aldehyde / ketone [33]. In typical catalytic reactions involving vanadium or molybdenum oxides (as well as oxides of other transition metal elements of groups V, VI, and VII) the oxide surfaces undergo reduction and oxygen vacancies are formed [1-3, 12, 14, 34]. When the concentration of these point defects surpasses a certain critical value at the surface they propagate into the substrate bulk. This may give rise to crystallographic shear (CS) planes [1, 14] where the linkage between polyhedra is changed towards higher connectivity, e.g. from comer- to edgesharing, which results in effect in annihilation of vacancies. The CS planes constitute extended two-dimensional defects in the bulk. As a consequence, the crystal can be considered as being composed of slabs of comer-sharing polyhedra of more or less undistorted parent stmcture, separated by CS planes. If the spacing between the CS planes is regular, the parent stmcture is transformed into a mixed valency oxide. Depending on the actual spacing different intermediate oxides may be formed which are usually described as homologous series. As an example we mention the series McnOsn-p appearing in ReOs like stmctures. Here, n is the width of the slab of parent stmcture as given by the number of coordination octahedra in unbroken rows between CS planes, and p measures the number of anionic sites eliminated by the crystallographic shear. Despite the enormous importance of vanadium and molybdenum oxides as catalysts many details of their microscopic behavior are still under debate [4, 12, 16]. This is due to the large structural complexity of these materials as well as to their diverse electronic behavior that makes quantitative studies, both experimental and theoretical, rather challenging. In the present chapter, which deals with theoretical concepts applied to vanadium and molybdenum oxide surfaces, we will restrict the discussion to binary oxide systems. So far, mixed metal oxide systems have not been studied by quantitative theory. Theoretical methods that have been used to study oxide surfaces can be classified according to the approximations made in the system geometry where two different concepts are applied at present, local cluster and repeated slab models. Local cluster models are based on the assumption that the physical/chemical behavior at selected surface sites can be described by finite sections cut out from the oxide surface. These sections (surface clusters) are treated as fictitious molecules with or without additional boundary conditions to take the effect of environmental coupling into account. Therefore, their electro-

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nic structure can be calculated by modem quantum chemical methods. Here methods using various ab initio density functional theory approaches [35] have proven to be particularly successful in recent work. Repeated slab models [36] rely on an exact two-dimensionally periodic arrangement of all atoms and molecules at the oxide surface. Within this approach, a slab of full translational periodicity and finite thickness (surface slab) describes the surface system. For computational convenience surface slabs are repeated perpendicular to their surface with vacuum separating adjacent slabs. This yields an altogether threedimensionally periodic system with a large supercell which can be studied by modem bulk methods of solid state theory. Here ab initio density functional theory methods, such as the full potential linear augmented plane wave (FPLAPW) method [37-39] have been applied exclusively. In the following, we discuss recent concepts and theoretical results concerning microscopic properties of vanadium and molybdenum oxide surfaces. While the two elements form different classes of oxides their surfaces exhibit numerous stmctural and electronic similarities, such as microscopic surface binding, adsorption, or oxygen vacancies, which we will point out accordingly.

2. VANADIUM OXIDES Binary oxides of vanadium, VxOy, form a large family of materials with different stoichiometry where vanadium ions differ by their formal valence charge as well as by their local coordination. Single valency oxides, where all vanadium ions occur with the same valence charge, include four oxides, VO, V2O3, VO2, and V2O5. The monoxide VO appears as a cubic (rocksalt) lattice [40, 41] and vanadium appears with a formal charge +2. The sesquioxide V2O3 forms an antiferromagnetic monoclinic phase [42] below 150 - 170K and a paramagnetic trigonal comndum phase [43-45] yielding in both phases a formal charge +3 for vanadium. The dioxide VO2 crystallizes in a monoclinic (distorted rutile) phase [46-50] below 337K and in a tetragonal mtile phase [47, 51-55] resulting always in a formal charge +4 for vanadium. In the pentoxide V2O5, which forms an orthorhombic layer-type lattice [56-58], vanadium assumes its highest formal valence charge, +5. In addition to single valency oxides there are many phases of intermediate composition where vanadium ions of different and even fractional valence charge occur. These mixed valency oxides may have well defined composition and bulk crystal structure. Examples are Magneli phases, which form homologous series Vn02n-i or Vn02n+i and have been associated with shear plane formation [59, 60]. Single crystals of Vn02n-i have been observed for 2 < n < 9. In these systems the average valence charge of vanadium is +(4-2/n), ranging

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between +3 and +4, which is described by ion ratios V'^/V^"^ = (n-2)/2. Apart from the single valency oxide V2O3 (n = 2) examples are monoclinicVsOs [61, 62], triclinic V4O7 [63-65], V5O9 [66, 67], V6O11 [68], V7O13 [68], and V9O17 [69]. Single crystals of Vn02n+i oxides have been observed for 2 < n < 6. Here the average valence charge of vanadium is +(4+2/n), ranging between +4 and +5, which is described by ion ratios V'^/V^'^ = (n-2)/2. Apart from the single valency oxide V2O5 (n = 2) examples are monoclinicVsOv [70], orthorhombic V4O9 [71], and YeOu where a monoclinic low temperature phase [72-74] and an orthorhombic phase at high temperatures [75] is observed. Further, single crystals of monoclinic V12O26 [76] and of monoclinic V14O6 [77] have been identified in the experiment. Point defects in vanadium oxides, such as oxygen vacancies, can order themselves and give rise to superstructures or complexes. This has been observed in several oxides (VO with about 20% oxygen vacancies being a good example [1]). Further, superstructures, as e. g. V52O64 or V244O320 [1], are formed under oxygen-rich conditions where the oxygen sublattice is almost complete. However, some vanadium ions occupy tetrahedral sites where each tetrahedrally coordinated cation has four vacant octahedral sites as nearest neighbors and the cluster is topologically similar to the rocksalt structure. In addition to bulk material, vanadium oxide cluster ions, such as V20(4.6/, V30(6V. V40(8-ii)\ VsOdi.n/, V60(i3-i5)\ and YjO^.e-m" [78, 79], can be prepared in gas phase. The chemical reactivity of these species shows a distinct dependence on cluster size. While the smaller clusters are reacting quite easily with other molecules, the reactivity of the larger systems is decreased with the exception of oxygen-rich clusters, which can release molecular oxygen upon collision with reactant gas. All of the vanadium oxides described above are interesting in their physical/chemical properties and may be used in many applications. However, the following discussion will be restricted to single valency oxides that have been dealt with in quantitative theoretical studies. 2.1. Vanadium oxide bulk systems 2.1,1. Geometric structure Single crystals of vanadium oxide can be classified in their geometric structure, apart from their lattice definition by lattice and lattice basis vectors, by the occurrence of specific elementary VOx building units. All single valency oxides are characterized by a common octahedral V06 unit. The different crystal lattices differ by the distortion of the octahedral units (as a result of actual V-0 distances and 0-V-O angles) as well as by the connectivity of the units in the

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crystal (linking of comers, edges, or faces). This becomes evident in the ballsand-sticks models of the four oxides, VO, V2O3, VO2, V2O5, shown in Fig. 1. In the monoxide VO, which forms a cubic rocksalt lattice [40, 41] with two atoms (1 V + 1 O) in the crystal unit cell, the VOg octahedron is regular with all distances dv-o = 2.06 A and angles Z(OVO) = 90°, see Fig. la. Here all adjacent VOe units are connected by their edges where oxygen centers share six octahedra each and are six-fold coordinated with respect to their vanadium neighbors. The dioxide VO2 exits as a tetragonal rutile crystal above 337K [47, 51-55J with six atoms (2 V + 4 O) in the crystal unit cell. This structure is distorted slightly by shear, becoming monoclinic below 337K [46-50, 80] and doubling its crystal unit cell (twelve atoms, 4 V -h 8 O). In the tetragonal lattice the regular V06 octahedron is elongated sightly along its 4-fold axis with four (two) distances dv-o = 1-90 A (1.95 A) and all angles Z(OVO) = 90°. In the monochnic lattice the VOg octahedron is distorted yielding three distances dy-o each in the range 1.76 - 1.87 A and 2.01 - 2.05 A, respectively and angles Z(OVO) varying between 78° and 99°. In both crystal lattices adjacent VOg units are

Fig 1 Geometric structure of (a) cubic bulk VO ((100) netplane stacking), (b) monochmc bulk VO2 ((Oil) stacking), (c) trigonal bulk V2O3 ((0001) stacking), and (d) orthorhombic bulk V2O5 ((010) stacking). Vanadium and oxygen centers are shown as light large and dark small balls, respectively. Octahedral VOg and bipyramidal VzOg units are indicated by black lines.

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connected by their edges or comers, see Fig. lb. The oxygen centers share three octahedra each and are three-fold coordinated with respect to their vanadium neighbors with distances varying by 3% (tetragonal) and 10% (monoclinic) respectively. The sesquioxide V2O3 crystallizes as a monoclinic phase [42] below 150 170K and as a trigonal corundum phase [43-45] above this temperature with ten atoms (4 V + 6 O) in the crystal unit cell. In the trigonal lattice the VOe octahedron is distorted with three distances dv-o each at 1.96 A and 2.06 A, respectively and angles Z(OVO) varying between 82° and 97°. Here adjacent VOe units are connected by their edges or comers, see Fig. Ic. The oxygen centers share four octahedra each and are four-fold coordinated with respect to their vanadium neighbors with distances varying by 5% in the trigonal lattice. The pentoxide V2O5 forms an orthorhombic layer-type lattice [56-58] with 14 atoms (4 V + 10 O) in the crystal unit cell. In this crystal lattice the WOe octahedra are distorted substantially with one small distance dy-o = 1.59 A and one very large value, dy-o = 2.79 A, while the other V-0 distances range beo

o

tween 1.78 A and 2.02 A. Angles Z(OVO) involving the vanadyl oxygen at short distance amount to 105° while those including the farthest oxygen lie between 73° and 77°. As in V2O3, adjacent V06 units are connected by edges or comers. However, in V2O5 three oxygen centers share three and two octahedra, respectively. This results in oxygen centers being singly, two-, and three-fold coordinated. The large distortion of the octahedra actually suggests the use of V2O8 bipyramids as building units to describe the lattice stmcture which reflects the layer-type lattice more appropriately. The oxide layers of bulk V2O5 shown in Fig. Id extend parallel to the (010) netplanes of the lattice. This experimental stmcture of bulk V2O5 [56-58] has been verified recently in theoretical total energy [81-83] as well as molecular dynamics studies [84, 85]. Note that depending on the choice of the orthorhombic crystal axes the layer net plane orientation may also be denoted by (001). The latter corresponds to an interchange of lattice constants b and c as proposed in Ref. [86]. In this chapter we have adopted an altemative nomenclature used e.g. in Ref. [87]. 2.7.2. Electronic properties The electronic stmcture of bulk vanadium oxides is determined to a major extent by the amount of d electron occupation in the vanadium ions. In the ideal three-dimensional periodic bulk, electrons are described by Bloch states with energy dispersions reflected in band stmctures and corresponding densities of states (DOS). These quantities can be calculated with high accuracy by modem band stmcture and total energy methods based on the density functional theory (DFT) method.

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The vanadium atom is described in its ground state by an electron configuration [Ar] 3d^ 4s^. Therefore, vanadium cations V^^ for 2 < p < 5 are given by a configuration [Ar] 3d^"^ resulting in an empty (p = 5) or partially filled 3d shell (p < 5). This result of the free ions is also reflected in the electronic ground states of the bulk oxides where vanadium possesses a formal charge +p. As a consequence, the amount of occupied V 3d electron states is expected to increase in going from V2O5 (corresponding formally to V^^ ions) to VO (formally V^"^ ions). This effect becomes obvious in results from recent band structure studies on the single valency oxides of vanadium discussed in the following. The electronic structure of bulk V2O5 has been studied extensively by both semi-empirical and ab initio techniques. This includes periodic Hartree-Fock (HF) pseudopotential calculations [88], tight binding studies within a perturbative approach [89], non-empirical atomic orbital method studies [90], ab initio orthogonalized linear combination of atomic orbitals (OLCAO) work [91], and self-consistent augmented spherical wave (ASW) [82, 83] as well as ab initio full-potential linear augmented plane wave (FP-LAPW) calculations [81]. These studies yield overall good agreement with the measured photoemission and optical properties. As an example, Fig. 2 shows the band structure and total density of states (DOS) of geometry optimized bulk V2O5 obtained by ab initio FPLAPW calculations [81]. The electron states with negative energies between 0.0 and -5.5 eV describe the occupied valence band region, characterized by dominant oxygen 2p and rather little vanadium contributions. The integration of the DOS over this valence band region yields close to 60 electrons per unit cell (containing 10 oxygen ions) which is consistent with a 2p^ valence configuration of ionic oxygen O^". The conduction band region starts at about 2 eV above the highest occupied electron level (defined by 8 = 0 in Fig. 2) and is determined dominantly by (empty) vanadium 3d type electron states. This indicates a 3d^ valence configuration of ionic V^"^ and confirms altogether the ionic nature of a (V^"^)2(0^")5 compound suggested by formal valence charges. However, a detailed binding analysis based on the electronic density and wavefunctions [81] shows clearly that the actual amount of atom charging is smaller than the formal valence charges indicate. As a consequence, V-O binding in bulk V2O5 is determined not exclusively by ionic coupling but contains also sizeable covalent contributions. This will be discussed in more detail in Sec. 2.2.2. The energy region of the calculated V2O5 valence bands is found to have a width of 5.5 ± 0.5 eV, cp. Fig. 2. This matches very well with values from angle-integrated [92] as well as angle-resolved ultra-violet photoemission spec-

144 6.04.0\^

2.0-

£

.0-

z

LU

-2.0-4.0rx -6.0-

uz

r

u R

Y

r

0.0

15.0

k point path

orthorhombic BZ

Fig. 2. Band structure and total DOS of bulk V2O5. The energy bands are shown for characteristic paths connecting high symmetry points of the irreducible part of the orthorhombic Brillouin zone (BZ) which is included at the bottom. All energies e(k) are taken with respect to that of the highest occupied state. The DOS is given in states per unit volume and per eV.

-7.0

-5.0

-3.0

1.0

Energy [eV]

Fig. 3. Angle resolved He-II ultraviolet photoemission spectrum of a ViOsCOlO) surface sample taken at normal incidence from Refs. [93, 94].

145

troscopy (ARUPS) [93, 94] experiments on V2O5. As an example, Fig. 3 shows recent ARUPS data for single crystal V2O5 using He-II light at normal incidence with respect to the (010) oriented surface [93, 94]. The emissionspectrum has an overall width of 6 eV and can be described by a superposition of mainly three peaks. A major central peak lies at about -3.2 eV below the onset of the spectrum (taken as the energy zero in Fig. 3) and two satellite peaks are found at -1.2 and -4.3 eV, respectively. The nature of these peaks is connected with the different coordination of the oxygen ions in V2O5 as will be discussed in detail in Sec. 2.2.2. From photodesorption experiments V2O5 is known to be a semiconductor with a visible band gap of Eg = 2.35 eV [95] at low temperatures. This is confirmed by values obtained from other measurements, for example, optical absorption (Eg = 2.30 eV) [96] and optical reflectance measurements (Eg = 2.38 eV) [97]. The FP-LAPW calculations [81] yield a direct band gap of 2.3 eV at F and an indirect gap (corresponding to a transition R to F, see Fig. 2) of 1.9 eV which seems to be in good agreement with the experimental gap values obtained for V2O5. However, this agreement, found also in other theoretical studies [82, 83, 98], is considered fortuitous. Bandstructure methods, using DFT schemes within the local density approximation (applied in the theoretical studies), are not expected to describe electronic excitations to a high accuracy [99]. There are two narrow split-off conduction bands at about 2 eV above the highest occupied electron level and about 0.5 eV below the upper conduction band range which have been discussed in the literature [83, 91, 100]. These bands are characterized by rather small dispersion widths corresponding to large effective electron masses and localized band states. The latter are described by dominant V 3dxy character (t2g symmetry) with small O 2px,y admixing. The presence of these split-off conduction bands hints at possible excitations to localized electron states and is of interest for conductivity as well as optical absorption experiments on bulk V2O5 [83, 91, 100]. The electronic structure of bulk VO2 has been examined in several theoretical studies [101-105] in connection with a structural phase transition at 337K (monoclinic to tetragonal rutile). This is accompanied by a semiconductor-tometal transition. As an example. Fig. 4 shows the total densities of states (DOS) of bulk VO2 for the monoclinic (left diagram) and tetragonal rutile (right diagram) lattice geometry obtained by ab initio FP-LAPW calculations [105]. Here the experimentally known geometries of the two bulk phases have been used, see Sec. 2.1.1. The electron states with negative energies (between -1.5 and -7.5 eV in the monoclinic and between -1.9 and -8.0 eV in the tetragonal rutile geometry) represent an occupied valence band region described by dominant oxygen 2p and almost no vanadium contributions. This is completely analogous to the valence band region found for V2O5 and reflects a 2p^ valence

146 10.0

5.0

10.0

DOS [a.u.:

15.0

5.0

DOS [a.u.]

10.0

Fig. 4. Total DOS of bulk VO2 for the monoclinic (left diagram) and tetragonal rutile (right diagram) lattice geometry. The energy zero coincides with the energy of the highest occupied state. The DOSs are given in states per unit volume and per eV.

configuration of ionic oxygen O^'. However, in contrast to the V2O5 system, in VO2 the energy region about the highest occupied electron levels (e « 0) is characterized by (occupied and empty) vanadium 3d type electron states. This is consistent with a 3d^ valence configuration of ionic V"^ and with an overall (V^^)(0^")2 compound suggested by formal valence charges. The DOS curves show subtle differences between the two different phases near (8 = 0). The rutile phase shows a continuous increase of the DOS from 8 = -0.5 eV with increasing energy that evidences a metallic state. In contrast, the monoclinic phase yields finite DOS values above 8 = -0.4 eV with a 0.1 eV wide deep minimum just above 8 = 0 indicating a small gap semiconductor. These findings can give a qualitative picture of the electronic consequences (semiconductor-to-metal transition) of the structural phase transition observed for bulk VO2. However, the actual size of the semiconductor band gap, 0.6 eV found in the experiment [92], cannot be reproduced by the band structure calculations. This has been explained by limitations of the local density approximation used in the theoretical treatment [104]. Further, the band structure calculations cannot uniquely identify the electronic origin of the first order phase transition for bulk VO2 nor the nature of the monoclinic phase (Peierls vs. Mott-Hubbard insulator) which is still under debate [104]. The electronic structure of bulk V2O3 has been studied by different theoretical methods [106-109]. The main goal was to explain the physics of the Mott transition at Tc= 150 - 170K where the material transforms from an antiferromagnetic insulating phase with monoclinic structure below Tc to a paramagnetic

147 10.0

5.0 10.0 DOS [a.u.]

15.0

Fig. 5. Total DOS of bulk V2O3 for the trigonal corundum lattice geometry. The energy zero coincides with the energy of the highest occupied state. The DOS is given in states per unit volume and per eV.

metallic phase with trigonal corundum structure. While the monoclinic low temperature phase is quite involved and not fully understood in its electronic properties, the paramagnetic metallic phase seems quite clear and can be accounted for by band structure calculations. As an example, Fig. 5 shows the total DOS curve of bulk V2O3 for the trigonal corundum lattice geometry obtained by ab initio FP-LAPW calculations [105] where the experimentally known geometry has been used, see Sec. 2.1.1. The characterization of the two main band regions shown in Fig. 5 is completely analogous to what has been found for the V2O5 and VO2 bulk systems and differs only in the relative position of the highest occupied electron levels. Electron states with negative energies between -3.5 and -8.3 eV define an occupied valence band region determined by dominant oxygen 2p and negligible vanadium contributions, thus, reflecting a 2p^ valence configuration of ionic oxygen O^". In addition, there is a valence band region between -1.6 and 0.0 eV, which originates from occupied vanadium 3d type electron states. This region is larger than that of bulk VO2 and is consistent with a 3d^ valence configuration of ionic V^^ and with an overall (V^"')2(0^-)3 compound suggested by formal valence charges. The electronic structure of bulk VO has been calculated by different band structure methods [110-114] and using correlated electron procedures [115]. This Mott-Hubbard metal, which forms a rocksalt type lattice, is the simplest of all single valence oxides of vanadium and has been treated theoretically already a long time ago. As an example, Neckel et al. [114] have published results from self-consistent APW calculations for the experimentally known lattice geometry

148

of VO, which continue the trends in the electronic parameters found for the previously discussed vanadium oxides. The VO bandstructure reveals a valence band region between 10.3 and 5.7 eV below the Fermi level described by oxygen 2p contributions. In addition, there is another valence band region between 4.1 eV below the Fermi level and the Fermi level where the band states are vanadium 3d type. The region is larger than that of bulk VO2 and V2O3 reflecting a larger V 3d occupation. This is confirmed by data from charge analyses which yield a V 3d contribution of 2.8 electrons in the oxide, somewhat larger than suggested by a (V^'^)(0^') configuration from formal valence charges.

2.2. Vanadium oxide surfaces 2.2.1. Geometric structure So far, theoretical studies on vanadium oxide surfaces have focused exclusively on single crystal surfaces which are described by low Miller indices and are believed to be energetically favorable. These surfaces are most easily accessible by theory due to their relatively simple geometry although their relevance as to catalytic activity has been doubted. Single crystal surfaces of the cubic monoxide VO, see Fig. 6, include the

Fig. 6. Geometric structure of the (100), (111), and (110) oriented single crystal surfaces of cubic VO. Vanadium and oxygen centers are shown as light large and dark small balls, respectively, with connecting sticks indicating nearest neighbor relations.

149

non-polar (100) surface described by planes containing square lattices of both V and O ions in a checkerboard arrangement. The polar (111) surface is determined by alternating planes of either V or O ions only in hexagonal lattice positions where the surface can be V or O terminated. Further, the open (110) surface can be characterized by planes with rectangular lattices containing chains of alternating V and O ions. So far, no theoretical studies involving VO surfaces have been published. The only single crystal surface of the trigonal corundum phase of the sesquioxide V2O3, which has been examined so far in preliminary cluster studies, is the (0001) surface, see Fig. 7. This surface is described by alternating rows of two hexagonal V planes (A, A') very close together and quasi-hexagonal O planes (B, whose atom density is three times as large as that of the V planes) in a stacking sequence ...BAA'BAA'B... This allows three different ideal surface terminations, ...BAA', ...AA'B, and ...A'BA where the latter reflects the ion arrangement with the smallest average Madelung potential at the surface. This corresponds to a surface arrangement of one hexagonal V plane on top of a quasi-hexagonal O plane (containing 3-fold coordinated oxygen 0(3)) as shown in Fig. 7. This surface geometry is very similar to that found in experimental LEED studies on the (0001) surface of trigonal Cr203 where in addition major relaxations of the top layers along the surface normal have been observed [116].

0(3)

Fig. 7. Geometric structure of the (0001) oriented single crystal surface of trigonal V2O3. for a surface termination A'BA. Vanadium and oxygen centers are shown as light large and dark small balls, respectively, with connecting sticks indicating nearest neighbor relations.

150

0(3)

Fig. 8. Geometric structure of the (Oil) oriented single crystal surface of the monoclinic (distorted rutile) phase of the VO2 for a surface termination A'BA. Vanadium and oxygen centers are shown as light large and dark small balls, respectively, with connecting sticks indicating nearest neighbor relations.

Theoretical studies on single crystal surfaces of the monoclinic (distorted rutile) phase of the dioxide VO2 have been restricted to the (Oil) surface (in the crystal definition according to Andersson [80]), see Fig. 8. This surface is described approximately by alternating rows of two quasi-rectangular O planes (A, A') and one VO plane (B, whose atom density is twice as large as that of the O planes) in a stacking sequence ...BAA'BAA'B... This allows three different ideal surface terminations, ...BAA', ...AA'B, and ...A'BA where the latter, see Fig. 8, is expected to be the energetically lowest in analogy with the surface structure of rutile type Ti02(l 10) [117]. This leads to 5- and 6-fold coordinated vanadium centers, V(5), V(6), and two- as well as three-fold coordinated oxygen, 0(2) and 0(3), respectively, at the ideal bulk terminated VO2(011) surface as indicated in Fig. 8. The (Oil) surface of monoclinic VO2 is, in its structure, very close to the (110) surface of tetragonal (rutile) VO2 with identical local surface elements such as rows of bridging oxygen and 5- as well as 6-fold coordinated vanadium ions. However, no theoretical studies on the high temperature rutile phase have been published so far. In the pentoxide V2O5, forming a rather open orthorhombic layer-type lattice, the only single crystal surface which has been examined in detail is the (010) surface, see Fig. 9. (It should be noted again that in this chapter the definition of Miller indices is based on orthorhombic crystal axes as introduced in Ref. [87], see also Sec. 2.1.1.). This surface, which represents the preferred cleavage plane of the crystal, is determined in its structure by local V-0 binding

151

0(1)

0(2)

0(2')

Fig. 9. Geometric structure of the (010) oriented single crystal surface of orthorhombic V2O5. Vanadium and oxygen centers are shown as light large and dark small balls, respectively, with connecting sticks indicating nearest neighbor relations. Differently coordinated surface oxygen centers, 0 ( l - 3 , 2',3'), are labeled accordingly.

rather than by stacking of the crystal planes (discussed above for the compact oxides) resulting in an altogether rough surface shape. Here six planar atom layers (4 oxygen, 2 vanadium layers) form a crystal slab which is repeated along to the (010) netplane direction. These crystal slabs can also be described by periodic arrangements of edge and comer sharing VO5 pyramids sticking out at both sides of the slab as discussed in Sec. 2.1.1. In each V2O5 slab there are three structurally different oxygen centers, terminal (vanadyl) oxygen, 0(1), coordinated to one vanadium atom through a short bond and bridging oxygen, 0(2) / 0(3), coordinated to two or three vanadium atoms. This leads to five different oxygen centers at the ideal bulk terminated ¥265(010) surface as indicated in Fig. 9. Terminal oxygen centers 0(1) are located above vanadium centers. Further, oxygen centers 0(2), 0(2') bridge two vanadyl groups pointing into the bulk and sticking out of the surface respectively, while oxygen centers 0(3), 0(3') are connected to three vanadyl groups pointing into different directions. In addition, the surface exposes bare vanadium centers connected to 0(2) andO(3, 3') centers.

152

2.2.2, Physical and chemical properties Theoretical studies on various physical and chemical parameters of vanadium oxide surfaces have been performed using both repeated slab as well as local cluster models. However, the vast majority of studies has focused on the (010) surface of the pentoxide, V2O5. This is a result of rather little experimental information on details of VxOy surface systems other than ¥205(010) that has attracted great interest due to its possible importance in catalytic applications. Further, most of the theoretical work has been based on cluster type studies where local surface behavior, in particular near surface oxygen sites, is discussed in detail. This will be reflected in the following discussion.

(a) Properties of clean surfaces The electronic structure and binding at vanadium oxide surfaces can be quantified within the DFT framework by analyses of the Kohn-Sham orbitals and corresponding electron densities resulting from periodic slab or local cluster calculations. In particular, cluster studies offer all traditional quantum mechanical tools, such as population analyses [118] or bond orders [119, 120], to examine inter-atomic binding at the surface. As an example, Table 1, lists atom charges from Mulliken populations and Mayer bond orders of substrate clusters V10O31H12 and V20O62H24, modeling one and two layers of the ¥265(010) surface, see Fig. 10. The data are obtained from DFT calculations [121-123] using the revised Perdue-Burke-Emzerhof (RPBE) functional [124, 125]. A compariTable 1 Atom charges q and bond orders p of substrate clusters V10O31H12 and V20O62H24, modeling the ¥205(010) surface, see text. All data refer to V, O surface species near the cluster center. The differently coordinated oxygen species 0(1), 0(2), 0(3) are indicated in Fig. 9. The two entries of q(V) for V20O62H24 correspond to vanadium of the first and second layer respectively. All values are given in atomic units. V10O31H12

V20O62H24

q(V)

1.59

1.66,1.78

q(0(l))

-0.34

-0.34

q(0(2))

-0.70

-0.70

q(0(3))

-0.87

-0.88

P(0(1)-V)

2.04

2.05

P(0(2)-V)

0.82

0.82

P(0(3)-V)

0.49

0.44

153

Fig. 10. Geometric structure of the clusters V10O31H12 (a, one layer) and V20O62H24 (b, two layers) representing a section at the ¥205(010) surface. The V (O) atoms are shown as large (small) shaded balls while very small balls refer to hydrogen atoms used to saturate oxygen atoms at the cluster periphery. For a definition of the differently coordinated surface oxygen centers, 0(1) to 0(3), see Fig. 9.

son of the calculated atom charges and bond orders shows very close similarity between the one- and two-layer clusters, which indicates that electronic interlayer binding is rather weak confirming the experimental result of a layer-type substrate material. All vanadium centers are positively charged and oxygen centers are negative. Vanadium ions are described as V^'^"^ to V^'^"^ where the variation reflects the location inside the cluster. Further, the negative oxygen charges scale with coordination, O^"^' for singly coordinated oxygen 0(1), O^"^' for doubly coordinated oxygen 0(2), O^'^" for triply coordinated oxygen 0(3). This indicates for the ¥205(0x0) surface that bridging oxygen sites are more nucleophilic than terminal vanadyl sites, which becomes important in view of the reactivity of the different sites with respect to surface chemical reactions. Altogether, the cluster data yield charging of the different cluster ions which is much smaller than formal valence charges, V^"^ and O^", would suggest. This reflects the well-known fact that MuUiken charges and formal valence charges can yield the same qualitative picture but may not be compared on a quantitative basis due to the absence of a unique definition of atom charges in compound material. In addition, the small MuUiken charges indicate sizeable covalent contributions to interatomic binding inside the surface layers of V2O5(010). This is confirmed by the bond order results of Table 1, which reveal the general picture based on atom coordination. Bonds between terminal oxygen, 0(1), and vanadium yield bond order values close to 2 which suggests double bonds. Bonds between bridging oxygen, 0(2), and each of their two vanadium neighbors result in bond order values close to 1, corresponding to two single bonds per oxygen. Further, V - 0 bond orders involving bridging atoms, 0(3), give values of 0.5 - 0.6 per bond revealing weaker than single bonds.

154

The electronic structure of the surface clusters in the valence region is determined by occupied Kohn-Sham valence orbitals being dominantly O 2sp type with very little V 3d admixture. Their energy range corresponds to valence energy widths A of about 5.5 eV. For comparison, very recent FP-LAPW band structure calculations [81] yield A = 5.35 eV for the V2O5 bulk and A = 5.05 eV for ¥205(010) single layer slabs, which is rather close to the cluster results. This indicates that the two different model concepts, periodic slabs and local clusters, yield the same electronic structure of the ¥205(010) surface. The energetic distribution of the Kohn-Sham valence levels in the clusters can be described by a total density of states ntot(e) (DOS) while its atom decomposition is accounted for by corresponding partial (atom projected) densities of states nA(e) (PDOS) [122, 126]. Fig. 11 shows DOS and PDOS curves for the V10O31H12 cluster, see Fig. 10, where the vanadium contributions as well as those from the differently coordinated oxygen atoms, terminal 0(1) and combined bridging 0(2) + 0(3), are included [126]. The curves have been smoothed by gaussian level broadening (FWHM of 0.4 eV) and the energy region of occupied levels is separated from that of the empty levels (gray background) by a thin vertical line. The total DOS in the energy region between -13 and -7 eV shows a multi-peak structure reflecting the above described dominantly O 2sp derived valence band region, in very good agreement with DOS results from larger clusters as well as from bulk V2O5 and ¥205(010) single layer slabs [81-83, 127]. (Note that the additional peaks below -13 eV correspond to split-off cluster levels arising from 12.0

3

CO

O Q

-16.0

-13.0

-10.0 -7.0 Energy [eV]

-4.0

Fig. 11. Total DOS and atom projected PDOS curves of the V10O31H12 cluster. For further details see text.

155

bond saturation of peripheral oxygen atoms by hydrogen terminator atoms and have to be considered cluster artifacts.) The PDOS curve referring to terminal oxygen 0(1) (dotted line in Fig. 11) is described by an overall confined (~ 3 eV wide) distribution that is concentrated near the center of the valence region with smaller contributions above the center. In contrast, the PDOS's of bridging oxygen 0(2,3) yield a broad distribution covering the full energy range of the valence region. Thus, the 0(1) derived cluster contribution shows a dispersion width smaller than that of bridging 0(2,3) species. This may be explained by the spatial distribution of the different oxygen centers in the crystal and their effective inter-atomic binding [94]. Recent angle resolved UPS experiments on a freshly cleaved ¥265(010) surface sample [93, 94] yield spectra at normal emission which exhibit three major peaks in the O 2sp valence region, see Fig. 3. This is consistent with the overall shape of the calculated total DOS curve of the V10O31H12 cluster shown in Fig. 11 and suggests that the most prominent central peak in the experimental spectrum may be assigned to emission from mainly terminal oxygen, 0(1), while the two peripheral peaks at the top and bottom of the valence band region are characterized as mixtures of vanadium with 0(2) and 0(3) induced emission intensity. This interpretation can help to identify the type of surface oxygen that participates in specific reaction steps in combined reaction and photoemission experiments. Table 2 Atom charges q and bond orders p of substrate cluster V17O42H16 modeling the monoclinic VO2(011) surface, see text. All data refer to V, O surface species near the cluster center. The differently coordinated oxygen species 0(2), 0(3) as well as the 5- and 6-fold coordinated vanadium species V(5), V(6) are indicated in Fig. 12a. The two and three entries in the bond orders reflect the variation between slightly different V-O distances. All values are given in atomic units. V17O42H16

q(V(5))

1.49

q(V(6))

1.57

q(0(2))

-0.46

q(0(3))

-0.71

p(0(2)-V)

0.95, 0.99

p(0(3)-V)

0.55,0.65,0.70

156

Fig. 12. Geometric structure of (a) the V17O42H16 cluster representing a section at the VO2(011) surface and (b) the V11O33H33 cluster representing a section at the V2O3(0001) surface. The V (O) atoms are shown as large (small) shaded balls while very small while balls refer to hydrogen atoms used to saturate peripheral oxygen atoms. The differently coordinated surface oxygen centers, 0(2), 0(3), and the 5- and 6-fold coordinated vanadium centers V(5), V(6) are labeled accordingly.

The electronic structure of the monoclinic VOiCOl 1) surface has been examined in preliminary studies [128] using clusters of different size and shape. As an example, Table 2, lists atom charges from MuUiken populations and Mayer bond orders for the substrate cluster V17O42H16, see Fig. 12a. The data are obtained from DFT calculations [128] using the RPBE functional [124, 125]. These results reveal local charging and V-O binding at the VO2(011) surface which is quite similar to that found for ¥205(010). Positive vanadium ions are described as V^'^"^ (5-fold coordinated) and V^"^"^ (6-fold coordinated) and the negative oxygen charges scale with coordination, O^'^" for bridging oxygen 0(2), O^*^" for bridging oxygen 0(3). As for the ¥205(010) surface clusters, charging of the different cluster ions in the V17O42H16 cluster is found to be much smaller than formal valence charges of the VO2 substrate, V"^ and O^", suggest. This has implications identical to those discussed before. In particular, the data evidence that the ion charges of vanadium differ only little (by 0.1 electrons) between the 5- and 6-fold coordinated surface species. Therefore, discrimination between the two must be based on covalent rather than ionic binding behavior. Furthermore, the vanadium charges determined from poputions in the VO2(011) surface clusters are decreased only slightly compared to those of the ¥205(010) surface clusters. This is in contrast to the standard chemical characterization in terms of formal valence charges +4 and +5 in the two substrates and confirms the importance of covalent binding in the substrates. The bond order results of Table 2 reflect again the general picture based on atom coordination. Bonds between doubly coordinated oxygen, 0(2), and each of their two vanadium neighbors result in bond order values close to 1, sugges-

157

Table 3 Atom charges q and bond orders p of substrate cluster V11O33H33 modeling the trigonal ¥203(0001) surface, see text. All data refer to V, O surface species near the cluster center. The 3-fold coordinated oxygen species 0(2) is indicated in Fig. 12b. The three entries in the bond orders reflect the variation between different V-0 bonds. All values are given in atomic units. V11O33H33

q(V)

0.82

q(0(3))

-0.67

P(0(3)-V)

0.50, 0.58, 0.95

ting two single bonds per oxygen. Further, V-O bond orders involving triply coordinated atoms, 0(3), give values of 0.5 - 0.7 per bond revealing weaker than single bonds. However, the V-O bond orders in the VO2(011) surface clusters are, altogether, somewhat larger than corresponding values found for the ¥205(010) surface clusters, see Table 1. The suggests that covalent binding becomes more important at the expense of ionicity at VOiCOll) that is consistent with V and O charges of the dioxide being smaller than those of the pentoxide. Preliminary cluster studies [129] on the electronic structure of the trigonal ¥203(0001) surface substantiate the above concepts. As an example. Table 3, lists atom charges from MuUiken populations and Mayer bond orders for the substrate cluster V11O33H33, see Fig. 12b. The data are obtained from DFT calculations [129] using the RPBE functional [124, 125]. Here positive vanadium ions closest to the surface are described as V^"^^, and the negative oxygen as O^'^". The vanadium charges determined from populations are the smallest of the three oxide surfaces, which goes along with the formal valence charge +3 of vanadium in V2O3 and emphasizes the increased importance of covalent binding in this substrate. In addition, the V-O bond order results of Table 3 reflect the coordination of the oxygen ions as discussed before.

(b) Surface oxygen vacancies Catalytic properties of vanadium oxide based catalysts depend strongly on their ability to provide surface oxygen as a reactant. Therefore, theoretical studies of physical and chemical (catalytic) properties of the structurally different surface oxygen centers can help to elucidate details of the reaction steps. Moreover, calculations of electronic, structural, and energetic properties of oxygen vacancies created by oxidation reactions are of vital importance as will be discussed in the following.

158

It should be noted that oxygen vacancies at oxide surfaces are of major importance for the physical and chemical behavior of these systems, apart from consequences for catalytic reactivity. As examples we mention point defects at silicon dioxide [130] or the existence of color centers at MgO surfaces which are due to localized electron states at oxygen vacancy sites and have been discussed extensively in the literature [131]. Surface oxygen vacancies have been studied extensively for the ¥205(010) surface by cluster calculations [122-123, 126]. Due to the existence of five structurally non-equivalent oxygen sites, 0(1-3, 2',3'), at ¥205(010), see Fig. 9, this surface may exhibit five different types of surface oxygen vacancies. These can be calculated by removing corresponding oxygen species from the surface clusters and re-evaluating the electronic structure of the resulting vacancy clusters. Here frozen vacancy states are obtained from calculations where all atom positions are kept frozen at the initial cluster geometry. In addition, relaxed vacancy states are determined by geometry optimizations allowing all cluster atoms to rearrange according to lowest total energy. As an illustration, Table 4 lists results from calculations based on a V10O31H12 substrate cluster [126], see Fig. 10, where oxygen has been removed from sites 0(l-3). Oxygen vacancy energies, ED^(0) (frozen geometry) and ED^''^(0) (relaxed geometry), are determined by appropriate cluster total energy differences (based on DFT calculations using the RPBE functional) using the ground state of atomic oxygen as a reference. All vacancy energies are found to be rather large. Frozen values range between 7 and 8 eV and vacancy induced relaxation decreases these energies by only 0.5 to 0.8 eV. This suggests that oxygen is bound very strongly to its substrate environment and may not be removed in one single step during an oxidation reaction. However, pre-adsorbed atoms or molecules may weaken surTable 4 Oxygen vacancy energies, EB^'^\0), for the surface sites 0(1-3) at the ViOsCOlO) surface with and without surface relaxation. All data are obtained for the V10O31H12-O (= V10O30H12) model cluster [126]. In addition, the table contains values of charges q^^'^^^V) and relaxation displacements Ar(V) of the central V atom(s) closest to the vacancy. Energies are given in eV and lengths are in A. Vacancy site

ED'(0)

ED'(O)

q'(V)

q^(V)

Ar(V)

0(1)

7.16

6.47

1.38

1.51

0.50

0(2)

7.95

7.19

1.26

1.36

0.70

0(3)

7.01

6.49

1.28

1.33

0.08

159

face binding of the oxygen and hence may make it available for oxidation processes as will be discussed later on. Figs. 13a-d visualize geometric consequences of substrate relaxation as well as charge transfer due to vacancy formation at the ¥205(010) surface where V10O31H12 based vacancy clusters are used as models of different vacancy sites 0(l-3). All cluster ions are displayed by shaded balls where ball radii represent atom charges as obtained by populations. As a main result, the relaxation effect is always found to be locally confined. If oxygen is removed from a (singly coordinated) terminal 0(1) site, see Fig. 13a, b, relaxation causes the vanadium center below the vacancy site (hatched ball) to move perpendicular to the surface towards the substrate by about 0.5 A. Thus, the vanadium penetrates further into the substrate and can interact with the nearest terminal oxygen of the

Fig. 13. Relaxed geometry of the V10O31H12-O ( = V10O30H12) cluster with (a, b) an 0(1) vacancy, (c) an 0(2) vacancy, and (d) an 0(3) vacancy. The views are perpendicular (plots a, c, d) and parallel (plot b) to the ¥205(010) surface, respectively. Cluster atoms are shown by shaded balls where ball radii represent atom charges. Dark (light) shading refers to negative (positive) charge while the radii give the amount of charge. The central V species next to the 0(1) vacancy is emphasized by hatching in plots a, b. The white balls behind the relaxed cluster describe the system without vacancy.

160

second substrate layer to form a V - O - V bond bridge between the two crystal layers. This effect has been confirmed by calculations using a vacancy cluster V20O62H24-O ( = V20O62H24) modeling two adjacent crystal layers [5, 126]. Here the V - O - V inter-layer bond bridge is found to be very similar to bridge bonding at the 0(2) surface site as confirmed by bond order and charge analyses [132]. Thus, vacancy formation at the first crystal layer of the ¥205(010) surface can increase the electronic coupling with the second layer. This may be the starting point of major surface relaxation and may result eventually in a reconstruction of the whole surface region. If oxygen is removed from a (doubly coordinated) bridging 0(2) site, see Fig. 13c, the strongest relaxation shifts occur for the two vanadium centers adjacent to the vacancy which move laterally by 0.6 A such that the vacancy opening is enlarged. However, the lattice topology of the ¥205(010) surface is conserved suggesting that a single 0(2) vacancy will not introduce major surface restructuring. Finally, the relaxation geometry of the 0(3) vacancy, given in Fig. 13d, can be understood by the same mechanisms discussed for the other vacancies. A closer inspection of the ball radii in Figs. 13a-d reveals that the positive charges of the vanadium ions closest to each oxygen vacancy are smaller than corresponding values of the cluster without the vacancy. This is also clear from a comparison of the charge values q(V) given in Tables 1 and 4 and evidences chemical reduction of the metal sites. A more detailed description of the effect becomes possible by (P)DOS and orbital analyses of the vacancy clusters. As an example. Fig. 14 shows (P)DOS curves obtained for the V10O31H12 based vacan12.0

3 C/)

O Q

-16.0

-13.0

-10.0 -7.0 Energy [eV]

-4.0

Fig. 14. Total DOS and atom projected PDOS curves of the valence band region for the V10O31H12 based vacancy cluster modeling an 0(2) surface vacancy, see text.

161

cy cluster modeling an 0(2) surface vacancy. A comparison of these curves with those of the initial substrate cluster, see Fig. 11, shows clear differences in the energy region of the highest occupied cluster levels. While in the initial cluster the highest levels are at the top of the O 2sp valence region and refer to oxygen type orbitals, those of the vacancy cluster are located at about 2 eV above the O 2sp valence region and are characterized by vanadium 3d orbitals. This shows that the oxygen vacancy leads to increased local vanadium 3d occupation, substantiating the microscopic picture of chemical reduction of the metal sites observed also in photoemission experiments [133]. Extended studies of local electronic states of 0(2) and 0(3) vacancies at the ViOsCOlO) surface [134] based on the V10O31H12 cluster have been carried out in search of possible color center states. These calculations do not yield any low-lying electron states localized at the vacancies and reminiscent of color centers as found at MgO surfaces [131] in accordance with experimental evidence [135]. This can be rationalized by simple geometric arguments [126] based on the size of the surface vacancies at ¥205(010) and by the strength of the electrostatic Madelung potential near the vacancies. Details of possible diffusion of surface oxygen vacancies are of great importance for a microscopic understanding of structural transformations at the V2O5 surface which happen as a consequence of oxidation reactions. Diffusion processes, where oxygen, residing at a surface site, diffuses to a nearby oxygen vacancy site, have been studied in some detail by cluster models [126]. As an illustration, we mention results from V10O31H12 based vacancy clusters with a model path 0(1) -^ 0(2')vac at the ¥205(010) surface. This corresponds either to diffusion of oxygen species from an 0(1) site filling a nearby vacancy at 0(2') or to diffusion of an oxygen vacancy at 0(2') to a nearby 0(1) site. From calculations of total cluster energies for oxygen moving along the path its transition state geometry T and the corresponding diffusion energy barrier EB can be evaluated. Fig. 15 visualizes the 0(1) -^ 0(2')vac diffusion path by small dark balls with the transition state geometry T indicated by a large dark ball. This path is found to yield the smallest barrier, only 0.5 eV, which can be easily released by typical surface reactions. Thus, 0(2^'^) vacancies, created by oxidation reactions at the ¥205(010) surface may be annihilated quite easily by oxygen transfer from terminal 0(1) sites. This result is rather interesting in connection with the identification of reactive surface sites. Depending on the experiment, one may not be able to distinguish between an oxidation reaction step occurring at an 0(1) site and one occurring at an 0(2) site followed by diffusion of oxygen from 0(1) to the vacancy at 0(2). However, more detailed calculations are needed to substantiate these findings.

162

Fig. 15. Diffusion of surface oxygen filling a neighboring vacancy at the ¥205(010) surface. The model path 0(1) -^ 0(2')vac is visualized for a V10O31H12 substrate cluster and connects sites 0(1) (geometry A) with 0(2') (geometry B). Oxygen positions along the path are shown by small black dots with the transition state geometry (T) emphasized by a larger dark ball.

(c) Adsorption Atomic and molecular adsorption at vanadium oxide surfaces have been studied theoretically using both periodic slab and cluster models where so far studies are restricted to the pentoxide, V2O5, as a substrate due to its possible importance in catalytic applications as mentioned before. Further, adsorbate species include in all cases atoms (H [122-123, 126, 136-142], O (see below)) or rather small molecules (O2 (see below), H2O [143-144], NH3 [145-147], NO [146, 148], C2H4 [149], propene (CsHg) [140], toluene (C6H5CH3) [140]) that are of catalytic interest but also small enough to make meaningful calculations feasible. Hydrogen adsorption at the ¥265(010) surface has been studied extensively by both cluster [122-123, 126, 136-140] and slab models [141] due to its catalytic relevance. Under reaction conditions hydrogen will always be present at vanadium oxide surfaces giving rise to surface hydroxyl groups. These groups may desorb from the surface or recombine forming water species that desorbs. Here it is interesting to find out which of the oxygen sites are involved in the adsorption and formation of surface OH or H2O species and which of these species will desorb most readily. These questions have been addressed in previous theoretical DFT studies [122-123, 126, 140] based on a V10O31H12 substrate cluster. Here one or two hydrogen atoms are added to one of the five different surface sites, 0(l-3, 2',3'), and local geometries of the resulting surface OH and H2O groups are optimized in total energy calculations. Fig. 16 visualizes resulting equilibrium geometries for both species and for different

163

OH, H20atO(1)

OH, HgO at 0(3) OH, H2O at 0(2) Fig. 16. Sketch of computed equilibrium geometries of hydrogen adsorption at different oxygen sites of the ViOsCOlO) surface. Single hydrogen adsorption (surface OH) as well as adsorption of two hydrogen atoms (surface H2O) are included. The results are obtained from optimizations of V10O31H12+H and V10O31H12+2H cluster models, respectively. The surface species, OH or H2O, are shown by haded balls while the surface lattice is sketched by light balls.

surface sites. In addition, Table 5 lists energetic results from the calculations. Here EB(H) denotes the binding energy of a single hydrogen adsorbed at a given oxygen site 0(l-3, 2', 3') while EB(H') gives the binding energy of a hydrogen adsorbing at an existing surface OH group to yield surface H2O. Further, EB(H) refers to the combined adsorption energy of two hydrogen species to form surface H2O. All energies are determined from total energy differences of the appropriate clusters where one or two isolated hydrogen atoms in their ground states are taken as reference. The EB(H) values of Table 5 show that hydrogen Table 5 Hydrogen adsorption energies, EB(H), EB(H'), and EB(2H), for the V10O31H12 substrate cluster with H, 2H adsorbed at different oxygen sites 0 ( l - 3 , 2', 3')- For definitions of the adsorption energies see text. All energies are given in eV. Site

EB(H)

EB(H')

EB(2H)

0(1)

2.34

1.56

3.90

0(2)

2.21

0.66

2.87

0(3)

1.88

1.08

2.96

0(2')

0.54

--

--

0(3')

0.76

--

--

164

can adsorb at all oxygen sites. Binding is found to be rather strong at the 0(1), 0(2), and 0(3) sites, with EB(H) values between 1.88 eV and 2.34 eV, which suggests rather stable surface hydroxyl groups at these sites. In contrast, the small EB(H) values found for the 0(2') and 0(3') sites reveal fairly weak adsorptive binding. The latter sites are located close to two vanadyl groups that obviously makes surface OH formation less likely to happen. The equilibrium geometries of surface OH resulting from the calculations, see Fig. 16, do not reflect the symmetry of the corresponding oxygen sites before adsorption. The 0-H distances of all surface groups are quite close to those of typical OH containing systems [150] while the distances between oxygen and its nearest vanadium neighbors at the V2O5 surface are enlarged. Thus, hydrogen adsorption, which results in strong O-H binding, weakens the V-0 bonds near the adsorption site. This is confirmed by bond order analyses [126], which show that for example the V=0 double bond at the 0(1) site is reduced to a single bond and the weaker V-0 bonds at 0(2, 3) are reduced even further. At all oxygen sites, hydrogen adsorption leads to the accumulation of additional negative charge at the corresponding surface oxygen and the surface OH species becomes slightly negative, OH^^" to OH^ '^' from population analyses. In addition, the positive charges at vanadium neighbors are decreased slightly which can be understood by hydrogen induced chemical reduction of the metal species. This is in complete analogy with the reduction effect caused by oxygen vacancies discussed above. Densities of states of the adsorbate clusters reveal additional occupied states, at about 2 eV above the oxygen 2sp valence region, which are characterized as V 3d type. The hydrogen induced reduction effect has been confirmed by recent photoemission experiments [151]. The EB(H') values of Table 5 show that a second hydrogen can adsorb at the three surface OH groups involving the 0(l-3) sites which results in surface H2O. However, corresponding O-H binding energies of the second hydrogen are always found to be smaller than those of the first. The largest EB(H') value, 1.56 eV, is obtained for the terminal 0(1) site yielding a combined adsorption energy EB(2H) = 3.9 eV. The EB(2H) values of the other sites are considerably smaller which is explained by the fact the surface oxygen in two- and three-fold coordinated sites is less flexible in its ability to rehybridize compared to singly coordinated terminal oxygen 0(1). It should be noted that the present definitions of hydrogen adsorption energies refer to adsorption starting from atomic species. If molecular hydrogen gas is exposed to the ¥205(010) surface the energetics of corresponding adsorption processes has to include H2 -^ H + H dissociation. This reduces the computed adsorption energies EB(H) for single H adsorbates by half of the dissociation energy of free H2 (V2 D(H2) = 2.25 eV [150]). As a consequence, positive adsorption energies (stable adsorption states) are obtained only for the 0(1) and

165

0(2) sites while 0(2'), 0(3), and 0(3') site adsorption states are metastable. Further, energies EB(2H) for adsorption of two H species at a given oxygen site would have to be reduced by D(H2) = 4.5 eV which, based on the present results, yields only metastable adsorption states. Surface oxygen which, as a result of catalytic oxidation at the V2O5 surface, is incorporated into a hydrocarbon will be removed from the surface leaving a vacancy behind. Here a possible removal scenario could consist of a combined process where hydrogen is adsorbed at the surface and the resulting surface OH or H2O group desorbs. The energetics of such processes can be studied within the cluster approach by combining the results obtained for the vacancy clusters with those describing hydrogen adsorption. Within the surface cluster approach, desorption processes can be described in their energetic consequen(V205)surf ^

( V 2 0 5 ) s u r f / 0 ^ " + Ogas

(A)

or (V205)surfHads ^

(V205)surf/0^^^ + (OH)gas

(B)

or (V205)surf2Hads " ^ (V205)surf/0^^^ + (H20)gas

( Q

ces by desorption energies E D ( 0 ) , ED(OH), and ED(H20), respectively, which are given by corresponding cluster total energy differences. Based on the V10O31H12 substrate cluster [122, 123, 126] desorption energies of the different processes are listed in Table 6. Note that process (A) describes oxygen vacancy formation and, therefore, desorption energies ED(0) of Table 6 are identical to oxygen vacancy formation energies EDÔ) of Table 4. As discussed before, desorption energies ED(0) are rather large (6.5 eV to 7.2 eV) for all surface oxygen sites, which shows that it is quite difficult to remove oxygen by itself from Table 6 Desorption energies E D ( 0 ) , E D ( O H ) , and E D ( H 2 0 ) for different oxygen sites at the V2O5(010)

surface. All data are obtained from calculations based on the V10O31H12 substrate cluster [122, 123, 126]. For definitions of the different quantities see text. All energies are given in eV. site

ED(0)

ED(OH)

ED(H20)

o(i)

6.47

3.98

0.48

0(2)

7.19

4.57

0.16

0(3)

6.49

3.54

-0.44

0(2')

7.19

2.90

—

0(3')

6.49

2.43

—

166

the ideal ¥205(010) surface. However, corresponding energies ED(OH), for desorption of OH and H2O, leaving an oxygen vacancy behind, are always found to be smaller. For OH removal, ED(OH) values vary between 2.4 eV and 4.6 eV and for H2O removal, ED(H20) values range even between -0.5 eV (metastability at the 0(3) site) and 0.5 eV. These decreased desorption energies originate obviously form weakening of V-0 bonds at the V2O5(010) surface caused by hydrogen adsorption. Altogether, the cluster results suggest that the presence of pre-adsorbed hydrogen at the vanadium oxide surface facilitates oxygen removal substantially and can, thus, contribute to an enhanced yield of oxygenated products near vanadia based surfaces. Atomic oxygen adsorption at the ¥205(010) surface has been studied recently by cluster models [134] in an attempt to describe details of healing of surface oxygen vacancies. These surface vacancies can be annihilated by either diffusion in the bulk or by binding of adsorbing oxygen, where in the latter case both atomic and molecular oxygen are possible candidates. The healing process ED(H20)

(V205)surf/0^^^ + Ogas " ^ (V205)surf

(A)

is the inverse of vacancy formation and, therefore, its energetics are determined by quantities ED(0) discussed above. Alternatively, one could envision a twostep process (V205)surf/0^^^ + (02)gas " ^ ( V 2 0 5 ) s u r f / 0 ^ " + (02)ads ^ (V205)surf + Oads " ^ (V205)surf + Ogas

(B)

where molecular oxygen adsorbs at the surface followed by desorption of atomic oxygen leaving a healed surface behind. Here the intermediate state corresponds to either O2 filling an oxygen vacancy or to O being adsorbed at an oxygen surface site. The latter healing process (B) has been considered in recent theoretical DFT studies [134] based on a V10O31H12 substrate cluster. The intermediate state describing atomic oxygen adsorption deserves special attention. In calculations, atomic oxygen is added at different oxygen sites of the V10O31H12 cluster and the geometry of the resulting local O2 surface group is optimized in total energy calculations. The resulting equilibrium geometries of the O2 surface species, visualized in Fig. 17, yield O2 lying flat on the surface at the 0(1, 2, 2') sites and protruding out of the surface at the 0(3, 3') sites. Its inter-nuclear distance is increased at all surface sites by about 0.2 A compared to the distance in gas phase O2, thus suggesting a per-oxo type species. Table 7 lists energetic results from the calculations. Here EB(0) denotes the binding energy of a single oxygen adsorbed at a given oxygen site 0(1-3, 2', 3') while EB(02/vac) gives

167

02 at 0(3') 02atO(3) \ r . r ^

O2 at 0(2')

02atO(1)

/

O2 at 0(2) Fig. 17. Sketch of computed equilibrium geometries of atomic oxygen adsorption at different oxygen sites of the ViOsCOlO) surface reflecting also molecular oxygen adsorption at respective oxygen vacancies. The results are obtained from optimizations of the local O2 surface species on a V10O31H12 substrate cluster. The O2 surface species is shown by shaded balls while the surface lattice is sketched by light balls.

binding energies of an O2 molecule at corresponding oxygen vacancy sites. The data are determined by total energy differences of the appropriate clusters. Obviously, atomic oxygen can adsorb at all oxygen sites where strong binding is found for the 0(1, 2, 2') sites. However, at all sites EB(0) values are smaller than corresponding EB(02/vac) values. This reflects the fact that oxygen vacancy energies at the V2O5(010) surface are, for all sites, larger than the dissociation energy of gas phase O2. It suggests further that the healing scenario (B) is an exothermic process. Table 7 Binding energies E B ( 0 ) , EeCOi/vac) for different oxygen sites at the ViOsCOlO) surface. All data are obtained from calculations based on the V10O31H12 substrate cluster [134]. For definitions of the different quantities see text. All energies are given in eV. site

EB(0)

EeCOi/vac)

0(1)

1.82

2.59

0(2)

1.78

3.33

0(3)

0.35

0.98

0(2')

1.78

3.34

0(3')

0.36

0.98

168

Water adsorption at the ViOsCOlO) surface has been studied by semi-empirical cluster models [143] as well as by periodic DFT slab models [144]. The DFT studies show that molecular adsorption of H2O can happen at all surface oxygen sites as well as at the 5-fold coordinated vanadium site. At the vanadium site hydrogen bonding is found to be most important for the H2O adsorption while at the oxygen sites donation from oxygen to the water molecule seems to dominate the adsorbate binding. Dissociation of H2O is found to be rather unfavorable at the ¥205(010) surface which is explained by pronounced Coulombic effects. For both molecular and dissociative water adsorption the vanadyl oxygen 0(1) is identified as the most favorable site. This interpreted as the 0(1) site being of highest catalytic activity in contrast to other theoretical work [121, 136-140, 152-153]. Ammonia adsorption at the ¥205(010) surface has been examined in DFT studies using very small cluster models [146, 147] as well as using periodic slabs [145]. In the slap studies it is assumed that surface oxygen sites are covered with hydrogen and adsorption of NH3 at the resulting surface OH groups (Bronsted acid sites) is considered. The calculations show clearly that NH3 is bound to the Bronsted acid sites fairly strongly forming an N H / ion where binding is the strongest at the terminal 0(1) site. On the other hand, NH3 can stabilize above the 5-fold coordinated vanadium site where a small binding energy is combined with rather large adsorbate - substrate equilibrium distances. This suggests polarization and hydrogen binding while dative binding involving the NH3 lone pair orbital seems unlikely in contrast to the claim made by the authors [145]. 3. MOLYBDENUM OXIDES Molybdenum oxides form a large group of materials resulting from the ability of molybdenum to possess different formal oxidation states as well as different local coordination [2-3, 12, 34]. Of particular importance is the coexistence of different oxidation states of Mo that leads to mixed valency oxides. Amongst the binary oxides MOxOy there are only two single valency oxides, M0O3 (Mo^"^) (low-pressure orthorhombic a-phase [154] as well as high-pressure monoclinic P phase [155, 156]) and monoclinic 6-M0O2 (Mo"^) [157]. Successive reduction of M0O3 to M0O2 gives rise to intermediate mixed valency oxides which are based on the formation of crystallographic shear (CS) planes [12, 34, 154, 158160]. These oxide have been suggested to play an important role in catalysis [161, 162]. During high-temperature reduction, removal of oxygen from M0O3 creates oxygen vacancies that can interact to order into equally spaced parallel

169

CS planes across which the M0O6 octahedra share edges rather than comers. These extended defects exist even at modest temperature due to the large mobility of oxygen vacancies, Further, the defects are stabilized by relaxation as a result of Mo cation displacement from their centers of symmetry to interstices. Homologous series of molybdenum oxides [1, 3, 12, 34, 154, 158, 160, 163, 164], which are based on CS plane formation starting from a ReOs-type structure, include MOnOsn-i and M0n03n-2 systems. Single crystals of oxides MOnOsn-i have been observed for 4 < n < 9 where n determines the thickness of Re-Os-type slabs separated by CS planes. Here the CS planes are perpendicular to lattice direction a with inter-planar distances corresponding to the height of n/2 MoOe octahedra [158]. In these systems the average valence charge of molybdenum is +(6-2/n), ranging between +4 and +6, which is described by ion ratios MO^VMO"^ = (n-l)/l. Apart from the single valency oxide M0O2 (n = 1) examples are M09O26 [158, 165] (with low-temperature triclinic ^ phase and high-temperature monoclinic p' phase), monoclinic M08O23 [154, 158, 164, 166-171], M07O20, MoeOn (not prepared as single crystal), tetragonal M05O14 [154], and M04O11 [154, 158, 167, 171-175] (with low-temperature monoclinic r| phase and high-temperature orthorhombic 7 phase). Both M04O11 structures are ReOa-based but adjacent slabs of octahedra are connected by M0O4 tetrahedra and the orientation varies between slabs. Electronic properties of MonOsn-i bulk systems are studied experimentally in [176-186]. Single crystals of oxides Mon03n-2 contain CS planes where M0O6 octahedra share faces rather than edges [154, 159, 187]. In these systems the average valence charge of molybdenum is +(6-4/n), which is described by ion ratios MO^VMO"^ = (n-2)/2. The only example observed so far is triclinic Moi8052[154, 167, 188-189]. In addition to these intermediate oxides, reduced phases built of M0O6 octahedra and pentagonal columns have been characterized. Examples include the low-temperature phase K-M017O47 [154] and M010O28 [154, 158]. Further, metallic molybdenum oxide M03O with cubic structure has been described [190]. Apart from bulk material, evaporation of M0O3 can produce cyclic species like M03O9, M04O12 or M05O15 with the structures based on comer-shared M0O4 tetrahedra [191-194]. Finally, molybdenum oxide clusters have been prepared as neutral and ionic species in gas phase and studied by quantum chemical methods at various levels of theory. This includes MoO [195-198], MoOn\ n=l-3 [199], M0O3, Mo04^", M0O4H2, Mo205^"^, and M02O6, [200] where the calculations provide a satisfactory description of stmctural, energetic, and spectroscopic properties. The following discussion will be restricted to the two single valency oxides of molybdenum which have been accounted for by quantitative theoretical work.

170

3.1. Molybdenum oxide bulk systems 3.1.1. Geometric structure Single crystals of the two single valency molybdenum oxides, M0O2 and M0O3, can be characterized both by arrangements of distorted M0O6 octahedra. The corresponding crystal lattices differ basically by the connectivity of these octahedra and by their Mo-0 distances as well as 0-Mo-O angles. This is obvious from the balls-and-sticks models of the two oxides shown in Figs. 18,19. The trioxide M0O3 forms an orthorhombic crystal lattice where the elementary cell contains four chemical M0O3 units [154, 201-202]. It has a layered structure with weakly binding bilayers parallel to the (010) crystal plane. Each bilayer consists of two interleaved planes of comer-linked distorted MoOe octahedra (see Fig.l, one bilayer is marked by shaded atom balls) where octahedra from adjacent planes share edges. The octahedra are described by Mo-0 distances dMo-o = 1-67, 1.73, 2*1.95, 2.25, and 2.33 A where the largest distance refers to oxygen and molybdenum of different planes. Alternatively, each bilayer can be discussed in terms of comer-sharing sheets are composed of ribbons of octahedra sharing edges with two adjacent M0O4 tetrahedra along the (001) direction. The tetrahedra form infinite strings that are combined to increase

Fig. 18. Geometric structure of orthorhombic M0O3 with netplane stacking along the (010) direction. Molybdenum (oxygen) centers are shown by large (small) balls where those of the topmost bi-layer are shaded in light/dark gray. Non-equivalent oxygen centers, 0 ( 1 , 2, 3), are labeled accordingly. An elementary unit cell of the crystal is included to the right.

171

the coordination of Mo to 6. The resulting octahedra and comers with octahedra of parallel ribbons of both sides. In this description the layer structure of M0O3 results from the existence of free unshared comers (one per octahedron) in one layer that are pointing between those of neighboring layers. Bulk molybdenum trioxide contains three structurally different oxygen sites, terminal oxygen 0(1) coordinated to only one Mo center, and two bridging oxygen sites, doubly 0(2) and triply 0(3) coordinated to molybdenum centers, see Fig. 18. These sites can be identified by infrared and Raman spectroscopy which allow a quantitative assignment by corresponding group frequencies for stretching vibrations and can correlate bond lengths with force constants [203, 204]. The dioxide M0O2, crystallizes in a monoclinic lattice that deviates only slightly from the tetragonal rutile structure and is characteristic for several early transition metal dioxides like VO2. Its elementary cell contains four chemical M0O2 units [157] and almost equal Mo-O bond lengths describe the octahedral Mo environments. The M0O6 units form rows in which they are connected by edges and the octahedra of adjacent rows are linked by comers. Further, every second row is rotated by 90° about the (Oil) direction, see Fig. 19, lb. Molybdenum dioxide contains only one type of oxygen sites, 0(3) linking three metal centers.

Mo

0(2)

0(3)

^-.JL^/A

unit cell

Fig. 19. Geometric structure of monoclinic M0O2 with netplane stacking along the (Oil) direction. Molybdenum (oxygen) centers are shown by large (small) balls where those of the topmost layer are shaded in light/dark gray. Non-equivalent oxygen centers, 0(2, 3), are labeled accordingly. An elementary unit cell of the crystal is included to the right.

172

3.1.2. Electronic properties The electronic structure of bulk molybdenum oxides is determined by the amount of 4d electron occupation of the Mo ions analogous to vanadium oxides discussed above. The molybdenum atom is described in its ground state by an electron configuration [Kr] 4d^ 5s\ As a consequence, molybdenum cations Mo^"*^ for 4 < p < 6 are characterized by configurations [Kr] 4d^"P with a partially filled (p < 6) or an empty (p = 6) 4d shell. The 4d occupation is reflected in the electronic ground states of the bulk oxides where Mo possesses a formal charge +p. As a result, the amount of occupied Mo 4d electron states is expected to increase in going from M0O3 (corresponding formally to V^"^ ions) to M0O2 (formally V"^ ions). Although modem total energy methods allow us to obtain band structures and corresponding densities of states (DOS) for molybdenum oxides, the electronic structure of these materials has been examined in only very few theoretical studies. Electronic ground state as well as structural properties of bulk M0O3 have been studied by ab initio periodic Hartree-Fock (HF) methods [205, 206] with different electron corrections. Here geometry optimizations yield the experimental bulk structure and confirm the weak attractive interaction between adjacent bilayers. Further, interatomic Mo-O binding was found to depend on bond distance, ranging from strongly covalent binding for the shortest distances to predominantly ionic for the longest. In addition to the electronic structure also the formation of shear structures in M0O3 was studied theoretically by the SCF-SW-Xa method using very small model clusters [196, 207]. Here a mechanism of oxygen removal and change in linkage of M0O6 octahedra from comer to edge linked was assumed and an interpretation of the corresponding core level photoemission spectra of M0O3 was given. Bulk M0O2 displays a striking metal-insulator stmctural transition that was the subject of extensive theoretical as well as experimental work. The electronic bulk stmcture was studied by periodic [208] as well as cluster DFT methods [209, 210]. The electronic properties of molybdenum dioxide were found to be dominated by strong hybridization of O 2p and crystal-field split Mo 4d states with bands near the Fermi energy originating almost exclusively from Mo 4d t2g-type orbitals.

3.2. Molybdenum oxide surfaces 3.2.1. Geometric structure Theoretical studies on molybdenum oxide surfaces have focused exclusively on single crystal surfaces of M0O3 and M0O2 that are described by low Miller indices. This is due to their relatively simple geometry considering the

173

M0-15O56 cluster

Fig. 20. Geometric M0O3. Molybdenum oxygen sites, 0 ( 1 , 2, model calculations is

Mo

0(1)

0(2)

0(3)

structure of the (010) oriented single crystal surface of orthorhombic (oxygen) centers are shown by large (small) balls and non-equivalent 3), are labeled accordingly. The M015O56 surface section used in cluster shown by shaded balls.

very complex bulk structures of molybdenum oxides. Single crystal surfaces of orthorhombic M0O3, considered by theory, are of (010) and (100) orientation. The (010) surface, which represents the preferred cleavage plane of the crystal and is chemically quite inert, is determined in its structure by three differently coordinated oxygen ions while molybdenum ions as such are not exposed, see Fig. 20. Oxygen ions appear as terminal (molybdenyl) oxygen species, 0(1), coordinated to one molybdenum center at a distance dMoo = 1-67 A. Bridging oxygen, 0(2), is coordinated asymmetrically to two Mo centers at distances dMo-o = 1-74, 2.25 A and bridging oxygen, 0(3), is coordinated to three Mo centers with two equal (dMo-o = 1-95 A) and one longer distance (dMo-o = 2.33 A). The (100) surface is rather different in its structure from the (010) surface, see Fig. 21. It is much rougher and less compact and there are many structurally different metal as well as oxygen sites. Molybdenyl oxygen ions 0(1) are similar to those of the (010) surface. However, they do not cover all Mo ions leaving bare metal sites at the (100) surface. There are also symmetrically bridging oxygen centers that are coordinated to two Mo centers of the surface and couple weakly with a third Mo surface center. These centers correspond structurally to triply coordinated oxygen 0(3) at the (010) surface. Finally, there are terminal oxygen centers which correspond to molybdenyl 0(1) of the (010)

174

(010)

M0H5O56 cluster

^(100) -^— M05O24 cluster

Fig. 21. Perspective view of the (010) and (100) surfaces of orthorhombic M0O3. Molybdenum (oxygen) centers are shown by large (small) balls. The M015O56 and M06O24 surface section used in cluster model calculations are shown by shaded balls.

MO11O39 Cluster

Mo(6) Mo(5)

0(3)

0(2)

Fig. 22. Geometric structure of the (Oil) oriented single crystal surface of monoclinic M0O2. Molybdenum (oxygen) centers are shown by large (small) balls and non-equivalent sites, V(5, 6), 0(2, 3), are labeled accordingly. The M011O59 surface section used in cluster model calculations is shown by shaded balls.

175

surface but point parallel to the surface. Theoretical studies on single crystal surfaces of monoclinic M0O2 have been restricted to the (Oil) surface, see Fig. 22. This surface contains 5- and 6fold coordinated molybdenum centers, Mo(5), Mo(6), and two sets of two- as well as three-fold coordinated oxygen, 0(2) and 0(3), as indicated in Fig. 22. The surface does not contain singly coordinated oxygen.

3.2.2. Physical and chemical properties Up to now, theoretical work on molybdenum oxide surfaces has been restricted to only few periodic slab as well as local cluster studies on the single valence oxide surfaces MoO3(010), (100) and MoO2(011). As for vanadium oxide, there is very little experimental information on quantitative microscopic details of these complex surfaces despite their importance concerning catalytic applications. The following discussion of theoretical results will be based exclusively on cluster studies that have yielded rather detailed microscopic insight.

(a) Properties of clean surfaces Electronic properties of the MoO3(010) surface have been examined by cluster model studies [211] using a M015O56H22 cluster, see Fig. 23a, which describes a bilayer section of the bulk, cp. Fig. 18. This cluster has been proven to yield electronic properties that are size converged. In fact, previous studies have

Mo(6)

0(3)

Fig. 23. (a) Geometric structure of the M015O56H22 cluster modeling a bi-layer section at the MOOBCOIO) surface, (b) Geometric structure of the M011O56H34 cluster modeling a one-layer section at the MoO2(011) surface. The central molybdenum atom and its neighboring nonequivalent surface oxygen centers, 0(l-3), are labeled accordingly.

176

shown [212] that clusters as large M07O30H18 are already sufficient to account for local electronic properties of the surface. The M015O56H22 cluster contains all structurally different oxygen centers in their complete surface environment. Further, hydrogen terminators are placed at the cluster periphery to saturate dangling oxygen bonds, thus accounting for electronic embedding of the surface cluster. The representation of the local MoOsCOlO) surface geometry by the M015O56H22 cluster can be tested by geometry optimizations in which all cluster atoms are allowed to rearrange except the peripheral hydrogen terminators. These optimizations yield equilibrium positions for the surface atoms that are very close to those of the bulk termination with inter-atomic distances varying by less than 0.05 A. The electronic structure of the clusters is determined by ab initio DFT methods [35, 213-214] using both local [215] and gradient corrected (RPBE [124, 125]) functional with model core potentials for Mo [216]. In addition to total energies and equilibrium geometries, detailed analyses of the electronic structure in the clusters are performed using Mullilcen populations [118] and Mayer bond order indices [119, 120]. Further, total density of states (DOS) and atom projected partial densities of states (PDOS) are evaluated for the interpretation of experimental photoemission spectra of M0O3 surface systems as discussed below. Table 8 lists geometric and electronic parameters of the M015O56H22 cluster with its geometry taken from the experimental bulk geometry [154] as well as Table 8 Selected atom charges q, bond orders p and distances do-Mo, (A) of the M015O56H22 cluster representing MoOsCOlO) surface. For a definition of the two sets of data, "bulk termination" and "optimized cluster", see text. Results refer to DFT calculations using the RPBE approximation. bulk termination

optimized cluster

q(Mo)

2.23

2.16

q(0(l))

-0.48

-0.48

q(0(2))

-0.74

-0.72

q(0(3))

-0.99

-0.97

1 p(0(l)-Mo) 1 do(i)-Mo

1.66 1 1.67

1.7411.69

p(0(2)-Mo) 1 do(2)-Mo

1.05 1 1.73 0.33 1 2.25

1.13 11.75 0.3112.33

p(0(3)-Mo) 1 do(3)-Mo

0.26 1 1.94 0.44 1 1.94 0.31 1 2.33

0.28 11.97 0.45 1 2.03 0.32 1 2.28

177

for the optimized geometry (all atom positions except for the hydrogen terminators are optimized according to the lowest energy). Results of atom charges from MuUiken populations and Mayer bond orders are given for selected atoms closest to the cluster center. All data are found to differ only little between the bulk terminated and optimized cluster calculations. The populations yield positively charged molybdenum ions, described as Mo^'^"^ with minor charge variation (Aq = 0.2) inside the cluster, and negative oxygen ions. The negative oxygen charges scale with coordination, O^'^" for singly coordinated oxygen 0(1), O^^" for doubly coordinated (asymmetrically) bridging oxygen 0(2), and O^^' for triply coordinated (symmetrically) bridging oxygen 0(3). This result indicates that the largest ionic contributions to binding at the MoO3(010) surface involve symmetrically bridging 0(3) sites, which is of importance in view of the reactivity of the different sites with respect to chemical reactions. Altogether, local charging of the different cluster atoms is found to be much smaller than formal valence charges suggest. This indicates that inter-atomic binding in M0O3 is described by both ionic and sizeable covalent contributions. Covalent contributions to Mo-O binding can be estimated roughly from corresponding bond order indices. The data of Table 8 confirm the general picture based on simple valence concepts revealing three types of Mo-O bonds. Bonds between terminal oxygen, 0(1), and Mo yield bond order values close to 2, which suggests double bonds and is consistent with the single coordination of 0(1). The asynunetrically bridging oxygen 0(2) binds to two neighboring Mo centers with bond order values of 1.0 for the center at closer and 0.3 for that at further distance, corresponding to one single and one much weaker bond that is reasonable in view of the environmental geometry of the 0(2). The symmetrically bridging oxygen 0(3), which is coordinated to three neighboring Mo centers (two of the upper and one of the lower part of the bi-layer), yields bond order values of 0.3 - 0.4 indicating three bonds which are similar in strength and considerably weaker than single bonds. If the total amount of covalent binding of each oxygen center with its Mo neighbors is defined by the sum of bond orders, the highest value is obtained for 0(1) (sum = 1.7), followed by 0(2) (sum = 1.4) and 0(3) (sum = 1.0). This correlates nicely with the increase in negative charge from 0(1) to 0(2) to 0(3). Altogether, the results suggest for the different oxygen centers at the MoO3(010) surface that an increase in coordination with respect to Mo neighbors leads to an increased ionic and decreased covalent binding character. The calculated energy of the highest occupied orbitals (HOMO), 8HOMO, of the M015O56H22 cluster amounts to 6.1 eV. Based on general DFT theory [35] the quantity -8HOMO represents the first cluster ionization potential which converges with increasing cluster size towards the workfunction of the MoO3(010) surface. The theoretical value of -8HOMO is reasonable in view of typical work-

178

function values of transition metal oxide surfaces (e. g. 6.4 ± 0.7 eV for ¥205(010) [94, 217]), while experimental workfunction values for MoOsCOlO) do not seem to exist in the literature. The character of the HOMO and lowest unoccupied orbital (LUMO) of the M015O56H22 cluster becomes obvious from population results. The analysis shows that the HOMO is described by dominant 2sp contributions of oxygen at the cluster periphery while the LUMO can be characterized by 4d contributions of molybdenum. This indicates that optical absorption leading to electronic HOMO-LUMO excitations is combined with an oxygen to metal charge transfer at the MoO3(010) surface. Further, adding electronic charge to the surface will result in an increased 4d occupation of the metal ions at the surface. This can be connected with chemical reduction and is important later on in connection with oxygen vacancy formation. The electronic structure of the M015O56H22 cluster can also be characterized by energy level distributions of the occupied Kohn-Sham valence orbitals. Fig. 24 shows the total DOS curve for the valence band region of the cluster together with its PDOS decomposition into Mo and O contributions. The DOS of the valence band region extends from -12.5 to -6.1 eV and is described by a multi-peak structure (regions A-D). The decomposition shows clearly that dominant contributions are due to oxygen while molybdenum contributions are smaller and confined to the lower part of the valence band region. Population analyses identify the oxygen contributions as O 2p and those of molybdenum as Mo 4d that is to be expected based on the valence configuration of the atoms

(a) total Mo 0

4.0

B . 0

CO

02.0

Q

A

1

0.0 -18.0

^

E]

/"' '

-14.0 -10.0 Energy [eV]

-6.0

-18.0

-14.0 -10.0 Energy [eV]

Fig. 24. Total DOS and atom projected PDOS curves of the M015O56H22 cluster, (a) Total DOS (thick solid) with decomposition into Mo (thin solid) and O (dashed) contributions, (b) PDOSs of the differently coordinated oxygen centers 0(1) (solid), 0(2) (dotted), and 0(3) (dashed). A gaussian level broadening of 0.4 eV is applied and the energetic position of the highest occupied cluster orbital EHOMO at -6.1 eV is marked by a thin vertical line.

179

and their formal valence in M0O3 bulk. There is an additional set of levels between -15.0 and -12.8 eV (region D, below the valence band region) which is due to split-off orbitals as a consequence of the bond saturation of peripheral oxygen by hydrogen in the clusters and therefore has to be considered a cluster artifact. Altogether, the valence band region assumes a width of A = 6.4 eV in the present cluster. Width A is expected to converge with increasing cluster size towards the total valence band width of the extended MoOsCOlO) surface. This is consistent with recent MoOsCOlO) slab and bulk [206, 218] calculations as well as experimental Hell photoemission studies [206]. The valence band region can be characterized in more detail using PDOS curves shown in Fig. 24b. Here the oxygen PDOS is decomposed into contributions of the differently coordinated surface oxygen sites 0(l-3), allowing a distinction of the three main peak regions A, B, C in the total DOS curve. The energetically highest region A (about 2 eV below the upper valence band edge, £HOMO) is described by contributions from oxygen of all three coordination types with that of 0(3) being most pronounced. The central region B (~ 4 eV below £HOMO) is characterized by oxygen with dominating 0(1) contributions in addition to small molybdenum ingredients. The lowest region C (about 5.5 eV below 8HOMO) has dominant 0(3) as well as molybdenum character. Region E, denoting unoccupied levels and simulating the conduction band region of the substrate within the cluster approach, starts at 1.8 eV above the upper valence band edge. This is consistent with the well known experimental result of bulk M0O3 forming a small band gap insulator. Fig. 24a shows in addition that the lower part of region E is dominated by molybdenum contributions that are described as Mo 4d type based on orbital analyses. This confirms the qualitative picture of the electronic structure of M0O3 being described roughly by a Mo 4d type conduction band region in addition to the O 2sp type valence band region. The present conclusions about electronic properties of the MoO3(010) surface are confirmed by other cluster model studies using ab initio HartreeFock methods [219] together with full geometry optimization. The results from the M02O11H10, M03O16H14, and M07O32H22 surface clusters confirm mixed ionic and covalent Mo-0 binding and suggest symmetric bridging oxygen as a probable active site for insertion into propylene. Further, the electronic structure of the MoO3(010) surface has been treated by periodic DFT calculations using plane-wave pseudopotential methods in a repeated slab model [218]. Here the binding analysis based on crystal orbital overlap populations indicates also that Mo-O binding is a combination of ionic and covalent contributions for all oxygen species. However, the analysis suggests that covalent Mo-O binding is stronger for the terminal oxygen that for either of the two bridging oxygens [218].

180

Relationships between chemical composition, surface geometry, and electronic properties of molybdenum oxide surfaces can be discussed by comparing surfaces of different orientation. As an example, electronic properties of the MoOsClOO) surface have been calculated for appropriate surface clusters, such as M06O24H12 shown in Fig. 21, using DFT methods [220]. A comparison of these results with those of MoOsCOlO) derived surface clusters [220] shows that electronic differences between both surfaces are mainly due to different atom arrangement while local charging and binding properties seem surface independent. In particular, geometrically equivalent oxygen centers at the MoOsCOlO) and (100) surfaces are similar in their electronic behavior. Therefore, the electronic structure at the two surfaces is determined mainly by their detailed atom arrangement and is not influenced by major charge redistributions due to substrate surface binding. As mentioned above, the MoOsCOlO) surface has been found to be chemically inert at room temperature. Valence band photoemission spectra for this surface [201] show negligible emission in the bulk band gap region. However, when the surface is exposed to electron or ion bombardment, even of relatively low energies, or to ultraviolet photons [221] the surface becomes reduced. Many reduced oxides of molybdenum have been observed and found to exhibit significant emission in the bulk band gap region. The band gap emission must involve electron states of predominant Mo 4d character. Angle-resolved UPS studies on band gap defect states conclude [202] that these states are associated with extended defects rather than with point defects, and are presumably due to the shear-plane structure of the reduced phases. One example of these phases is M0O2 resulting from deep reduction of molybdenum trioxide and being always present as domains at M0O3 surfaces [222, 223]. The reduction effect can be seen in model studies using surface clusters which simulate local sections of the MoO2(011) surface. As an illustration. Table 9 lists atomic charges from Mulliken populations and Mayer bond orders for the Moii039H34cluster cut out of the MoO2(011) surface (see Figs. 22, 23b). The data are compared with those of a M015O56H22 cluster modeling the MoO3(010) surface. The comparison suggests local charging and Mo-O binding which is quite similar between the MoO2(011) and MoO3(010) surface. In the MoO2(011) surface cluster positive molybdenum ions are described by Mo^'^"^ (6-fold coordinated) and Mo''^^ (5-fold coordinated) while negative oxygen charges scale with coordination, O^^' for doubly and O^'^" for triply coordinated oxygen. (Note that at the monoclinic MoO2(011) surface there are two very slightly differing 0(2) and 0(3) species and the results in Table 9 refer to averages.) As for the MoO3(010) surface clusters, charging of the different cluster ions in the M011O39H34 cluster is much smaller than formal valence charges of the M0O2 substrate, Mo"^"^ and O^", suggest. In addition, charging of the 5- and 6-fold coordinated Mo ions differs only by little. There-

181

Table 9 Comparison of selected atom charges q, bond orders p, and distances do-Mo, (A) of two surface clusters representing MoOiCOll) and MoOsCOlO) surfaces, respectively. The results refer to DFT calculations using the LDA approximation. MOii039H34for

M015O56H22 for

MoO2(011)

MoOsCOlO)

q(Mo(6))

1.28

1.89

q(Mo(5))

1.54

1.89 -0.43

q(0(l)) q(0(2))

-0.52

-0.62

q(0(3))

-0.79

-0.87 1.73 1 1.67

p ( 0 ( l ) - M 0 ) 1 do(2)-Mo p(0(2)-M0) 1 do(2)-Mo

0.87 1 2.06 0.92 1 2.07

1.21 1 1.73 0.42 1 2.25

p(0(3)-M0) 1 do(3)-Mo

0.64 1 2.06 0.54 1 2.07 0.59 1 1.98

0.38 1 1.94 0.50 1 1.94 0.38 1 2.33

fore, discrimination between the two should be based on covalent rather than ionic binding behavior. A comparison between the Mo ion charges (determined from populations) of the MoOaCOl 1) and MoOsCOlO) surface clusters shows that those of the reduced oxide (M0O2) are decreased by only 0.3 electrons. This is in contrast to the standard chemical characterization in terms of formal valence charges and confirms further the importance of covalent binding in the substrates. The general picture based on atom coordination at the MoO2(011) surface is reflected in the bond order results listed in Table 9. Bonds between doubly coordinated oxygen and each of its two molybdenum neighbors result in bond orders close to 1 indicating single bonds. Further, Mo-O bonds involving triply coordinated oxygen yield bond orders of 0.5-0.6 revealing weaker than single bonds. A comparison of bond orders in the M0O2 and M0O3 surface clusters suggests that in the case of M0O2 covalent binding becomes more important at the expense of ionicity which it consistent with molybdenum and oxygen charges of the dioxide being smaller than those of the trioxide. Additional studies have been published to explain factors that influence solid acidity in pure molybdenum oxides and those supported on silica-alumina [224]. These studies are based on cluster models of respective acid sites where electronic properties are evaluated by self-consistent HF, electron-correlated

182

M0ller-Plesset (MP2) and local density functional (LDA) methods. Here it was possible to identify relationships between structural transformations in the molybdenum trioxide tetrahedra and changes in Bronsted-Lewis acidity. In addition, the experimentally observed influence of the support composition on the acid strength of the supported molybdenum oxide catalyst was interpreted by the calculated charge redistribution and by molecular orbital energies. The proposed mechanism of acidity changing with molybdenum loading was supported by the agreement between calculated and measured infrared vibration frequencies.

(b) Surface oxygen vacancies The oxidation of organic molecules at the surface of a molybdenum oxide catalyst, which involves surface oxygen, creates surface oxygen vacancies. Here the initial state of the catalyst may be restored by molecular oxygen from gas phase (reoxidation) or the density of oxygen vacancies may grow and give rise to extended (rather than point) defects [14, 154, 158]. These defects may lead to the annihilation of oxygen vacancies by formation of CS planes. For M0O3 surfaces this process is postulated to be essential for the ability of the oxide to in(a)

0(1) vacancy

0(2) vacancy

0(3) vacancy

(b) 0(1) vacancy

0(2) vacancy

0(3) vacancy

Fig. 25. (a) Sketch of vacancies at different surface oxygen sites 0(l-3) in the M015O56H22 cluster. Cluster atoms of the top (bottom) part of the bi-layer are shown as shaded (white) balls where the ball size decreases from Mo to O to H. The vacancies are sketched as black dots, (b) Relaxed geometry of the M015O56H22-O (=Moi5056H2i ) clusters with different vacancy sites.

183

sert oxygen into the adsorbed species in a selective oxidation process [12, 33, 225]. This is supported by quantum chemical calculations on M02O10 clusters [196] where changes in topology by transforming M0O6 octahedra from comersharing (with a vacancy) to edge-sharing favor the edge-linked geometry. However, formation of CS planes cannot exclude the existence of oxygen vacancies at M0O3 surfaces [226]. Therefore, studies on electronic properties in the presence of surface oxygen vacancies are of a great importance. In recent cluster studies [212] the electronic and geometric structure of the three different types of oxygen vacancies, that can exist at the ideal MoOsCOlO) surface, have been examined by DFT calculations using the RPBE approach. These studies are based on a M015O55H22 substrate cluster where oxygen is removed from one of the three sites 0(l-3), see Fig. 25a. In a first set of calculations the electronic structure of the indented cluster is evaluated with its atom centers kept fixed at their positions of the bulk termination (frozen geometry). In subsequent calculations all atoms of the indented cluster (except the peripheral hydrogen terminators) are allowed to relax in response to the oxygen removal (relaxed geometry). As a result. Table 10 lists oxygen vacancy energies, ED^(O), EDXO) corresponding to frozen and relaxed geometries. The vacancy energies are determined by appropriate cluster total energy differences using the ground state of atomic oxygen as a reference. Obviously, oxygen vacancy energies are rather large for all surface oxygen sites. This indicates that it is quite difficult to remove oxygen by itself from the ideal MoO3(010) surface. The frozen geometry value ED^(0) is largest, 7.6 eV, for the singly coordinated terminal 0(1) site, while for two bridging sites 0(2), 0(3) smaller values, 7.1 and 6.8 eV, are obtained. Table 11 lists selected atom charges from MuUiken populations and Mayer bond orders of the M015O56H22 based vacancy clusters representing the different oxygen vacancies 0(1-3) at the MoO3(010) surface (frozen geometries). For comparison, corresponding results for the vacancy free M015O56H22 cluster with bulk termination (denoted "substrate") are included. Table 10 Results of vacancy energies, E/'^(0), of the different oxygen surface sites 0(l-3) with and without surface relaxation, for definitions see text. The data are obtained from DFT calculations on the M015O56H22 cluster representing a local section of the MoO3(010) surface. All energies are given in eV. Site 0(1) 0(2) 0(3)

ED'(0)

EDXO)

7.64 7.09 6.81

4.84 5.10 6.35

184

Table 11 Selected atom charges q, bond orders p, and distances do-Mo (in A) of the M015O56H22-O (= M015O55H22) clusters representing oxygen vacancies 0(1-3) at the MoO3(010) surface. The results refer to DFT calculations using the RPBE approach and frozen cluster geometries. The fourth column, denoted "substrate", gives data of the M015O56H22 cluster for comparison 0(1) vac.

0(2) vac.

0(3) vac.

substrate

q(Mo)

1.57

1.32

1.70

2.23

q(0(l))

--

-0.46

-0.48

-0.48

q(0(2))

-0.66

~

-0.73

-0.74

q(0(3))

-0.96

-0.96

--

-0.99

p(Mo-0(l))ldMo-o(i)

--

1.9711.67

1.75 1 1.67

1.6611.67

p(M0-0(2)) 1 dMo-0(2)

1.45 11.73

~

1.11 11.73

1.05 11.73

p(Mo-0(3)) 1 dMo-o(3)

0.4611.94

0.3611.94

-

0.2611.94

The atom charges evidence that the existence of a surface oxygen vacancy always reduces the central Mo atom(s) closest to the vacancy (i. e. positive Mo charges are decreased) where the effect amounts to 0.5 - 0.9 electrons. In contrast to the sizeable charge reduction at the metal ions, the oxygen ions near the vacancy experience only small changes in their charge state. In all cases, oxygen ions become less negative with charge differences below 0.1 electrons. A comparison of the bond order results for the different oxygen vacancy clusters with those of the vacancy-free cluster shows that the largest effects occur when the 0(1) or 0(2) vacancies are created. For the 0(1) vacancy the bond order between the central Mo center of the asymmetric bridging oxygen 0(2) increases considerably, suggesting a Mo-0(2) bond strength near the 0(1) vacancy similar to that of the original Mo-O(l) bond. Thus, in terms of Mo-0 binding the asymmetric bridging oxygen 0(2) takes the place of the original terminal 0(1) atom, which has been removed in the vacancy creation process. On the other hand, the existence of an 0(2) vacancy results in strengthening of a nearby Mo-O(l) bond. These findings become important for the vacancy properties when surface relaxation is included in the vacancy calculations [211-212, 227]. Geometric consequences of substrate relaxation due to vacancy formation at the 0(1-3) sites are sketched in Fig. 25b for the M015O56H22 based vacancy cluster. A comparison with the initial cluster geometry, shown in Figs. 23a, 25a, evidences that in all cases relaxation is locally confined. The response of the substrate to the creation of a triply coordinated 0(3) vacancy is rather small. In contrast, removing an oxygen from the terminal 0(1) or doubly coordinated 0(2) sites leads to significant changes, where for both sites the same relaxed

185

substrate geometry is achieved. Here a terminal oxygen species, which originates from either the 0(1) or 0(2) site (complementary to the site where the oxygen has been removed), stabilizes in the interstitial region between the two sites. This suggests that a clear distinction between the two types of oxygen vacancies may not be possible in the experiment. An alternative view of the chemical reduction effect at the metal sites near oxygen vacancies becomes possible by orbital analyses and atom projected partial densities of states (PDOS) of the vacancy clusters. As an example. Fig. 26 shows DOS and PDOS curves of the valence band region obtained for the M015O56H22 based vacancy cluster with a triply coordinated 0(3) vacancy. A detailed comparison of Fig. 26 with Fig. 24 (giving the (P)DOS results of the vacancy-free cluster) shows only very small differences over the whole range of the oxygen 2sp dominated valence band region except for its upper edge. In the region near the HOMO at -6.1 eV the occupied orbitals of the vacancy-free cluster are determined by O 2sp contributions whereas in the vacancy cluster the highest occupied orbitals (below 5.3 eV) are characterized as Mo 4d type. An analysis of these metal derived valence orbitals (unoccupied in the vacancy-free cluster) reveals further that their charge distribution is located mainly at the three Mo sites adjacent to the vacancy. Thus, the calculations show that the oxygen vacancy leads to increased local molybdenum 4d occupation. This result is also found in the (P)DOS and orbital analyses of the vacancy clusters for 0(1) and 0(2) sites. It confirms the picture of chemical reduction of adjacent metal

-14.0 -10.0 Energy [eV]

14.0 -10.0 Energy [eV]

Fig. 26. Total and atom projected (P)DOS curves of the valence band region fo the M015O56H22 based vacancy cluster with an 0(3) surface vacancy, (a) Total DOS (thick solid) with decomposition into Mo (thin solid) and O (dashed) contributions, (b) PDOSs of the differently coordinated oxygen centers 0(1) (solid), 0(2) (dotted), and 0(3) (dashed). A gaussian level broadening of 0.4 eV is applied and the energetic position of the highest occupied cluster orbital £HOMO at -5.3 eV is marked by a thin vertical line.

186

centers induced by oxygen vacancies and observed already in the population results. The existence of additional occupied states of Mo character, located above the O 2sp derived valence region, is relevant for the interpretation of experimental photoemission spectra of molybdenum oxide surfaces. According to the results of the cluster studies additional photoemission intensity above the valence band region may be indicative of chemical reduction of the metal centers, leading to lower oxidation states, where the effect can be introduced by oxygen vacancies or by different chemical composition of the oxide. This has been verified in UPS experiments on differently prepared MoOsCOlO) surfaces in comparison with measurements of other single and mixed valency molybdenum oxide samples [212].

(c) Adsorption Atomic and molecular adsorption at molybdenum oxide surfaces have been studied theoretically using both periodic slab and cluster models where so far studies are restricted to the trioxide, M0O3, as a substrate. Further, adsorbate species include in all cases atoms (H [138, 218, 227-229]) or small molecules (methanol [230, 231], CO [206], H2O[206], CH2 [232], CH3 [232], CH4 [233], OCH3 [234-236], C3H5 [228, 229]) which are of catalytic interest. Hydrogen adsorption at the MoO3(010) surface has been examined extensively by both cluster [138, 227-229] and slab models [218]. Under reductive or oxidative conditions of catalytic reactions, the surface/bulk structure of M0O3 or M0O2 can transform from one oxide to the other rapidly [237]. There are several possible processes that lead to reduction of M0O3 to lower valency oxides. One process starts with adsorption of hydrogen at surface oxygen sites and results in surface hydroxyl or water species. Under reaction conditions these surface species may desorb leaving oxygen vacancies behind. Here it is interesting which of the differently coordinated oxygen sites at the M0O3 surface is dominantly involved in hydrogen adsorption and OH/H2O desorption, see e. g. [238]. This question has been addressed in cluster model studies of hydrogen adsorption at 0(1-3) sites of the MoO3(010) surface, see Fig. 18. Atomic hydrogen is added to one of the different surface oxygen sites, 0(1-3), of a M015O56H22 model cluster, see Fig. 21, and the local geometry of the resulting surface OH species is optimized in DFT total energy calculations [227] using the RPBE functional. Fig. 27 visualizes calculated equilibrium geometries of the different surface OH species. In addition, Table 12 lists electronic and energetic parameters from the calculations. Here EB(H) denotes the adsorption energy of hydrogen at a given

187

0(2)H

(010) view

0(3)H

0(1 )H

(001) view

Fig. 27. Sketch of calculated equilibrium geometries of hydrogen adsorption at different oxygen sites 0(1-3) of the MoOsCOlO) surface, yielding surface OH. The results are obtained from DFT optimizations using M015O56H22H adsorbate clusters. Darker shaded balls show the surface OH species while light balls refer to the surface lattice environment.

oxygen site 0(l-3), calculated from corresponding total energy differences, and Table 12 Selected atom charges q from populations and Mayer bond orders p of the M015O56H22H clusters modeling adsorption of hydrogen at surface sites 0(l-3) at the MoO3(010) surface. The table includes energies EB(H) for H adsorption and ED(OH) for OH desorption, based on relaxed cluster geometries and obtained by DFT calculations using the RPBE functional. Further, data for the substrate M015O56H22 cluster as well as for isolated OH are included. 0(1)H

0(2)H.

0(3)H

Surface / OH

Q(Mo)

2.17

1.84

2.11

2.23

Q(0(1))

-0.81

-0.47

-0.46

-0.48

Q(0(2))

-0.74

-1.04

-0.73

-0.74

Q(0(3))

-0.99

-0.98

-1.18

-0.99

Q(H)

0.49

0.56

0.59

0.42/0.24

P(Mo-0(l))

1.60

1.67

1.75

1.66

1 P(Mo-0(2))

1.05/0.32

0.98/0.86

0.99/0.35

1.05/0.33

0.25/0.43/0.3

0.26/0.45/0.3

0.12/0.15/0.3

0.26/0.44/0.3

0.71

0.54

0.43

0.80/0.91

-

P(Mo-0(3)) P(H-O) EB(H)

(eV)

E D ( O H ) (eV)

1.62

1.43

1.34

3.92

3.74

3.32

188

is the desorption energy of OH (leaving a surface oxygen vacancy). Table 12 shows that hydrogen can adsorb at all surface oxygen sites where binding is found to be strong (with binding energies of 1.3 - 1.6 eV) suggesting rather stable surface hydroxyl species. The equilibrium geometries of the surface OH groups, see Fig. 27, do not reflect the symmetry of the corresponding oxygen sites. The O-H distances are always quite close to those of typical OH containing molecules (similar to hydrogen adsorption at vanadium pentoxide discussed above). Further, the distances between oxygen and its nearest vanadium neighbors at the M0O3 surface are enlarged. This indicates that hydrogen adsorption weakens the Mo-0 bonds near the adsorption site. The effect is most pronounced for the triply coordinated oxygen site 0(3) where hydrogen is tilted lying almost flat at the surface. This may be an initial step for penetration of the adsorbate into the substrate leading to molybdenum oxide bronzes, see for example [239]. At all oxygen sites, hydrogen adsorption leads to accumulation of negative charge at the corresponding site and surface OH appears as a negative species, see Table 12. In addition, the positive charges at vanadium neighbors are decreased slightly, which can be described as hydrogen induced chemical reduction. The reduction effect is also obvious from (P)DOS curves of the adsorbate cluster where additional occupied Mo 4d type states occur [227]. In analogy to the previous discussion for vanadium pentoxide surface oxygen vacancies at the M0O3 surface may be formed in a two-step process where, in the first step, hydrogen is adsorbed and, in the second, OH is desorbed. Table 12 includes calculated desorption energies ED(OH) for the different oxygen sites. These energies range between 3.3 and 3.9 eV and are all considerably smaller than corresponding oxygen vacancy energies of 4.8 to 6.4 eV, see Table 10. The energy decrease is obviously due to weakening of Mo-0 bonds at the MoO3(010) surface caused by hydrogen adsorption. This suggests that the presence of pre-adsorbed hydrogen facilitates oxygen removal, completely analogous to the findings for vanadium pentoxide. The interaction of atomic hydrogen with the MoO3(010) surface was also studied using periodic DFT methods in pseudopotential plane-wave calculations 218]. Here the inclusion of full surface relaxation is found to be important even for a qualitative description of adsorbate bonding. The results confirm that hydrogen is most strongly bound at the 0(1) site followed by 0(2) and 0(3). Similar to the discussion above, hydrogen was predicted to diffuse into bulk forming molybdenum oxide bronzes, HxMo03 [1, 239]. Methanol adsorption at MoO3(010), (001), and (100) surfaces has been studied in cluster models using the semi-empirical ENHT method [230]. The calculations suggest that both molecular and dissociative adsorption may occur at the different surfaces. The stability of molecular adsorption complexes increases from the (100) surface to (001) and (010) if the latter contains oxygen ED(OH)

189

vacancies. At the (100) and (001) surface, which contain unsaturated oxygen, dissociative adsorption is energetically preferred over molecular adsorption. The oxidative dehydrogenation of methanol to formaldehyde on M0O3 surfaces was also studied by quantum chemical ab initio methods [231]. It was concluded that dual dioxo Mo=0 sites (existing at MoO3(010)) are necessary for the reaction to happen where each of these sites extracts one hydrogen in a sequence of steps. The importance of 0x0 Mo=0 groups for the olefin methathesis reaction as well as for hydrocarbon oxidation was also pointed out by quantum chemical studies on small molecular models, such as CI2M0O2 [240, 241]. Theoretical studies on adsorption and oxidation reactions of CH3 and CH2 fragments at the MoO3(100) surface [232] have also been connected with methanol adsorption. The adsorption energy and binding of these fragments was computed using a methodology based on the semi-empirical atomic superposition and electron delocalization molecular orbital theory (ASED-MO). Here both homolytic and heterolytic mechanism of C-H bond breaking are discussed. On the MoO3(010) surface, exposing mainly oxygen ions, the interactions between the CH3/CH2 fragments and the surface is found to be weak. In contrast, on other M0O3 surfaces, containing bare molybdenum ions, lone pair orbitals of the fragments can hybridize with the metal ions to form stronger C-Mo bonds. The adsorption of H2O and CO has been studied using periodic ab initio Hartree-Fock and correlated methods [206] where the five-fold coordinated molybdenum centers at the MoO3(100) surface were chosen as adsorption sites. The calculations show that adsorption of both molecules is rather weak and determined by electrostatic binding. Numerous theoretical studies are devoted to details of C-H activation at M0O3 surfaces. In particular, C-H activation and bond breaking in adsorbed CH4 at the MoO3(010) surface was considered in periodic DFT studies using pseudopotential plane-wave methods [233]. The calculations indicate that methyl is adsorbed most strongly above terminal oxygen sites 0(1) at MoO3(010) resulting in a surface methoxy species. On the other hand, methyl adsorbing above the asymmetric bridging 0(2) site yields formaldehyde with hydroxyl binding near the 0(1) site. Finally, methyl above the symmetric bridging 0(3) site reacts to form a formyl species with water near the 0(1) site. The understanding of C-H bond activation was also the goal of theoretical studies on methoxy adsorption at MoO3(100) surfaces based on a M03O11 cluster model using the serm-empirical ASED-MO method [234, 235]. In particular, the mechanism of C-H bond breaking during the reaction of methane near a surface oxygen vacancy is considered where methane is found to adsorb dissociatively. Further, photon assisted C-H activation in methoxy, adsorbed near bare molybdenum centers at M0O3 and M08O24 surfaces, has been treated by ASEDMO calculations on a Mo30i3^" model cluster [236].

190

Adsorption of allyl, C3H5, species at the MoOsClOO) and (100) surfaces has been examined in DFT cluster studies [228, 229]. Altogether, the adsorptive interaction is found to be rather weak at both surfaces. At the (010) surface, the adsorbate binds most strongly above asymmetric bridging oxygen sites 0(2) where the allyl plane extends perpendicular to the surface. This yields a weakening of both C-H and adjacent C-C bonds while the C-C bond further away becomes stronger. The asymmetric C3 skeleton is qualitatively similar to that in acroleine and can be interpreted as an initial step in the oxidation process. At the (100) surface, the adsorbate binds most strongly above bare molybdenum sites with the allyl plane parallel to the surface. As a result, there is no trend towards asymmetric C-C bonds which makes the allyl oxidation less likely to happen. The different activity of the MoO3(100) and (100) surfaces with respect to allyl adsorption was confirmed further by charge sensitivity analyses [242].

SYNOPSIS Vanadium and molybdenum oxides, which represent commercially important materials, form a fascinating class of compounds concerning their crystal structure as well as their electronic properties. The knowledge of these properties is a necessary prerequisite to understand their physical and chemical behavior. In single valency oxides of both vanadium and molybdenum the building units are distorted Me06 octahedra linked by comers or/and edges. The richness of arrangements of these octahedra is a source for the existence of differently coordinated oxygen ions and has an important impact for surface properties of the oxide systems. Despite the importance of vanadium and molybdenum oxides as technologically relevant materials many details of their microscopic behavior are still under debate. According to theoretical studies, single valency oxides are compounds described by mixed ionic and covalent character of metal - oxygen bonds. Charging of metal and oxygen ions is smaller than expected from their formal valence charges and oxygen ionicity is found to scale with the coordination numbers of the corresponding sites. The character and energetic distribution of the highest valence orbitals of these oxides, characterized also by densities of states (DOS), are correlated with the formal oxidation state of the metal ions. The analysis reveals differences in electronic properties as well as transformations from metal to semiconducting and insulating behavior for different oxidation states. Theoretical studies on surface oxygen vacancies show that the oxygen species is bound quite strongly to the surface. Thus, it is rather difficult to remove surface oxygen by itself. However, pre-adsorbed atoms or molecules, atomic hydrogen serves as an example, can facilitate oxygen removal resulting in surface vacancies. Further, surface oxygen vacancies are able to diffuse at the

191

oxide surface or move into the bulk and they can be easily refilled by adsorbing oxygen from gas phase. In conclusion, the present theory on vanadium and molybdenum oxide surfaces, while it is not yet fully developed, can already give important input towards a microscopic understanding of their physical and chemical behavior.

ACKNOWLEDGEMENT This work has been supported by the Deutsche Forschungsgemeinschaft through its Sonderforschungsbereich 546, and by grant No. 3T09A 14615 of the State Committee for Scientific Research in Poland.

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Chapter 5

Geometry of adsorbates on metal oxide surfaces R. Lindsay, B.G. Daniels and G. Thornton Surface Science Research Centre and Chemistry Department, Manchester University, Manchester Ml3 9PL, United Kingdom 1. INTRODUCTION Currently, considerable research effort, both basic and applied, is being devoted to understanding surfaces of metal oxides on an atomic scale. This work is partly motivated by the huge significance of these surfaces in a variety of commercially important applications which can potentially benefit from such detailed knowledge. These include heterogeneous catalysis, corrosion inhibition, electronic component manufacture and the emerging fields of nanoand bio- technologies. Along with other chapters in this volume, this contribution is concerned with fundamental studies on such surfaces, performed on single crystal metal oxides (or highly oriented thin films). Here we focus on quantitative structural determinations of adsorbates on these surfaces, knowledge of which is a necessary prerequisite for comprehending, and eventually predicting various surface phenomena e.g. surface reaction mechanisms. There is a wide variety of experimental surface science techniques for quantitative structural determinations (see for example Refs. 1 and 2), and many of them have been applied to studying adsorbate/metal oxide systems, as will be seen below. It should be noted that we have excluded work involving scanning probe microscopy (SPM) which normally provides only a qualitative or at best a semi-quantitative picture of adsorption geometry. The extent to which adsorption geometries have been elucidated varies widely, ranging from merely extracting the angular orientation of an adsorbate to precise determination of the atomic coordinates. The amount of structural information obtained depends both upon the experimental technique employed, and the depth of analysis. Compared to the sizeable volume of quantitative structural determinations of adsorbates on metals and semiconductors (see for example Ref 3) the number

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of similar studies on metal oxides is somewhat restricted. In our opinion this limited archive of literature is not due to a lack of interest in this area but rather the following: (i) (ii) (iii) (iv)

Insulating nature of metal oxides Difficulties with surface preparation Degradation of surface during measurement Lack of structural data for the clean surfaces

The insulating nature of metal oxides (point (i)) has restricted the number of such studies since a majority of the relevant experimental techniques involve the incidence and/or emission of charged particles, usually electrons. In an insulator this loss/gain of charged particles normally results in significant surface charging due to the low electrical conductivity of the solid, and this severely distorts the data, making structure determination usually an intractable problem. For some oxides the problem can be overcome simply by increasing electrical conductivity via reduction of the bulk through thermal treatment (e.g. Ti02) [4], whereas for other metal oxides the problem is less easily solved. One way around it is to employ photon in/photon out techniques such as surface x-ray diffraction (SXRD) or fluorescence yield surface extended x-ray absorption fine structure/near edge x-ray absorption fine structure (SEXAFS/NEXAFS). Another option is to conduct studies on highly ordered ultra thin metal oxide films grown on conducting substrates (e.g. NiO(100)/Ni(100) [5]). This approach allows one to obtain surfaces which can mimic, if prepared correctly, the geometric and electronic structure of a single crystal metal oxide surface rather well, whilst avoiding charging problems. Difficulties with surface preparation (point (ii)) mean that for many metal oxides it has been, and still is in many cases, a significant task just to prepare clean metal oxide surfaces reproducibly and reliably. In fact, currently one could state that the preparation of most metal oxide surfaces, other than those of some well studied binary metal oxides, largely remains something of a black art. Commonly, in situ surface preparation involves cycles of inert gas ion sputtering and high temperature annealing, perhaps followed by annealing in a partial pressure of O2 to restore surface stoichiometry. However, for compositionally more complex metal oxides, including high temperature superconductors, the most promising techniques are either in situ cleavage or thin film growth. An excellent illustration of the difficulties associated with the surface preparation of metal oxides is provided by the (110) surface of rutile Ti02, which has long been considered a prototypical metal oxide surface. From qualitative low energy electron diffraction (LEED) data it was believed.

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for a number of years, that one could relatively easily prepare the unreconstructed (1x1) surface phase [4]. Recent scanning tunneling microscopy (STM) studies (e.g. Refs 6,7), however, have demonstrated that LEED alone cannot be relied upon to confirm unequivocally surface quality. STM images show that surface geometry depends strongly upon both in situ preparation conditions, and the thermal history of the sample, in particular the extent of bulk reduction of the single crystal. Degradation of metal oxide surfaces during measurement (point (iii)), is seemingly rather a common problem, in particular for molecular adsorbate covered surfaces. The damage occurs due to impinging (or outgoing) photons/electrons/ions destroying the surface through one or more of three processes: desorption, dissociation and disordering. For example the (2x1) LEED pattern formed following acetic acid adsorption on Ti02(l 10)1x1 reverts to (1x1) within 30 seconds using standard LEED optics [8]. Destruction on this timescale normally renders the collection of appropriate data almost impossible, since data acquistion times are normally significantly longer. There are, however, methods to circumvent this problem. For example for LEED studies, like the one just described, one may utilise a modified LEED optics, which operates with electron beam current densities several orders of magnitude lower than normal, thus exposing the sample to a much lower dose of electrons during the measurement. The viability of this approach has been demonstrated for weakly bound molecular adsorbates on MgO(lOO) [9,10]. Lack of structural data for clean metal oxide surfaces (point (iv)), has also hindered quantitative structural determinations of adsorbates on these surfaces. This is particularly true for those techniques providing a rather complete description of the adsorption geometry (e.g. quantitative LEED (LEED-IV)), since at least some knowledge of the clean surface geometry is usually necessary for structure elucidation. The reasons for the lack of structural data for clean metal oxide surfaces are basically identical to the previous three points (i-iii) for adsorbate covered surfaces. As the existence of this review demonstrates, the various experimental problems, described above, have not discouraged everybody from attempting to perform structural determinations of adsorbates on metal oxide surfaces. The following sections of this paper detail such work. It is structured such that there is a separate section on each metal oxide for which quantitative structural determinations of adsorbates have been carried out. We have restricted ourselves to only those studies involving monolayer/submonolayer adsorbate coverages, but have included atoms, molecules and metal adsorption in this coverage regime.

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2. MAGNESIUM OXIDE - MgO Under normal conditions MgO exhibits the rocksalt structure. Cleavage of MgO occurs preferentially along the non-polar (100) face (see Fig. 1), revealing a relaxed surface which is not far removed from the ideal bulk termination. The other low Miller index surfaces, (111) and (110), have significantly higher surface energies, and have a tendency to reconstruct to display (100) facets, at least in vacuum [11]. Only the structure of adsorbates on the (100) surface have been investigated quantitatively so far. The primary hurdle to overcome for such studies is that MgO is highly insulating. Although thin films of MgO(lOO) have been utilised to circumvent this problem in a number of other surface studies, most notably by Goodman et al (see for example Ref. 12), somewhat surprisingly nobody has exploited this technique for quantitative structural determinations of adsorbates. The most probable reason for this concerns the degree to which the films are representative of the single crystal (100) surface [13]. 2.1. MgO(lOO) - H2O Water adsorption on MgO(lOO) has been subject of two recent quantitative structural determinations [9,14]. The first of these is a LEED-IV study of a p(3x2) phase [9] formed following exposure of a vacuum-cleaved surface at 200 K to 5x10'^ mbar of water for 30 min, and then maintaining a water partial pressure of 1x10"^ mbar. This procedure resulted in a stable water overlayer of coverage 1 monolayer (ML) (where 1 ML is equal to the number of surface

Fig. L A schematic plan view of the (100) surface of MgO (or NiO).

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Fig. 2. Theoretically derived minimum energy configuration for a (3x2) unit cell of H2O on MgO(lOO) [9].

Mg cations) suitable for LEED-IV measurements, which were performed using low current LEED optics. The automated tensor LEED method was utilised to determine the surface geometry [15]. Both the atomic coordinates of the water molecules and the atoms in the top MgO layer were varied during the optimisation, a total of 54 structural parameters. The results of a semiempirical potential calculation, described in the same paper, were used as the initial geometry for the structural optimisation. A schematic diagram of the minimum energy configuration obtained from this calculation for the (3x2) unit cell is displayed in Fig. 2. The water molecules lie approximately flat (i.e. the H-O-H molecular plane is approximately parallel to the MgO(lOO) surface) with the oxygen atoms sitting almost directly above surface Mg cations at a perpendicular height of 2.11 A. The authors state that this overlayer geometry is adopted as it favours hydrogen bonding between molecules adsorbed in consecutive rows. The optimum, LEED-IV derived, coordinates of the oxygen atoms in the water monolayer are listed in Table 1 along with the coordinates obtained from the calculation (scattering from hydrogen atoms was ignored as is typical in LEED-IV studies of H-containing molecules, since it is a very weak electron scatterer). From this table it can be seen that the experimentally determined structure is very similar (including the estimated errors) to that obtained from the calculation: perpendicular heights of the oxygens range from 2.05 ± 0.05 A to 2.22 ± 0.05 A. The surface MgO layer, which was not optimised in the theoretical calculation, exhibits only small displacements from bulk termination: less than 0.05 ± 0.05 A along the surface normal and of the order of 0.1 ± 0.15 A in the surface plane. One final point to note is that the

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Table 1 LEED-IV derived coordinates of oxygen atoms in water molecules (Mj) for MgO( 100)3x2H2O. The error bars are estimated to be ± 0.15 A for the in-plane x,y distances (see Fig. 2 for definition of x and y axes) and ± 0.05 A for the surface-atom distances z. The positions obtained from semi-empirical calculations are in parentheses (from Ref 9). Molecule

x(A)

Ml

0.19(0.00) 2.88 (2.89) 2.90(3.10) 0.06 (0.00) 5.81 (6.05) 5.89 (5.84)

M2 M3 M4 M5 M6

y(A) 0.20 (0.21) 0.23 (0.00) 2.85 (2.92) 3.28(3.19) 2.85 (2.98) 5.94 (5.90)

z(A) 2.22(2.11) 2.05(2.11) 2.14(2.11) 2.07(2.11) 2.11(2.11) 2.10(2.11)

authors suggest that the results of their structural optimisation should be regarded with some caution due to the limited data set and the large number of parameters. The other structural determination following water adsorption on MgO(lOO) was performed employing normal incidence x-ray standing waves (NIXSW) by Liu et al [14]. Again cleaved MgO(lOO) surfaces were employed. The water dosing conditions were, however, completely different from the study described above. Measurements were recorded following (i) exposure to a partial pressure of 1x10"-^ Torr of water for 3 min at 300 K, and (ii) immersion in water at 300 K for 10 min. Only (1x1) LEED patterns were observed after either preparation, although the intensity variation of the diffraction spots from the sample immersed in water were apparently somewhat different, possibly suggesting a modified surface structure. From previous photoemission data [16], recorded by the same group from similarly prepared surfaces, it is concluded that the adsorbed moiety resulting from both preparations is surface hydroxyl, present at coverages of 1.2 ML and 2 ML for preparations (i) and (ii), respectively (the greater than 1 ML coverages arise simply due to fact that besides the binding of OH" species, formed following H2O dissociation, to surface Mg^"^ sites, the H^ counter ions also form hydroxyl species by attachment to oxygen anions in the surface layer (see Fig. 3). NIXSW data were obtained from the MgO(200) diffraction planes, thus providing information about the position of atoms along the surface normal. The locations of both oxygen and magnesium atoms were elucidated using the O KLL and Mg KLL Auger signals, respectively. A method, developed by Woicik et al [17], was employed to extract the surface contribution from the measured NIXSW profiles, and thus determine the surface geometry. For both

205

d200 "^^ Fig. 3. A cross section of the fully hydroxylated MgO(lOO) surface proposed in Ref 14.

surface preparations it was concluded that the perpendicular separation between the oxygen atom of the hydroxyl and the surface layer is 2.13 ± 0.04 A. Assuming, as Liu et al have done, that the hydroxyls are bonded to the surface Mg cations so that the Mg-0 bond Hes along the surface normal, then this distance of 2.13 ± 0.04 A represents the Mg-OH bond length (Fig. 3). Additionally, they found that surface hydroxylation does not induce any vertical shift in the top layer Mg cations. Somewhat surprisingly, at least initially, the derived Mg-OH distance is equal, within experimental uncertainty, to the Mg-0 bond length in the bulk oxide (2.106 A). The authors note, however, that the Mg-O bond lengths in bulk MgO and Mg(0H)2 have been shown to be the same from diffraction studies [18]. Besides the semi-empirical potential calculations discussed above a number of other theoretical methods have also been used to investigate the adsorption geometry of water on MgO(lOO). These include two ab initio calculations, one employing the Car-Parinello approach [19] the other Hartree-Fock methodology [20], performed on isolated water molecules on the surface. Both conclude that adsorption atop a surface magnesium is favourable, with an adsorbatesurface distance of about 2.1 A. They disagree, however, about the orientation of the H-O-H plane, and neither predict the flat geometry suggested by the semi-empirical calculations; a set of Monte Carlo and molecular dynamics simulations [21] does support this flat geometry. Furthermore, neither of these ab initio calculations find that dissociative adsorption is favourable on ideal terrace sites. In contrast, more recent ab initio calculations (e.g. Refs. 22 and 23), which employ larger surface unit cells involving more than one water molecule, indicate that dissociation can occur on the perfect surface to produce hydroxyl species.

206

2.2. MgO(lOO) - C2H2 Only one quantitative structural determination has been performed for acetylene (C2H2) on Mg(lOO), which employed LEED-IV [24]. This work is particularly significant as it represents the first detailed structure determination of a molecular adsorbate, in this case a physisorbed species on an oxide substrate. Data were recorded from a (2x2) overlayer of acetylene, prepared by adsorption onto a cleaved MgO(lOO) surface maintained at 88 K. The (1/2,0) and (0, 1/2) diffraction spots were missing from the LEED pattern at normal incidence, indicating the presence of two orthogonal glide planes. Structure determination involved the same methodology as described in the last subsection for MgO(100)3x2-H2O. The structure predicted by semi-empirical potential calculations [24,25] was used as initial input for the automated tensor LEED structural optimisation. Fig. 4 shows the most favourable structure arising from the theoretical calculations. This has a (2x2) surface unit cell with two glide planes, and contains two acetylene molecules, the centre of mass of each being directly above a surface Mg cation. In the LEED simulations the positions of the carbon atoms and the surface layer Mg and O ions were all optimised (36 structural parameters). The coordinates of the centres of mass and the orientations of the two molecules obtained from the best fit between theory and experiment (Pendry R-factor = 0.14) are displayed in Table 2. Also shown in this table are data from the theoretical calculations and an earlier neutron diffraction experiment [26]. It can be seen that the fully optimised

Fig. 4. The lowest energy structure for the (2x2) acetylene overlayer on Mg(lOO) obtained from semi-empirical potential calculations [24,25].

207

Table 2 Coordinates of the centres of mass and orientations of the two acetylene molecules (1 and 2) in the (2x2) unit cell on MgO(lOO). The values obtained from LEED-IV data [24], neutron diffraction (ND) [26], and semi-empirical potential calculations [24,25] are tabulated. The definitions of x, y, and O are given in Fig. 4. 6 is the polar tilt of the molecule away from the surface plane, and z is with respect to the surface MgO layer. The two calculated z values are for effective charges of ± 1.2 and ± 2, respectively. Molecule/ Method

x(A)

y (A)

z (A)

O

6

1/LEED 1/ND 1/Theory

0.22 ±0.10 ?

-0.16 ±0.10 ?

2.50 ± 0.05 ?

0.00

0.00

2.49/2.39

67° ±10° 45° ±15° 60°

108° ± 5 ° 90° ±15° 90°

2/LEED 2/ND

2.79 ±0.10 7

3.06 ?

2.50 ± 0.05 ?

118° ±10° 135° ±15°

89° ± 5° 90° ± 5°

2/Theory

2.98

2.98

2.49/2.39

120°

90°

LEED-IV structure does not deviate greatly from the theoretically determined structure. The most obvious difference is the tilt of the C-C axis away from parallel to the surface plane (18° ± 5°) for molecule 1. HC-CH bond lengths of 1.2 ± 0.05 A are deduced from the carbon atom positions. The authors note that within experimental error this distance is identical to the C-C bond length of the free molecule (1.18 A), which is as expected for a physisorbed molecule. Surface Mg and O ions were only displaced by amounts less than the experimental uncertainty. 2.3. MgO(lOO) - CO Quantitative structural information pertaining to the orientation of CO molecules adsorbed on MgO(lOO) has been deduced from polarisation infrared spectroscopy (PIRS) [27]. For this study MgO(lOO) surfaces were created via in situ cleavage, and measurements were performed ki a partial pressure of 5x10"9 mbar of CO at various substrate temperatures in the range 32 K to 56 K. IR spectra were acquired in transmission geometry at an angle of incidence of 45° ±1°, with p- and s-polarisation data being obtained alternately. The spectra indicate that the MgO(100)-CO system undergoes a reversible phase transition at approximately 45 K. This is consistent with helium atom scattering (HAS) work, which shows that the surface symmetry changes from c(4x2) to (1x1) as the substrate temperature is increased from below to above 45 K [28]. The higher temperature phase exhibits only one absorption band,

208

detected only with p-polarisation. From this information it was concluded that CO molecules in the (1x1) phase are either oriented perpendicular to the surface plane or are uniformly inclined with random azimuthal orientation. Three absorption features are present in the spectra recorded below 45 K, one being observed with both s- and p-polarisations, the other two with the latter only. Given these data and the HAS results [28], the authors concluded that there are three CO molecules per c(4x2) surface unit cell, and that the molecular axis of one CO molecule lies along the surface normal, whilst the other two are tilted antiparallel to each other. Simulation of the IR spectra produced a tilt angle of 29° away from the surface normal. Theoretical calculations support the conclusion that there are both tilted and upright CO molecules in the c(4x2) phase (e.g. Refs. 29-31). 2.4. MgO(lOO) - CO2 CO2 adsorption on MgO(lOO) has been the subject of both PIRS [32] and C AT-edge NEXAFS [33] measurements. For the PIRS study, preparation of the adsorbate overlayer involved exposure of an appropriately cleaved MgO crystal at 82 K to a CO2 partial pressure of 1x10"^ mbar, which was maintained throughout the measurements. Interestingly, anisotropic broadening of the diffraction spots in the (1x1) clean surface LEED pattern evidenced preferential step orientation, giving rise to a predominantly single domain adsorbate induced (2V2xV2)R45° LEED pattern, at least on freshly cleaved samples. Orientational information was extracted from the dependence of the intensities of the vibrational bands upon the angle of incidence and polarisation of the incoming IR beam. From these data, in combination with LEED, it was concluded that CO2 forms a herring bone structure on the surface with the molecular tilt angle being determined to be 27° ± 10° with respect to the surface plane, and the twist to be 22° ± 20° out of the direction. A contemporary ab initio cluster calculation predicts [34] that an isolated CO2 molecule is adsorbed atop a Mg cation aligned perpendicular to the surface plane, although a flat lying geometry between two neighbouring cations was only slightly less favourable. Overlayer preparation for the C A^-edge NEXAFS study [33] was radically different to that described above. In this work cleaved MgO(lOO) was exposed to 260 Torr of CO2 for 15 min at room temperature. This procedure resulted in the formation of surface CO32-, as evidenced by a sharp resonance in the NEXAFS spectra at 290.2 eV, which was assigned to the Is - 7C* transition of 003^" on the basis of previous work. The angular orientation of the surface CO32- was examined by recording NEXAFS spectra as a function of polar photon incidence angle. The intensity of the 7C* feature varied very little with photon incidence angle, which indicates that either the overlayer is

209

orientationally disordered, or that the tilt angle of the molecules is close to the magic angle [35]. 2.5. MgO(lOO) - Ca Both ion scattering spectroscopy (ISS) with a neutral incident beam [36] and surface x-ray diffraction (SXRD) [37,38] have been applied to the investigation of the structure of Ca segregated onto the MgO(lOO) surface. For the ISS experiment a clean and ordered MgO(lOO) surface was obtained by cleaving in air a MgO single crystal and then repeatedly flash annealing in vacuum to 1073 K. To induce Ca surface segregation, the sample was further annealed up to 1573 K. The LEED pattern was observed to be a clear (1x1) both before and after Ca segregation. A neutral beam of He atoms rather than a beam of He"^ ions was used for the ISS measurements to minimise sample charging problems. Structural information was obtained from the intensity variation with scattering geometry of the spectral peaks due to scattering from Ca and Mg. It was determined, in agreement with theoretical calculations [3942], that Ca substitutes for Mg in the surface layer, and that the surface Mg and O ions are coplanar within experimental error (±0.1 A), whereas Ca is located 0.4 ± 0.1 A above this layer. It is suggested in the paper that both the protrusion of the surface Ca cations, and the Ca surface segregation are driven by the reduction of local strain, which is induced by the ionic radius of Ca2+ being 1.5 times greater than that of Mg2+.

Fig. 5. The optimum structure of the Ca-segregated MgO(lOO) surface determined from SXRD [37,38]. The numerical labels are employed in Table 3.

210

Table 3 z displacements of atoms in MgO(100)(V2xV2)R45°-Ca from analysis of SXRD data [37,38]. The numerical labelling of the atoms corresponds to that in Fig. 5. Both absolute displacements (Az (A)) and shifts relative to the MgO bulk nearest neighbour distance (Az/d) are listed.

l:Ca

2:Mg

Az/d

0.3 ±0.10

-0.032 ± 0.066

Az(A)

0.63 ±0.03

-0.066 ±0.14

3:0

4:0

0 (fixed) 0 (fixed) ±4 ±4 0 ±9

0 ±9

5:Mg

6:Mg

7:0

8:0

0.095 ±0.014

-0.2 ±0.08

-0.2 ±0.14

0.2 ±0.06

0.20 ±0.03

-0.42 ±0.18

-0.42 ±0.32

0.42 ±0.12

The SXRD study [37,38] was performed by Renaud's group. For these measurements the segregation was induced by annealing a sample at 1773 1873 K in air for several hours - the first step in a procedure developed for preparing high quality (100) surfaces of MgO. To remove both carbon contamination and oxygen vacancies, the MgO(100)-Ca surface was further annealed at approximately 1173 K for 30 min in 1x10-4 Torr of O2 in the UHV measurement chamber. In contrast to the ISS work in which the LEED pattern remained (1x1), diffuse, half-order spots were observed, corresponding to a (V2xV2)R45° surface unit cell. This difference probably arises because there was significantly more Ca on the surface in the SXRD study, on account of the significantly higher sample anneal temperature. The surface coverage of Ca was, in fact, estimated to be of the order of 1 ML from Auger electron spectroscopy. Surface structure was elucidated by fitting the (11/) and (20/) crystal truncation rods. Here it was assumed that (i) Ca is in substitutional sites, and (ii) Ca is present only in the topmost atomic plane. The best fit to the experimental data was obtained for a structure having the observed (V2xV2)R45° surface unit cell, in which every other surface Mg cation is replaced by a Ca cation. A schematic diagram of the structure is displayed in Fig. 5, and the vertical displacements of the atoms in the first two layers are listed in Table 3. Qualitatively similar atomic positions were obtained from theoretical calculations performed on the (V2xV2)R45° unit cell, which was found to be energetically favourable [39-42]. 2.6. MgO(lOO) - Ag The atomic scale structure of the MgO(100)/Ag interface has been examined experimentally by SEXAFS [43] and SXRD [37,44-46]. For the SEXAFS experiment an air-cleaved substrate was cleaned in situ by annealing

211

Table 4 Results of best fits to SEXAFS modulations, recorded at normal and grazing incidence, from MgO(100)-Ag at 1 ML equivalent coverage [43]. Also listed are the theoretically expected effective coordination numbers (N*theo) ^^^ ^be adsorption geometry discussed in the text. N*expt (N*theo) Ag-Ag Ag-0 Ag-Ag Ag-0

R (A)

Normal Incidence 8.3 ± 1 (8.5) 3.02 ±0.05 1 ± 1 (0.25) 2.55 ± 0.05 Grazing Incidence 7.3 ± 1 (7) 3.03 ±0.05 2±1(1) 2.53 ±0.05

at 1200 K for 10-20 min in an oxygen ambient of 5x10'4 Pa, followed by cooling at the same oxygen pressure to room temperature. Silver deposition and measurements were undertaken at 300 K. The estimated coverage of silver was 1 ML equivalent (i.e. if continuous over the surface the film would be one atom thick). Ag Ljjj edge SEXAFS data were collected at normal and grazing (15°) photon incidence angles to the surface plane. The results of the best fits to the SEXAFS spectra are Usted in Table 4. The Ag-Ag distance in the overlayer is significantly longer than the nearest neighbour distance in bulk silver (2.89 A). It is, however, within the experimental uncertainty, identical to the lateral separation (2.98 A) of equivalent sites on the MgO(lOO) surface. As to the adsorption site of the silver atoms, the effective coordination numbers (N*exp) ^^^ consistent with adsorption atop a surface ion, fitting favouring oxygen anions over magnesium cations. It should be noted that the authors mention that, on the basis of other work [47], the 1 ML equivalent Ag film is not actually one atomic layer thick, but rather it consists of mostly 2 ML thick islands covering about 60% of the surface area. They took account of this morphology in their structural determination. An SXRD study of this system investigated both the interfacial geometry, and details of the film morphology from submonolayer coverage up to many monolayers thickness [37,44-46]. Their preparation of the MgO(lOO) surface involved firstly, as described in the last subsection, annealing a polished sample in air at 1773 K for 3 h, and then in situ Ar"^ bombardment whilst annealing the sample at 1823 K. The resulting defected surface was improved by oxygen annealing. Silver was deposited at room temperature. Scattered x-ray intensity measured during scans along {hhl = 0.1), and (hOl ==0.1) demonstrated that up to 0.5 ML equivalent coverage the Ag film has, on average, its bulk lattice

212

parameter. This demonstrates that there is only weak coupling between Ag and the MgO(lOO) surface. Between 0.5 ML equivalent and 1 ML equivalent (and beyond) the islands begin to become strained by interaction with the MgO(lOO) substrate. Details about the adsorption site of on site Ag ( i.e. Ag atoms that sit on a lattice having the same in-plane lattice parameter as the substrate) in the first layer, which comprises only a small fraction of the total amount deposited, and the interfacial distance were obtained from analysis of (11/) and (20/) crystal truncation rod scans. It was determined, in agreement with the SEXAFS study [43], that Ag preferentially sits atop oxygen anions. The interfacial distance was found to be on average 2.52 ±0.1 A, again consistent with the SEXAFS result. Ab inito calculations of the Mg(100)/Ag interface agree that the most probable adsorption geometry is Ag atop a surface oxygen [48-51]. However, a variety of Ag-0 bond lengths, ranging from 2.34 A to 2.70 A, are obtained. The most recent and sophisticated calculation provides a value of 2.55 A [51]. 2.7. MgO(lOO) - Pd The Renaud group have also investigated details of the MgO(100)/Pd interfacial structure using SXRD [37,52-54]. Prior to Pd deposition the substrate was cleaned and ordered using the method described above for the MgO(100)-Ag SXRD study. Diffraction measurements were made following Pd deposition at room temperature for a wide range of coverages. At 1 ML equivalent coverage it was concluded from analysis of the (20/) and (31/) crystal truncation rods that 50 % of the Pd film is in registry with the substrate, and that in registry first layer Pd atoms are located above oxygen anions, as was found for silver, at a distance of 2.22 ± 0.02 A. This adsorption site contradicts a surface electron energy loss fine structure (SEELFS) study which concluded that for 3-D Pd clusters of diameter 2 nm on MgO(lOO), first layer Pd is located atop Mg cations [55]. The SXRD results are, however, in rather good agreement with several recent ab initio calculations which show that adsorption of Pd atop a surface layer oxygen is favourable, and predict Pd-O bondlengths ranging from 2.11 - 2.23 A [51,56-58]. 2.8. MgO(lOO) - Ni The MgO(100)-Ni interface was the subject of another SXRD study [59]. Surface preparation replicated that described above and SXRD data were recorded for Ni coverages from 0.2 to 125 ML equivalent. From simulations of the (11/) crystal truncation rod it was again concluded that part of the Ni film adsorbs on site, and that the epitaxial site is above a surface oxygen. For coverages up to 1 ML equivalent the distance between the MgO surface and

213

first Ni layer is approximately 1.82 A. Theoretical calculations are once more in rather good agreement with the experimental results [e.g. Refs. 51 and 60]. 2.9. MgO(lOO) - Fe Urano and Kanaji have utilised LEED-IV to examine the geometric structure of an approximately 1 ML equivalent film of Fe on MgO(lOO) [61]. The MgO(lOO) substrate was prepared by cleavage in air followed by annealing up to 1073 K in a partial pressure of oxygen (lxlO"3 Pa) inside the UHV measurement chamber. Iron film growth was performed with the sample surface at room temperature. A sharp 1x1 LEED pattern was observed at a film thickness of around 1 ML, which the authors interpret as evidence for the film being pseudomorphic. LEED-IV curves were acquired with an electron beam incidence angle of 6.1°. To determine the Fe adsorption site and height of the Fe above the surface layer, experimental data were compared with simulated data, calculated for three different adsorption sites: atop oxygen, atop magnesium and bridge. It should be noted that the position of the substrate ions were not optimised, but rather remained fixed in their bulk terminated positions. From this analysis it was determined that Fe adsorbs above a surface oxygen anion at a distance of about 2.0 A. A theoretical calculation for this system also finds that adsorption atop an oxygen is energetically preferred, but at a rather larger interfacial spacing of 2.3 A [62]. In contrast to the pseudomorphic relationship proposed in the LEED study, SXRD data [37,63] from the MgO(100)-Fe system indicate that for a 1 ML equivalent film the Fe lattice parameter is approximately 2.89 A, which is close to that of bulk Fe. For the latter study, however, a rather different preparative procedure was employed. Firstly, the MgO substrate was cleaned in solvents ex situ and only heated to 633 K in UHV. Then the sample was maintained at 633 K, whilst Fe was sputtered onto it, using a planar magnetron sputter gun. 3. NICKEL OXroE - NiO Nickel oxide, like MgO, usually adopts the relatively simple rocksalt lattice. The natural cleavage plane of NiO is (100), and studies have shown that the resulting surfaces are of high quality, relaxing only slightly away from the ideal bulk terminated (100) surface (see Fig. 1). Structural determinations of adsorbates have been performed on both this surface and the polar (111) surface. To circumvent surface charging problems almost all of these studies have been performed on highly oriented NiO thin films.

214

3.1. Ni(lOO) - NO N K-QdgQ NEXAFS [64] and scanned-energy mode N Is photoelectron diffraction (PhD) [65,66] have been employed to investigate the adsorption geometry of NO on NiO(lOO). Both of these studies utiUsed a NiO(lOO) thin film rather than the surface of a single crystal. Film preparation involved oxidation of a Ni(lOO) surface at elevated temperature, a procedure used extensively by Freund's group, who performed the NEXAFS measurements. Such NiO(lOO) thin films are known to contain a high density of surface defects, which could drastically affect adsorption properties. However, it has been clearly demonstrated that for NO the adsorption behaviour is dominated by ideal (100) surface sites [64,67]. The N ^-edge NEXAFS data were recorded from a NO saturated substrate, by monitoring the N KLL Auger yield as a function of photon energy. Spectra were recorded at a range of photon incidence angles between normal and grazing. The angular dependence of the intensity of the leading 7i* resonance was extracted and fitted with the appropriate trigonometric expression to determine the angular orientation of the N-O bond axis. The angle deduced was 45°. The PhD investigation [65,66] was performed on a NiO(lOO) thin film dosed with a saturation exposure of NO at 135 K. A partial pressure of NO of approximately 7x10"9 mbar was maintained in the experimental chamber during PhD data collection. This was done to overcome problems with photon beam induced NO dissociation/desorption (for more details see Ref 66). Two peaks, separated by approximately 4.4 eV, were observed, as had been done previously [64], in the N Is core level region. Comparing the diffraction displayed by each of these two peaks confirmed Freund et aFs assertion that the two peaks arise due to two different core hole states of the same surface species [64]. Quantitative elucidation of the geometry of this adsorbate involved a two-step process. Initially, a visual inspection of the PhD spectra was carried out to ascertain the approximate position of the N atom relative to its nearest neighbours below. Next, this geometry was used as an initial guess for the generation of simulated PhD data via electron multiple scattering calculations, which were compared to the experimental data set. Agreement between experimental and theoretical data was iteratively improved by varying structural parameters until the best fit, and therefore the optimised structure, was obtained (R-factor = 0.09: this R-factor is defined differently to the Pendry LEED Rfactor, see Ref 68 for details). NO is found to adsorb N atom down atop a Ni atom, and in agreement with the NEXAFS data [64], the N-O bond axis is tilted away from the surface normal. This optimised adsorbate geometry is shown ia Fig. 6 and the values of the optimised parameters are listed in Table 5. Interestingly, although theoretical calculations [64,69,70] favour this adsorption

215

Fig. 6. Schematic of the optimum NO adsorption geometry on NiO(lOO) determined from PhD [65,66]. The geometrical parameters varied during the structural optimisation are displayed. For the adsorbate, the darker shaded circle is nitrogen, and the lighter oxygen.

Table 5 Parameter values obtained from the best fit between experimental N Is PhD data and theoretical simulations for NiO(100)-NO [65,66]. The definitions of the first five parameters are given in Fig. 6. and are the mean square vibrational amplitudes of the nitrogen atom parallel and perpendicular to the surface, respectively. The positive error for Y^Q is replaced by an asterisk due to the fact that all N-O bond lengths greater than the optimum lie within the estimated error. The double asterisk next to the error for 02, the N - 0 tilt, indicates that the error given is only for N-0 bond lengths [HCOO]- [87].

[001]

[110] • Fig. 19. Two suggested adsorption geometries for rhodium gem-dicarbonyl (Rh(C0)2) on Ti02(l 10)1x1 [104].

231

4.10 TiO2(100) - SO2 SO2, chemisorbed on either the (1x1) or (1x3) phases of TiO2(100), can undergo the following reaction via thermal activation [105]: SO2 -> SO32- -^ SO42-. Raza et al have performed S AT-edge NEXAFS measurements to obtain the surface orientation of the reactant, intermediate, and product of this reaction [105]. Surfaces were prepared by Ar"^ sputtering and annealing, with either a final anneal in 1x10"" mbar of O2, or a 10 min anneal in vacuum at 1100 K to produce the (1x1) and (1x3) surfaces, respectively. SO2 exposure was carried out at 110 K to achieve -0.5 ML coverage (where 1 ML is defined as occupation of all of the O and Ti surface atom sites by sulfur-containing species). To investigate molecular orientations, NEXAFS spectra were recorded over a range of polar photon incidence angles with the x-ray electric vector in the [001] and [010] sample azimuths. These measurements were performed at three substrate temperatures, namely 130 K, 200 K, and 500 K, in

® Ti

Fig. 20. Models of TiO2(100)lx3 (top) and TiO2(100)lxl (bottom) along with the proposed SO2 adsorption sites corresponding to, from left to right, chemisorbed SO2, 803^"", and SO42- -like species [105].

232

order to determine the orientations of SO2, SO32- and SO42-, respectively. For SO2 from analysis of the 7i* resonance it was found that the C2 molecular axis is oriented 20° ± 2° away from the (100) surface normal on the (1x3) surface, and 26° ± 3° on (1x1), with no azimuthal alignment on either surface. It should be noted that at 110 K SO2 displayed no specific orientation on either surface. Similar analysis for SO32-, which adopts a trigonal-based pyramidal structure, again indicates no azimuthal alignment and tilts of the C3 molecular axis of 24° ± 3° and 25° ± 3° away from the surface normal on the (1x3) and (1x1) surfaces, respectively. Data for S04^" from either surface at 500 K show little photon incidence angle dependence. This is consistent with a sulfate species having T j symmetry, although there is evidence from the spectral features of some distortion away from this highly symmetric configuration. This distortion is more pronounced for the S04^" already present on the surfaces at lower temperatures, specifically between 130 K and 220 K. As the results from the (1x1) and (1x3) surfaces are so similar, the authors use these structural data to argue that the reaction of SO2 with Ti02 is driven by local geometry effects rather than longer range surface structure. Proposed adsorption structures for the various surface species on the two surfaces are shown in Fig. 20. 4.11. TiO2(100) - K The local geometry of the K adsorption site in TiO2(100)c(2x2)K has been probed using K i^-edge SEXAFS [106]. To form the c(2x2) overlayer 0.5 ML of K was deposited onto a stoichiometric TiO2(100)lxl surface, cleaned as described above. SEXAFS measurements were performed at normal and C(2x2)

atop

K inclined

atop bridge

'°^°1

Fig. 21. Schematic of TiO2(100)c(2x2)K derived from SEXAFS data [106], with K atoms bonded to two bridging oxygens (left). The altemative K adsorption sites considered are also shown (right).

233

grazing x-ray beam incidence with the electric vector aligned to the [001] azimuth. Successful modelling of the data required a single shell of oxygen neighbours at 2.62 ± 0.03 A, the analysis rejecting models which contained K or Ti backscatterers. Adsorption site determination entailed a comparison of the experimentally determined oxygen effective coordination numbers with ones calculated for four plausible sites: angled bridge, atop bridge, atop, and Kinclined (see Fig. 21). From this process it was concluded that the angled bridge is the K adsorption site, in which the K-O2 plane is at 45° to the surface normal. It is noted in the paper that the K-0 bond length is close to the sum of the appropriate ionic radii. The structure predicted by a recent ab initio Hartree-Fock calculation is in excellent agreement with the experimental data [107]. 5. ALUMINIUM OXIDE - AI2O3 A number of different bulk phases of AI2O3 exist, including a-Al203 (corundum) and Y-AI2O3. Quantitative structure determinations of adsorbates have been performed on the (0001) face of an a-Al203 single crystal, and a (0001) oriented thin film of y-like AI2O3. a-Al2O3(0001) displays a variety of surface periodicities [37]. Of interest, as regards this review, are the (1x1) and (V31xV31)R±9'^ unit cells, both of which have been investigated by SXRD [37]. For the (1x1) surface it was concluded that the surface is terminated by a single-aluminium layer, whilst a model consisting of a tiling of domains bearing a close resemblance to that of two metal Al(l 11) planes separated by a hexagonal network of domain walls is proposed for the (V31xV31)R±9° termination. The y-like Al2O3(0001) thin film, grown on NiAl(llO) via preferential oxidation, exhibits a complex LEED pattern [108], and its atomic scale structure is the subject of some debate [109]. 5.1. a-Al2O3(0001) - Cu Gota et al have employed SEXAFS at the Cu ^-edge to probe the structures of a-Al2O3(0001)lxl-Cu and a-Al2O3(0001) (V31xV31)R±9°-Cu, following deposition of 0.5 ML equivalent of Cu [110]. Preparation of the surfaces involved Ar+ bombardment, with subsequent annealing at 1073 K in 5x10-7 Torr of O2, and at 1673 K in UHV for the (1x1) and (V31xV31)R±9° terminations, respectively. SEXAFS spectra were obtained at normal photon incidence. On both surfaces the majority of Cu is present in the form of small 3D clusters, and best fits to the SEXAFS data have a significant contribution

234

Table 7 Results of best fits to SEXAFS modulations, recorded at normal incidence, from Al2O3(0001)lxl-Cuand Al2O3(0001)(V31xV31)R±9°-Cu at 0.5 ML equivalent coverage of Cu [110]. N* is the effective coordination number of a shell of scatterers, R is its distance, and Aa^ the mean square relative displacement. N* (±10%) Cu-Cu Cu-Cu Cu-Al

0.7 4.3 2.0

Cu-Cu Cu-Al

4.7 5.7

R(A) (± 0.02 A) (1x1) 2.0 2.37 3.16 (V31xV31)R±9° 2.42 3.24

Aa2 (A2) (±0.5x10-3) 0.010 0.010 0.010 0.012 0.012

from Cu, although information about the interface was also extracted. Specifically, including a shell of Al scatterers was found to significantly improve the fits. It should be noted that no evidence for oxygen scatterers was found. Table 7 lists the structural parameters obtained from analysis of the SEXAFS. The authors interpret the rather long Cu-Al bond distances as possibly being a consequence of the very weak interaction between Cu and the AI2O3 surfaces The relatively high effective coordination number of Al for the (V31xV31)R±9° reconstruction (5.7) is ascribed to Cu adsorption at defects. 5.2. y-like Al2O3(0001) - di-tert-butyl nitroxide The orientation of di-tert-butyl nitroxide (CgHigNO), the structure of which is illustrated in Fig. 22, on a thin film of (0001) oriented y-like AI2O3 has been deduced from N AT-edge NEXAFS measurements [111]. The film was grown on a clean and well-ordered NiAl(llO) substrate, via oxygen exposure and annealing (see Ref. 108 for details). NEXAFS spectra were recorded from a monolayer of di-tert-butyl nitroxide, which had been dosed at a substrate temperature of 100 K. No 71* resonances were observed in these spectra, although such a resonance was present in multilayer spectra. This difference was explained as being due to electron donation from a surface oxygen to the adsorbate's nitrogen atom. Given the lack of TT* resonances, molecular orientation information was obtained from the variation in intensity of the a*(N-C) and a*(N-0) resonances in the spectra. A tilt angle of about 70° for the N-O bond axis away from the surface normal was concluded. From complementary electron spin resonance (ESR) measurements, it is known that

235

Fig. 22. Ball and stick model of di-tert-butyl nitroxide.

the NEXAFS spectra actually arise from two differently bound di-tert-butyl nitroxide species. However, the majority species comprises more than 90% of the adsorbate coverage, and so it is presumed that the contribution of the minority species to the NEXAFS spectra is not significant and that the tilt angle extracted is essentially that of the majority species. 6. CHROMIUM OXIDE - Cr203 Cr203 has the same bulk crystal structure as a-Al203, namely corundum. Of its several low Miller index surfaces only one, (0001), has been employed for adsorbate structural determinations so far. To overcome sample charging problems a thin film has been utilised for these studies, rather than a single crystal. The surface structure of this (0001) oriented thin film has been investigated by LEED-IV [112]. Simulations of the experimental data evidence a chromium terminated surface with large vertical interlayer relaxations, reaching down five or six layers. 6.1. Cr2O3(0001) - CO2 Freund and coworkers have used C ^-edge NEXAFS to examine the angular geometry of the CO32- moiety resulting from CO2 adsorption on Cr2O3(0001) [113]. As indicated above, the substrate was a highly oriented

236

thin film rather than a single crystal. It was grown by first annealing the sample at 500 K in a partial pressure of oxygen for 3 min and then annealing in vacuum at 1000 K to remove excess oxygen. CO2 was adsorbed with the substrate at approximately 100 K. Two NEXAFS spectra were recorded, one at normal photon incidence the other at grazing incidence, to elucidate the orientation of the CO32-. For these measurements the sample was maintained at approximately 170 K to avoid physisorbed CO2. Analysis of the carbonate 7C* resonance at 290.2 eV, which is significantly more intense at normal incidence, indicated that the molecular plane of the CO32- is approximately parallel to the surface normal. 6.2. Cr2O3(0001) - NO NEXAFS, at the N iC-edge, has been performed by Freund's group to obtain the tilt angle of the N-0 bond axis relative to the Cr2O3(0001) surface [114]. A thin film, prepared as above, was employed. NO dosing and NEXAFS measurements were carried out at substrate temperatures of between 100 K and 105 K. The variation in the relative intensities of the NO 71* and a* resonances with polar photon incidence angle was used to obtain the molecular orientation. Fitting to an appropriate trigonometric function [35] provides an angle of 50° between the N-0 bond axis and the surface normal. The authors note that there is a significant error for this tilt angle. 7. IRON OXIDE - FexOy There are three stable bulk forms of iron oxide: FcxO, Fe304, and Fe203. Single crystal faces of each of these three phases have been the subject of surface studies. However, adsorbate geometry determination has only been performed on (111) oriented thin films of FeO and Fe304 grown on Pt(lll). Details of the structure of both of these films have been revealed by the employment of a range of surface science techniques (see Ref 115 and Refs. therein). The FeO (111) film is comprised of a single bilayer of iron and oxygen on the Pt(lll) substrate, with oxygen being uppermost. It has a lateral lattice constant of 3.11 A (the lattice constant in bulk FeO is 3.04 A), and an ironoxygen interlayer spacing of 0.68 A, which is much reduced compared to the (111) interlayer separation in bulk FeO (1.25 A). The Fe304 film is substantially thicker, of the order of 100 A - 200 A. At its surface 0.25 ML of iron cations sit upon a close packed oxygen layer, giving rise to a hexagonal surface unit cell of side 5.92 A

237

Fig. 23. Schematic diagram of stryrene. Only the carbon atoms are shown.

7.1. FeO(lll) - H5C6CHCH2 Wiihn et al have elucidated the angular geometry of styrene (H5C6CHCH2: see Fig. 23 for schematic diagram) on a FeO(lll) thin film using C X-edge NEXAFS [116]. The methodology for substrate preparation involved cycles of iron deposition onto a clean and ordered Pt(lll) surface at room temperature, followed by oxidation for 2 min at approximately 1000 K in a partial pressure of oxygen (10"6 mbar). Adsorbate overlayers of coverage 0.6 ML and 1 ML (where 1 ML is defined as the saturation coverage of styrene on FeO at 200 K) were produced by exposure to 0.7 L and 1 L of styrene, respectively, at 200 K. NEXAFS spectra were recorded at normal and grazing photon incidence for both coverages. Comparison with multilayer/gas phase NEXAFS data, along with results from other techniques, demonstrated that the styrene molecule remains intact upon adsorption. From analysis of the variation in intensity of the strongest 7C* resonance (285 eV), it was concluded that the angles between the surface and the molecular plane of the styrene are 41° and 43° for coverages of 0.6 ML and 1 ML, respectively (the ethylenic and benzenoid fragments of the styrene molecule are believed to be coplanar). An error of ±2° is attributed to these tilt angles by the authors. 7.2. Fe304(lll) - H5C6CHCH2 The orientation of styrene on Fe304(lll) is also derived from C i^-edge NEXAFS in Ref 116. A thin film of (111) oriented Fe304 was prepared using a procedure identical to that described in the last subsection, except that more deposition/oxidation cycles were required. NEXAFS spectra were obtained in the submonolayer regime for styrene coverages of 0.4 ML and 0.8 ML, with the substrate at 240 K and 225 K, respectively. At the lower coverage the tilt of the styrene plane away from the surface was found to be 28° ± 2°. This angle increases to 42° ± 2° at 0.8 ML coverage. On the basis of TPD data [115, 116], this change in geometry is attributed to there being more than one adsorption state for styrene in the submonolayer regime. It is suggested that initially styrene binds via its TT system to surface iron cations in a near flat

238

geometry, then once these iron sites are saturated styrene bonds more weakly to other surface sites giving a tilted geometry. It should be noted that on FeO(lll) ( subsection 7.1) there are no iron cations in the surface layer, and only the weakly bound, more tilted styrene species is observed. 8. ZINC OXIDE - ZnO ZnO adopts the wurtzite structure in which Zn and O are both tetrahedrally coordinated to their counter ions. The O-terminated (0001) and Zn-terminated (0001) basal faces, as well as the non-polar (1010) prism face have all been the subject of adsorbate structure determinations. As regards their clean structures, LEED-IV was utiUsed as a probe over twenty years ago [117,118], and very recently they have been reexamined with SXRD [119-121]. From these measurements it was found that neither polar face relaxes greatly, a feature ascribed to surface enhanced covalency [121]. For the (1010) face, SXRD [120] indicates only small displacements of the surface Zn and O away from their bulk terminated positions. 8.1. ZnO(lOlO) - HCOOH Davis et al have determined the orientation of [HCOO]" on ZnO(lOlO) with C AT-edge NEXAFS [122]. For these measurements the substrate was prepared by cycles of Ar"^ bombardment at 800 K, followed by annealing for 10 min at 800 K in 1x10"6 mbar O2 with subsequent cooling in O2 to 550 K. The formate overlayer was formed by exposure of the substrate at 295 K to 8 L of formic acid, to give a coverage of about 0.2 ML. The NEXAFS data indicate that formate is oriented with its molecular axis tilted towards upright, and that there is no azimuthal alignment. Quantitative analysis, employing the leading 71* resonance, provides a tilt of the formate C2 axis of about 25° away from the surface normal. A possible bond geometry is shown in Fig. 24, in which formate is coordinated in a bidentate configuration to one Zn atom. This structure is consistent with the lack of azimuthal ordering. A recent energy minimised calculation of formate on ZnO(10 10) predicts that a bidentate configuration is more stable than monodentate [123]. However, the suggested structure has formate bridging two Zn atoms along the [1120] direction. Unless the experimental data is very strongly affected by dynamical effects, such ordering would give rise to a clear azimuthal polarisation dependence, which is not apparent. 8.2. ZnO(lOTO) - CO2 Exposure of ZnO(10 10) to CO2 leads to a surface CO32- species. The

239

Fig. 24. Model of ZnO(lOTO) surface, with the bond geometries of formate and carbonate proposed in Ref 122. The tilts out of high symmetry sites have been neglected, and for formate only one azimuthal orientation is shown.

orientation of this CO32- has been determined, using C AT-edge NEXAFS [122]. For these measurements the substrate was prepared as described above (subsection 8.1), and exposed to 45 L of CO2 as it was cooled to the measurement temperature of 180 K. This achieved a coverage of about 0.2 ML. NEXAFS spectra were recorded at a number of photon incidence angles with the x-ray electric vector in either the [000T] or [1120] azimuth The angular dependence of the NEXAFS data indicates that the plane of CO32- lies in the [000 1 ] azimuth, tilted away from the surface normal by approximately 15°. Fig. 24 shows an adsorption geometry suggested in Ref. 122. 8.3.ZnO(1010)-C6H6 The angular geometry of the prototypical aromatic molecule benzene (C6H6) on ZnO(lOTO) has been investigated with C A'-edge NEXAFS [124]. The substrate was cleaned and ordered as outlined above (subsection 8.1), and the overlayer was formed by exposure to 5 L of benzene at the measurement temperature of 180 K. The coverage of benzene was about 0.3 ML. From the intensity variation with photon incidence angle of the leading 71* resonance in the NEXAFS spectra, it was determined that the benzene ring is tilted by 10°

240

away from the surface plane, i.e. benzene lies more or less flat on the surface. This geometry mirrors that found on metal surfaces, which arises due to benzene coupling with the surface through its n system [35,125]. 8.4. ZnO(lOTO) - C5H5N NEXAFS, both at the C ^-edge and N ^-edge, has been used to examine the orientation of another aromatic molecule, pyridine (C5H5N), on ZnO(10 10) [124]. A preparative procedure similar to that used for the studies discussed in the previous three subsections was performed for the substrate. The pyridine adlayer, of coverage approximately 0.1 ML, was formed by exposure to 1 L at the measurement temperature of 295 K. NEXAFS data were recorded at a series of polar photon incidence angles, with the electric vector of the linearly

290 310 Photon Energy (eV) Fig. 25. C AT-edge NEXAFS spectra of ZnO(lOTO) after exposure to 1 L of pyridine at the measurement temperature of 295 K [124]. The peak labels are discussed in Ref. 124. The inset shows the experimental variation in the intensity of the 7i* feature (peaks al and all) with angle of incidence in the two measurement azimuths (data points). This is compared with the best fit to the data (solid lines).

241

polarised photons either parallel or perpendicular to the surface Zn-O dimer rows (see Fig. 24). The NEXAFS spectra show that pyridine stands up on the surface with the ring plane in the [000 1 ] azimuth. Two C K-edgQ spectra, recorded at normal and grazing incidence with the electric vector in [1120] azimuth, are displayed in Fig. 25. The inset in this figure shows the intensity variation of the strongest 7i* feature (a doublet, with components at 285.0 eV and 285.5 eV) with angle of photon incidence in the two measurement azimuths, along with the best fits to these data using an appropriate trigonometric expression [35]. These fits correspond to a tilt of the molecular plane towards the surface of 28° and a 27° twist out of the [11 20] azimuth. A consistent result was obtained from the N ^-edge data. The authors note that this apparent distortion out of a high symmetry geometry should be viewed with some caution, since dynamical effects are known to significantly influence NEXAFS results obtained at room temperature.

8.5. ZnO(lOTO) - C6N3H5, C7N2H6, C7N2H6, C7N3H7 The orientation on ZnO(lOlO) of benzotriazole (C6N3H5) and related molecules (indazole (C7N2H6), benzimidazole (C7N2H6), and 1methylbenzotriazole (C7N3H7)), which are of interest in connection with corrosion protection, has been studied using C and N K-QdgQ NEXAFS [126]. The substrate was prepared as above and the gas phase molecules used to dose

Benzimidazole

1 -Methyl-Benzotriazole

Fig. 26. Ball and stick models of the adsorbates studied in Ref 126.

242

ZnO(10 10) were obtained by evaporation of the solids. Ball and stick models of these molecules are displayed in Fig. 26. At submonolayer coverage, benzotriazole is found to adsorb in an upright geometry with the molecular plane within 35° of perpendicular to the substrate, as indicated by the polarisation dependence of the 7i* resonances. Seven slightly different models of the bond geometry are consistent with the data; two of the most likely models are shown in Fig. 27. Indazole (C7H6N2) was found to be oriented similarly. Benzimidazole (C7H6N2) and 1-methyl benzotriazole (C7H7N3) were found not to be have a preferred orientation at sub-monolayer coverage. 8.6. Z n O ( 0 0 0 1 ) - C O Thornton and coworkers have employed C A^-edge NEXAFS measurements to investigate CO adsorption at 130 K on ZnO(OOOT) [127,128]. The substrate was prepared by cycles of Ar"^ bombardment and annealing in vacuum to about 750 K. This was followed by a final anneal in IxlO"^ mbar O2 and cooling in this ambient to 550 K. CO adsorption gave rise to two K* resonances in the C ^-edge NEXAFS spectra at 287.7 ± 0.2 eV and 290.4 ± 0.2 eV. These were assigned to molecular CO, which is metastable under the prevailing

[1120]

Fig. 27. Model depicting two NEXAFS compatible adsorption geometries for benzotriazole onZnO(10T0)[126].

243

experimental conditions, and CO32-, respectively. XPS/x-ray absorption step edge measurements indicated that the coverage of CO32- was 0.1 ± 0.05 ML, whilst that of CO was 0.04 ML at most. From the dependence of the CO3271* intensity on photon incidence angle it was deduced that the molecular plane of CO32- is inclined away from the surface normal by 32° ± 20°. The authors state that the error bar is large due to the limited dataset, and also uncertainties in normalisation of the data arising from the variation in CO surface concentration. It should be noted that CO was also dosed onto a ZnO(000 1) surface precovered with approximately 0.55 ML of Cu. In this case only one intense 7C* resonance, at 287.7 ± 0.2 eV, was observed. This was again attributed to adsorbed molecular CO. The polarisation dependence of this resonance indicated that the CO molecular axis is tilted by 17° ± 10° away from the surface normal. 8.7. ZnO(OOOT) - CO2 Besides CO adsorption, Ref. 127 also presents C K-edge NEXAFS spectra obtained following exposure at 130 K of ZnO(OOOT) to CO2. Surface CO32again forms, as evidenced by a 7i* resonance located at 290.4 ± 0.2 eV. The estimated CO32- coverage was 0.1 i 0.05 ML, identical to that of the CO32species formed following CO adsorption, as discussed in the last subsection. Analysis of the intensity variation of the 7C* resonance with measurement geometry produced a value of 25° ± 5° for the tilt angle of C03^" away from perpendicular. This value is similar to that for CO32- produced from CO exposure. 8.8. ZnO(OOOT) - HCOOH The orientation of [HCOO]" on ZnO(000 1) has been investigated using C AT-edge NEXAFS [129]. Prior to the measurements, the sample was cleaned in situ by cycles of Ar+ bombardment and annealing in UHV to between 800 K and 1000 K, and ending with an anneal in a partial pressure of O2 (lxl0"6 mbar). Surface formate was created by exposure of the substrate to 20 L of formic acid at 110 K and heating to 340 K. This procedure gave a formate coverage of 0.2 ± 0.05 ML. Two NEXAFS spectra were recorded, one at normal incidence and the other at grazing. From quantitative analysis it was concluded that the tilt angle of the molecular plane of the formate is 55° ± 10° with respect to the surface normal. The authors speculate that the formate may be located at step edges, on the basis that earlier work has demonstrated that formate can only be located at defects on this surface [130]. It must be noted, however, that currently the actual identity of the defects remains unknown.

244

NEXAFS spectra were also acquired from [HCOO]' adsorbed onto a ZnO(000 1) surface, which had previously been exposed to Cu vapour at 140 K such that approximately half of its surface area was covered by 2D Cu islands. The formic acid was again dosed at a substrate temperature of 110 K, but in this case no annealing was necessary to dissociate molecularly adsorbed formic acid. The tilt of the formate in this case was 32° ± 5°, which is close to the magic angle [35]. It is suggested in the paper that this result is due either to buckling of the 2D Cu islands giving rise to a molecular tilt of about 32°, or randomly oriented formate. These two possibilities are indistinguishable by NEXAFS. 8.9. ZnO(000T)-C5H5N Hovel et al have investigated the adsorption of C5H5N (pyridine) on the O terminated basal ZnO face with several techniques, including NEXAFS at the C i^-edge, to determine its angular geometry [131]. The in situ substrate preparation involved Ar"*" bombarding with the sample at 650 K, heating in vacuum at 800 K, and frequent anneals in IxlO"^ mbar of O2, with subsequent cooling to 550 K in the O2 ambient. An adsorbate overlayer was prepared by exposing the surface to 1 L pyridine at the measurement temperature of 110 K. Adsorption was concluded to be molecular in the monolayer regime, and an angle of 66° ± 5° between the molecular plane of pyridine and the surface was derived from the angular dependence of the NEXAFS spectra. 8.10.ZiiO(000T)-K The structure of a p(2x2) overlayer of K on the O terminated basal face of ZnO has been studied with Auger and fluorescence yield polarisation dependent SEXAFS at the K A^-edge [132]. Substrate preparation involved cycles of Ar"^ bombardment and annealing in vacuum to 1100 K. A final anneal was Table 8 The effective coordination numbers resulting from single shell fits to K X-edge SEXAFS data recorded from ZnO(000 I)2x2-K [132]. Also listed are the calculated effective coordination numbers for the three-fold hollow, bridge and atop adsorption sites. Effective coordination numbers Angle of Incidence

Experimental

3-fold hollow

Bridge

Atop

90° 20°

3.8 ±1.5 4.7 ±1.5

2.2 4.4

1.1 3.5

0.0 2.6

245

performed in O2, with subsequent cooling in O2 to 500 K. The p(2x2)K overlayer was formed by evaporating K onto the substrate at room temperature and annealing to 900 K. SEXAFS data were recorded at normal and grazing photon incidence with the substrate maintained at a temperature of 125 K. Successful modelling of the grazing incidence data required a single shell of oxygen neighbours at 2.70 ± 0.04 A, the analysis routine rejecting models containing K or Zn scattering atoms. The K adsorption site was extracted by comparing the experimental effective coordination numbers with those calculated for high symmetry sites (see Table 8). From this comparison it was concluded that K adsorbed in a three-fold hollow site. On this surface there are two different hollow sites, the open and cave (see Fig. 28). On the basis of a lack of any Zn scatterer contribution to the SEXAFS oscillations, the authors argue that of the two sites the cave site is more probable. This is consistent with the prediction of Madelung potential calculations [133,134]. Models of the structure are shown in Fig. 28. [1120] [1010]

M v;i>^^^^^^ :.:mm'm^^m. ^-^f^'n^^m 'imîsw's^^^^,;,,.,.,,;

Fig. 28. Ball and stick models of ZnO(000 1 )2x2-K derived from SEXAFS data, with K atoms bound to three oxygen atoms in a three-fold cave site [132]. The K atoms form the comers of the p(2x2) unit cell outlined. The alternative K adsorption sites considered in modelling the SEXAFS are also shown. The bridge and atop sites are shown in the top and side view, with the open hollow site shown only in the top view.

246

c 3

7C*(CO) [0001]^

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°

• o • 1.5

Fig. 10. Adsorption energy of CO on NiO(lOO) cleaved in vacuum as a function of coverage. The data have been determined from TDS spectra like those shown in Fig. 6 (upper part) using the leading edge method and complete analysis. TDS data taken with heating rates of 0.1, 0.2, 1 and 2 K/s have been used.

0.0

Complete analysis 1 K/s Leading edge 1 K/s Leading edge 0.2 K/s 0.5

1.0

1.5

Relative coverage [©/©monolayer]

Fig. 11. Adsorption energy of NO on NiO(lOO) cleaved in vacuum as a function of coverage. The date have been determined from TDS spectra like those shown in Fig. 7 (upper part) using the leading edge method and complete analysis. TDS data taken with heating rates of 0.2 and 1 K/s have been used.

0.30 and 0.57 eV for CO and NO, respectively. The low coverage adsorption energies for CO and NO on NiO(lOO) and MgO(lOO) are compiled in Table 2. According to theory the interaction of the adsorbates with MgO(lOO) and NiO(lOO) are expected to be similar since the bonding should be mainly electrostatic in nature [52] (the electric fields at the surfaces of NiO(lOO) and MgO(lOO) are similar). However, according to Table 2 the bonding energies are considerably different, with the higher values being obtained for NiO(lOO). Covalent interactions involving the Ni3d electrons may play a role for the adsorbate-substrate interaction which do not show up in the calculations published so far. As far as it concerns the basis set superposition error (ESSE) corrected calculations listed in Table 1 which are expected to yield qualitatively better results as compared to the non-corrected calculations, it appears that the theoretical results for adsorption on MgO(lOO) are in general in line with our experimental results, whereas a similarly favorable comparison can not be made for NiO(lOO). It appears necessary to re-investigate the role of the NiSd electrons in future theoretical studies. The interaction between the molecule and the d-electrons of the substrate system is more obvious in the case of NO interaction with NiO. The molecule is tilted with respect to the surface normal as is experimentally evident from X-ray absorption measurements [59] as well as photoelectron diffraction investigations

338

by the Woodruff group [69]. The tih angle is near 45-60° with the nitrogen end of the molecule pointing towards a Ni ion. The interatomic distance between N and O is near the one known from the free molecule, the separation from the surface Ni ion is determined to be 1.88 + 0.2 A. It was early pointed out by Staemmler and his group that this tilt is a consequence of the interaction of the unpaired electron on the NO moiety with the open shells on the Ni ion. A detailed discussion of the bonding mechanism can be found in ref [57]. On MgO(lOO) preliminary calculations have indicated that the NO molecule is oriented perpendicular to the (100) surface plane. This appears to be reasonable in light of the missing d-electrons of the Mg ion. There are further indications for weak chemisorptive interactions between NO and the NiO(lOO) surface: Although there is only one NO species on the surface, the Nls photoelectron spectrum exhibits a broad band with two maxima. Such phenomena are connected with the dynamics of the screening processes accompanying the formation of the core hole. Detailed studies on screening phenomena have been reported for molecules on metal surfaces, [70-72], and it is well known, that systems with weak chemisorptive interactions exhibit the most complex line shapes. For weak chemisorption systems there is a high probability to find the final hole, both on the adsorbed molecule itself or, through extra molecular screening, on the substrate, giving rise to several final states, and thus to several intense components in the spectrum. It was proposed some time ago that such phenomena are also responsible for the line shapes observed for NO/NiO(100) [59]. The above mentioned recent photoelectron diffraction investigations have clearly demonstrated that the observed components in the Nls spectrum stem from a single species. The adsorption of CO on MgO has been thoroughly investigated by Heidberg and his group using IR-spectroscopy [73] and by Weiss and Toennies using helium atom scattering spectroscopy (HAS) [74]. They have clearly demonstrated that CO develops ordered phases on the cleavage planes and that order and spectroscopic properties depend on the quality of the prepared surfaces. From their experiments the influence of the presence of surface defects on adsorption properties is very obvious but a quantitative evaluation based on the number and the nature of the defects has not been reported. The quantitative evaluation of defects is a well defined but hard to tackle problem for future studies that has to be taken on by our research community. Water adsorption is an example that lends itself to a study of the influence of defects, because at lower coverage the (100) cleavage planes of MgO and NiO do not dissociate water, while the presence of defects does induce water dissociation. Water adsorption on MgO(lOO) is the best studied system so far [64, 75-77]. The groups of Heidberg [75], Weiss and Toennies [76] have studied it with IR, LEED and HAS. While these groups have used surfaces prepared through in situ cleavage, Stimiman et al. [78] have employed bulk single crystals prepared via

339

MgO(100) sputtered 0,3Ar7Mg2kV 6,8, 17LH2O)

196 K 239 K 0.51 eV 0.66 eV (Redhead)| heating rate 0.5 K/s 150

200 250 temperature [K]

300

Fig. 12. Thermal desorption spectra of H2O on UHV-cleaved NiO(lOO) and MgO(lOO). A schematic representation of the c(4x2) structure is included (according to ref. [75]). For comparison a thermal desorption spectrum from MgO(lOO) after creation of defects via sputtering is shown.

sputtering and heating in oxygen. Their TD spectra correspond grossly to the TDS data published by Wichtendahl et al. on in situ cleaved MgO(lOO) surfaces, which are shown in Fig. 12 [64]. Data taken on a thin IVlgO film [79] reveal the essential features, although the peaks appear to be shifted significantly with respect to the temperature-calibrated data of Wichtendahl et al. The assignment of the feature at 196 K was unclear so far, although it was near at hand to assume that it is connected with the phase transition from the c(4x2) to the (3x2)pg phase which occurs at about this temperature and was reported by several groups [75, 76, 79]. It could be shown that the molecule density within the c(4x2) phase is higher than in the (3x2)pg phase (~ 1.3 atoms and -1.0 atoms per magnesium surface ion, respectively) [78] so that a transition between these phases must be accompanied by a desorption peak in TDS. The TDS data of H2O/MgO(100) in combination with data for H2O/NiO(100) taken in our group [80] shed new light upon this question. Fig 12 [80] reveals the data. The feature near 150 K is due to multilayer desorption and shall not be discussed here in more detail. The peak at 196 K, which corresponds to the transition from

340

the c(4x2) phase to the (3x2)pg phase, is observed in the spectra from both surfaces at the same temperature. This observation is a first hint that part of the molecules of the c(4x2) phase belong to the second layer and do not have contact with the oxide surface. The highest temperature peak, which corresponds to desorption of the (3x2)pg phase in the MgO(lOO) case, i.e. the monolayer, is found at slightly different temperatures, i.e. 225 K for NiO(lOO) and 239 K for MgO(lOO), because the direct interaction between water and the Ni and Mg sites, respectively, is different. This conclusion is supported by results of molecular dynamics simulation by the Besan9on group on H2O/MgO(100) [77]. Results are displayed in Fig. 13 [77]. The interesting result is that a monolayer of H2O forms with the H2O molecules oriented with the plane parallel to the oxide surface and with a network of Hbonds within that layer, but fully disconnected from the second and higher layers. This monolayer, which is to be identified with the (3x2)pg phase, completely shields the higher layers from the underlying oxide surface. Therefore, the molecules desorbing during the phase transition from the c(4x2) phase to the (3x2)pg phase may be identified as being part of the second layer, since these molecules do not feel the differences between NiO(lOO) and MgO(lOO) as documented by the data shown in Fig. 13. In closing this chapter we would like to mention that there are several studies in the literature where more complex molecules, which are also of technical interest, have been studied. These include methanol, formic acid and aromatic compounds [81-88]. We shall not dwell on these studies here but merely refer the reader to the literature.

1=150 K

T=300 K

Fig. 13. Simulation of water on MgO(lOO) at coverages higher than one monolayer at 150 K and 300 K, according to ref [77]. a) top and side views of map-shots at 150 K. b) top and side views of map-shots at 300 K. (The first layer molecules are not represented in the top view.) The side views show that the overlayer beyond the monolayer tends to aggregate.

341

3. SURFACES OF OXIDES WITH CORUNDUM STRUCTURE Fig. 14 shows the lattice structure of a non-magnetic system, i.e. a-alumina [15]. We note here in passing that there are antiferromagnetic corundum structures where the metal centers carry different spin-orientations which renders the size of the unit cell considerably larger [89-91]. The surface energies for a structure as shown in Fig. 14 has been calculated for a variety of orientations [93]. Some of these energies are compiled in Table 3. Surfaces with (0001) and (1120) orientations have been investigated with the large data set available for the (0001) orientation. We therefore restrict ourselves Corundum unit cell

Corundum(OOOI) surface

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a-Al203 4.7628 a-Cr203 4.954 V2O3 5.105 a-Fe203 5.035

20

13.003 13.584 14.449 13.72

40

line scan direction (A)

Fig. 14. Top: Schematic representation of the corundum unit cell and surface structures with different terminations. Bottom left: Scanning tunneling image of the Fe2O3(0001) surface. The darker area is believed to be metal terminated and the brighter area metal terminated according to corundum(0001)-Me [92]. Bottom right: List of lattice parameters for different oxides with corundum structure [15].

342

Table 3. Recent Surface Energies from First Principles [93]

Surface Orientations 0001 1011 1012 1120

Surface Energy 2.76 Jm"' 2.55 Jm"^ 1.97 Jm"^ 1.86 Jm"^

to a discussion of the latter case. The (0001) surface is an example for the socalled charge neutralized polar surfaces. These surfaces exhibit relatively strong relaxations remnant of the phenomena discussed in connection with the polar (111) rocksalt type surfaces, because if the (0001) surface were terminated between an oxygen and a metal layer, these surfaces would be polar with a diverging surface potential. However, the surface is terminated with only half of a metal layer present. Formally, this corresponds to a surface created by cutting the bulk structure within the buckled metal ion layers. Fig. 15 shows the results of structural determinations for the three related systems Al2O3(0001) [8], Cr2O3(0001) [94, 95] and Fe2O3(0001) [96]. In all cases a stable structure in UHV is the metal ion terminated surface retaining only half of the number of metal ions in the surface as discussed above [92]. The interlayer distances are very strongly relaxed down to several layers below the surface [10]. The perturbation of the structure of oxides due to the presence of the surface is considerably more pronounced than in metals, where the interlayer relaxations are typically of the order of a few percent. The absence of the screening charge in a dielectric material such as an oxide contributes to this efSide view Y-î

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"'/•

A

^ < ; ,-*% f^

^

"^^

#

" i , -'-^^ ^ ^

-

A2 A4 A6

^3 A5

AUO 2^3

CfoO 2'^3

Weak lateral relaxation

No lateral relaxation

FesOg

Fe term.

and

O term.

Fig. 15. Experimental data on the structure of corundum type depolarized (0001) surfaces (side and top views).

343

Fig. 16. Low energy electron diffraction diagram of ¥203(0001) grown on Au(l 11) [98]

feet eonsiderably. Another material belonging to this family is V2O3 which can also be prepared as a (0001) surface [97]. The structure of this system has, however, not yet been thoroughly investigated, although the LEED pattern, as shown in Fig. 16, is quite sharp [98]. In this case the V2O3 (0001) surface has been prepared as a thin film on a Au(lll) substrate by evaporation and oxidation. We assume that the surface structure is similar to the related cases discussed above. It has recently been pointed out that oxide structures may not be as rigid as one might think judged on the relatively stiff phonon spectrum in the bulk [99, 100]. In fact, at the surface the phonon spectrum may become soft so that the geometric structure becomes rather flexible, and thus also very much dependent on the presence of adsorbed species [100]. The vibration modes of a clean Cr2O3(0001) surface have been studied under ultrahigh vacuum conditions with high resolution electron energy loss spectroscopy [100]. Fig. 17 shows the spectrum of the clean surface at the bottom. A mode, which is confined to the first few atomic layers of the oxide, was detected, in addition to the Fuchs-Kliewer phonons, at 21.4 meV. This mode shows only a very small isotopic shift when the Cr2O3(0001) film is prepared from ^^62 instead of ^^02. In contrast to the Fuchs-Kliewer phonons which extend deeply into the bulk the 21.4 meV mode is very sensitive towards the presence of adsorbates, as is shown in Fig. 17. This can be seen by its attenuation upon exposure of the surface to CO, which in this case is only weakly adsorbed. Full-potential linearized augmented plane wave calculations [100] show that this mode is an in-phase oscillation of the secondlayer oxygen atoms and the Cr atoms of the two layers below. A schematic representation of the mode is shown in the inset of Fig. 17. Bulk oxide stoichiometrics depend strongly on oxygen pressure, a fact that has been recognized for a long time and we have alluded to above [10]. So do oxide surfaces, structures and stoichiometrics, a fact that has been shown again in a recent study on the Fe2O3(0001) surface by the Scheffler and Schlogl

344

HREELS

Cr2O3(0001)/Cr(110) at 90 K Ep=7.5 eV, FWHM=2.3 meV

CO exposure

1.2 L

0.6 L

.50 clean C r , 0 , 150 Energy [meV]

Fig. 17. Electron energy loss spectra in the vibrational regime of clean and CO-covered Cr2O3(0001)/Cr(110) surface. (Ep! 7.5 eV, specular scattering) The inset shows a schematic representation of the calculated normal mode of a Cr2O3(0001) surface at 21.4 meV, according to ref. [100].

groups [92]. In fact, if a Fe203 single crystalline film is grown in low oxygen pressure, the surface is metal terminated while growth under higher oxygen pressures leads to oxygen termination. This surface would be formally unstable on the basis of electrostatic arguments [10]. However, calculations by the Scheffler group have shown that a strong rearrangement of the electron distribution as well as relaxation between the layers leads to a stabilization of the system. STJVl images by Weiss and co-workers corroborate the coexistence of oxygen and iron terminated layers and thus indicate that stabilization must occur. Of course, there is need for further structural characterization. For the surface of Cr2O3(0001) infrared spectroscopic data suggest a different termination when the clean surface is exposed to oxygen [101] as will be discussed below. The electronic structure of corundum type oxide surfaces have been studied with photoelectron spectroscopy [102, 103] as well as with electron energy loss spectroscopy [104]. Henderson and Chambers [103] have recently reviewed the photoelectron spectroscopic evidences. Still, there seems to be an ambiguity about the oxidation state of the chromium ions in the immediate surface layer. XPS Cr 2p data may be to a large extent reconciled as being due to Cr(III) in the surface, because there is obviously no change in the spectra as a function of

345

emission angle. However, cluster calculations indicate that even though the number of d-electrons on the Cr ions is compatible with Cr^^, the total charge on the Cr ions is closer to Cr^^ because of a considerable hybridization between the chromium and oxygen wave functions [104]. In a previous review we have discussed the assignment of the local d-excitations on chromia surfaces and the reader is referred to this paper for details [7]. We would like to mention, however, a particular aspect of surface charge transfer excitations for the case of chromia(OOOl) which also touches upon the question of the oxidation state of surface chromium ions. The charge transfer states often mark the edges of the band gap in transition metal oxides, and they are also modified by the presence of the surface. Unfortunately, NiO does not show this effect very clearly. The electronic excitations at the chromia(OOOl) surface, on the other hand, are good examples to discuss the phenomenon. Fig. 18 shows the EELS spectrum [105-107] of this surface at low temperature. Again, a thin film has been used. The sharp features in the band gap are excitations within the manifold of d-orbitals. A detailed discussion [107] has shown that the excitations are characteristic of surface Cr ions with three delectrons. However, the Cr ions do not carry a net charge of 3+ as expected (and found for the bulk Cr ions) but rather of 2+ charge due to strong hybridization with the neighboring oxygen atoms. When we now perform EELS measurements after adsorption of CO2, not only the d-excitations are influenced but also the very intense feature near 4.8 eV. Again, on the basis of cluster calculations performed in the Staemmler group [104], it has been possible to assign this intense feature to a surface charge-transfer excitation at the band gap, which is shifted to lower energy as compared with the corresponding excitation in the bulk (see Fig. 18). The decrease in energy is reasonable because the Cr d-orbispecular I detection lE„=100eV|

ind gap

i 20 L CO2 IT=300 K

\ clean 1

2

'

I

4

'

I

6

I

'

I '

I ' I

10 12 14 16

Energy loss [eV] Fig. 18. EELS spectra of the clean and adsorbate covered Cr2O3(0001) surface. Peak assignments are indicated [105-107].

346

6 L ^''O2/Cr,O3(0001) oII o

^V^^n^^^.W^-^-.V.-.

..^.^-v.

415 K

£ 1 0)

o 150 K 90 K 1

800

'

1

'

1

'

1

'

1

'

1

'

r

1000 1200 1400 Wavenumbers [cm'']

Fig. 19. IRAS spectra after adsorption of a saturation coverage of O2 on Cr2O3(0001)/Cr(l 10). The spectrum at the bottom was taken at 90 K. Other spectra are taken at the temperatures indicated [101].

tals have been lowered, the more open surface structure (see Fig. 15) and the coordination of the Cr ions to only three oxygen ions allows for better charge separation than in the bulk. In NiO the effect is less pronounced because the surface Ni ions are still five-fold coordinated. The analysis of the electronic structure of Cr2O3(0001), as discussed so far, has been performed at 90 K. We note that upon increasing the temperature, the structure of the surface changes [107], and we have speculated that these changes are connected with changes in the magnetic structure of the surface. Oxide surface magnetism is a field that needs to be explored in the future. Before we discuss the adsorption and bonding of molecules such as CO and NO to the corundum type surfaces in order to accentuate the differences with respect to the rocksalt type surfaces, we would like to comment on the chemistry of molecular oxygen on the corundum type surfaces. As the (0001) corundum type surfaces are depolarized polar surfaces we always have to define whether we are talking about the metal terminated or the oxygen terminated surface. If not mentioned specifically we always assume metal termination. Chromia when exposed to oxygen leads to a very interesting chromyl terminated surface. What do we know about the chemistry between molecular oxygen and chromia(OOOl) that leads to this final result? Figure 19 shows the appearance of a sharp line in infrared spectra which develops out of a broad feature after heating the surface to room temperature. As has been discussed earlier and also supported by isotopic labeling experiments by Dillmann et al. [101] chromyl groups are formed on the surface and it was suggested that the chromium ions in the surface layer of the clean surface bind oxygen to form such chromium-oxygen double bonds. As is obvious from Fig. 19 there are adsorbed oxygen precursors to the chromyl

347

(TDSl m/e = 34

^^02, ""Ôs mixture (50:50) / Cr2O3(0001)/Cr(110) \

r-~JL \r~JX

'_ A _A

rv 100

1x10-"'^ mbar

\h

Z

1x10"^ mbar

I

3

°-^'"

0.1 L

10L

^ /V r^ 100

Temperature [K]

300

rn/e = 32 500

Fig. 20. Thermal desorption spectra of oxygen desorbing from Cr2O3(0001) [101].

formation. The width as well as frequency is indicative of a O2' species being adsorbed on the surface in a more or less upright position [101]. Fig. 20 is indicative of the thermal chemistry going on the surface. The TDS data reveal that around room temperature most molecular oxygen is lost from the surface. This molecular oxygen does not result from a recombinative desorption as can be seen directly from the isotopic labeling experiments. There is basically no isotopic mixing beyond the originally present contaminants. Above room temperature there is a very high temperature peak due to recombinative desorption stemming in all likelihood from the chromyl oxygen present on the surface at elevated temperatures, and observed in the infrared spectra. We will see further below examples of how these chromyl groups influence the chemistry of other adsorbed species on the surface. CO adsorption on corundum type (0001) surfaces is interesting because, opposite to checker board rocksalt type surfaces, on the (0001) surface of Cr203 CO has been shown by angle resolved photoemission to assume a strongly inclined almost flat adsorption geometry. Infrared spectroscopy and recent cluster calculations by Staemmler's group have revealed the reason for this [108]. Briefly, on the basis of TDS data the activation energy for the desorption of CO from chromia(OOOl) is estimated to be about 45 kJ/mol. So the appropriate state may be termed weakly chemisorbed. In addition, more weakly bound CO molecules (28 kJ/mole) which are present on the surface at lower temperature have been detected. Multilayer adsorption of CO does not occur on Cr2O3(0001)/Cr(110) at 90 K. IR-spectra confirm the presence of different CO species on Cr2O3(0001). Fig. 21 shows a series of IRAS-spectra as a function of initial CO dosage at 90 K. Using a CO-exposure of 0.05 L we detect a signal of the CO-stretching vibration at 2178 cm"^ and at 0.1 L a second signal appears at 2165 cm'\ With increasing coverage the high-frequency signal considerably

348

increases and slightly shifts to lower wave numbers. The weak signal at 2165 cm'^ does not gain further intensity. Instead it disappears in the shoulder of the main signal. This state is to be correlated with the chemisorbed state observed in TDS. Finally, at exposures above 3 L a third signal appears at 2136 cm"\ which is due to the more weakly held species. Cluster calculations and TDS measurements agree fairly well as far as the adsorption of CO at the Cr2O3(0001) surface at low coverages is concerned. The chemisorbed species corresponds to an undissociated CO molecule residing in the Os-hollow position with the CO axis strongly tilted against the surface normal, oriented along a line connecting two surface Cr ions and in a hollow position with respect to the O^" ions in the second layer. The more weakly held species corresponds to a shallow local minimum in the potential energy curve, in which CO lies nearly flat on the surface, also between two surface Cr ions, but on top an O^" in the second layer (0-ontop position). In contrast to the adsorption geometry there is some discrepancy concerning the adsorption energies as calculated by the cluster calculations. The values of 45 and 28 kJ/mol estimated for the two TDS peaks from the Redhead formula are considerably larger than the calculated values of about 27 and 15 kJ/mol. As has been pointed out in connection with NiO, it has to be expected that the calculated values are too low since neither adsorbate-induced relaxations of the IRAS I

CO / Cr2O3(0001)/Cr(110)

90 K

2178

1 1900

1 2000

1 2100

1

\

1

1

2200

2300

2400

2500

Energy [cm'']

Fig. 21. FTIR spectra of CO adsorbed on Cr2O3(0001) at 90 K as a function of exposure [108].

349

Cr2O3(0001) surface nor extra cluster van der Waals attractions are accounted for. Whether the inclusion of these effects will lead to a perfect agreement with the measured values is not clear. In any case it seems justified to call the O3hoUow minimum a „chemisorption" though the dominant contribution to the interaction energy has electrostatic origin and is characterized by the situation in which the quadrupole moment of CO fits into the strongly inhomogeneous electrostatic field above the Cr2O3(0001) surface. Good agreement between theory and experiment exists concerning the frequencies shifts of the CO stretching vibration. The calculations predict a blue shift of about +22 cm'^ for the CO in the Os-hollow position, which compares favorably with the observed shift of 30-35 cm'^ in the most prominent IRAS band. For the weakly bound CO in the O-ontop position a small red shift of about -5 cm"^ is obtained; the signal at 2136 cm'^ observed at higher coverages for the weakly bound CO species does indeed exhibit such a small red shift of7 cm"\ However, the ab-initio calculations cannot offer an explanation for the shoulder in the IRAS peak at 2165 cm'^ observed for dosages in the range between 0.1 and 0.8 L (Fig. 21). The frequency itself seems to correspond to that of CO in the Cr-ontop position, but this does not constitute a local minimum and will hence be not occupied. We also might speculate that the shift observed in the main IRAS peak from 2178 cm"^ at low coverage to 2170 cm"^ at high coverage is due to lateral CO interactions, most probably because each CO molecule is slightly squeezed out of the absolute minimum by neighboring CO molecules. The results of the present study should also be interrelated with UPS and NEXAFS experiments which have been performed on the CO/Cr2O3(0001)/Cr(110) system previously [105, 106]. Whereas the present data allow the identification of strongly and weakly bound CO molecules, only the more strongly adsorbed species has been observed in the photoelectron spectroscopic studies. The explanation may be found in the experimental conditions: At the slightly elevated surface temperature (105 K) used in those experiments and because of UV- and X-ray-induced CO-desorption the weakly bound CO species could not be detected. (In fact, the UPS and NEXAFS measurements had to be performed at a background CO pressure of 1 x 10"^ mbar.) A very interesting point is the adsorption geometry of the CO-species on Cr2O3(0001)/Cr(l 10). On the basis of the ARUPS and NEXAFS data [105, 106] the tilt angle between surface plane and molecular axis of the chemisorbed CO was estimated to be about 20°. The present IR-data are consistent with an adsorption model of a strongly tilted CO-chemisorbate, in line with the cluster calculations. On the other hand, the IR-signal of the weakly bound CO-species (at» 2136 cm'^) shows a very small intensity compared to the relative population of states as deduced from TDS (not shown) [108]. Neglecting possible differences in the strength of the dynamic dipoles we can conclude that for the

350

weakly bound CO molecules the tilt angle between surface normal and molecular axis is even larger than for the chemisorbed CO: the weakly bound CO lies almost flat on the surface. This conclusion is confirmed by the cluster calculation as well. It is very interesting to note that CO adsorption on ¥203(0001) seems to show a very much related behavior although the study of orientation of the molecule is not complete yet [98]. Water adsorption has been studied for Fe203 as well as for Cr2O3(0001) surfaces by Henderson and Chambers [103, 109, 110]. In general, both, dissociative adsorption as well as molecular water adsorption has been observed. Fig. 22 shows TDS data for chromia [103]. The low temperature peak corresponds to molecular physisorption, including multilayer formation (peak at 165 K), while the higher temperature peaks are consistent with, partially, recombinative desorption of chemisorbed water from the surface. The two states at 295 K and at 345 K are populated sequentially, starting with the high temperature state at the lowest exposure. Vibrational spectra revealed the presence of isolated hydroxyl groups. These are associated with the peak at 345 K. In fact, Henderson and Chambers [103] proposed that water in the dissociated state is comprised of two distinct hydroxyl groups, i.e. a terminal and a bridging hydroxyl group. Molecularly chemisorbed water desorbs at 295 K. These authors also mention the strong influence of a predosage of the surface by oxygen on the surface

100

200

300 Temperature [K]

400

Fig. 22. Thermal desorption spectra of water desorbing from CriOsCOOOl) after various exposures [103].

351

chemistry. Their observations are in line with results on the adsorption of CO2 described briefly below [111, 112]. Water adsorption on alumina leads to molecular adsorbates [110]. There are some reports on the formation of isolated hydroxyl groups on regular sites [113, 114]. Whether they are the results of dissociative water adsorption is not completely clear, although there is a recent theoretical paper that predicts the dissociation of water on Al2O3(0001) by protonation of a surface oxygen, while the molecule previous to dissociation resides on a cation site [115]. There is no experimental verification of this result as yet. There are, however, explored routes to hydroxylate alumina surfaces either by exposure to atomic hydrogen or by a combination of aluminium deposition and subsequent hydrolyzation with water [116]. Carbondioxide adsorption is another example of molecular adsorption which has been studied quite extensively in the past. While for some time is was generally accepted that CO2 forms carbonates with chromia surfaces very readily upon exposure, it was recently observed that carboxylate species may be formed. TDS spectra indicate [111, 112] that there are more weakly and less weakly bound CO2 species on the surfaces. We have studied the nature of those species by various techniques including infrared spectroscopy. Fig. 23 shows several sets of IR spectra. The pair of sharp bands around 2300 cm"^ can easily be assigned to the more weakly bound CO2 with only slightly distorted structure Identification of adsorbed species

Comparison with microcrystalline samples CO2 on Cr2O3(0001)/Cr(100)

'^COp

1282 1313

"'135]W1342" 190 K

1289

"Isoi^ 342

300 K

12691 ^2002"^

il

250 K 90 K

CO2 on polycrystalline a-chromia

^

^3002^"

15i3|^

90 K

1^89

«^„

825 ° ' ^

-

1242 16251695

2344I2358

D. Scarano, A. Zecchina, Universlta Degli Studi Di Torino

CO2 on O/Cr2O3(0001)/Cr(100) s^T k.

90 K

si

1

+ 4 L CO2 at90K

2352

1289 [2375 T — I — I — I — I — I — I — r 1000 1200 1400 1600 1800 2000 2200 2400 energy [cm'']

T CO2 on oxygen precovered chromia

" ~ ^

D. Scarano, A. Zecchina, Universita Degli Studi Di Torino

0.1 t o r r 0 2 at 100 K

—1—1—1—1—1—1—1—1—I—r1200 1600 2000 energy [cm''']

Fig. 23. IRAS spectra of CO2 adsorbed on Cr2O3(0001) surfaces and on polycrystalline chromia, left panel: IRAS spectra at different surface temperatures and with isotopically labeled CO2 as well as Cr203, right panel: adsorption of CO2 after pre-adsorption of oxygen [112].

352

as compared with the gas phase species. By a combination of isotopic labeUng the adsorbed CO2 (shift of frequencies) as well as the oxide layer (no shift of CO2 bands) we have demonstrated that the single band centered around 1300 cm'^ is due to the presence of a carboxylate species, i.e. a bent anionic CO2 species, and not, as perhaps expected, to a carbonate [105]. The bands between 1610 and 1700 cm"^ are missing because of the applicability of surface selection rules in thin film systems. This means, all nontotally symmetric bands are suppressed in intensity. A quick comparison with CO2 adsorption on chromia microcrystalline material as shown in Fig. 23 indicates the presence of the bands between 1610 and 1700 cm"^ as expected for adsorption on a bulk dielectric material. It is remarkable how similar the thin film data are in comparison with the microcrystalline material. This has been discussed in detail by Adriano Zecchina's group [117]. Also, the response of the two systems with respect to preadsorption of oxygen is very similar. In fact, as shown in Fig. 23, CO2 adsorption in form of the less weakly bound CO2" is fully suppressed on the thin film system and very strongly attenuated for the microcrystalline system. This indicates that CO2 occupies the chromium sites, because we know that oxygen from the gas phase adsorbs on the chromium ions. As we remarked above, ELS [106] and XPS [102] spectra of the Cr2O3(0001) surface have been used to deduce that the Cr-ions in the surface are in a low oxidation state, i.e. Cr^^ as opposed to chromium ions in the near surface and bulk regions. It is therefore not surprising that such a surface provides electrons to adsorbed molecules, leading to electron transfer as, for example, documented by the formation of O2" and CO2". The low valence state of the Cr surface ions also has consequences in other reactions, such as the polymerization of ethene which has been studied on Cr2O3(0001) [118], and in connection with other model studies [119]. A tri-atomic molecule that has so far not found wide attention is NO2, although the molecule is interesting with respect to environmental catalysis [120]. We have studied the interaction of NO2 with alumina films [121]. Fig. 24 shows TDS data for this system: The dosages given in Fig. 24 are expressed in terms of the rise of the base pressure for 10 s at a fixed distance between doser and surface. By multimass TPD measurements it has been checked that NO2 is the only desorbing species. For small coverages a single desorption peak is found at 135 K which saturates at higher exposures. This first desorption peak can be assigned to isolated NO2 molecules physisorbed to the oxide with a binding energy of 38 kJ/mol estimated according to Redhead's analysis [65]. At higher coverages a second desorption feature at nearly 10 K higher temperature appears. The observation of a TPD peak at higher temperature which does not saturate upon increase of exposure is quite unusual compared to the "normal" desorption behavior of, e.g., N02/Au(l 11) published by Bartram and Koel [122]. It indicates an intermolecular interaction energy larger than the binding energy from the

353

TPD

N02/Al203(l 1 l)/NiAl(l 10)

(m=46)

S-IO"'^ mbar A* 90 K [42]. Fig. 7A shows a waterfall plot of MIE spectra of NO exposed to a low-defect MgO thin film [grown on Mo(lOO)] at 90 K. In order to create a high density of defects, the MgO/Mo(100) film was annealed to 1100 K for ~1 min. and quenched to -100 K; this treatment has been shown in a previous HREELS study to generate defects at the near-surface [43]. The MgO/Mo(100) sample was thus similarly treated and NO was exposed to this surface. Fig. 8 shows a waterfall MIES plot on the defect-rich MgO thin film. A --0.6 eV work function increase upon saturation of the NO coverage (Figs. 7B and 7C) is indicative of charge transfer from the MgO to the adsorbate. There is a -25% attenuation of the O 2p intensity at the onset of the appearance of bands A, B and C, which we assign to the ionization of the 27i, 7a and \n states of N2O, respectively. The presence of N2O was verified by TPD, which showed the characteristic mass spectral cracking pattern for the molecule. Separate MIES/UPS and TPD measurements were carried out independently after N2O adsorption on the MgO/Mo(100) to further verify these peak assignments and preclude the

384 — I

'

1

'

1

'—

M g O ( 1 0 0 ) 95 K

as-prepared (O

c

(D

MgO 0 ( 2 p )

CO LU

with defects -O o

,

50

100

150

200

•

O—^

#—•

as-prepared

250

300

35

NO dose [L]

Fig. 8 Evolution of MgO 0(2p) and N2O 7a as a function of NO coverage. The open and filled symbols denote defect-rich and as-prepared Ti02, respectively.

possibility of N2O impurity in the NO gas. These peak positions match closely with reported gas-phase measured ionization energies for NjO [44]. The attenuation is more pronounced for the high-defect surface (Fig. 7B) where the O 2p peak is attenuated by 80%. The UPS waterfall plot (Fig. 7C) acquired simultaneously with MIES shows only minor changes (A, B and C bands not well-developed) during the adsorption, indicating that a relatively small fraction of the surface is covered with adsorbate. Using integrated TPD peak areas to calibrate the CH3OH coverage on MgO, Gtinster et al. [27] found that the fraction of adsorbate coverage could be quantified (approximately) via attenuation of the O 2p MIES signal. Based on the abatement of the O 2p intensities following NO adsorption on the MgO/Mo(100) systems, we estimate the NO saturation coverage to be 13±2 % and 35±5 % for the low- and high-defect MgO films (Figs. 7A and 7C), respectively [45]. Fig. 8 shows a plot of the MIES O 2p and N2O 7a peak intensities as a function of NO coverage [35]. Saturation occurs at varying coverages, depending upon the extent of thermal quenching. The as-prepared MgO films appear to contain a significant density of extended defects (i.e. steps, kinks and grain boundaries), which increased after thermally quenching the sample. A noticeable

385

broadening of the O 2p MIES peak is observed after the quench. This broadening, accompanied with a decrease in the band gap, is diagnostic of an increased extended defect density [20, 46]. The expansion of the extended defects coincide with the accumulation of N2O and appears to adsorb on the defect sites, perhaps near steps or grain boundaries. We postulate that dissociative adsorption of NO on MgO to produce N2O occurs via the following mechanism: N O ^ NOd -> Nd + 0 ->N20d

(1)

where the subscript "d" signifies a defect site, stimulating the production of N2O. Differences in the MIE spectra between the as-prepared MgO and the quenched MgO is strongly suggestive that defect density is an important factor controlling NO adsorption kinetics. 3. TITANIA Ti02 is an important reducible metal oxide support used for stabilizing nanosized noble metal clusters that catalyze a number of industrially relevant reactions in photocatalysts, chemical sensors and heterogeneous catalysis. These reactions include low-temperature CO oxidation [47-65], hydrogenation and partial oxidation of hydrocarbons [56, 58, 66, 67] and the selective oxidation of propylene to propylene oxide [66, 68]. Rutile TiO2(110) is the most thermodynamically stable form of Ti02 and has thus been the focus of numerous experimental and theoretical studies due to its well known electronic and geometric properties. In addition to using this single crystal, complementary experimental studies have also been performed using thin film Ti02 [69-72], which can be studied with the use of charged particle probes. 3.1. Structural and electronic properties of Au on TiO2(110) Single crystal TiO2(110) has been widely used for obtaining atomicallyresolved images via scanning probe microscopy. Images showing individual atom rows running along the [001] have been reported by various groups [7376]. The n-type semiconductor was found to be amenable for STM and charged particle spectroscopy after cycles of Ar^ bombardment and annealing to 700-1100 K [77]. This property makes the single crystal an ideal support for preparing and surface-analyzing supported metal cluster model catalyst systems. In particular, model planar catalysts consisting of finely dispersed Au and Ag clusters of varying size have been prepared on TiO2(110) for study. Since atomic-scale characterization of the unique physical and chemical properties of these nanosized supported metal cluster systems have been

386

discussed at length elsewhere [6, 7, 78-80], only one selected example will be illustrated in this section with emphasis on a recent landmark work, namely the finding that the onset of catalytic activity for a Au/Ti02 system coincides at a specific cluster size accompanying the appearance of non-metallic admetal clusters [81]. Fig. 9A shows a constant current topographic scanning tunneling micrograph (CCT-STM) of 0.25 monolayer equivalent (MLE) of Au deposited onto single crystal TiO2(110)-(lxl). The clusters were deposited at 300 K, followed by an 850 K anneal for 2 min [81]. The inter-atomic distances of the Ti02 rows are --0.65 nm. Three-dimensional clusters with dimensions of-2.6 nm X 0.7 nm (diam. and height, corresponding to a 2-3 atom thickness) are observed as bright protrusions and preferentially nucleate at the step edges. Fig. 9B shows STS taken at various points on the surface. The ST spectra had been acquired by positioning the STM tungsten tip at a predetermined point (over a cluster) and interrupting the feedback loop. The tunneling current (I) as a function of voltage bias (V) was measured, providing I-V curves measured for a variety of Au cluster sizes on the Ti02 surface. The "plateaus" between the inflection points of the I-V curves along the bias voltage axis (where the slope rapidly changes at the zero tunneling current) denotes the band gap. The STS lineshapes thus show that the electronic character of these deposits vary between that of a metal and non-metal depending upon its size. As the clusters enlarge, they gradually adopt metallic characteristics with an enhanced density of states (DOS) at the Ep, signified by more abrupt changes in the I-V curve for larger clusters. For example, the cluster with the 2.5 nm x 0.7 nm size has a larger band gap than that of the 5.0 nm x 2.5 nm cluster. Smaller Au clusters have non-metallic properties as signified by relatively large band gaps in the STS whereas larger clusters having bulk-like characteristics have virtually no band gap. Intermediate-sized 4.0 nm x 1.5 nm and 3.0 nm x 1.0 nm clusters also follow this cluster-size-band gap dependency. Interestingly, there is a correlation between both the Au cluster size and electronic structure with its catalytic activity for CO oxidation on Au/Ti02. Fig. lOA shows a plot of activity for the oxidation of CO [expressed as (product molecules)/((total atoms on surface sites) x (second)) or turnover frequency (TOF)] on Au clusters supported at 350 K on a Ti02 thin-film surface (prepared via "hot" filament evaporation on a Mo(lOO) substrate [82]) as function of cluster size. The product molecules (CO2) from the 2 CO + O2 -> 2 CO2 reaction was extracted from the reactor with a vacuum syringe, compressed and then analyzed using a gas chromatograph [81]. In a parallel experiment, a one-to-five C0:02 mixture had been reacted on the Au(0.25 MLE)/TiO2(110)-(lxl) surface at 40 Torr [61, 67, 81]. The cluster sizes obtained from STM of equivalent Au coverages on the respective Ti02 surfaces

387

Bias Voltage (V) Fig. 9 (A) A CCT-STM image of 0.25 MLE Au deposited onto Ti02(l lO)-(lxl) prepared just prior to a C0:02 reaction. The sample had been annealed to 850 K for 2 min.; (B) STS data acquired for Au clusters of varying sizes on the Ti02(l 10)(1x1). An STS of the Ti02 substrate, having a wider band gap than the Au clusters, is also shown as a point of reference.

vêre used to calculate the TOF. For each point in Fig. lOA, a particular Au cluster size was prepared and then subjected to the C0:02 reaction. The activity for the Au/Ti02 catalyst exhibits a maximum 1.90 TOF at an average cluster size of 3.5 nm x 1.0 nm (300 atoms/cluster). Fig. lOB shows a plot of the STS band gaps measured over the same Au cluster size range used in the C0:02 reactions. Maximum activity coincides with the size-dependent metalto-nonmetal transition, showing the first evidence of the reaction being sensitive to the electronic structure of the Au particles. Fig. IOC shows a histogram showing the relative distribution of the Au sizes ranging from 2.04.0 nm in diameter at the maximum catalytic activity conditions, having drastically reduced band gaps of 0.2-0.6 V. It appears that 2-atom-thick clusters (with a 2.5-3.0 nm diam.) are optimally active for CO oxidation. Noteworthy is the fact that maximum activity is also observed at the same Au cluster size (2.9±0.5 nm diam.) on the (practical) high-surface area Ti02 catalyst [53] as in the (model) planar metal oxide, thus illustrating the utility of a surface science approach to obtain insight into "real world" systems.

388 2.20

o

Au clusters with a band gap of 0.2-0.6 V measured by STS

45 30

0

o

15

OH

0 0.0

2.0

4.0

6.0

8.0

10.0

Cluster Diameter (nm) Fig. 10 (A) The activity for CO oxidation at 350 K as a function of Au cluster size supported on TiOjCl lO)-(lxl) assuming a total dispersion of the Au. A 1:5 CO: O2 mixture had been used at a total pressure of 40 Torr. Activity is expressed as (product molecules/ ((total Au atoms) x (second)); (B) cluster band gaps measured by STS as a function of Au cluster size supported on TiO2(110)-(lxl). The band gaps were obtained while the corresponding topographic scan was acquired on various Au coverages ranging form 0.20.4 MLE. ( • ) 2D clusters, (O) 3D clusters, two-atom layers in height, (A) 3D clusters, three-atom layers or greater in height; (C) relative population of Au clusters (two-atom layers in height) that exhibit a band gap of 0.2-0.6 eV, as measured by STS of Au/ TiO2(110)-(lxl).

3.2. Adsorption of CO on Au/TiOj The nature of CO bonding to Ti02 is an important consideration for understanding how CO oxidation reactions are governed, in particular with regard to activity and selectivity. Previous theoretical and experimental work has shown that CO binds weakly onto the Ti02 support. Ab initio molecular orbital calculations indicate that CO absorbs through the carbon atom to the lattice Ti sites on the defect-free TiO2(110) surface with an adsorption energy

389

of 17 kcal/mol (0.73 eV per CO molecule) [83]. A similar result was obtained using a full potential linearized augmented plane-wave (FLAPW) method and along with the local density approximation (LDA); an adsorption energy value of 15 kcal/mol (0.64 eV per CO molecule) was determined [84]. In particular, the presence of metallic clusters supported on the metal oxide could dramatically affect the CO bonding properties. In order to better understand the effects of nanosized Au particles on CO bonding behavior, a study was performed comparing CO adsorption onto single crystal Au(llO) and a (0.5 MLE)Au/Ti02 planar model catalyst. This 0.5 MLE Au coverage corresponds to the formation of a particular mean cluster size in which maximum catalytic activity is known to occur [61]. Studies of CO adsorption on these respective substrates employing IRAS as a function of CO pressure and sample temperature were performed. Au(llO)(1x2) serves as an excellent model for comparison with highly dispersed Au/Ti02 systems since this reconstructed surface has a regular array of (111) facets with step-edges consisting of 6-coordinate atoms for every fourth atomic row [85-88]. This high density of low-coordination sites on the Au(l 10)-(lx2) is comparable to that of practical catalyst systems. IRAS is a useful technique for probing the chemical and physical properties of Au supported metal oxide systems. Coverage dependent absorption intensities of the CO probe molecule [89-91] can be measured in order to obtain isosteric energies of adsorption (Eads) via the Clausius-Clapeyron treatment of the adsorbate absorbance intensities [92]. As expounded by Woodruff et al. [90], relative shifts in frequency can be used to determine the relative position of the CO 2n^ state with respect to the Ep and thus provide insight to its bonding behavior to the metal. The 2n^ level lies relatively close to Ep at the onset of adsorption to the surface. Upon broadening of this band with increased CO coverage, the state can become less or more filled, depending upon whether it lies below (blue shift) or above the Ep (red shift), and can be experimentally measured using the surface vibrational spectroscopy. Eight IRAS data sets were acquired varying the CO coverage at a range of surface temperatures at pressures from 1.3 x 10"^ to 1.0 x 10"^ Torr on Au(110)-(lx2). At 1.3 X 10~^ Torr, the low pressure limit, the CO absorbance (with sample temperatures in parantheses) is a single entity at 2118 cm"^ (170 K) at the onset of adsorption. As the temperature was decreased, the CO coverage increased, resulting in a red shift of this feature to 2108 cm~^ (110 K) when saturation was reached. According to the Blyholder model [93], the red shift can be explained by a chemical effect due to decreased n back-bonding to the substrate concomitant with increased donation of electron density to the substrate upon CO adsorption. Comparable red shifts were observed at the

390

range of pressures used for these data sets. For instance, at the upper pressure limit used (1.0 x 10^ Torr) a frequency of 2116 cm~^ (230 K) was observed at the onset of adsorption, red shifting to 2106 cm"^ (110 K) at the CO saturation coverage (Fig. 11). The adsorbed CO measurements were red-shifted relative to gas-phase CO at 2143 cm"^ [4]. These isobaric data sets are rearranged into isotherms in Fig. 12. Integrated IRAS peak areas correlate linearly with coverage, setting the saturation intensity approximately equal to a 50% CO coverage (0 = 0.5) [94]. A representative set of isosteres are plotted at every 1.6% coverage increment up to and including 35%) (Fig. 13). CO adsorption energies can then be determined by the slopes of these isosteres via the Clausius-Clapeyron relation: d(\nP) d{\IT)

(2)

ads

R

where R is the gas constant. E^ds values are found to be -12 kcal/mol below a 6%

70 1 ' 1 ' 1 i T i T - ^

f"

2106 cm

//

^ >-

.-^60

5%

=>55

2

130 K

F^^ V 45 f ^ -" 0 J 35% Q-

'< ^ ' • ^ J 32% 300 K facilitates disruption of Au clustersupport interactions. 5.2. Imaging solution-deposited Au6(PPh3)JBF4]2 deposited TiOjCllO) Iwasawa and co-workers have recently begun employing phosphinestabilized Au complexes, Au(PPh3)(N03) and Au9(PPh3)8(N03)3, for preparing highly dispersed, nanosized Au particles on high surface area and asprecipitated oxide supports for catalytic reactions [119-122]. The use of phosphine-stabilized Au clusters provides a means of controlling monodispersion via synthesis of various-sized Au cluster precursors as starting materials (e.g. AU4, Aun, AU55, ^^c). Upon deposition and calcination on high surface area Ti02, the Au aggregated into 30 nm diam. particles, inhibiting its catalytic properties. More success was obtained on the as-precipitated Ti(0H)4, which resulted in clusters 3 nm in diam. that exhibited high activity for CO oxidation [122]. These Au complexes, while showing promise as precursors for practical catalyst preparations, have yet to be prepared on model metal oxide surfaces, making them amenable for surface sensitive probes and hence atomic-scale surface structure characterization. To this end, a phosphine-stabilized Au cluster, Au6(PPh3)6[BF4]2 (Au^L^) analogous to those used by Iwasawa and co-workers, has recently been deposited onto a single crystal Ti02(l 10) via solution deposition under ambient conditions (1 atm, 300 K) and probed with STM [123]. During initial attempts in depositing the Au^L^ onto the TiO2(110) crystal in CH2CI2 solvent, large agglomerations due to Au cluster aggregation of the complex (50-100 A in diam.) were observed. To avoid this problem, it was necessary to pre-treat the crystal with acetone in order to obtain the finely dispersed complex on the TiO2(110) crystal. Measured surface charge

403

Fig. 22 (A) 100 nm X 100 nm CCT-STM image of a freshly sputter-cleaned Ti02(l 10) surface. (B) 100 nm x 100 nm CCT-STM view of homogene -ously distributed AugLg clusters deposited onto TiO2(110). ( C ) 3 0 0 n m x 3 0 0 nm CCT-STM wide-range view of the Au6L6/Ti02(l 10).

100

150

200

250

300

404

densities and zeta potentials of colloidal mineral oxide suspensions, such as yAI2O3, Ti02 and Si02, have been reported to be strongly influenced by organic solvents [124-126]. A study by Rodrigues et al. [126] has shown that the presence of acetone lowers the net proton surface charge density (G^) of silica colloids as measured by potentiometric titration (at pH = 8). We postulate an analogous phenomenon occurring with the Au6L6/TiO2(110) system enabling the Au^L^ (which has a net charge of+2 without the BF4~ counterions) to better disperse and adhere onto the substrate. Fig. 22 shows a series of reproducible CCT-STM images of (A) a freshly, sputter-cleaned TiO2(110) to serve as a reference point (i.e. a blank), (B) Au^L^ deposited on the Ti02(l 10) surface and (C) a wide-range view of the Au6L6/TiO2(110). All images were obtained ex-situ using an Omicron UHV STM-1 scanner. The hexagold particles were widely and homogeneously distributed over the 1 cm^ TiO2(110) surface area. Multiple areas of the Au6L6/TiO2(110) that were sampled showed the same structural morphology (Fig. 22C). These particles are remarkably stable after prolonged exposure to air for 72 hours. Surface roughening is evident in Figs. 22B and 22C. However, the diagonal TiO2(110) terraces are still apparent. Using the integrated areas of the STM (assuming a one-atom-layer thick coverage), 10% of the TiO2(110) substrate was covered. XPS showed ample signal from the Au 4f core level, indicating its presence, and no signal from either B or F, indicating that the counterions were apparently removed by solvent evaporation. STM line profiles of these clusters showed an average height of 8±1 A (n = 10), which is approximately the same volume as expected for the Au^L^ entity. Experiments are currently under way for optimizing ligand removal in order to prepare these surfaces for various catalytic reactions (e.g. CO oxidation, propylene oxidation, etc.) in hopes of monitoring atomicscale structure of these practical admetal clusters during their performance. 6. SUMMARY Recent significant strides have been made on the development of available techniques that can be used to investigate the interfacial activity of reactions occurring at well-defined metal oxide interfaces. MIES/UPS has been an especially useful technique for elucidating geometric molecular orientation of molecules as they adsorb onto the metal oxide; the high surface sensitivity of the probe (in conjunction with TPD) can be exploited to quantitate near-surface defects and its effects on molecular adsorption. STM and STS has been shown to be particularly effective in monitoring the admetal cluster size and electronic structure effects on model planar oxide catalyst systems, showing a correlation of relative activity with these properties. Further advances have been made in gaining insight on the nature of the CO bond to these highly reactive systems

405

and further extended into the study of mixed metal oxide systems using IRAS, LEED and TPD. In the area of scanning probe microscopy, progress has been made in bridging the pressure gap for obtaining atomic-resolution images at realistic catalytic reaction conditions. These tools will play an important role in the frontier areas of surface science, permitting us to monitor molecularlevel interactions in order to control the size, structure and spatial distribution of adsorbed admetal particles and/or reactants (e.g. NO, CO) on oxides for the rational design of chemically active surfaces. ACKNOWLEDGMENTS We acknowledge with pleasure the support of this work by the Robert A. Welch Foundation, the Department of Energy, Office of Basic Sciences, Division of Chemical Science and the Dow Chemical Company. We also thank Jeffrey A. Stultz for proofreading the manuscript. REFERENCES 1. D. P. Woodruff and T. A. Delchar, Modem Techniques of Surface Science, 2nd ed.; Cambridge University Press: Cambridge, 1994. 2. G. A. Somorjai, Introduction to Surface Chemistry and Catalysis; Wiley: New York, 1994. 3. D. W. Goodman, J. Vac. Sci. Technol. A 14 (1996) 1526. 4. S. C. Street and D. W. Goodman, Annu. Rev. Phys. Chem. 48 (1997) 43. 5. S. C. Street and D. W. Goodman, in: Growth and Properties of Ultrathin Epitaxial Layers; (Elsevier, Amsterdam, 1997) 375. 6. D. R. Rainer and D. W. Goodman, J. Mol. Catal. A 131 (1998) 259. 7. T. P. St. Clair and D. W. Goodman, Top. Cat. 13 (2000) 5. 8. A. J. Martin and H. Bilz, Phys. Rev. B 19 (1979) 6593. 9. G. Lakshmi and F. W. de Wette, Phys. Rev. B 22 (1980) 5009. 10. G. Lakshmi and F. W. de Wette, Phys. Rev. B 23 (1981) 2035. 11. V. E. Henrich, G. Dresselhaus and H. J. Zeiger, Phys. Rev. B 22 (1980) 4764. 12. T. Urano, T. Kanaji and M. Kaburagi, Surf Sci. 134 (1983) 109. 13. M.-C. Wu, J. S. Comeille, C. A. Estrada, J.-W. He and D. W. Goodman, Chem. Phys. Lett. 182(1991)472. 14. J.-W. He, J. S. Comeille, C. A. Estrada, M.-C. Wu and D. W. Goodman, J. Vac. Sci. Technol. A 10 (1992) 2248. 15. E. Lindholm, Faraday Disc. Chem. Soc. 54 (1972) 200. 16. H. Kubota, T. Hirooka, T. Fukuyama, T. Kondow, K. Kuchitsu and A. J. Yencha, J. Electron Spectrosc. Relat. Phenom. 23 (1981) 417. 17. J. Giinster, G. Liu, V. Kempter and D. W. Goodman, Surf Sci. 415 (1998) 303. 18. K. Y. Yu, J. C. McMenamin and W. E. Spicer, Surf Sci. 50 (1975) 149. 19. S. Hiifner, Photoelectron Spectroscopy; Springer-Verlag: Berlin, 1995.

406 20. D. Ochs, W. Maus-Friedrichs, M. Brause, J. Giinster, V. Kempter, V. Puchin, A. Shluger and L. Kantorovich, Surf. Sci. 365 (1996) 557. 21. J. Gunster, G. Liu, V. Kempter and D. W. Goodman, J. Vac. Sci. Technol. B 16 (1998) 996. 22. L. H. Tjeng, A. R. Vos and G. A. Sawatzky, Surf. Sci. 235 (1990) 269. 23. S. Masuda, M. Aoyama, K.Ohno and Y. Harada, Phys. Rev. Lett. 65 (1990) 3257. 24. S. C. Street, Q. Guo, C. Xu and D. W. Goodman, J. Phys. Chem. 100 (1996) 17599. 25. M. Xi, M. X. Yang, S. K. Jo and B. E. Bent, J. Chem. Phys. 101 (1994) 9122. 26. W. Ueda, H. Ohowa, K. Iwasaki, T. Kuwabara, T. Ohshida and Y. Morikawa, Stud. Surf Sci. Catal. 90 (1994) 35. 27. J. Gunster, G. Liu, J. Stultz and D. W. Goodman, J. Chem. Phys. 110 (1999) 2558. 28. J. Gunster, G. Liu, J. Stultz, S. Krischok and D. W. Goodman, J. Phys. Chem. B 104 (2000) 5738. 29. H. Yamakado, M. Yamauchi, S. Hoshino and K. Ohno, J. Phys. Chem. 99 (1995) 17093. 30. M. C. Gallagher, M. S. Fyfield, J. P. Cowin and S. A. Joyce, Surf. Sci. 339 (1995) L909. 31. D. W. Turner, C. Baker, A. D. Baker and C. R. Brundle, Molecular Photoelectron Spectroscopy; John Wiley and Sons, Ltd.: London, 1970. 32. M. S. Banna, B. H. McQuaide, R. Malutzki and V. Schmidt, J. Chem. Phys. 84 (1986) 4739. 33. V. Cermak and A. J. Yencha, J. Electron Spectrosc. Relat. Phenom. 11 (1977) 67. 34. J. H. Lunsford, Catal. Today 6 (1990) 235. 35. A. Kolmakov, J. Stultz and D. W. Goodman, J. Phys. Chem. B, in press. 36. T. Gotoh, Y. Fukunaga and S. Takagi, Surf Sci. 357-358 (1996) 690. 37. L. N. Kantorovich, J. M. Holender and M. J. Gillan, Surf Sci. 343 (1995) 221. 38. L. N. Kantorovich, A. L. Shluger, P. V. Sushko, J. Gunster, P. Stracke, D. W. Goodman and V. Kempter, Faraday Discuss. 114 (1999) 173. 39. P. R. Underhill and T. E. Gallon, Solid State Comm. 43 (1982) 9. 40. V. E. Henrich and P. A. Cox, The Surface Science of Metal Oxides; Cambridge Univ. Press: New York, 1994. 41. X. Zhang, A. B. Walters and M. A. Vannice, J. Catal. 146 (1994) 568. 42. R. Wichtendahl, M. Rodriguez-Rodrigo, U. Hartel, H. Kuhlenbeck and H.-J. Freund, Phys. Stat. Sol. (a) 173 (1999) 93. 43. M.-C. Wu, C. M. Truong and D. W. Goodman, Phys. Rev. B 46 (1992) 12688. 44. K. Kimura, S. Katsumata, Y. Achiba, T. Yamazaki and S. Iwata, Handbook of Hel Photoelectron Spectra of Fundamental Organic Molecules; Halsted Press: New York, 1981. 45. A. Kolmakov and D. W. Goodman, Catal. Lett., in press. 46. L. N. Kantorovich, A. L. Shluger, P. V. Sushko and A. M. Stoneham, Surf Sci. 444 (2000)31. 47. M. Haruta, S. Tsubota, T. Kobayashi, H, Kageyama, M. J. Genet and B. Delmon, J. Catal. 144(1993)175. 48. S. D. Lin, M. Bollinger and M. A. Vannice, Catal. Lett. 17 (1993) 245. 49. F. Boccuzzi, A. Chiorino, S. Tsubota and M. Haruta, J. Phys. Chem. 100 (1996) 3625. 50. M. A. Bollinger and M. A. Vannice, Appl. Cat. B 8 (1996) 417. 51. N. W. Cant and N. J. Ossipoff, Catal. Today 36 (1997) 125. 52. K. Fukushima, G. H. Takaoka, J. Matsuo and I. Yamada, Jpn. J. Appl. Phys., Part I 36 (1997)813.

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53. G. R. Bamwenda, S. Tsubota, T. Nakamura and M. Haruta, Catal. Lett. 44 (1997) 83. 54. S. Minico, S. Scire, C. Crisafulli, A. M. Visco and S. Galvagno, Catal. Lett. 47 (1997) 273. 55. Z. M. Liu and M. A. Vannice, Catal. Lett. 43 (1997) 51. 56. M. Haruta, Catal. Today 36 (1997) 153. 57. Y. lizuka, H. Fujiki, N. Yamauchi, T. Chijiiwa, S. Arai, S. Tsubota and M. Haruta, Catal. Today 36 (1997) 115. 58. M. Okumura, S. Nakamura, S. Tsubota, T. Nakamura, M. Azuma and M. Haruta, Catal. Lett. 51 (1998) 53. 59. M. A. P. Dekkers, M. J. Lippits and B. E. Nieuwenhuys, Catal. Lett. 56 (1998) 195. 60. S. Tsubota, T. Nakamura, K. Tanaka and M. Haruta, Catal. Lett. 56 (1998) 131. 61. M. Valden, S. Pak, X. Lai and D. W. Goodman, Catal. Lett. 56 (1998) 7. 62. J.-D. Grunwaldt and A. Baiker, J. Phys. Chem. B 103 (1999) 1002. 63. J.-D. Grunwaldt, C. Kiener, C. Wogerbauer and A. Baiker, J. Catal. 181 (1999) 223. 64. E. D. Park and J. S. Lee, J. Catal. 186 (1999) 1. 65. H. Liu, A. I. Kozlov, A. P. Kozlova, T. Shido, K. Asakura and Y. Iwasawa, J. Catal. 185 (1999) 252. 66. T. Hayashi, K. Tanaka and M. Haruta, J. Catal. 178 (1998) 566. 67. M. Valden and D. W. Goodman, Isr. J. Chem. 38 (1998) 285. 68. M. Haruta, in: 3rd World Congress on Oxidation Catalysis; (Elsevier Science, Amsterdam, 1997) 123. 69. D. Martin, F. Creuzet, J. Jupille, Y. Borensztein and P. Gadenne, Surf. Sci. 377-379 (1997) 958. 70. D. Abriou, D. Gagnot, J. Jupille and F. Creuzet, Surf. Rev. Lett. 5 (1998) 387. 71. K. Luo, T. P. St. Clair, X. Lai and D. W. Goodman, J. Phys. Chem. B 104 (2000) 3050. 72. Q. Guo, W. S. Oh and D. W. Goodman, Surf Sci. 437 (1999) 49. 73. D. Novak, E. Garfunkel and T. Gustafsson, Phys. Rev. B 50 (1994) 5000. 74. S. Fischer, A. W. Munz, K.-D. Schierbaum and W. Gopel, Surf. Sci. 337 (1995) 17. 75. A. Szabo and T. Engel, Surf Sci. 329 (1995) 241. 76. P. W. Murray, N. G. Condon and G. Thornton, Phys. Rev. B 51 (1995) 10989. 77. C. Xu, X. Lai, G. W. Zajac and D. W. Goodman, Phys. Rev. B 56 (1997) 13464. 78. X. Lai, T. P. St. Clair, M. Valden and D. W. Goodman, Prog. Surf. Sci. 59 (1998) 25. 79. X. Lai, T. P. St. Clair and D. W. Goodman, Faraday Discuss. 114 (1999) 279. 80. C. C. Chusuei, X. Lai, K. Luo and D. W. Goodman, Top. Cat., in press. 81. M. Valden, X. Lai and D. W. Goodman, Science 281 (1998) 1647. 82. W. S. Oh, C. Xu, D. Y. Kim and D. W. Goodman, J. Vac. Sci. Technol. A 15 (1997) 1710. 83. H. Kobayashi and M. Yamaguchi, Surf Sci. 214 (1989) 466. 84. Z. Yang, R. Wu, D. W. Goodman and Q. Zhang, J. Chem. Phys., submitted. 85. W. Moritz and D. Wolf, Surf Sci. 88 (1979) L29. 86. M. Garofalo, E. Tosatti and F. Ercolessi, Surf Sci. 188 (1987) 321. 87. G. Binnig, H. Rohrer, C. Gerber and E. Weibel, Surf Sci. 131 (1983) L379. 88. C. Hofner and J. W. Rabalais, Surf Sci. 400 (1998) 189. 89. P. Hollins and J. Pritchard, Surf Sci. 89 (1979) 486. 90. D. P. Woodruff, B. E. Hayden, K. Prince and A. M. Bradshaw, Surf Sci. 123 (1982) 397. 91. P. Hollins, K. J. Davies and J. Pritchard, Surf Sci. 138 (1984) 75. 92. C. M. Truong, J. A. Rodriguez and D. W. Goodman, Surf Sci. 271 (1992) L385.

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93. G. Blyholder, J. Phys. Chem. 68 (1964) 2772. 94. R. Ryberg, Surf. Sci. 114 (1982) 627. 95. D. C. Meier, V. Bukhtiyarov and D. W. Goodman, submitted. 96. L. C. Feldman and J. W. Mayer, in: Fundamentals of Surface and Thin Film Analysis; (North-Holland, New York, 1986) Chapter 6. 97. B. Ealet, M. H. Elyakhloufi, E. Gillet and M. Ricci, Thin Solid Films 250 (1994) 92. 98. X. Lai, C. C. Chusuei, K. Luo, Q. Guo and D. W. Goodman, Chem. Phys. Lett. 330 (2000) 226. 99. Y. Wu, E. Garfunkel and T. E. Madey, J. Vac. Sci. Technol. A 14 (1996) 2554. 100. Y. Wu, E. Garfunkel and T. E. Madey, Surf. Sci. 365 (1996) 337. 101. P. J. Chen and D. W. Goodman, Surf Sci. Lett. 312 (1994) L767. 102. M.-C. Wu and D. W. Goodman, J. Phys. Chem. 98 (1994) 9874. 103. C. Xu, X. Lai and D. W. Goodman, Faraday Discuss. 105 (1996) 247. 104. D. R. Jennison, C. Verdozzi, P. A. Schultz and M. P. Sears, Phys. Rev. B 59 (1999) R15605. 105. C. Verdozzi, D. R. Jennison, P. A. Schultz and M. P. Sears, Phys. Rev. Lett. 84 (1999) 799. 106. J.-W. He, C. A. Estrada, J. S. Comeille, M.-C. Wu and D. W. Goodman, Surf Sci. 261 (1992) 164. 107. C. M. Truong, M.-C. Wu and D. W. Goodman, J. Chem. Phys. 97 (1992) 9447. 108. M.-C. Wu, C. M. Truong and D. W. Goodman, J. Phys. Chem. 97 (1993) 4182. 109. J. S. Comeille, J.-W. He and D. W. Goodman, Surf Sci. 306 (1994) 269. 110. C. Xu, W. S. Oh and D. W. Goodman, J. Phys. Chem. B, in press. 111. E. E. Platero, S. Coluccia and A. Zecchina, Surf Sci. 171 (1986) 465. 112. E. E. Platero, B. Fubini and A. Zecchina, Surf Sci. 179 (1987) 404. 113. G. A. Somorjai, Appl. Surf Sci. 121/122 (1997) 1. 114. R. A. Bennett, P. Stone and M. Bowker, Faraday Discuss. 114 (1999) 267. 115. A. Berko, G. Menesi and F. Solymosi, J. Phys. Chem. 100 (1996) 17732. 116. M. Li, Wilhelm, Hebenstreit, L. Gross, U. Diebold, M. A. Henderson, D. R. Jennison, P. A. Schultz and M. P. Sears, Surf Sci. 437 (1999) 173. 117. R. A. Bennett, P. Stone, N. J. Price and M. Bowker, Phys, Rev. Lett. 82 (1999) 3831. 118. J. A. Jensen, K. B. Rider, Y. chen, M. Salmeron and G. A. Somorjai, J. Vac. Sci. Technol. B 17 (1999) 1080. 119. Y. Yuan, K. Asakura, H. Wan, K. Tsai and Y. Iwasawa, Catal. Lett. 42 (1996) 15. 120. Y. Yuan, K. Asakura, H. Wan, K. Tsai and Y. Iwasawa, Chem. Lett. (1996) 755. 121. Y. Yuan, A. P. Kozlova, K. Asakura, H. Wan, K. Tsai and Y. Iwasawa, J. Catal. 170 (1997) 191. 122. Y. Yuan, K. Asakura, A. P. Kozlova, H. Wan, K. Tsai and Y. Iwasawa, Catal. Today 44(1998)333. 123. C. C. Chusuei, X. Lai, K. A. Davis, E. K. Bowers, J. P. Fackler and D. W. Goodman, submitted. 124. M. E. Labib and R. Williams, J. Colloid Interface Sci. 97 (1983) 356. 125. M. Kosmulski, J. Colloid Interface Sci. 135 (1990) 590. 126. F. A. Rodrigues, P. J. M. Monteiro and G. Sposito, J. Colloid Interface Sci. 211 (1999) 408.


409

Chapter 10

Principles of Reactivity from Studies of Organic Reactions on Model Oxide Surfaces A. B. Sherrill and M. A. Barteau Center for Catalytic Science and Technology, Department of Chemical Engineering University of Delaware, Newark, Delaware, 19716 USA

1. INTRODUCTION AND SCOPE The surface science of metal oxides has come of age over the past decade, stimulated by a variety of technological challenges. Examples include the needs to understand metal-metal oxide interfaces in semiconductor devices, sensors, and catalysts. Consequently, new techniques for surface analysis and modeling have been exploited to probe the complex surface-adsorbate interactions that govern chemical processes on oxide surface. Several recent reviews have explored relationships between the surface and bulk characteristics of metal oxide materials, and the ways in which these characteristics subsequently influence the surface chemical pathways [1-5]. Instead of simply adding a new edition to that list, two "case-studies" are presented, each highlighting the critical roles of surface anion-cation pairs and the redox properties of the surface in determining the product slate. The reactions of small organic molecules on titanium dioxide surfaces provide the central focus of these examples. Comparisons with reactions on single crystal surfaces of other oxides illustrate some of the key structural and electronic issues in oxide surface reactivity and the importance of local interactions between surface sites and adsorbed molecules. 2. TITANIUM DIOXIDE AS A MODEL FOR TRANSITION METAL OXIDES Transition metal oxides exhibit a number of properties that are conducive to catalytic applications, including thermal and mechanical stability needed to survive severe reaction conditions. More importantly, transition metal cations can typically exist in several different valence states. Titanium dioxide has a bulk band gap energy of about 3.2 eV, but electrons can be placed in (3d) gap states

410

about 0.7 eV below the Fermi level created by vacuum annealing. In contrast, for main group metal oxides such as MgO (bulk band gap ca. 7.7 eV), the d orbitals are not accessible [4]. Further, the lattice oxygen anions electronically isolate cations, thus surface-adsorbate bonding on oxides is highly localized. Consequently, transition metal oxides are often compared to transition metal complexes in solution. This so-called "cluster-surface analogy" has been used to advance molecular approaches to heterogeneous catalysis [6]. The surface science literature of titanium dioxide is replete with studies of the thermodynamic stability of different surface structures, as well as studies of the adsorption of metals, light gases, light and condensable organic molecules to the those surfaces. The full arsenal of surface analytical techniques has been applied to study Ti02 surfaces, including electron and vibrational spectroscopies, scanning probe microscopies, desorption spectroscopies, and computational models of varying complexity [1,3-5]. Experimental studies of single crystal Ti02 surfaces have utilized almost exclusively the rutile crystal structure, which is more stable than the anatase or brookite structures. Most of the work reviewed here was carried out on the (110) or (001) crystal planes, so a short introduction to the structures and coordination environments of those planes is presented first. The Ti02(l 10) crystal face (figure 1) is the most stable face of rutile titanium dioxide as well as the most thoroughly characterized. TiO2(110) has been imaged with low-energy electron diffraction (LEED), scanning tunneling microscopy (STM) and atomic force microscopy (AFM), and the electronic states have been characterized with X-ray photoelectron spectroscopy (XPS) and ultraviolet photoelectron spectroscopy (UPS). Polished TiO2(110) exhibits a (1x1) LEED pattern, although the surface can reconstruct into a (1 x 2) pattern during high temperature annealing in vacuum. The TiO2(110) surface is characterized by alternating parallel rows of bridging oxygen anions (and corresponding six-fold coordinate titanium cations beneath) and five-fold coordinate titanium cations [4,7-15]. The TiO2(001) surface (figure 2) is comparatively less stable, as all of the surface cations are four-fold coordinate in the ideal structure [4]; because of this, the TiO2(001) surface facets upon annealing to high temperatures. Upon annealing to 750 K, the surface facets into a {011} reconstruction with all of the cations in a five-fold coordinate arrangement. Further annealing to 900 K creates {114} faceting of the surface [4]. This structure ideally produces equal populations of four, five, and six-fold coordinate cations (thus retaining an average coordination number of five). A third reconstruction into a (7V2 x V2) R45° arrangement has recently been reported for (OOl)-oriented crystals annealed to 1473 K and rapidly quenched. The TiO2(001) surface has been imaged with AFM and STM, and has been characterized using Auger electron

411

Fig. 1. Structure of the unrelaxed Ti02(l 10) surface. Black and grey spheres are Ti and O atoms, respectively. H. Idriss and M.A. Barteau, Advances in Catalysis, v. 45 (2000). Reproduced by permission of Academic Press.

Fig. 2. Structure of the unreconstructed TiO2(001) surface. Black and grey spheres are Ti and O atoms, respectively. H. Idriss and M.A. Barteau, Advances in Catalysis, v. 45 (2000). Reproduced by permission of Academic Press.

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spectroscopy (AES), XPS, LEED, UPS, and near-edge X-ray absorption fine structure (NEXAFS) [4,16-22]. 3. CASE STUDY I: FORMIC ACID AS A PROBE OF SURFACE PROPERTIES Formic acid is a popular molecule for probing the catalytic properties of metal oxides [23-28]. The selectivity of its decomposition has frequently been used as a measure of the acid-base properties of oxides. This is a tempting generalization to make; oxides that produce dehydration products (H2O and CO) are described as acidic oxides, while their basic counterparts produce dehydrogenation products (H2 + CO2). It has been shown that in many cases the product selectivity is better connected to the surface redox behavior of the oxide [29]. Thus, more reducible surfaces produce higher yields of CO2. Consequently, particular attention has been paid in surface science studies to the interaction between adsorbed formate ions (the primary reaction intermediate) and surface metal cations, as well as to the participation of lattice oxygen anions in the surface reaction mechanism. 3.1. Formic Acid Decomposition on Metal Oxides As discussed in detail below, several reaction pathways are available to formic acid in addition to the "simple" dehydration and dehydrogenation reactions noted. Interactions between neighboring molecules and between the adsorbed molecules and the surface both influence the product selectivity. Formic acid may adsorb molecularly on oxide surfaces as shown in reaction (1). Formic acid also adsorbs by proton transfer to surface oxygen anions (reaction 2). The resulting intermediate is a surface formate. Details about the structure of this intermediate on different crystal planes will be presented below. The formate intermediate may decompose by dehydration, producing carbon monoxide, or by dehydrogenation, producing CO2. Gas phase molecules are labeled with "(g)" (e.g., HCOOH(g)) Adsorbed molecules are marked with asterisks (e.g., HCOOH*), surface lattice oxygen atoms are denoted by Os, and hydroxyl groups formed by protonation of them are designated with OsH. (1) (2) (3) (4)

HCOOH(g) ^ HCOOH* HCOOH(g) + O s ^ HCOO* + OsH HCOO* -^ CO(g) + Os + H* HCOO* -> C02(g) + H*

(dehydration, e.g. MgO [26]) (dehydrogenation, e.g. ZnO [25])

Formic acid decomposition has been examined on a number of metal oxides using a wide assortment of surface analytical techniques to probe the decomposition pathways. One might expect the most facile adsorption to occur

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on the most basic surfaces. Indeed, formic acid adsorbed on thin films of MgO (grown on Mg single crystals), and on stoichiometric and defective single crystals, easily dissociates into surface formate, even at 150 K and below [26,3033]. The formate intermediate dehydrates on MgO(lOO) surfaces exclusively at 520 K. The amount of formic acid adsorbed on MgO does not appear to be structure sensitive. The MgO(lll) and (100) surfaces adsorbed the same amount of formic acid as measured by XPS [30]. On both surfaces thermal decomposition produced exclusively CO, the nominal dehydration product, demonstrating that dehydration is not an indication of surface acidity [26,30,32]. Magnesium oxide is widely considered to be a strong base, but the formation of surface formates also occurs readily on bare surfaces of oxides that have been typically considered to be less basic (e.g. ZnO) or even acidic (e.g. Ti02). Such observations point out the limited utility of these labels. The zinc-terminated ZnO(OOOl) surface also readily dissociates formic acid to form surface formates; the subsequent decomposition of these yields both dehydration and dehydrogenation products at 575 K [25,34-39]. The dehydrogenation reaction results in the reduction of surface zinc cations as evidenced by desorption of Zn metal atoms coincident with CO2 desorption [39]. Pretreating the ZnO(0001)-Zn surface with chlorine atoms prior to formic acid exposure blocks Zn^"^ sites (each Zn cation on the surface has a single coordination vacancy, and roughly three zinc cations are blocked by each chlorine atom) and enhances the selectivity of the dehydrogenation reaction to form CO2. Coincidentally, chlorine was observed to promote the abstraction of lattice oxygen anions by the surface hydrogen produced by surface formate decomposition [37]. Earlier work on the ZnO(0001)-Zn surface demonstrated a correspondence between the extent of surface reduction produced by the adsorption and reaction of different molecules, e.g., formic acid, formaldehyde, and methanol, from which formates were synthesized. The greater the extent of surface reduction produced by the formate generation reaction, the greater was the selectivity to CO observed in formate decomposition [25]. The effect of chlorine adsorption on formate chemistry is consistent with this trend. Chlorine is electron withdrawing; consequently, chlorine adsorption causes the energy of the valence band maximum to shift away from the vacuum level and create a depletion layer. Thus, the chlorine adsorption may be thought of as increasing the extent of surface oxidation [4,37]. The change in surface electronic character and the corresponding increase in the formate dehydrogenation selectivity following chlorine adsorption are consistent with the selectivity trends for formate reactions on the clean ZnO (0001) surface: the more highly oxidized the surface, the higher the dehydrogenation selectivity. Formic acid decomposition has been studied on the (110), (001), and (100) surfaces of Ti02 [23-25,40-51]. The degree to which surface reducibility influences the reaction paths (e.g., dehydrogenation vs. dehydration; unimolecular reactions vs. bimolecular ones) will be explored in more detail

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below. Here it is sufficient to understand that the primary reaction of formic acid in temperature programmed desorption experiments on titanium dioxide single crystals is via dehydration at 550 K [43]. The TiOiCllO) surface facilitates both molecular and dissociative adsorption of formic acid. The dissociative adsorption of formic acid (to form surface formates) is mediated by surface oxygen anions; low-energy electron diffraction studies have shown that formate is ordered into ( 2 x 1 ) domains with formates bridging the surface titanium cations. This ordered layer is disrupted on heating; formate may recombine with surface hydroxyl groups to desorb formic acid (reverse of reaction 2) or it may decompose to form the dehydration products CO and H2O, as well as small amounts of CO2 and H2 [43]. Net dehydration reactions dominate the surface decomposition chemistry of formic acid on {011}-TiO2(001) faceted surfaces. Temperature programmed desorption experiments on these surfaces revealed that formic acid desorption resulted from both recombinative (390 K) and decomposition (560 K) processes. X-ray photoelectron spectroscopy measurements demonstrated that formic acid adsorbed on the {011}-TiO2(001) surface at 165 K both as formate and formic acid. This reconstruction of the TiO2(001) surface exposes titanium cations in a five-fold coordinate manner only, leaving one free coordination site per cation for interaction with adsorbates. Formic acid adsorbed on TiO2(110) at 180 K also formed a layer containing both molecularly and dissociatively adsorbed species; heating to 270 K left only adsorbed formate ions in a (2 x 1) pattern. This ordered arrangement vanished on heating to 350 K, coincident with the desorption of water and formic acid; further heating decomposed the remaining formate species, forming CO2 and CO as the principal carbon-containing products [43,52]. In both cases the reaction products are generated from exclusively unimolecular decomposition [45]. Surfaces with four-fold coordinate Ti"*"^ cations are capable of bimolecular coupling of surface formate to form formaldehyde; these sites can be created by more severe thermal faceting [1,3,23]. Reaction (5) illustrates this, where R may be a hydrogen atom or an alkyl group, and where Os denotes surface lattice oxygen. R

(5)

R

co,^ 0/ -Ti— / \

^

o R-c-R

+ CO.2(g)

This bimolecular coupling is a structure-sensitive reaction, and it illustrates a key characteristic of metal oxides: multiple coordinative unsaturation of surface metal cations may facilitate coupling of ligands in a manner similar to that for unsaturated metal complexes in solution. Examples of other coupling reactions

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observed on TiO2(001) surfaces include alkyne cyclotrimerization (on reduced surfaces) and alkoxide etherification reactions (on oxidized surfaces). Alkynes are readily cyclized into aromatics on reduced surfaces of TiO2(001) exposing Ti^^ cations on the surface; this chemistry is analogous to catalysis of alkyne cyclization by low-valent mononuclear transition metal complexes in solution. In both cases, the metal center must be able to complete a two-electron oxidation and reduction cycle in order to form a metallacycle intermediate and eliminate the aromatic product [53,54]. A reaction more closely related to the bimolecular ketonization reaction described above is alcohol etherification via coupling of surface alkoxides. It was observed that {114}-faceted TiO2(001) surfaces containing adsorbed methoxide species produced dimethyl ether, while the {011}-faceted TiO2(001) surface produced no detectable ether products [55]. Titania surfaces with oxygen vacancies reduce oxygen-containing intermediates such as formates and methoxides in the process of the filling the oxygen vacancy [41]. Vacancy healing is also important in the reductive coupling of aldehydes and ketones to form symmetric olefins on reduced TiO2(001) surfaces. This reaction is analogous to the well known McMurry reaction, wherein slurries of low-valent titanium carry out the coupling process [20]. Thus, not only do coordinatively unsaturated metal cations affect the local electronic structure of the oxide surface, but they also serve as active sites for a variety of bimolecular coupling reactions with potential synthetic utility. The formation of formaldehyde from formic acid is not expected to occur on stoichiometric TiO2(110) because the surface cations are exclusively fivecoordinate, and formaldehyde formation in the absence of reduced surface sites has been shown to be dependent on the presence of low-coordinate metal cations on titanium dioxide surfaces. Iwasawa et al. noted the absence of formaldehyde from the product slate of formic acid reaction products in their studies on the TiO2(110) surface [43]. 3.2. The Participation of Oxygen Vacancies in Formic Acid Decomposition The extent of lattice oxygen incorporation into the decomposition products is linked to the surface redox properties of the oxide. Recent SSIMS and FTRAIRS studies have provided initial evidence for the transient formation of oxygen vacancies on the surface of TiO2(110). Henderson has proposed that surface oxygen vacancies may explain the formation of trace amounts of formaldehyde from formic acid on TiO2(110) (figure 3) [41]. Both reactions (6) and (7) are proposed to occur below 500 K; the water produced from formic acid exposure desorbed at 475 K, before the onset of formate decomposition to form CO. Formaldehyde and CO2 were produced in minor quantities relative to the production of CO [41]. Formaldehyde was formed from formic acid on reduced

416

Oxygen vacancy formation on TiO 2(110) during bridging hydroxyl combination water desorption

bridging ^'''''P^

\ > J % L

A v X

^ - ^ vacancy newly formed five-coordinate Tl'^+ cation sites (one cation hidden from view)

Fig. 3. Schematic model for the creation of oxygen vacancies on Ti02 (110) by combination of bridging hydroxyl groups to form water. M.A. Henderson, Journal of Physical Chemistry B, V. 101 (1997). Reproduced by permission of the American Chemical Society.

surfaces of TiO2(001), and the amount of oxygen deposited on the surface was consistent with that required to convert the available Ti^"^ sites to Ti"^^ [23]. Consequently, the low yield of formaldehyde and CO2 likely represents the small surface population of oxygen vacancies on the crystal surface in the work by Henderson [41]. Experiments with ^^0-enriched surfaces did demonstrate incorporation of lattice oxygen into the CO product of formic acid decomposition, illustrating the relative ease of lattice oxygen exchange and abstraction [40,41]. (6) (7) (8)

HCOOH* + Os ^ HCOO* + OsH 2 OsH -> H20(g) + Os + vacancy 2 HCOO* + vac. -> Os + H2C0(g) + C02(g)

The formation of formaldehyde at reduced titanium surface cation sites can also occur, but via formate reduction rather than via bimolecular formate coupling. (9)

HCOO* + H* -> H2CO + Os

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Idriss et al noted that the decomposition temperature of surface formate is nearly constant across a wide range of different TiO2(001) faceted and reduced surfaces, suggesting that this reaction occurs via a conmion intermediate, whether the products are those of dehydration or reduction. These workers proposed scission of the formate C-H bond as the rate-determining step [23]. Catalytic experiments on single crystals of TiO2(110) suggested that transportlimited reaction conditions favor unimolecular dehydration, and that lower surface temperatures favor bimolecular dehydrogenation [43]. Thus, it is important to realize that interactions with neighboring molecules may influence the surface chemistry as much as surface-adsorbate interactions. The case for oxygen vacancies is supported by recent work with Fourier Transform Reflection-Absorption Infrared Spectroscopy (FT-RAIRS) by Hayden and coworkers that indicated the presence of two types of surface formate on the TiO2(110) surface [49]. Adsorption of formic acid at 300 K yielded a surface coverage of about 0.6 ML (from XPS) [50,51]. One type of surface formate (accounting for about 2/3 of the surface coverage) has the OCO plane aligned in the direction. The less populous type of surface formate is oriented orthogonally to the first, in the direction (figure 4). FT-RAIRS indicated that this second variety of formate interacts with oxygen vacancies on the surface created during the formation of formate along the direction in the manner

6.49 Ang. 0^.oxygan(toprow)^ I formate

titanium

hydrogen

Fig. 4. Schematic showing the possible adsorption sites for formate species on TiOa (110). B.E. Hayden, A. King, and M.A. Newton, Journal of Physical Chemistry v. 103 (1999). Reproduced with permission of the American Chemical Society.

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of reactions (6) and (7) above. The population of the second variety of formate was increased when formic acid was adsorbed at 450 K, where vacancy creation by hydroxyl recombination would be more energetically favorable [49]. Impingement on the surface of low-density electron beams can also be used to form Ti^"^ sites on the crystal surface. Recent work has shown that additional surface capacity for bridging formate created by these new surface sites corresponds to the amount of additional formic acid uptake by the surface [50,51]. The adsorbed formate did not heal the oxygen vacancy until the formate decomposed, lending additional credence to a reaction step other than defect healing controlling the formation of formaldehyde at Ti^"^ sites. Sputter reduction of the TiO2(110) surface greatly increased the surface capacity (to ca. 0.92 ML), and subsequent exposure of a formate-saturated, sputtered surface to O2 produced almost no defect healing [49]. The formation of formaldehyde and the attendant formation of CO2 reveals a likely interaction with oxygen vacancies, and when contrasted with decomposition chemistry on the MgO(lOO) surface, highlights the importance of surface redox properties in carboxylic acid chemistry on metal oxide surfaces. 3.3. Tracking Formic Acid Decomposition with Scanning Probe Microscopy Appealing as the traditional probes for formic acid chemistry may be, recent work with temperature programmed scanning tunneling microscopy (STM) and non-contact atomic force microscopy (NCAFM) has generated new interest in oxide surface science by dynamically illustrating surface diffusion and reaction of formates [8,24,56-60]. The strong interaction of formate with the metal oxide surface helps make the surface reaction visible to STM at elevated temperatures; in contrast, surface diffusion of adsorbates on metal surfaces above cryogenic temperatures is often fast on the time scale of the microscope, making identification of important reaction intermediates difficult [61]. Iwasawa and coworkers have published several accounts of STM and NCAFM investigations of formic acid decomposition on the stoichiometric surface of TiO2(110), and have imaged surface titanium cations, oxygen anions, and oxygen vacancies [8,24,52,56-59,61]. Formate ions on the surface were observed in ordered (2x1) domains with STM, in agreement with earlier LEED and FT-RAIRS work. Surface diffusion and reaction pathways were probed with STM as well. The migration of formate ions was observed by using the microscope tip to clean off a portion of the scanned area and monitoring the back-filling of the bare patch from the surrounding domains. The bare area was initially roughly square; subsequent images indicated that surface migration was anisotropic, favoring transport along the rows of titanium cations between the oxygen ridges. When viewed under catalytic conditions, Iwasawa et al.

419 Step

/

(001]

Ti203 rowj Upper terrace

Step Lower terrace

Fig. 5. A post-reaction STM image of the Ti02 (110) surface heated to 450 K and exposed to 1.3 X 10'^ Pa HCOOD gas for 10 min. One product particle is marked by the white circle. One ( 2 x 1 ) unit of the formate monolayer is shown with the solid rectangle. Y. Iwasawa, H. Onishi, K. Fukui, S. Suzuki, and T. Sasaki, Faraday Discussions, v. 114 (1999). Reproduced by permission of the Royal Society of Chemistry.

observed the formation of bright particles across the surface in addition to the (2x1) array of formate (figure 5). The bright features were located atop oxygen anion positions and were about the same size as formate ions. Based on the size of the bright particles and the preference for dehydrogenation on surfaces with high formate coverage, surface carbonate ions were proposed. Such species may be formed as intermediates in the production of CO2, e,g. via reactions such as the bimolecular ketonization of acetate [46,47]. Thus, surface formate and hydroxyl species may not be the only participants in formic acid decomposition [8,24,56-59]. Now consider the fate of formic acid adsorbed on two (1x2) reconstructions of TiO2(110). When annealed in vacuum at temperatures near 900 K, doublerow strands of Ti203 were formed on the surface of the crystal. Formic acid dosed onto these surfaces did not make formate [8,25]. The inability of the acid to undergo dissociation is reminiscent of the oxygen-polar ZnO(OOoi) surface; on that surface, the overlayer of oxygen anions obscured the zinc cations. Thus, by extension, there must not be Ti cations accessible to the adsorbed formic acid. Recent work by Bowker on TiO2(110) surfaces annealed to higher temperatures (ca. 1200 K) has revealed the cross-linked structures on the surface [62]. The surface included added rows of Ti02 running in the direction.

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Temperature programmed STM experiments under oxygen led to new growth of (1 X 1)-Ti02 on the crystal surface. Formic acid dissociated on the cross-linked surface to make formate ion. As the temperature ramp progressed, the population of formate diminished while small islands of (1 x 1) termination formed on the terraces of the crystal. Bowker and coworkers proposed that the decomposition of formate ions supplies oxygen atoms to interact with interstitial titanium cations and generate the Ti02 islands [60]. The nucleation of new (1x1) titanium dioxide is not based on the healing of oxygen vacancies, but it does highlight the role that the surface redox environment plays in formate decomposition. The change in reactivity of formate on (1 x 2) reconstructed TiO2(110) shows that although the coordination environment surrounding the cation occupies a central role in determining product selectivity, the surface electronic structure should not be ignored. 3.4. Catalytic Reactions of Formic Acid on Titanium Dioxide (110) Iwasawa and coworkers conducted molecular beam experiments to investigate the catalytic decomposition of formic acid on single crystal TiO2(110) surfaces [43]. These experiments revealed competing reaction channels (figure 6) for formate dehydration (rO and dehydrogenation (r2) on the single crystal surface under catalytic conditions that could not be accounted for by the chemistry observed in TPD experiments alone. At surface temperatures below 650 K, the predominant catalytic pathway was concluded to be dehydrogenation of formic acid through a bimolecular reaction between an adsorbed formate and a molecule of formic acid. An activation-energy of 15 ± 10 kJ/mol and a reaction order of 0.9 ±0.1 with respect to the gas phase formic acid pressure (at 650 K) were estimated from catalytic rate studies. The surface coverage of adsorbed formate at saturation was estimated at 0.5 ML based on C (Is) and O (Is) XPS measurements; this value was consistent with the (2x1) LEED pattern observed when formic acid was adsorbed on the clean surface at 300 K. The adsorbed formate ions were assumed to align themselves between top-layer oxygen anion rows in a bidentate fashion. Thus, each bridging formate is coordinated to two five-fold coordinate Ti"^^ cations. UHV TPD data indicated that surface formate decomposes at 570 K. The first-order rate dependence mentioned above was also observed at 600 K, suggesting that formic acid must interact with formate ions on the near-saturated surface [43]. A mechanism for catalytic dehydrogenation was proposed to explain these observations. The researchers used deuterium-labeled formic acid to discriminate against background water and hydrogen. In the scheme below, Os represents surface lattice oxygen, while oxygen originating from formic acid is denoted by Of*. Adsorbed species are labeled with (*).

421

H0.5

0

Fig. 6. Reaction rates simulated as a function of reaction temperature with a collision frequency of 0.5 s'^ site'\ ri: dehydration; T2: dehydrogenation; 0: the coverage of formates. H. Onishi, T. Aruga and Y. Iwasawa, Journal of Catalysis, v. 146 (1994). Reproduced by permission of Academic Press.

(10)

(11) (12)

-OsD DCOOD(g) + Os ^ D C O O * + ( D C O O * -> CO(g) + OfD D C O O * + DCOOD(g) -> C02(g) + D2(g) + D C O O *

The break-point temperature in dehydration (above which the rate was temperature insensitive) matched the maximum temperature for dehydrogenation, suggesting that a conmion intermediate exists for each reaction, and that the product selectivity is determined by interactions with other molecules and the surface. Above 650 K, the catalytic dehydration channel dominates, but the rate-determining step changes above 700 K. Below 700 K, the reaction rate is nearly independent of the partial pressure of formic acid (ca. 0.2 order). Above 700 K, the rate of the reaction is essentially independent of temperature, implying that reaction is limited by formic acid adsorption and dissociation; thus, above 700 K, the rate becomes first-order with respect to the partial pressure of formic acid. Higher pressures of formic acid over the crystal surface should therefore increase the transition temperature - this behavior was observed by Iwasawa and coworkers, and the turnover frequency for catalytic dehydration approached the collision frequency of formic acid at high

422

temperatures. Consequently, a mechanism was proposed to account for the catalytic dehydration of formic acid on the titania surface. Reaction (13) is the decomposition reaction of formate ions on titania, and is rate-determining below 700 K. Reaction (14) is a proton transfer reaction that forms water and formate ions, and reaction (15) highlights the scrambling reaction possible between lattice oxygen anions and formate-derived oxygen atoms [43]. (13) DCOO*-^ CO(g) + OfD (14) DCOOD(g) + OfD -> D20(g) + DCOO* (15) OfD + Os-^ OsD + Of* It should not be surprising that, except in those cases where the coordination environment around the cation is suitable for bimolecular coupling reactions, UHV decomposition reactions tend toward unimolecular processes. The bimolecular dehydrogenation process that occurs during steady-state reaction on the single crystal TiO2(001) surface should serve as a reminder that steady-state processes involve interaction between the surface, surface species, and vapor phase components, and that it is not appropriate to attribute product selectivity exclusively to oxide characteristics. 3.5. Extension to Higher Carboxylic Acids The decomposition of formate on titanium dioxide is a model for the decomposition of higher carboxylic acids. Adsorption and reaction of acetic acid on TiO2(110) demonstrated behavior similar to formic acid [63,64]. Experiments with LEED showed that acetate ion on the (110) surface formed ( 2 x 1 ) overlayers that vanished on heating to 450 K. The return of the (1 x 1) LEED pattern agreed with the disappearance of long-range order as acetic acid and water desorbed from the surface. Williams et al. argued, based on LEED data, for a shortening of the Ti-Ti bond distance because of the formation of bridging acetate [63]. Heating to 450 K did not change the angular distribution of H"^ ions desorbing from the surface during ESDIAD experiments, which implied that acetate ions retained a bridging configuration [63,64]. Non-contact AFM of (2 X 1) acetate overlayers on TiO2(110)-(l x 1) has illustrated electronic interactions of acetate species with neighboring acetate ions and hydroxyl groups [65]. Scanning tunneling microscopy has also been used to study acetate orientation on the (110) surface. Temperature programmed STM demonstrated the formation of (2 x 1) ordered domains of acetate and subsequent disappearance of acetate from the surface on heating to 540 K. From this, Iwasawa and coworkers estimated a first-order rate constant of about 4 x 10"^ s'^ at 540 K, in reasonable agreement with prior TPD experiments [61]. ESDIAD, STM and LEED investigations of the orientation of benzoic acid on the (110) surface again focused on the formation of the adsorbed carboxylate

423

ion. ESDIAD confirmed that the molecule stood upright on the surface of the crystal, agreeing with earlier studies with formate and acetate. Benzoate also formed a (2 x 1) LEED pattern at saturation. STM of benzoate on the (110) surface detected (2 x 2) patterns; this discrepancy was attributed to intermolecular interactions between phenyl rings of adjacent intermediates forming dimers on the surface [64]. Thermal reaction of larger carboxylic acids on titanium dioxide surfaces traced reaction pathways similar to those of formic acid. The {110}-faceted surface of TiOiCOOl) selectively decomposed acetate and propionate via unimolecular dehydration to form ketene and methyl ketene. The dehydration yields from both reactants were quite similar, as were the desorption temperatures of the products [46]. Acetic acid decomposed on the {114}-faceted of the TiOa (001) surface to produce ketene as well as acetone [44]. The acetone generated arose from bimolecular coupling of pairs of surface acetates at four-fold coordinate cations; this is analogous to the production of formaldehyde from surface formate on identically prepared surfaces. The reaction of propionic acid corresponded directly to the reaction of acetic acid, producing methyl ketene and 3-pentanone [46]. Adding bifunctional character to the molecule adds another dimension to the overall reaction on the surface. For example, acrylic acid (which contains both carboxyl and vinyl groups) adsorbed on {110}-faceted TiO2(001) to form surface acrylates. The acrylate ion not only decomposed via unimolecular dehydration to form ethylene and CO, but was also reduced to form acrolein. Coupling of acrylates to form divinylketone was not observed on this surface, but was observed on the {114}-faceted surface, where divinylketone was formed at the expense of acrolein formation. Two mechanisms are at work here. On annealing the crystal to higher temperatures, the surface becomes slightly more oxidized, decreasing the acrylate reduction. Simultaneously, surface restructuring to form {114}-TiO2(001) exposes four-fold coordinate titanium cations, providing sites for the bimolecular coupling reaction [44]. Reactions of maleic anhydride on {Oil }-faceted TiO2(001) produced ketene, ethylene, acetylene, vinylacetylene, benzene, butene, CO and CO2 [66]. The reaction was sensitive to the population of surface defects; increased surface reduction (via sputtering) produced a corresponding increase in production of hydrocarbon coupling products (benzene, vinylacetylene, and butene). Reaction of maleic anydride was proposed to occur at five-coordinate Ti"^"*" cations via a ring opening process leading to a six-membered metallacycle incorporating a lattice oxygen anion. This intermediate was proposed to form acetylene by decarboxylation; subsequent cyclotrimerization of acetylene produced benzene [66].

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The most complicated carboxylic acid examined to date is likely biisonicotinic acid (2,2'-bipyridine-4,4'-dicarboxylic acid), adsorbed on the TiOiCllO) surface by Patthey et al. [67] Their investigation with NEXAFS experiments and INDO calculations revealed that both carboxylate groups were bound to the surface in a bridging fashion [67]. DeSegovia and coworkers examined glycine (comprising both carboxyl and amino substituents) on the (1 X l)-TiO2(110) surface with photoelectron spectroscopy (PES). Their experiments suggested that multilayers of glycine on the (110) surface have a zwitterionic form [68]. Glycine adsorbed on the (1 x 2) reconstructed surface at submonolayer coverage predominantly formed adsorbed carboxylates and quenched Ti(3d) features in the PES spectra [69]. The reaction chemistry of carboxylates depends on many factors, including the surface coordination environment, as demonstrated by the bimolecular coupling of formate ions to form formaldehyde or acetate ions to form acetone on the {114}-reconstruction of TiO2(001). The oxidation state of the cation center is critical as well, as the reduced-surface chemistry of formic acid (forming formaldehyde) attests. Steady-state catalytic experiments with formic acid demonstrated that gas phase interactions with adsorbates lead to an unexpected dehydrogenation pathway, and that it is important to consider vapor phase interactions. In UHV experiments, the product selectivity from formic acid decomposition across a wide range of oxides is most often attributed to the ease with which the oxide surface is reduced. Considering the important role of oxygen vacancies in the decomposition process, it is unsurprising that surface redox properties affect reaction selectivity. Consequently, the reactions of carboxylic acids (even multifunctional ones) on metal oxide catalysts are dominated by carboxylate-surface interaction. 4. CASE STUDY II: THERMAL AND PHOTOREACTIONS OF METHANOL ON TiOi Like formic acid, methanol decomposition has also been used to probe the acidbase properties of metal oxides [70]. However, methoxide decomposition is dependent on surface structure in much the same way as formate decomposition. For example, methanol undergoes parallel dehydration and dehydrogenation reactions on the same crystal surface of zinc oxide [25]. Once again, product selectivity ratios may not necessarily serve as a diagnostic of acid-base properties alone. Alcohol decomposition does provide additional insight into the interaction of adsorbates with metal oxide surfaces. The reactions of alkoxides on titanium dioxide have been used to probe thermal and photo-reactivity of powder and single crystal samples.

425

4.1. Thermal Reactions of Methanol on TiOi Single Crystals Methanol decomposes on titanium dioxide surfaces by mechanisms that are similar to those by which formic acid decomposes. Methanol can reversibly adsorb on single crystal surfaces of titania (reaction 16) in a molecular state, or it may dissociatively adsorb by interaction with surface lattice oxygen anions, forming a surface methoxide (reaction 17). Reaction (18) represents the disproportionation reaction of hydroxyl groups on the surface of the metal oxide. (16) (17) (18)

CHsOHcg) ^ CH3OH* CH30H(g) + Os-> CH3O* + OsH OsH + OsH -^ H20(g) + Os

It is important to note that the production of a molecule of water via reactions (17) and (18) creates an oxygen vacancy on the surface. This vacancy can be filled by decomposition of a methoxide leaving a methyl group bound to the surface. This methyl group can decompose to leave carbon and hydrogen on the surface, or it can incorporate hydrogen (from the decomposition of other methyl groups or from the dissociation of methanol to form methoxide) from the surface to produce methane. Methoxide may also combine with surface hydrogen to form methanol. This step is similar to the production of formic acid at higher temperatures from adsorbed formate. (19) (20) (21) (22)

CH3O* ~> CH3* + Os CH3* -> C* + 3H* CH3* + H * - > CH4(g) CH3O* -^ CH20(g) + H*

Methanol desorbs from the (110) surface in low and high temperature states, as recently verified in separate studies by Henderson et al. [71,72] and Vohs et al [73] The low temperature desorption state has been assigned to desorption of molecularly adsorbed methanol. The higher temperature states have been assigned as recombinative desorption of methanol with surface hydroxyl groups based on SSIMS and HREELS identifications of surface species [52,71-73]. The TiO2(110) surface can be modified to alter the surface chemistry of adsorbed methoxides. Recently Vohs et al. [73] have reported temperature programmed desorption studies of methanol on TiO2(110)-supported V2O5. Deposited monolayer films of vanadia converted some of the adsorbed methanol to formaldehyde and water, while multilayer films of vanadia on the TiO2(110) surface were found to be inactive for methanol oxidation. Furthermore, adsorption studies of formaldehyde indicated that formaldehyde production from

426

methanol is reaction-limited and that there is little further decomposition of formaldehyde to CO or CO2. Consequently, the authors proposed that methoxide is the key intermediate in the oxidation of formaldehyde on V2O5 monolayers [73], although spectroscopic confirmation is lacking to date. Henderson et al. [71,72] have recently used SSIMS and HREELS to study methanol adsorbed on hydroxylated, clean, and oxygen-precovered vacuumannealed surfaces of TiOiCllO). The vacuum-annealed surface contains oxygen vacancies at a surface concentration of between five and ten percent. Coveragedependent TPD of methanol on these defective surfaces demonstrated the presence of both molecularly and dissociatively adsorbed methanol molecules on the surface of the oxide; these give rise to methanol desorption between 165 and 480 K. HREELS experiments suggested that both molecularly and dissociatively adsorbed methanol molecules were present on adsorption at 135 K. SSIMS experiments showed that the low-temperature desorption event at 290 K involved molecularly adsorbed species, while the high temperature feature at 480 K involved recombination of surface methoxide and hydroxyl groups. The absence of surface hydrogen atoms may explain why methane is not detected among the products desorbing from the (110) surface; the methoxides present on the surface may be isolated, with other surface species unavailable for reaction. Methanol adsorbed on an ^^0-enriched surface of TiOiCHO) did not produce CHs^ÔH; thus it was proposed that methoxy groups bound at vacancy sites were immobile [71,72]. Methanol co-adsorbed with water displaced most of the water on the surface; methanol co-adsorbed with oxygen formed surface methoxides stable to 625 K. Oxygen pretreatment of the surface did lead to the formation of a species assigned as formaldehyde, which Henderson proposed to be formed via a disproportionation reaction between two methoxides. Spectroscopic probes of the controlling intermediate were inconclusive [71,72]. Methanol decomposition has also been studied on sputtered, {Oil}-, and {114}-faceted crystal faces of TiO2(001). Chemisorbed methanol formed surface methoxide species; about half of the methoxides formed recombined with surface hydroxyl species to form methanol at 365 K. The remainder of the methoxide species decomposed at higher temperatures to form some combination of methanol, methane, dimethyl ether, formaldehyde, and CO depending on the preparation of the crystal surface [74]. The sputtered (001) surface produced methanol as the primary hightemperature desorption product - more than half of the methoxy species present on the surface at 400 K recombined to form methanol at 580 K. The balance of surface methoxides decomposed into methane and CO in a 2:1 ratio between 590 and 600 K. X-ray photoelectron spectroscopy measurements for methanol adsorbed on the sputtered (001) surface at 300 K demonstrated that two types of carbon-containing species were present on the surface; subsequent flashing to

427

450 K desorbed molecularly bound methanol and shifted the binding energy to 286.8 eV, characteristic of surface methoxy groups [74]. The {011}-faceted (001) surface generated methanol, methane, and formaldehyde between 610 and 640 K. Again, methanol accounted for about half of the high temperature products, with the balance of products identified as formaldehyde and methane in a roughly 1:5 ratio. Deuterium labeling of the hydroxyl group illustrated the formation of surface methoxy groups; CH3OD desorbed at 370 K, while CH3OH desorbed at 620 K. Further confirmation of surface methoxide was provided by XPS. Heating the methanol-dosed surface to 300 K desorbed the molecular state, leaving only one type of carbon-containing intermediate (FWHM = 1.9 eV) at 286.8 eV, as observed on sputtered (001) surfaces. [74] The identification of surface methoxide as a precursor for decomposition products is in agreement with methanol adsorption experiments on ZnO[25], Ti02 (110) [71,72], Zr02 [75] and UO2 [76]. Formaldehyde production was associated with a net dehydrogenation of surface methoxides [74]. Methanol decomposition on the {114}-faceted surface of TiO2(001) produced methanol, methane, and formaldehyde, as well as a product not demonstrated on the other crystal faces, dimethyl ether. Methanol desorption on the {011}-faceted and sputtered crystal surfaces typically accounted for about half of the high-temperature products, and the {114}-faceted surface gave rise to a similar result. Formaldehyde and methane were produced in a 1:1 ratio on the {114}-faceted surface; while formaldehyde production increased slightly on the more oxidized surface, production of methane dropped to less than half of the production on the {011}-faceted surface. Generation of dimethyl ether, the product of a bimolecular reaction, occurred at 400 K and accounted for about eight percent of the total carbon yield. No CO was produced on the {114}faceted surface. X-ray photoelectron spectroscopy experiments revealed that the total methoxide coverage was about sixty-five percent of the methoxide coverage on the {011}-faceted surface. This corresponds to the population decrease of coordinatively unsaturated cations on the surface. One-third of the five-fold coordinate titanium cations on the {011} reconstructed surface are converted to six-fold coordinate sites after reconstruction to the {114}-faceted surface (another third are converted into four-fold coordinate sites, preserving the fivefold surface average). It is likely that dissociative adsorption depends on coordinatively unsaturated titanium cations. In fact, the relative coverage change matches the relative population change in available unsaturated centers (down by about a third in both cases), demonstrating the importance of cation coordination vacancies in surface reactions on oxide surfaces. Water desorption occurs concurrently with methanol desorption on the oxidized surfaces, albeit in low yields (ca. 2-5%). Production of water requires dissociation of two methanol molecules; consequently, there is a small net reduction in the population of

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surface oxygen anions that ultimately affects subsequent thermal reactions of methoxide. Methane selectivity increased as the sputtered surface was oxidized to form the {Oil}-faceted surface, and then decreased upon further oxidation to the {114}-faceted surface [74]. The ideal {011}-faceted surface contains exclusively five-coordinate titanium cations, so a bimolecular disproportionation reaction (see reaction (23)) to produce the ether product is precluded. Although the sputtered surface contains low-coordinate titanium cations and ample oxygen vacancies, the formation of dimethyl ether is exclusive to the {114}-faceted surface. In fact, the defect density drives all of the methoxides to fill those oxygen vacancies, and the resulting carbon is deposited on the surface, only to be burned off as CO, as illustrated in reaction (24) [74]. (23) (24)

2 CH3O* -> HsCOCHscg) + Os C* + Os ^ CO(g)

4.2. Thermal Reactions of Methanol on TiOi Powders As evidenced by the chemistry on the (001) surface, the activity and selectivity of methanol decomposition is dependent on defect and coordination vacancy populations; thus, application of single crystal studies to polycrystalline powders requires some caution. In fact, single crystal surface chemistry has been scaled up to powders before. For example, aldol condensation of acetaldehyde to crotonaldehyde on the {114}-faceted Ti02 (001) crystal face, while not a structure sensitive reaction, has been shown to be similar to acetaldehyde aldolization on rutile powders [77]. The presence of four-, five-, and six-coordinate titanium cations on the {114}-faceted surface has been suggested to provide a good model for polycrystalline materials; this conclusion is based on experimental studies of a variety of small oxygen-containing molecules that illustrate the utility of carefully selected model surfaces for understanding more complex, high surface area materials [55,78]. Methanol decomposition has been studied on the anatase and rutile polymorphs of titanium dioxide [55,78]. Methanol uptake experiments on anatase TiOi demonstrated a molecular coverage of about 3.2 molecules-nm"^ [78]. Both reversibly adsorbed methanol and irreversibly adsorbed methoxides were present on the surface at low temperatures. Infrared spectroscopy showed that surface methoxides were responsible for desorption of the parent alcohol and the decomposition products. Surface methoxy groups decomposed into both dehydration and dehydrogenation products at 650 K. Temperature programmed desorption experiments demonstrated a net weight loss for the powdered sample at high temperatures; the starting weight could be recovered by oxidizing the sample at 400 K. The weight loss was directly related to the product distribution since oxidative dehydrogenation to form formaldehyde and water involved the

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abstraction of lattice oxygen. Major reaction products at 650 K included formaldehyde, dimethyl ether, methane, CO, and H2O. Small amounts of H2 and methanol also desorbed at 650 K [78]. Methanol reacted on rutile Ti02 to form a similar product distribution. Methanol uptake on rutile was about 1.7 times higher per unit area than on anatase Ti02. Earlier work has suggested that the estimated ratio of exposed titanium cations on rutile and anatase polycrystalline surfaces is about 1.5, in good agreement with the methanol uptake ratio [55]. The selectivity to formaldehyde was approximately the same on both anatase and rutile, although the yield was higher on rutile, reflecting greater methanol uptake. The selectivity ratio for dehydration products (methane and dimethyl ether) was also about 1:1 [55]. 4.3. Thermal Reactions of Methanol on SnOi - Another Rutile Oxide Because titania and tin oxide have the same (rutile) bulk structure, comparison of their surface chemistry permits other influences on reactivity to be examined. One important difference is that the metal-oxygen bond energy is lower for Sn02 than Ti02. Titania is often compared to Sn02, although differences in the electronic structure of Sn02 result in weaker bonding of lattice oxygen anions to the surface. Consequently, it is considerably easier to remove bridging oxygen anions from the (110) surface of Sn02 by annealing or sputtering than from Ti02. Thus, one might expect that surface methoxides would form more oxidized products on tin oxide than on titania [4]. Consistent with the connection between the ease of reduction and product selectivity of surface oxygenates, recent work on SnO2(110) [79] has demonstrated that methanol decomposition produced formaldehyde exclusively. The formation of methoxide was determined to be dependent on the availability of bridging oxygen vacancies on the surface; stoichiometric surfaces converted only about ten percent of adsorbed methanol. It was demonstrated that it was also possible to over-reduce the surface as well. Surfaces characterized by fourcoordinate Sn^"^ cations created by vacuum annealing at 700 K converted as much as half of the adsorbed methanol, while ion-bombarded surfaces annealed to 950 K converted only five percent of methanol to formaldehyde. This disparity in conversion was attributed to in-plane oxygen vacancies created by the more severe annealing (figure 7). These in-plane defects were responsible for a new, higher temperature desorption state for formaldehyde that did not exist for surfaces containing only bridging oxygen vacancy sites. The presence of this additional high-temperature state implies greater stabilization of methoxide on the surfaces annealed to high temperature, but the decreased yield of formaldehyde from these surfaces illustrates the difficulty of dissociatively adsorbing methanol at in-plane sites [79].

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(a) [110]

(b)

^Bridging vacancy

(c)

•in-plane vacancy

Fig. 7. Ball model illustrations of: (a) the ideal, stoichiometric surface of SnOi (110), (b) the surface missing all bridging oxygen anions and (c) the surface with in-plane oxygen vacancies. The small solid circles represent Sn cations while the large open circles represent O ^" anions. V.A. Gercher, D.F. Cox, and J.-M. Themlin, Surface Science, v. 306 (1994). Reproduced by permission of Elsevier Publications.

Methanol decomposition on {114}-faceted TiO2(001) was accompanied by the formation of dimethyl ether attributed to the presence of four-coordinate Ti"^"^ surface cations, yet no ether was formed on the vacuum annealed SnO2(110) surface, which also contains four-coordinate metal cations. However, the tin cations present on the surface are present as Sn^"*^; SnO has a bulk fourcoordinate structure, hence there is only one coordination vacancy, again precluding bimolecular coupling reactions [79]. 4.4. CH3OH Decomposition on Fluorite Metal Oxides Methanol decomposition has also been studied on fluorite metal oxides. In this bulk structure each metal cation is octahedrally coordinated to oxygen anions, and each anion is tetrahedrally coordinated to metal cations. Methanol decomposition on Zr02 has been the subject of considerable surface analysis largely because the decomposition pathway has demonstrated structure sensitivity [75]. The ZrO2(100) crystal face is characterized by square arrays of zirconia bonded to four subsurface oxygen anions. Methanol adsorbed on ZrO2(100) at 125 K decomposed into methoxide and hydroxyl groups. Seventy-five percent of the adsorbed methanol desorbed as the parent molecule at 230 K through a recombinative pathway. The balance of surface methoxide reacted between

431

ZrOa (100)

•

:Zr*'^ in 1st layer

® . iZr'"'^ in 2nd layer

ZrOa (110)

Q .

-O'^ •" Ist layer lO'"^ in 2nd layer

Fig. 8. Models of the (100) and (110) surfaces of cubic zirconia. Both top and side views are presented. P.A. Dilara and J.M. Vohs, Surface Science, v. 321 (1994). Reproduced by permission of Elsevier Publications.

610 K and 630 K to liberate CO and methane. Comparison between lightly sputtered and stoichiometric Zr02 surfaces demonstrated that oxygen vacancies were not essential for the formation of methoxide [75]. The ZrO2(110) surface, characterized by rectangular arrays of zirconium cations and oxygen anions in the same plane, exhibited substantially different surface chemistry in the high temperature desorption state. Figure 8 illustrates both the (110) and (100) basal planes of zirconia. Although nearly three-quarters of the adsorbed methanol desorbed at 240 K via recombinative generation of the parent molecule, the remaining methoxides primarily formed formaldehyde and methane in a 2.5:1 ratio. Approximately equal amounts of methane and CO were formed on the Zr02 (110) surface. Temperature programmed HREELS recorded the dissociative adsorption of methanol to form methoxide. However, HREELS also indicated the formation of a new intermediate after heating to 300 K, identified as dioxymethylene (CH2O2), which could decompose to produce formaldehyde [75]: (25) (26)

CH3O* + Os ^ CH2O2* + H* CH202*^CH20(g)+Os

432

The production of dioxymethylene from methoxides competed with dehydrogenation to form CO and methane [75]: (27) (28)

CH3O* ^ CO(g) + 3 H* CH3O* + H* -^ CH4(g) + Os

Interaction of methoxide with the in-plane oxygen anions on the (110) surface is necessary to form the dioxymethylene intermediate [75]. CH 3

A1

O—Zr-0

H3C^

p I O—Zr-0

CH,

•

o

o ^ "^ *—Zr-O

However, the methoxide would be expected to have a much more difficult time forming the dioxymethylene intermediate on the (100) surface, as the oxygen anions are below the surface plane (see figure 8). Consequently, this surface produced methane instead. The reaction to produce methane is also driven by the coordinative unsaturation around the zirconium cations on the ZrO2(100) surface. The Zr"^"^ cations are four-fold coordinate on ZrO2(100) as compared to six-fold coordinate on ZrO2(110), giving the ZrO2(100) surface a greater affinity for oxygen and therefore those site may be better able to cleave the carbon-oxygen bond in the methoxide [75]. Methanol reactions have also been studied on polycrystalline wafers of UO2 [76]. Two parent desorption states existed for methanol adsorbed at 90 K. Molecularly adsorbed methanol desorbed at 110 K, and methanol generated by surface recombination of methoxides and protons desorbed at 180 K. Carbon (Is) XPS demonstrated that methanol dissociatively adsorbed on the urania surface and that methoxide was the only surface intermediate present above 150 K. Primary reaction products were methane and carbon monoxide at 480 K. Oxygen atoms not removed from the surface as CO were incorporated into the oxide surface; isotopically labeled U^^02 surfaces did not exchange oxygen with methoxide to produce C^Ô [76]. 4.5. Decomposition of Higher Alcohols on Titanium Dioxide Single Crystal Surfaces Decomposition reactions of larger aliphatic alcohols have been examined in detail on the {Oil}-faceted TiO2(001) surface [80]. Ethanol adsorbed at 300 K exhibited a low temperature desorption peaks for ethanol and water at 365 K and a high temperature desorption state for decomposition products at 588 - 595 K. Half of the ethanol adsorbed on the surface desorbed as ethanol at 365 K. Half of the remaining surface ethoxide groups desorbed as ethanol at 588 K. The

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remaining twenty-five percent of carbon-containing products desorbed as ethylene and acetaldehyde in a 5:1 molar ratio between 588 - 595 K. Temperature programmed XPS measurements for ethanol adsorbed at 300 K revealed no intermediates distinguishable from surface ethoxide (e.g., no carboxylates), and measured peak area changes corresponded well to quantitative changes in TPD desorption events. High-temperature production of olefins from ethoxide is consistent with hydride elimination from the p-position observed on polycrystalline anatase powders (reaction (29)), while the concomitant production of ethanol must occur via reaction of from the released hydrogen with other ethoxides (reaction (30)). Acetaldehyde is also produced on the (001) surface (reaction (31)), the result of ethoxide dehydrogenation via hydride elimination from the a-position [80]. (29) (30) (31)

CH3CH2O* -> H2C=CH2(g) + Os + H* CH3CH2O* + H* -^ CH3CH20H(g) CH3CH2O* -^ CH3CH0(g) + H*

Qualitatively similar results were obtained for reaction and desorption of normal and iso-propanol on the {Oil}-faceted TiO2(001) surface. In the case of normal propanol, almost half of the molecules initially adsorbed desorbed as the parent molecule at 370 K, while half of the remaining surface species reacted to form propanol at 580 K. The ratio of propene to propionaldehyde generated at 580 K was 10:1. Desorption of isopropanol quantitatively mirrored the desorption of normal propanol in two desorption states at 365 and 512 K. Isopropanol did not generate any dehydrogenation products (e.g., acetone), and the surface did not generate any bimolecular coupling products for any of the probe alcohol molecules. The absence of ether formation on the {011}-faceted surface is consistent with the need for double-coordination vacancies to facilitate that reaction, and the absence of such sites on this surface of titanium dioxide [80]. Campbell et al. have recently studied the reactions of surface ethoxy groups on the TiO2(110) surface [81]. Temperature programmed desorption studies of isotopically labeled ethanol were used to probe the surface intermediates in ethanol decomposition. As in the other work discussed here, low temperature desorption features (< 250 K) were associated with the desorption of molecularly adsorbed ethanol. Between 250 and 400 K, ethanol desorbed as a result of the recombination of surface ethoxy groups with surface hydroxy] groups. Seventyfive percent of the surface ethoxy groups produced by saturation exposure of ethanol to the (110) surface generated ethanol below 400 K. The remaining onequarter of the initial ethoxy population desorbed at 650 K as ethanol and ethylene. The high temperature desorption channel was thought to involve

434

hydrogen elimination from the p-position (earUer proposed on {Oil }-TiO2(001)) as a result of further isotopic labeling experiments. The CD3CH20H-dosed TiO2(110) surface produced primarily CD2CH2 [81]. Had the reaction proceeded through hydrogen elimination from the aposition and subsequent hydride shift, the expected product would have been CD2CHD. A kinetic isotope effect of about 2 kJ/mol was detected for ethylene production in TPD of CD3CH2OH versus CH3CH2OD, consistent with C-D vs. C-H scission as the rate determining step in ethoxide decomposition. Furthermore, TPD analysis suggested that the dehydrogenation reaction was not concerted with ethylene desorption. To explain these observations, Campbell proposed that initial ethanol exposure to the TiO2(110) surface generates surface hydrogen bound to bridging oxygen anions and surface ethoxides bound to titanium cation centers [81]: (32)

CH3CH2OH* -^ Ti-OCH2CH3 + Obridging-H

Seventy-five percent of the ethoxy groups recombine with the surface hydrogen to reform ethanol as the reverse of reaction (32). The remaining twenty-five percent are stable to 500 K. In fact, these ethoxides were irreversibly formed; when exposed to water at low temperature, they did not

Ti02 (110) Surface

(^J) Bridging Oxygen ((V) I Surface Oxygen (QJ

#

Ti''

V^r = Bridging Oxygen Vacancy

Bulk Oxygen (O^)

Fig. 9. Representation of the TiOi (110) surface. Binding of-OR groups (to represent either OH or OEt) is shown at both Ti ^'"*" sites and bridging oxygen vacancies. L. Gamble, L.S. Jung, and C.T. Campbell, Surface Science, v. 348 (1996). Reproduced by permission of Elsevier Publications.

435

react to form any ethanol. This suggests that the surface ethoxides converted to a different surface binding site. Examination of the CH3CH2OD TPD revealed coverage-dependent production of HDO and D2O in addition to backgroundadsorbed water. Consequendy, Campbell proposed a water-generation step (reaction (33)) stoichiometrically analogous to reaction (18) that involved the creation of surface oxygen vacancies. Figure 9 illustrates the two binding sites proposed [81]. (33)

2 Obridging-H - > H20(g) + Obridging + ^ bridging

(34)

T i - O C H 2 C H 3 + Vbrldging - ^ Obridging-CH2CH3

Reaction (34) represents surface diffusion of the ethoxide to fill in oxygen vacancy. The resulting Obridging-CH2CH3 remains on the surface until the Phydrogen is eliminated, forming ethylene [81]. 4.6. Higher Alcohols: Extension to Poly crystalline TiOi and Other Materials Ethanol decomposition has been studied on anatase and rutile powders of titanium dioxide. Much like methanol decomposition, it was demonstrated that the bulk form of titania did not influence the surface chemistry significantly [55]. Furthermore, TPD studies of ethanol decomposition on clean and hydroxylated titania powders have also demonstrated the insensitivity of high-temperature reaction channels to the initial presence of surface hydroxyl groups, while lowtemperature desorption modes of ethanol and water are affected by pre-exposure of the oxide surface to water, where it acts as a site-blocker [82]. These results were in good agreement with similar experiments on the (110) single crystal face conducted by Campbell et al, and together they suggest that there are two separate adsorption modes for ethanol on the surface of titanium dioxide [81]. Ethanol TPD on anatase powders revealed the existence of two desorption states (390 and 590-620 K). Approximately sixty percent of the adsorbed ethanol desorbed as the parent molecule via the low-temperature state in powder TPD experiments. The remainder decomposed into acetaldehyde, ethylene, diethyl ether, and water. Small amounts of ethanol and butene were produced as well. Reaction and decomposition of n-propanol and isopropanol on anatase titanium dioxide powders were also characterized by two desorption states; much like ethanol and methanol, the low temperature state was characterized by desorption of the parent molecule, while the high temperature state contained decomposition products. Major decomposition products from normal propanol reaction included propylene, di-n-propyl ether, propionaldehyde, and water, while isopropanol decomposition produced mostly propylene and water. It was suggested that steric effects inhibited the bimolecular coupling of isopropoxides to form the corresponding ether [78].

436

Recent work by Idriss and Seebauer has focussed on using thermal reactions on ethanol on many oxide powders to correlate material properties with oxide reactivity [83]. In steady-state catalytic experiments at low conversion over Ti02, ethanol product selectivity heavily favored the production of acetaldehyde (selectivity ca. 93%) and its derivatives, ethyl acetate and acetone. Acetaldehyde is produced via oxidative dehydrogenation of surface ethoxides, the formation of which Idriss and Seebauer determined was related to the degree of coordinative unsaturation of surface cations. Furthermore, the dehydrogenation reaction requires that the surface be reducible and reoxidizable due to the participation of lattice oxygen in the surface reactions. Acetaldehyde reacted further to produce ethyl acetate via a proposed Tischenko reaction involving hydrogen transfer between two aldehydes and subsequent coupling of the resulting aldehyde/alkoxide pair [83]. (35)

H3C(CH)=0* + H3C(CH)=0* ^

H3C(CO)-0-CH2CH3(g)

Acetone was suggested to be formed by oxidation of acetaldehyde followed by coupling of the resulting acetate intermediates [83]: (36) (37)

CH3CHO* + Os ^ CH3COO* + H* 2 CH3COO* -> C02(g) + (CH3)2C=0(g) + Os

Two conditions must be satisfied for acetone formation from acetaldehyde on the metal oxide surface. The cation must be reducible, as lattice oxygen must be abstracted to form the surface acetate group. Si02, which will dehydrogenate ethanol to form acetaldehyde, is not capable of this, as the oxygen-silicon bond is too strong. The metal cation must also possess two coordination vacancies. For example, MgO will also produce acetaldehyde but not acetone in UHV. Magnesia possesses a rocksalt structure with six-coordinate occupation around the cation, and cleavage of the (100) surface simply truncates the bulk lattice, leaving a single coordination vacancy at each magnesium cation. Single crystal work with MgO(lOO) has demonstrated that acetates on this surface will not couple, but instead undergo a Cannizzaro-like disproportionation reaction [83]. There is a rich array of chemistry centered on alkoxides and carboxylates such as the reaction of acetic acid to form ketene on titania [45,46] and silica [84], and aldol condensation of acetaldehyde to form on titania [77]. In fact, both of these reactions represent cases where single crystal work ultimately provided the basis for understanding more complex powder-catalyzed chemistry.

437

4.7. Photoreactions of Alcohols on TiOi Photoreactions on titanium dioxide have been the focus of considerable interest for some time. Titania offers the opportunity to oxidize organic compounds in polluted environments, and has also been exploited to generate titania-supported nanoparticles of metals (e.g., silver) via photoreduction reactions [85]. While there is not enough room here to thoroughly treat photocatalytic processes, a brief introduction to the subject is presented below. Readers seeking detailed treatments of this subject are referred to a recent review by Yates et al. on titania-facilitated photocatalysis [86]. The utility of titania for photocatalytic applications is based on the band gap of the semiconductor. By definition, there are no states in the band gap to facilitate the recombination of electron-hole pairs; thus, excitation across the bandgap may allow charge transfer to an adsorbate and therefore its activation. Studies have concluded that the lifetime of excited states across the bandgap is on the order of nanoseconds, which is long enough to facilitate charge transfer. Two competing processes determine the fate of the electron-hole pair. The first is simple electron-hole recombination, leading to no reaction and a return to the initial state. The second route allows the semiconductor to donate the electron to an electron-accepting adsorbate. The hole may then combine with an electron from an oxidizable adsorbate. The efficiency of this process is conventionally referred to as the quantum yield, which is defined as the number of reaction events per photon adsorbed. If all incident light is assumed to be absorbed, it is more effectively defined as the ratio of the rate of charge transfer to adsorbates to the sum of all electron-hole consuming events. On titanium dioxide, the concentration of electrons tends to exceed that of holes, because the transfer of electrons to adsorbed oxygen is relatively slow [86]. Early studies on titania powders showed that methanol generated methyl formate as the principle photooxidation product. Molybdena- and vanadiamodified Ti02 catalysts demonstrated at least an eighty percent drop in activity relative to pure titania, although selectivity to dimethoxymethane (and thus suppression of further oxidation products) was almost total [87]. Recently, Lin et al. have studied methanol photochemistry on Ti02 powders with IR in order to spectroscopically identify the important surface intermediates [88]. Investigation of methanol reactivity in the absence of oxygen indicated the presence of surface methoxides which decomposed to form methanol, formaldehyde, water and hydrogen. No additional IR features were detected during UV exposure reactions, although the surface concentration of methoxide decayed on consumption during the reaction. In the presence of oxygen, UV exposure caused features in the IR spectrum associated with surface formate to grow at the expense of methoxide. The rate of methoxide decomposition was initially slower in the oxygen atmosphere than without oxygen, although methoxide concentration monotonically decreased throughout the experiment.

438

The surface methoxide concentration approached a stable population when oxygen was absent from the system (figure 10) [88]. Electron spin resonance studies of aqueous methanol reactions with colloidal titania have suggested the appearance of methoxy radicals [89]; the participation of similar radicals was suggested for methanol decomposition in the absence of vapor phase oxygen. Lin suggested that methoxy groups absorb holes and generate methoxy radicals on the surface. (38) (39)

Ti^^-0-Ti^^-0CH3 + h^ -> Ti^^-0-Ti^^-0CH2- + [Hi* Ti^^-0-Ti^^ + CH2O,'(g) Ti^^-0-Ti'^-0CH2

The rapid decrease in rate shown in figure 10 was likely due to electron/hole recombination at the surface as the concentration of surface electrons increased with continued irradiation [88]. The apparently slower rate of methoxide decomposition in the presence of oxygen is likewise dependent on the photoexcitation process. As stated above,

(A) -oin the absence of oxygen

2000 4000 6000 8000 Photo-irradiation Time (sec)

10000

Fig. 10. Relative concentration of CH30(a) as a function of UV irradiation time in the absence of (A) O2 and in the presence of (B) 10 Torr of O2. C.-C. Chuang, C.-C. Chen, and J.-L. Lin, Journal of Physical Chemistry B, v. 103 (1999). Reproduced by permission of the American Chemical Society.

439

charge transfer to adsorbed oxygen is a relatively slow process. Lin postulated that the methoxy radicals incorporated photoexcited oxygen in the form of peroxide radicals [88]: (40) (41) (42)

O2* + H* + e- -^ HOO* CH3O* + hv + h"^ + 02(g) -^ *0CH200» HOO* + *0CH200* -> [OCH2OOOOH]* -^ HCO2* + H2O + O2*

Thus, the formation of the organoperoxide transition state is dependent on the availability of the peroxide radical formed by photoexcitation of adsorbed oxygen, the rate limiting step of the process [88]. Methyl formate photodecomposition has also been studied on titanium dioxide powders. In the presence of oxygen, methyl formate decomposed to form CO, CO2, formaldehyde, and water [90]. 5. SUMMARY Selectivity on metal oxide catalysts is ultimately determined by complex intermolecular and surface-adsorbate interactions. Competing reaction channels are facilitated or hindered by the coordination geometry around metal cations, the ease of reduction of the surface, and the resulting stabilization of surface intermediates. The decomposition of relatively simple organic molecules like methanol and formic acid can be surprisingly complex, but attention to a few concepts may help to understand the reaction processes: 1) The population of surface defects and coordination vacancies drives alkoxide and carboxylate formation and decomposition. When cations have at least two coordination vacancies, bimolecular reactions are possible (e.g., acetone from acetate ions, dimethylether from methoxy groups). 2) The ease of reduction of the metal oxide is often coupled to the product selectivity due to the involvement of oxygen vacancies in the reaction. For example, MgO(lOO) exclusively dehydrates formic acid to form CO, while the more reducible ZnO(0001)-Zn surface readily produces CO2 by dehydrogenation of formate. 3) Reduced cation oxidation states can drive reduction reactions of oxygenated adsorbates. Sputter-reduced surfaces of TiO2(001) readily converted formic acid into formaldehyde and methanol into methane and CO. 4) Interactions with other molecules can affect the ultimate reaction selectivity of the process. For example, the catalytic dehydration of formic acid on TiO2(001) occurred as a unimolecular process at high temperatures and low formate coverage, while a bimolecular dehydrogenation process dominated at near-saturation coverage of the titania crystal.

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Even though many of these principles appear to be applicable to for multifunctional carboxylates and alkoxides, it is important to recognize that more complex molecules may be more or less influenced by differences in surface conditions than others. Small, monofunctional molecules ably serve to highlight site requirements for some reactions, but there is no reason to expect that a large, multifunctional molecule will necessarily interact with a metal oxide catalyst as a mere combination of its functionalities. Consequently, it is important to characterize the adsorption and synthesis of larger molecules where possible, to determine the limitations of the principles explored here, and to develop an understanding of adsorption and reaction characteristics that will lead to more selective catalysts. REFERENCES [I] [2] [3] [4] [5] [6]

[7] [8] [9] [10] [II] [12] [13] [14] [15] [16] [17] [18] [19] [20]

H. Idriss and M.A. Barteau, Advances in Catalysis, 45 (2000) 261. M.A. Barteau and J.M. Vohs in Handbook of Heterogeneous Catalysis, G. Ertl and H. Knozinger, VCG, Vêinheim, (1997) Vol. 2, pp. 873. M.A. Barteau, Chemical Reviews, 96 (1996) 1413. V.E. Henrich and P.A. Cox, The Surface Science of Metal Oxides, Cambridge University Press, Cambridge (1994). H.-J. Freund, Faraday Discussions, 114 (2000) 1. J.-M. Basset, B.C. Gates, J.-P. Candy, A. Choplin, M. Leconte, F. Quignard and C. Santini, Eds., Surface Organometallic Chemistry: Molecular Approaches to Surface Catalysis, Kluwer Academic Publishers, Dordrecht (1986) Vol. 231. M.A. Henderson, Surface Science, 419 (1999) 174. H. Onishi, K. Fukui and Y. Iwasawa, Bulletin of the Chemical Society of Japan, 68 (1995) 2447. M. Li, W. Hebenstreit, L. Gross, U. Diebold, M.A. Henderson, D.R. Jennison, P.A. Schultz and M.P. Sears, Surface Science, 437 (1999) 173. M. Ashino, T. Uchihashi, K. Yokoyama, Y. Sugawara, S. Morita and M. Ishikawa, Applied Surface Science, 157 (2000) 212. S. Petigny, H. Mostefa-Sba, B. Domenichini, E. Lesniewska, A. Steinbrumm and S. Bourgeois, Surface Science, 410 (1998) 250. M. Brause, S. Skordas and V. Kempter, Surface Science, 445 (2000) 224. S. Tanaka, K. Mase, M. Nagasono, S. Nagaoka and M. Kamada, Surface Science, 451 (2000) 182. M. Li, W. Hebenstreit, U. Diebold, A.M. Tyryshkin, M.K. Bowman, G.G. Dunham and M.A. Henderson, Journal of Physical Chemistry B, 104 (2000) 4944. M. Li, W. Hebenstreit and U. Diebold, Surface Science, 414 (1998) 1951. V.S. Lusvardi, M.A. Barteau, J.G. Chen, J. J.J. Eng, B. Fruhberger and A. Teplyakov, Surface Science, 397 (1998) 237. B.A. Watson and M.A. Barteau, Chemistry of Materials, 6 (1994) 771. G.S. Herman, Y. Gao, T.T. Tran and J. Osterwalder, Surface Science, 447 (2000) 201. H. Norenberg, F. Dinelli and G.A.D. Briggs, Surface Science, 446 (2000) 183. H. Idriss and M.A. Barteau, Catalysis Letters, 26 (1994) 123.

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R.A. Bennett, P. Stone, R.D. Smith and M. Bowker, Surface Science, 454-456 (2000) 390. H. Onishi, Y. Yamaguchi, K. Fukui and Y. Iwasawa, Journal of Physical Chemistry, 100 (1996) 9582. P. Stone, R.A. Bennett and M. Bowker, New Journal of Physics, 1 (1999). Q. Guo, I. Cocks and E.M. Williams, Journal of Physical Chemistry, 106 (1997) 2924. Q. Guo and E.M. Williams, Surface Science, 433-435 (1999) 322. K. Fukui and Y. Iwasawa, Surface Science Letters, 464 (2000) L719. J.N. Wilson, D.J. Titheridge, L. Kieu and H. Idriss, Journal of Vacuum Science and Technology A, 4 (2000) 1887. L. Patthey, H. Rensmo, P. Persson, K. Westermark, L. Vayssieres, A. Stashans, A. Petersson, P. Bruhwiler, H. Siegbahn, S. Lunell and N. Martensson, Journal of Chemical Physics, 110 (1999) 5913. E. Soria, E. Roman, E.M. Williams and J.L. deSegovia, Surface Science, 433-435 (1999)543. E. Soria, I. Colera, E. Roman, E.M. Williams and J.L. deSegovia, Surface Science, 451 (2000) 188. J.M. Tatibouet, Applied Catalysis A, 148 (1997) 213. M.A. Henderson, S. Otero-Tapia and M.E. Castro, Faraday Discussions, 114 (2000) 313. M.A. Henderson, S. Otero-Tapia and M.E. Castro, Surface Science, 412/413 (1998) 252. G.S. Wong, D.D. Kragten and J.M. Vohs, Surface Science, 452 (2000) L293. K.S. Kim and M.A. Barteau, Surface Science, 223 (1989) 13. P.A. Dilara and J.M. Vohs, Surface Science, 321 (1994) 8. J.A. Lloyd, W.L. Manner and M.T. Paffett, Surface Science, 423 (1999) 265. J.E. Rekoske and M.A. Barteau, Langmuir, 15 (1999) 2061. K.S. Kim, M.A. Barteau and W.E. Fameth, Langmuir, 4 (1988) 533. V.A. Gercher, D.F. Cox and J.-M. Themlin, Surface Science, 306 (1994) 279. K.S. Kim and M.A. Barteau, Journal of Molecular Catalysis, 63 (1990) 103. L. Gamble, L.S. Jung and C.T. Campbell, Surface Science, 348 (1996) 1. V.S. Lusvardi and M.A. Barteau, Journal of Physical Chemistry, 100 (1996) 18183. H. Idriss and E.G. Seebauer, Journal of Molecular Catalysis A, 152 (2000) 201. R. Martinez, M.C. Huff and M.A. Barteau, Applied Catalysis A, 200 (2000) 79. W.E. Farneth, R.S. McLean, J.D. Bolt, E. Dokou and M.A. Barteau, Langmuir, 15 (1999) 8569. A.L. Linsebigler, G. Lu and J. J.T. Yates, Chemical Reviews, 95 (1995) 735. T. Carlson and G.L. Griffin, Journal of Physical Chemistry, 90 (1986) 5896. C.-C. Chuang, C.-C. Chen and J.-L. Lin, Journal of Physical Chemistry B, 103 (1999) 2439. O.I. Micic, Y. Zhang, K.R. Cromack, A.D. Trifunac and M.C. Thurnauer, Journal of Physical Chemistry, 97 (1993) 13284. C. Chuang, W. Wu, M. Huang, I. Huang and J. Lin, Journal of Catalysis, 185 (1999) 423.

Oxide Surfaces D.P. Woodruff, editor © 200J Elsevier Science B. V. All rights reserved.

443

Chapter 11

The Structure of TiOi surfaces Ulrike Diebold Department of Physics, Tulane University, New Orleans, LA 70118, U.S.A. 1. INTRODUCTION Unraveling the relationship between the atomic surface structure and other physical and chemical properties is probably one of the most important achievements of surface science. Because of the mixed ionic and covalent bonding in metal oxide systems, the surface structure has an even stronger influence on local surface chemistry as compared to metals or elemental semiconductors [1]. A vast amount of work has been performed on Ti02 over the years, and this is certainly the best-understood surface of all the metal oxide systems. This chapter starts with a brief description of the bulk structure of titanium dioxide crystals, and their stable crystal planes. Because bulk nonstoichiometries influence the surface properties of Ti02 in a variety of ways, we include a short discussion of bulk defects. The longest part of the chapter is devoted to the rutile (110) surface. The (bulk-truncated) (1x1) surface is known with a very high accuracy from experimental as well as theoretical studies. Even more importantly, there has been given a reason on why theory and experiment disagree in some aspects [2]. Surface defects are categorized as step edges, oxygen vacancies, line defects (closely related to the (1x2) reconstruction), common impurities, and the manifestation of crystallographic shear planes on surfaces. The long-standing argument of the structure of the (1x2) phase seems to be settled, as discussed in section 3.2.2. Scanning Tunneling Microscopy and, more recently. Atomic Force Microscopy, studies have revealed the complexity of the seemingly simple rutile (110) surface, hence the section on Ti02(l 10) commences with a recommendation on how to best to prepare this surface. The two other low-index planes, rutile (100) and (001) are described in sections 4 and 5, respectively. Until fairly recently the (1x3) reconstruction of the rutile (100) seemed well understood, but inconsistencies in theoretical calculations as well as new interpretations of xray diffraction data show that closer look on the structure of this phase may be

444

needed (4.2.2.). Very recent developments on structural investigations of anatase surfaces are included at the end.

2. BULK STRUCTURES Titanium dioxide crystallizes in three different structures; rutile (tetragonal; D4hi4-P42/mnm, a = b = 4.584 A, c = 2.953 A, ), anatase (tetragonal; D4hi9I4i/amd, a = b = 3.733 A, c = 9.37A), and brookite (rhombohedrical, D2h^^Pbca, a = 5.436 A, b = 9.166 A, c = 5.135 A) [3]. Only rutile and anatase play any role in the applications of Ti02 and have been studied with surface science techniques. Their unit cells are shown in Fig. 1. In both structures, the basic building block consists of a titanium atom surrounded by six oxygen atoms in a more or less distorted octahedral configuration. In each structure, the two bonds between titanium and the oxygen atoms at the aspices of the octahedron are slightly longer. A sizable deviation from a 90° bond angle is observed in anatase. In rutile, neighboring octahedra share one corner along type directions, and are stacked with their long axis alternating by 90°. This results in three-fold coordinated oxygen atoms. Rutile Ti02 single crystals are widely available. They can be bought in cut and polished form from companies such as Commercial Crystal Laboratories, U.S.A.; Kelpin Kristallhandel, Germany; Goodfellow Ltd., U.K.; or Earth Jewelry, Japan; and many others. A very small roughness is achieved by grinding the sample, and then polishing the surface for many hours with a chemo-mechanical treatment. This is also referred to as epitaxial polish. Practical aspects of surface preparation and handling are discussed in ref [4].

Anatase

Rutile 1.946 A

titanium

[010] ygen

1.983

[001] [010] [100]

1.937 A

Fig. 1. Unit cells of rutile and anatase. The tetragonal bulk unit cells (rutile a = b = 4.584 A, c = 2.953 A; anatase a = b = 3.733 A, c = 9.37A) are indicated. In both structures, slightlydistorted octahedra are the basic building units. The the bond lenghts and angles of the octahedrally-coordinated Ti atoms are indicated.

445

Ramamoorthy et al. [5] calculated the total energy of periodic Ti02 slabs using a self-consistent ab initio method. The (110) surface has the lowest surface energy, and the (001) surface the highest. This is expected from the considerations of surface stability, based on electrostatic and dangling-bonds arguments discussed in section 3.1.1. below. The thermodynamic stability of the (100) surface was also considered, and was found to be stable with respect to forming (110) facets. The experimental results on the three lowindex rutile surfaces discussed below fit rather well with the stability expected from these calculations. For rutile, the (110), (001), and (100) surfaces have been studied, with (110) being the most stable one. These three surfaces are discussed in this chapter. The two simple approaches that are commonly used to estimate the stability and structure of oxide surfaces are exemplified in detail for the rutile (110) surface. For anatase, the (101) and the (100)7(010) surface planes are found in powder materials, together with some (001). The (101) surface was calculated to have the lowest surface energy, even lower than the rutile (110) surface [6]. The first experimental results on this surface appear to confirm this theoretical prediction [7,8]. 2.1. Bulk defects The titanium-oxygen phase diagram is very rich with many stable phases with a variety of crystal structures, see Fig. 2 (ref [3]). Consequently, Ti02 can be reduced easily. Bulk reduction and the resulting color centers are reflected in a pronounced color change in Ti02 single crystals from initially transparent to light and, eventually, dark blue [9]. These intrinsic defects result in n-type doping and a high conductivity. This high conductivity makes Ti02 single crystals such a convenient oxide system for experimentalists. As has been pointed out recently [10], bulk defects play a major role in a variety of surface phenomena where annealing to high temperature is necessary , e.g. during the encapsulation of Pt [11-13], in bulk-assisted re-oxidation [14,15], in restructuring and reconstruction processes [9,16], and in gas adsorption [17]. The relationship between crystal color, conductivity, bulk defects as characterized by EPR measurements, and surface structure of rutile (110) has been investigated systematically by Li et al. [10]. The bulk structure of reduced Ti02-x crystals is quite complex with a various types of defects such as doubly charged oxygen vacancies, Ti3+ and Ti4+ interstitials, and planar defects such as crystallographic shear (CS) planes. The defect structure varies with oxygen deficiency which depends on temperature, gas pressure, impurities, etc. Despite years of research, it is still subject to debate which type of defect is dominant in which region of oxygen deficiency [18,19]. It was shown that the dominant type are Ti interstitials in the region from TiOi.9996 to TiOi.9999 (from 3.7x10^8 to 1.3x1919 missing O atoms per cm^) [19]. CS planes precipitate on cooling crystals across the Ti02-x (0 < X < 0.0035) phase boundary [20]. They show a very strong

446

1200

0.8

1.2

1.6

O/Ti ratio

Fig. 2. Phase diagram of the Ti-0 system re-drawn from ref. [3]. The region Ti203-Ti02 contains Ti203, Ti305, seven discrete phases of the homologous series Tin02n-l (Magneli phases), and Ti02. See ref [3] for a more detailed description.

dependence on the cooling history and are absent in quenched specimen. The formation mechanism was reviewed by Smith et al. [20-22]. Such CS planes may extend all the way to the surface [23-27] and their appearance is discussed in section 3.1.4 below. The diffusion mechanism for the various types of defects is quite different; oxygen migrates via a site exchange (vacancy diffusion) mechanism, while excess Ti diffuses through the crystal as interstitial atoms. The interstitial diffusion happens especially fast through the open channels along the (001) direction (the crystallographic c-axis) [28,29]; see Fig. 3a. A Ti interstitial located in these channels is in an octahedral configuration, similar to the regular Ti sites. Consequently, the diffusing species in oxidation reactions of reduced TiaOb surfaces (where a > b/2 but probably less than b) produced by sputtering and/or Ti deposition is the Ti atom and not the O vacancy, as has been shown in a series of elegant experiments with isotopically labeled 'Ô and 46Ti by Henderson [14,15].

3. THE STRUCTURE OF THE RUTILE TiO2(110) SURFACE The rutile (110) surface is the most stable crystal face and simple guidelines can be used to essentially predict the structure and estimate the stability of the TiO2(110)(lxl) surface. Because these concepts are very useful for the other crystal faces of Ti02 as well other oxide materials, they are exemplified for this surface. The relaxations from the bulk-terminated coordinates are reviewed as

447

well as the types and shapes of defects. Although the TiO2(110) surface is very stable, it nevertheless reconstructs and restructures at high temperatures under both oxidizing and reducing conditions. 3.1. The rutile TiO2(110)(lxl) surface 3.1.1. Bulk truncation Two concepts have been introduced to predict the stability of oxide structures. Tasker [30] discussed the stability of ionic surfaces based on purely electrostatic considerations. The second concept, autocompensation, was originally developed for surfaces of compound semiconductors and applied to metal oxide surfaces by LaFemina [31]. The most stable surfaces are predicted to be those which are autocompensated, which means that excess charge from cation-derived dangling bonds compensates anion-derived dangling bonds. The net result is that the cation- (anion-) derived dangling bonds are completely empty (full) on stable surfaces. This model allows for the partially covalent character found in many metal oxides, including Ti02. Both concepts are used in a complementary way, and represent a necessary (but not sufficient) condition for stable surface terminations. With very few exceptions, stable metal oxide surfaces for which the structure is known are non-polar [30]a^(i fulfill the autocompensation criterion [31]. Tasker's and LaFemina's approaches are exemplified in creating a stable (110) surface (Fig. 3). In Tasker's concept, the dipole moment of a repeat unit perpendicular to the surface must be zero in order for the surface energy to converge. He introduced three categories for ionic (or partially ionic) structures. Type 1 (neutral, with equal numbers of cations and anions on each plane parallel to the surface) is stable. Type 2 (charged planes, but no dipole moment because of the symmetrical stacking sequence) is stable as well. Type 3 surfaces (charged and a dipole moment in the repeat unit perpendicular to the surface) will generally be unstable. Consider, for example, the rutile structure as being composed of (110)oriented planes such as drawn in Fig. 3 a. The top plane in Fig. 3a consists of the same number of Ti and O atoms. In a purely ionic picture, the titanium and oxygen atoms have formal charges of 4+ and 2-, respectively. Hence, the top layer has a net positive charge. The next two layers consist of oxygen atoms, hence both of them have a net negative charge. A type-2 repeat unit is outlined by the dashed lines A and B in Fig. 3a. It consists of a mixed Ti, O layer, sandwiched between two layers of oxygen atoms. The total unit does not have a dipole moment. A crystal, cut or cleaved^ to expose a (110) surface, will naturally terminate with the surface created by cutting along line A (or B) in Fig. 3a. In Fig. 3b, the top of the model is shifted along the (110) direction (cutting* the crystal in a 'Gedankenexperiment'). The resulting surface is very

Unfortunately, Ti02 fractures and does not cleave well.

448

corrugated because one 'layer' of oxygen atoms is left behind. As shown in below, there is overwhelming evidence that the (1x1) surface of TiO2(110) closely resembles the 'bulk-terminated' structure depicted in Fig. 3b. The same surface structure is also predicted using the rules of autocompensation. In Fig. 3b, the same number of oxygen-to-titanium bonds are broken as titanium-to-oxygen. Transferring electrons from the dangling bonds on the Ti cations will just compensate the missing charge in the dangling bonds on the O anions. Hence, the surface is autocompensated [31]. Note that only the longer (and, consequently, weaker) bonds are broken when the crystal is sliced in this way. This is yet another indication that this surface termination is a likely one. The rutile (110)(lxl) surface in Fig. 3b contains two kinds of oxygen atoms. Along the [001] direction, rows of 6-fold coordinated Ti atoms (as in the bulk) alternate with 5-fold coordinated Ti atoms with one dangling bond perpendicular to the surface. Two kinds of oxygen atoms are created as well. Within the main surface plane, they are three-fold coordinated as in the bulk. The so-called bridging oxygen atoms miss one bond to the Ti atom in the removed layer and are two-fold coordinated. These bridging oxygen atoms are

;S[v W ^M y^

w'Jk ^^^ ^ ^ ^

[110]^

ii [001] 1]

^^T^^'r

^l^^^îH^^/^^l^^^îl^^^^îi

3.

B

X^=(?>^m, F^^^^vJinl

Y^^

^

uO ^" j l

r^ ^^^^^jX

^rp\^ ~^

î7!iS

l)Lny

TZ)

/rS

Y\_^^^^^^^^fej J l [no]

Fig. 3(a) Ball-and-stick model of the rutile crystal structure. It is composed of slightly distorted octahedra, two of which are indicated. Along the [110] direction these octahedra are stacked with their long axes alternating by 90°. Open channels are visible along the (001) direction. The dashed lines A and B enclose a charge-neutral repeat unit without a dipole moment perpendicular to the [110] - direction (a 'type 2' crystal plane [30]). (b) The crystal is 'cut' along line A. The same number of Ti -> O and O -> Ti bonds are broken, and the surface is autocompensated [31]. Note that only the longer bonds are broken. The resulting (110) surface plane is stable and overhelming evidence exists for such (1 x 1 )-terminated TiO2(110) surfaces.

449

subject to much debate. Because of their coordinative undersaturation, atoms from these rows are thought to be removed relatively easily by thermal annealing. The resulting point defects (section 3.1.4.) affect the overall chemistry of the surface, even in a macroscopic way [32]. A (1x1) LEED pattern is generally observed upon sputtering and annealing in UHV. To my knowledge no quantitative LEED study has been reported, probably because of the difficulty of creating defects when the sample is bombarded with electrons (see section 3.1.4). A medium-energy electron diffraction study (MEED) study of TiO2(110) employed an ESDIAD optics with a channelplate; this setup is more sensitive than a conventional LEED apparatus, and allows for very small electron currents to be used. The results of this study are basically consistent with the (1x1) structure depicted in Fig. 3b. X-ray photoelectron diffraction (XPD) spectra also fit the expected (1x1) termination [33], as do STM and AFM results (which are discussed in section 3.1.3 below). 3.1.2. Relaxations Every surface relaxes to some extent. In recent years, the geometry of the Ti02(l 10)(lxl) surface has been studied in some detail both experimentally and theoretically. The results of a surface x-ray diffraction experiment (SXRD) [34] and of several total-energy calculations are listed in Table 1. The experimentally-determined directions of atoms in the first layers are sketched in Fig. 4. As is expected from symmetry, the main relaxations occur perpendicular to the surface. Only the in-plane oxgyens ('4', '5' in Fig. 4) move laterally towards the 5-fold coordinated Ti atoms. The bridging oxygen atoms (labeled '3' in Fig. 4) are measured to relax downwards considerably, and the 6-fold coordinated Ti ('1') atoms upwards. The 5-fold coordinated Ti atoms ('2') move downwards and the neighboring 3-fold coordinated oxygen atoms ('4', '5') upwards, causing a rumpled appearance of the surface. The relaxations in the second Ti02 layer are roughly a factor of two smaller. The most striking feature is the large relaxation of the bridging oxygen atoms by -0.27 A. The measured geometry would indicate a very small bond length between the six-fold coordinated Ti atom ('1') and the bridging oxygens ('3') of only 1.71± 0.07A instead of the 1.95 A expected from the bulk structure. The relaxation results in vertical distances of 0.89 ±0.13A and 1.16 ±0.05A from the 6-fold ('1') and 5-fold coordinated ('2') Ti atoms, respectively. This is in agreement with ion scattering measurements, where vertical distances of 0.87 A and 1.05 ±0.05A were found [35,36]. (Another ion scattering study found the height of the bridging oxygen atoms comparative to that of the bulk structure but the interlayer distance largely relaxed with about -18 ± 4% [37].) Photoelectron diffraction results [38] are also in agreement with relaxations from the x-ray diffraction work given in Table 1.

450

A [110]

[110]

[001]

Fig. 4. Model of the TiO2(110) surface. The relaxations of surface atoms, determined with SRXD are indicated [34]. The labels refer to the relaxations listed in Table 1. Charlton

Harrison Harrison FPLAPW 7 layers

SXRD experiment

0.12*0,05 -0.16±0.O6

Tl(1)(6-fold) Ti(2)(5-fold) p(3)bridging) 0(4,5) [110] [1 1 0] 0(6)

ppiiii; 0{9) TK13) _ TK14)

,

0(15) O(16,17)[110] [110] 0(18) 0(19)

-

0,08 -0.23

0,23 /-0.17

-0.27±0.08 0.05±0.05 ±0.16±0.08 0.03±0.08

-0.16 0.09 ±0.06 -0.09

0.07*0.04 • and < 1 1 2 > with the expected step height of 1.6 A, as well as a strong variation in mesoscopic morphology [26]. The CS planes also act as nucleation sites for re-growth of new Ti02 layers during high-temperature oxidation (see section 3.2.2) [86]. 3.2. Reconstructions 3.2.1. Reconstruction under reducing conditions: The structure (s) of the (1x2) phase The most common reconstruction observed on TiO2(110) surfaces has a (1x2) symmetry with a doubling of the periodicity along the [110] direction. Various models have been suggested for this reconstruction and are depicted in Fig. 11 [50]. A (1 X 2) LEED pattern was originally observed after high-temperature annealing of a reduced Ti02(l 10) sample in ultrahigh vacuum (UHV). Based on Ti:0 AES ratios it was interpreted as alternate rows of bridging oxygen missing from the regular (1 x 1) surface ("missing-row moder'[48], Fig. 11a). One of the first atomically-resolved STM results on this surface was also interpreted as missing bridging oxygen rows [51,87]; however, the Ti atoms underneath the missing oxygen atoms had to be shifted by half a unit cell in [001] direction to account for the observed image contrast. A structure with a (3x2) symmetry was reported by one group [88]. A model for the this reconstruction was proposed where this symmetry is achieved by removing 1/3 or 2/3 of the oxygen in the bridging oxygen rows. However, such a reconstruction has not been reported by other groups. The simple missing row model for the (1x2) structure in Fig. H a has been discounted on the basis of more recent results. In STM the (1x2) reconstruction appears as a series of bright strings along the [001] direction [50,51,60,87,89-91]. At low coverage, the strings grow preferentially out of the upper terrace onto the lower one (Fig. 9 a) [60]. At first, they are scattered across the terraces with a minimum distance of 13A along the [001] direction. They consist of bright double strings (although the double-ridge structure is often not resolved), with a bright dot at the end (see Fig. 9 ). Antiphase boundaries are observed in high-resolution images of a fullydeveloped (1x2) reconstructed surface [50]. Higher periodicities, i.e., a local (1x3) reconstruction, have also been observed [90,92,93]. In STM images the (1x2) strands generally have an apparent height smaller than a regular Ti02 step edge of 3.2 A, and are in registry with the bright rows of the (1x1) substrate. Because most researchers report empty-state images and because these are dominated by the tunneling into mostly Ti 5(i-derived states bright

464 (a) missing row

(b) added TizOa' row

iMfiM (a) added row

Fig. 11. Models for the Ti02(l 10)(lx2) surface. Small white balls, Ti, large black balls, O. (a) The 'missing row' model, obtained by removal of one row of bridging oxygens, was originally proposed by Moller et al. [48]. This model is inconsistent with more recent STM images and first-principles calculations, (b) The 'added row model' has Ti203 stoichiometry as proposed by Onishi and Iwasawa [59]. (c) The 'missing unit' model was proposed by Pang et al. [92] . Recent evidence suggests that the structures in b) and c) might both be present at Ti02(l 10)(lx2) under different conditions. From ref [50], with permission.

strands centered on top of bright substrate rows imply that the (1x2) strands are at the position of 5-fold coordinated Ti atoms and not at the bridging oxygen atoms. STM images of a simple missing row structure are expected to show a bright feature above the missing bridging oxygen row (provided the STM tip is a reasonable distance from the surface [49]), inconsistent with the registry observed experimentally. The rows can be removed by tunneling under 'extreme conditions' (Vsampie = +1.5 V, I = 10 nA [50]). First-principles total energy calculations show that the missing row structure is energetically equivalent to a (2x1) structure (where every other bridging oxygen is removed) [41]. For all these reasons, the missing row reconstruction is no longer considered a viable model. Early on, Onishi and Iwasawa [59,60] suggested a quite different model. It consists of double rows of Ti cations in a distorted tetrahedral configuration

465

(Fig. 11). The structure has a Ti203 stoichiometry, and the model is often called 'added Ti203 rows'. However, it needs to be emphasized that the structure does not resemble the one found in corundum Ti203 crystals. Selfconsistent total-energy and electronic structure calculation found that this added 'Ti203' row structure has a lower surface free energy than the missing row structure and that it is consistent with the expected contrast in STM [94]. It is also supported by ESDIAD [95], high-resolution STM [50], and ion scattering [93] measurements. Based on the fact that (1x2) rows extend out of step edges, a modified model of the missing row structure has been proposed by Murray et al. [27] which involves narrow rows with missing bridging oxygens that are effectively part of the upper terrace. Lateral relaxations were also included [27]. This model was shown to be consistent with calculated surface charge densities [49]. Based on the same scheme, an 'added row model' was proposed by Pang et al. [92] for the fully-reconstructed surface. This consist of narrow, long regions of the regular Ti02 structure with all the atoms in bulk-like positions and with all bridging oxygen atoms missing (Fig. 11). The black grooves between the bright rows are then due to the missing Ti02 units separating the rows. (Consequently this model has also been called 'missing unit' structure [50].) The stoichiometry of the added rows is Ti305. The model was based on STM images with unusually high resolution and the observation of a (1x3) phase consisting of thicker rows. Charge-density calculations by Pang et al. supported the added row model, but were inconsistent with either the added Ti203 or the simple missing row model (Figs. 1 la and b, respectively). The off-normal lobes in ESDIAD images were supposed to stem from the O atoms at either side of the added Ti02 'rows', adjacent to the missing units. Pang's model has been questioned by Tanner et al. [50,96,97]. The 1x2 strands that are part of 'restructured' surfaces after low-temperature oxidation (see the next section) are consistent with the added Ti203 rather than the added row model [98]. Ion scattering measurements are also consistent with the added 'Ti203' model [93] although (somewhat surprisingly) extra oxygen atoms at the position of the 5-fold coordinated Ti atoms were postulated according to these measurements. To my knowledge, no total-energy calculations have been made that would support the added row structure. Such calculations could help to decide which of the added row models has a lower energy. As is discussed in the next section, the two added row models appear not to be mutually exclusive, and the formation of one or the other structure may just depend on the preparation and crystal parameters. 3.2.2. Restructuring under oxidizing conditions The first atomic-level investigation of the dynamic processes that occur when reduced Ti02 crystals are exposed to oxygen was reported by Onishi and Iwasawa [99]. These authors used a blue crystal with a resistivity of 2 Qm. They acquired STM images while the sample was kept at a temperature of 800

466

K and under an O2 background pressure of IxlO-^ Pa. Added rows (interpreted as Ti203 rows, Fig. lib) and 'hill-like features', appeared while imaging the surface, and disappeared when the oxygen was pumped out from the chamber. This effect was not tip-induced; the same structures were observed on areas of the sample that were not imaged during the hightemperature oxygen exposure. It is now established [9,16,100-103] that such an oxygen-induced surface restructuring effect is attributed to the re-oxidation of the reduced crystal, as already suggested by Onishi and Iwasawa [99]. As mentioned above, the bulk defects in sub-stoichiometric Ti02-x consist partially of Ti interstitials which show a high diffusitivity at elevated temperatures. When these interstitials appear at the surface, they can react with gaseous oxygen and form additional Ti02 (or TiaOb) structures. At high oxygen partial pressures, a complete reoxidation of the whole crystal can be achieved [9]. The reoxidation process has pronounced effects on the surface structure. The kinetics of the oxygen-induced restructuring process as well as the resulting surface phases depend on sample temperature, annealing time, gas pressure, and reduction state (i.e., 'age' or color) of the crystal. These parameters have been investigated in detail by Li et al. [9,98,100-102]. For example. Fig. 12 shows the effect of annealing in 1x10-^ mbar O2 at various

Fig. 12. STM images (500 A x 500 A) of a Ti02 (110) surface taken at room temperature. The surface was exposed to I8O2 (1 x 10-6 îbar) at (a) 500 K, (b) 520 K, (c) 550 K, (d) 660 K for 10 min, (e) 710 K for 15 min, and (f) 830 K for 20 min. Before each gas exposure the sample was sputtered and annealed in UHV at 880 K for 30 min which renders flat, (1x1) terminated surfaces. Fromref [102].

467

temperatures [102]. Before each gas exposure, a flat (1x1) terminated surface was prepared by sputtering and annealing in UHV. The morphology of the sample is very temperature dependent. For medium temperatures, surfaces are relatively rough with many small-scale features. Isotopically labeled 18O2 gas was used for the annealing excursions. In SSIMS and low-energy He+ ion scattering measurements, the signal from iÔ atoms can clearly be separated from the (naturally much more common) iÔ isotope in the crystal. All the surfaces in Fig. 12 showed an enrichment with iÔ, with a maximum at ^Ô surface concentration around 660 K. The structures that form for intermediate annealing temperatures (520 K - 660 K) are better seen in the small-scale image in Fig. 13a. They consist of small, (1x1) terminated islands and irregular networks of connected 'rosettes', i.e., six bright spots in a pseudohexagonal arrangement, as well as small strands. A model for the rosette network is shown in Fig. 13b. It consists of atoms in bulk-like positions with some atoms missing from the regular (1x1) structure. LDAbased first-principles calculations [102] have shown that such a rosette structure is stable. The same calculations also predict sizable relaxations. The rosette structure can be explained simply by the formation of (partially incomplete) Ti02 layers through a growth process where the Ti atoms come from the reduced bulk and the iÔ from the gas phase. The kinetics of the growth determines the relative concentration of the incomplete structures (the rosettes and strands) and the (1x1) islands on the surface. The surface is flat and (1x1) terminated when the growth is slow in comparison with surface diffusion processes. This can be achieved in various ways, either during annealing at high temperatures. Fig. 12f, or when the flux of one of the constituents is small (in very light samples with a small concentration of interstitials [9]) or at lower O2 background pressures. Conversely, on very dark crystals, or at intermediate temperatures, a substantial part of the surface can be covered with rosette networks. Note that 60% of the Ti atoms in the rosettes are 4-fold coordinated, whereas the Ti atoms exposed on the (1x1) surface are 5-fold coordinated. One should take this into account when preparing rutile (110) surfaces for surface chemistry experiments. The surface structure in Fig. 12e, obtained after annealing at 710 K in 1x10-6 mbar 1^02, shows the presence of (1x2) strands on otherwise flat, (1x1) terminated surfaces. This is in agreement with Onishi and Iwasawa's results described above [99]. From atomically-resolved images of strands connected to rosettes as well as UHV annealing experiments of restructured surfaces it has been shown that these strands also have the Ti203 structure depicted in Fig. 11 [98].

468

(b) Fig. 13. (a) An atomically resolved STM image (150 A x 150 A) of a surface prepared as in Fig. 1 Ic. Small (1x1) terminated islands and patches of connected pseudohexagonal rosettes are seen, (b) Atomic model (top and side view) for the oxygen-induced structure observed in (a). A bulk-terminated (1x1) island is shown on the right side and the unit cell is indicated. Small white balls are Ti atoms. Shadowed large balls represent oxygen atoms; darker shading indicates higher z-positions. The rectangle indicates the unit cell of the (1x1) structure. The network patch ('R') on the left side consists of an incomplete Ti02(l 10)(lxl) layer and contains only atoms at bulk position. The strands probably have a structure similar to the added Ti203 model in Fig. 1 Ic. From ref. [102] .

The dependence of the restructuring effect (as well as the type of (1x2) reconstruction) on the reduction state of the bulk was shown by Bennett et al. [16]. High-temperature STM studies were performed on two different Ti02 samples. On a dark blue/black crystal (that showed already evidence for CS plane formation) two different structures were observed (see Fig. 14.) The dark and bright strings in Fig. 14 were attributed to added Ti203 rows and added rows of a bulk-terminated Ti02 layer, but with bridging oxgyens at the center of the strands, respectively. The latter structure also appears crosslinked with partial 'rosettes' (Fig. 13b). Both, the added row structure and the rosettes are just incomplete Ti02 structures that form during the growth of additional Ti02(l 10)(lxl) layers. The authors have published impressive webbased STM 'movies' [104] (which can be viewed at http://www.njp.org/) of the growth process that results from the cyclic completion and new formation of the cross-linked added-row structures. In contrast, the dark rows (the Ti203 added rows) appeared relatively unreactive for additional growth.

469

148

^3

iLine Profile A-B

148

Line Profile C-D

A

D) 2

W^^ ^ 0

A

l'\

\d

0 10 20 30 40 60 60

0

20 40 60 80 100

Line Length (A)

Fig. 14. High temperature STM image of oxygen-induced features on a rather dark, nonstoichiometric TiO2(110) crystaL The crystal was exposed to oxygen (5.5 x 10"^ mbar, 833 K), stopped midreaction by removal of the oxygen overpressure, and imaged at the same temperature. The two types of strands with different apparent height (dark strands, DS and bright strands, BS) are attributed to the presence of different (1x2) structures. The added Ti203 rows account for the dark strings. The bright strands are interpreted as a slightly modified (addition of bridging oxgyens at the strands) added row structure (Fig. l i e ) . Additionally, bright rows and vacancies are visible on the (1 x 1 ) surface (marked BR and V, respectively). Line profiles are taken in the fast scan direction to minimize thermal drift and tip change problems. Reproduced from ref. [16], with permission...

3.3 Recommendations for surface preparation Although the Ti02(l 10)1x1 surface is considered the 'bestcharacterized', prototypical metal oxide surface, the above summary clearly shows that its atomic-level structure is significantly more complex than originally assumed. The recent STM results clearly indicate that both the oxidation conditions and the history of the TiO2(110) sample have significant influence on the morphology of the surface, the presence of strands, rosettes.

470

or CSP's. The variations in the surface structure with O2 pressure, crystal temperature and bulk defect density are so vast that one could suspect chemistry of the TiO2(110) surface could also be significantly variant for samples oxidized under different conditions. For example, the issue of whether water is molecularly or dissociatively adsorbed on TiO2(110) [63,64,72,105,106], may be significantly clouded in the literature because of studies in which the morphology of the surface was unknowingly disordered by the presence of the rosettes and/or strands observed recently by STM. This level of ambiguity may also permeate many other adsorption studies on TiO2(110). Guidelines of surface preparation of TiO2(110) can be extracted from this recent work [9,16,83,102,103]. If solely (1x1) terminated surfaces are desired, light blue crystals should be used. Annealing in oxygen will then result in stoichiometric, (1x1) terminated surfaces. On the other hand, more complex morphologies with a range of different coordination sites can be formed if dark crystals are annealed in oxygen. Oxygen plasma-treatment might be a good alternative to the more commonly used sputter/annealing cycles [83], but is not accessible in many standard surface science set-ups. The rich array of surface structures that can be produced on TiO2(110) may provide a playground for surface science experiments where the influence of different adsorption sites can be tested.

4. THE STRUCTURE OF THE RUTILE (100) SURFACE 4.1. The TiO2(100)(lxl) surface The rutile (100) surface has received considerably less attention than the (110) crystal face. The rules of autocompensation and creation of non-polar surfaces discussed above (3.1.1 allows a straightforward prediction of the stable surface termination (Fig. 15a). Again, the same number of Ti -> O as O -> Ti need to be broken, as indicated by the line in Fig. 15 b. This results in a strongly corrugated surface, with rows of bridging oxygen atoms at the outermost, [001]-oriented ridges (Fig. 15 a). Indeed a (1x1) terminated LEED pattern is observed on this surface after sputtering and annealing, and STM and noncontact AFM images are consistent with this model [107,108]. Several theoretical calculations have determined likely relaxations of the (1x1) surface [5,40,109,110]. In ref [109] different theoretical approaches and basis sets were tested. All these calculations agree in the general motions of the atoms, although the amount of relaxations differ somewhat. As expected from symmetry, no relaxations occur along [001]. In the [100] direction only the 5-fold coordinated Ti atoms show appreciable relaxations (downwards) [5]. Substantial relaxations occur along the [010] direction with the two-fold coordinated and the three-fold coordinated oxygen atoms moving in opposite direction of the five-fold and six-fold coordinated Ti atoms.

471 2-fold coordinated oxgyen s.foid coordinated titanium 3-fold coordinated ^^ _ , ^ oxgyen

A

(110) ^°S

f^ micro-

interstitial site: A

trigonal prismatic coordination: B

Fig. 15. (a) Geometry of the unreconstructed TiO2(100)(lxl) surface. This surface results when the same number of Ti -> O as O -> Ti bonds are broken in a bulk crystal, see dashed line in (b). The formation of the (1x3) microfaceted surface, originally proposed by Zschak et al. [118] is displayed in (b). Removal of the volume labeled y produces a stochiometric surface. In order to reconcile the reduced character of the (1x3) surface (observed in photoemission), the outermost bridging oxygen atoms (Y) are thought to be missing as well [119]. The transition from the (1x1) surface to the (1x3) microfaceted surface was proposed to occur via breaking the bonds a which leads to the relaxation indicated by the small arrow and the creation of an intermediate (1x3) phase [107]. (c) A surface x-ray structure analysis by Zajonz et al. has proposed a modified, heavily relaxed model (From ref. [124], with permission), (d) The glancing angle x-ray diffraction data that have lead to the original microfacet model have been re-evaluated by Landree et al. [125] and a new structure has been proposed. The octahedra model schematic representations in (d) show of the microfacet model (top) and the new model (bottom) (From ref [125], with permission).

In Fig. 15a, O atoms would move to the right and Ti atoms to the left. The net effect of these displacement is to increase the effective coordination of the 5-fold coordinated Ti atoms [109]. To my knowledge, no experimental data on relaxations of the TiO2(100)(lxl) surface exist. X-ray photoelectron and Auger electron diffraction were performed but are insensitive to the details of the surface structure [111].

472

4.2. Reconstructions 4.2.1. The microfacet model of the rutile TiO2(100)(lx3) surface In addition to the (1x1) terminated surface, a (1x3) reconstructed surface forms relatively easily upon annealing to high temperatures in UHV. It is partially reduced as is clearly evident from photoemission experiments [112115]. Initially, the observed reconstruction was interpreted as removal of every third row of the outermost, 'bridging' oxygen atoms [115]. Such a proposal was very much in line with the initially proposed [116] (and now largely abandoned) 'missing row' model for the (1x2) structure of the TiO2(110) surface (see Fig. 11a). The step structure on this surface may lead to a misinterpretation as additional (1x5) and (1x7) reconstructions as pointed out by Muryn et al. [115]. The first STM work on a reconstructed TiO2(100) surface was reported by Clark and Kesmodel [117]. A glancing angle X-ray diffraction and low energy electron diffraction study [118] suggested a 'microfaceted' model, see Fig. 15 b. Removal of the volume assigned with y creates facets of the lowest-energy, (110) plane. Note that, again, the same number of O -> Ti as Ti -> O bonds are broken. This results in a stoichiometric surface, with 2-fold coordinated bridging oxygen atoms at the outermost ridges, as is the case for the TiO2(110) surface. The unit cell in [010] direction is 3 times wider. Atomically-resolved STM images showed bright ridges with a geometry consistent with the microfaceted model [119,120]. The apparent height between top and bottom of the reconstruction was measured as 3 A with STM, instead of the expected 5A. This has been associated with tip effects [119,120]. Scanning tunneling spectroscopy (STS) measurements showed considerable difference between dl/dV spectra taken on and in between the bright ridges, respectively [119,121].. This was interpreted as missing oxygen atoms on the outermost ridges, (i.e., removal of the oxygen atoms labeled y in Fig. 15 c). This results in a surface termination with 3-fold coordinated Ti atoms in the outermost plane, consistent with the observation of a reduced surface in photoemission [115]. A photoelectron diffraction study of the Tiip level was performed by Hardman et al [114]. The PED curves of the (curve fitted) Ti3+ feature were evaluated. These Ti^"^ features are supposed to be come from Ti atoms located next to oxygen vacancies, and diffraction effects should give information about their surface geometry. Three different positions for vacancies on the microfaceted (1x3) surface were tested. The missing oxygen atoms at the outermost ridge of the (110) facets was determined as the most likely configuration [114]. In this context it is also interesting to note that the reduction state of the substrate plays a role in the formation of the (1x3) reconstruction. Almost stoichiometric, Nb-doped TiO2(100) films are thermally much more stable than reduced Ti02 substrates and do not reconstruct at temperatures where reduced Ti02 substrates already show a clear (1x3) structure [122].

473

A model involving 'discrete bond breaking' was proposed for the formation of the (1x3) surface [107,108]. Surfaces were prepared that exhibited both, the (1x1) termination and the ridges typical for the (1x3) surface. STM and non-contact AFM images showed an intermediate phase, which had a (1x3) symmetry, but did not possess the characteristics of the micro faceted structure. Raza et al. [107] suggested that the bonds labeled a in Fig. 15 are broken, which allows the Ti atoms to relax towards the other rows of bridging oxygen atoms. 4.2.2. Is the simple microfacet model valid? Despite the experimental results discussed in the previous section, it is presently not clear if the microfacet model for the (1x3) surface (Fig. 15b) is valid. In fact, recent evidence indicates that it may be a oversimplification. Theoretical calculations are not in agreement with the microfacet model. A tight-binding calculation derives a higher surface energy for the relaxed, micro faceted surface as compared to the (1x1) terminated one [123]. DFT ab initio calculation [110] also showed a considerably higher surface free energy. This implies that there would be no driving force for the reconstruction. These surprising results were attributed to the fact that stoichiometric surfaces has been considered in the calculations (i.e., the oxygen atoms y. Fig. 15c, were not removed). Experimental spectroscopic data show that the (1x3) surface is partially reduced. As was pointed out by Lindan et al., the Ti3+ state associated with the reduction should correctly be treated with spin-polarized DFT calculations [110]. A recent grazing incidence x-ray diffraction (GIXD) analysis shows strong lateral and vertical relaxations (1 A and more!) of the titanium and oxygen atoms in the top layer [124]. The coordination of the surface Ti atoms differ considerably from the simple microfacet model, see Fig. 15c. Threefold coordinated Ti atoms were found at the facet ridges, in agreement with previous work. In addition, various Ti atoms were found in different configurations, e.g., Ti in an oxygen bridge site, titanium B in a trigonal prismatic coordination, as well as an interstitial Ti site A. Obviously, this model deviates strongly from the simple microfaceted structure. The original GIXD data, which were the basis for the microfaceted model, were re-evaluated recently by Landree, Marks, Zschack, and Gilmore [125]. In ref [118] the data were interpreted using Patterson functions which show only interatomic vectors, not the 'true' atomic positions. The reevaluation was based on a technique known as 'direct method', in essence a Fourier transform of the measured data in connection with a search for possible (unknown) phases [126]. It was found that the microfacet model gave very poor agreement which the data as compared to a model that contained four Ti and six to eight O atoms in the surface unit cell. In this model, the Ti atoms reside in edge- and corner-sharing octahedral units, as opposed to the normal rutile structure which is composed of comer-sharing octahedra only

474

[001]

- • [010] - • [010] [010]

Fig.l6. Surface termination of the rutile TiO2(001) surface. Only one possibility exists to cut a Ti02 crystal in this direction, see the side view at the left side. Surface Ti atoms are 4fold coordinated and surface O atoms 2-fold coordinated.

(see Fig. 15d). The reconstructed structure is rationalized as a standard configuration for non-stoichiometric defects such as CS planes, see Fig. 10. The reduced states observed with spectroscopic measurements would then be accommodated similar as in the bulk. The surface unit cell of the TiO2(100)(lx3) structure is quite large with many atoms, hence it is no surprise that it is hard to determine conclusively the exact surface geometry. While the electron diffraction [114], STM and AFM results [107,108,119-121], together with ESDIAD measurements [127] are consistent with the microfacet model of the (1x3) surface, neither of these techniques gives direct evidence of surface geometries. The (1x1) and (1x3) reconstructions provide a convenient system for site-sensitive surface chemistry experiments, because one can reversibly cycle a TiO2(100) crystal between the (1x1) surface and the reconstructed one [128,129]. Hence, it would be quite important to have additional experimental as well as theoretical support for one of the models currently suggested and depicted in Fig. 15.

5. RUTILE (001) There is only one way to cut a rutile crystal in (001) direction (Fig. 16.) Although this creates a non-polar, autocompensated surface, it does not represent a low-energy configuration. This becomes clear immediately when reviewing the coordination of the surface atoms. All the Ti atoms are 4-fold coordinated, and all the O atoms 2-fold coordinated. Hence the number of broken bonds on this surface is higher than on the other low-index rutile surfaces discussed so far. Consequently, the (001) surface has a high surface

475

Fig. 17. A TiO2(001) surface after annealing at a very high temperature (1300°C, 1 h in UHV). The image has been taken with an optical microscope. It shows lines due to slip along certain crystallographic directions. From ref [133], with permission.

energy and tends to facet. Based on LEED studies, (Oil) and (114) facets have been identified by Tait and Kasowski [130] and Firment [131]. Several additional crystal planes were identified in an STM/LEED study by Porier et al. [132]. . The image in Fig. 17 shows the drastic changes that can occur upon very high-temperature annealing. The micrograph has been taken with an optical microscope after heating a TiO2(001) crystal for Ih at 1300°C in UHV [133]. The lines are caused by slip. A non-equilibrium structure was observed after rapidly quenching a TiO2(001) crystal from a similar high-temperature anneal [134]. 6. VICINAL SURFACES Vicinal surfaces of TiOi have not been studied extensively. A study of Na adsorption on a stepped (441) surface was reported by Onishi et al [135]. Unpublished experiments in my own laboratory with a similarly cut crystal showed macroscopic faceting upon annealing. The most detailed investigation was performed recently on a TiO2(210) surface [136]. In a formal sense, TiO2(210) lies midways between (110) and (100), and is the most simple vicinal surface. Atomistic simulations, based on Coulombic interaction between ions and a short-range repulsive interaction, predicted an asymmetric sawtooth-like structure of the surface, consisting of {110} nanofacets. The width of each nanofacet is 1.5 times the width of the surface unit cell of the (110)(lxl) structure (i.e., 3a/V2). The nanofacets terminate with a row of Ti

476

atoms carrying bridging oxygen atoms. The surface energy of this structure is predicted to be 2.07 J/m^. (This is to be compared by a surface energy of 1.78 J/m2 for (110) derived using a similar calculation [136].) STM images showed a (1x1) terminated surface that could be consistent with this structure, although the interpretation was again made difficult by balancing electronic effects with the very strong corrugations of this surface.

7. ANATASE SURFACES Commercial titania powder catalysts are a mixture of rutile and anatase (e.g., the most often-used Degussa P-25 contains approx. 80 to 90% anatase and 10 to 15% rutile [137]), and, for reasons that are not completely understood, anatase is the photocatalytically more active form. It also shows enhanced activity in (non-photoinduced) catalysis, and behaves different than rutile in gas sensing devices. High-purity titania powder catalysts are typically made in a flame process from titanium tetrachloride [137]. Many additional synthetic techniques processes have been applied [138-142]. The shapes of the crystallites vary with preparation techniques and procedures. Typically, the (101) and the (100)/(010) surface planes are found, together with some (001) [6]. Several theoretical studies have predicted the stability of the different low-index anatase surfaces [6,143,144]. The (101) face is the thermodynamically most stable surface. Calculated surface energies are 0.52 and 0.81 J/m^ for the (relaxed) anatase (101) and (001) surfaces. Compared with the (also GGA-calculated) value for rutile (110) surface [145] it appears that the anatase (101) surface is even the most stable surface of all the Ti02 polymorphs! Experimental investigations on single-crystalline anatase, are just starting. Meaningful surface science investigations necessitate singlecrystalline samples. While rutile crystals are readily available, sufficiently large and pure anatase crystals are more difficult to obtain. Because anatase is a metastable phase, it transforms into rutile at relatively low temperatures [139], with the transition temperature dependent on impurities, crystal size, sample history, etc. However the recent progress in synthesizing singlecrystalline anatase samples with high purity [138] as well as the successful growth of epitaxial thin-films on appropriate substrates [7,146,147] has rendered the first surface science investigations of this material. 7.1. Anatase (101) Two reports on the structure of anatase (101) surfaces have appeared very recently [7,8]. The (101) surface on an anatase sample grown by chemical transport [138] showed a (1x1) surface after mild sputtering and annealing [8]. A mineral sample was used in an STM study by Hebenstreit et al. [7]. In order to avoid the contaminations in this natural single crystal, a 700 A thick, epitaxial film was grown on the surface. Sputtering and annealing again

477

film was grown on the surface. Sputtering and annealing again produced a (1x1) termination in LEED. The surface has only pm symmetry, giving rise to a preferential orientation of step edges (Fig. 18a). Based on the rules of autocompensation for step edges a detailed model for steps is given in Fig. 18 [7]. Titanium atoms at the terraces they have five-fold and six-fold coordination, and titanium atoms at the step edges are four-fold coordinated. These have a higher reactivity against gas adsorption [7]. Two-fold coordinated oxygen atoms are located at the ridges of the saw-tooth like structure. According to ref. [6], these atoms relax. The tunneling site in the atomic-resolution images probably extends across both, the two-fold coordinated oxygen atoms and the five-fold coordinated Ti atoms [7]. In an image taken with a higher tunneling current (12 nA), where the tip was probably closer to the surface, the 2-fold coordinated oxygen atoms are distinguished as single features [7]. In correspondence to the rutile (110) surface, one might expect that the twofold coordinated oxygen atoms are removed easily upon annealing in UHV, giving rise to point defects. Several imperfections with atomic dimensions are identified in the atomic-resolution images [7]. At this point it is not clear which one of these features, if any, are indicative of oxygen vacancies. Their number density is very small, certainly smaller than the usually quoted 5 - 10 % [53,64] on rutile (110). This would be in agreement with the (calculated) low surface energy of the (101) surface [6]. Calculations of the electrostatic potential by Woning and van Santen also predict that the rutile (110) surface can be reduced easier than anatase [148]. 2-fold^

5-fold Ti

r

Fig. 18 (a) STM results of an anatase (101) single crystal. The monoatomic terraces terminate with step edges that run predominantly in certain preferred orientations, (b) Atomic models (side and top view) of the anatase (101) surface. From ref [7], with permission..

478

7.2. Anatase (001) As shown in Fig. 19, the stable, autocompensated anatase (001) surface exhibits exclusively 5-fold coordinated Ti atoms, as well as two-fold and threefold coordinated oxygen atoms. Calculations show that the corrugation increases somewhat upon relaxation, from 0.82 A to 0.92 A [6]. The most detailed structural investigations on this surface so far have been taken on thin films, grown epitaxially on the SrTiO3(001) substrate [146,147]. The SrTiO3(001) surface shows a very good lattice match with the anatase (001) surface, but a poor one with the rutile phase. Their anionic sublattices bear substantial resemblance despite their overall crystallographic dissimilarities [149] As Ti02 is deposited, heteroepitaxial growth can be regarded as a continuous formation and extension of the oxygen atom network from the substrate into the film. Within this oxygen sublattice the (relatively small) Ti cations become arranged in their appropriate sites. Formation of interfaces where the oxygen sublattice continues and the metal cation sublattice changes abruptly is often exploited for thin film heteroepitaxy of metal oxides [150]. On SrTiO3(001) epitaxial, stoichiometric anatase thin films of high crystalline quality have been grown by several techniques [149,151]. Recent, unpublished results show that MOCVD-grown films are stable up to 1000°C [152]. During the growth, the anatase (001) films are (1x1) terminated. The (1x1) termination has been confirmed ex-situ with LEED and XPD by Herman et al. [146]. The surface was only outgassed at 100°C and not sputtered. Hence, a carbon coverage of 0.6 monolayers was estimated based on XPS measurements. The surface was well ordered and the structure in XPD was fully consistent with the bulk-like termination depicted in Fig. 19. A twodomain (1x4) reconstruction formed on a similar sample after sputtering and annealing the (1x1) surface in UHV [147]. Based on angle-resolved mass5C-Ti

2C-0

(001) ,[010] • [100]

^ vv.^., V . - • /

I

^ ^^ QQj

anatase (001) (1 x4), microfacet model

Fig. 19 The surface structure of teh anatase (001)(lxl) surface and the proposed model for the (1x4) reconstruction found after annealing a clean anatase (001) surface in UHV (after ref [147], with permission).

479

spectroscopy of recoiled ions (AR-MRSI) a 'microfaceted' model was proposed. In this model (103) facets are exposed which contain two-fold oxygen and both four-fold and five-fold coordinated Ti atoms, see Fig. 19. Such a model resembles in many ways the microfacet model for the rutile (100) (1x3) surface discussed in section 4.2. It remains to be seen if this model will be confirmed by additional experimental and theoretical work. There is no reason to assume that epitaxial anatase films, provided they are thick enough, would have a different surface structure than an single crystals. In fact, LEED studies on single-crystalline anatase samples confirm the results by Herman et al. A (1x1) termination was observed after sputtering and annealing in oxygen at 400°C an iron contaminated mineral anatase sample [153 ]. A single crystal grown with chemical transport [138] showed a (1x1) termination initially and a (1x4) reconstruction after sputtering and annealing to 900K [154]. 8. CONCLUSION Much has been learned about the surface structure of the titanium dioxide system in recent years. Two somewhat contradictory lessons can be drawn from all this work: simple approaches work well for obtaining a first guess of surface structures and oxide surfaces are even more complicated than anticipated. It is comforting that the elemental rules for predicting surface terminations outlined in section 3.1.1. work so well for predicting the structure of the (1x1) terraces and the step edges of all the orientations of both rutile and anatase. The extensive theoretical work has helped to refine the understanding of surface relaxations, and the level of detail on the atomic geometry of the TiO2(110) surface is certainly comparable to that of certain elemental semiconductors or metals. On the other hand, scanning probes techniques have unraveled a very rich picture of surface structures. One interesting theme is the interplay between surface structure and bulk defects; the reduction state of the crystal is quite important for the presence of different structural features under exactly the same preparation conditions. It remains to be seen if such a behavior is also present on other metal oxides. Total-energy calculations have helped enormously to confirm models for surface reconstructions, and have acted as a warning sign when models derived from experimental results were too naive. It appears that, again, bulk structures should act as a guideline for devising models for new surface reconstructions. The expanding data base has made rutile the model system for metal oxides. Nevertheless, there are still many open questions concerning the crystal structure of rutile surfaces as pointed out throughout this Chapter. One interesting aspect is the advent of surface studies on anatase. From the recent

480

progress in synthesizing single-crystalline anatase samples with high purity as well as the successful growth of epitaxial thin-films on appropriate substrates, I would expect a rapid increase in the interest in anatase in the near future.

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Chapter 12

The Anisotropy of Metal Oxide Surface Properties G.S. Rohrer Department of Materials Science and Engineering, Carnegie Mellon University, Pittsburgh PA, 15213-3890, U.S.A. 1. INTRODUCTION The surface properties of metal oxides influence the rates of heterogeneous chemical reactions, the growth of heteroepitaxial films, and the sintering of particles during the consolidation of ceramics. Just as certain bulk properties of crystalline materials are anisotropic, surface properties also depend on orientation. The dependence of surface properties on orientation can be rationalized by recognizing that the atoms on crystallographically distinct facets have different coordination environments, as illustrated schematically in Fig. 1. This same figure also illustrates that in binary and more complex materials, surfaces with identical orientations can be terminated by different atomic layers with distinct compositions. Throughout this chapter, we shall take the "character" of a surface to be defined by its orientation, {hkl}, and its atomic termination layer. The relationship between surface properties and surface character allows the macroscopic properties of surface active materials to be manipulated by controlling the distribution of the exposed surfaces. For particulate materials, this distribution is determined by the crystal habit. The character of surfaces exposed by a dense polycrystal can be controlled by the introduction of some crystallographic texture and thin film materials can be oriented through heteroepitaxial seeding. While techniques for structural control are generally available to the material scientist, what is less frequently known is how the physiochemical properties of a material vary with surface character. In this chapter, efforts to measure the surface properties of oxides as a function of character will be described. Section two of this chapter reviews some specific examples of the relationship between surface character and properties that have been derived from studies of particulate oxides, single crystals, and thin films. In section three, the factors that determine surface morphology and the range of achievable orientations are described. The factors that influence surface stoichiometry and the composition of the termination layer are described in

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section four. Two specific examples illustrating anisotropic reactivity are reviewed in sections five and six, which deal with partial oxidation reactions on M0O3 and photochemical reactions on Ti02, respectively. 2. SURFACE CHARACTER.PROPERTY RELATIONSHIPS 2.1 Particulate oxidation catalysts One clear piece of evidence for surface character-property relationships lies in the observation that particulate catalysts with the same composition, but different crystal habit, have different reactivities. This phenomenon, sometimes referred to as a catalytic anisotropy or a structure-sensitive catalytic reaction, has been observed in a variety of Mo, V, and W containing materials [1]. Because the catalytic anisotropy of a-Mo03 has been studied in relatively greater depth, this exemplary case is described below [1-10]. The orthorhombic (a) phase of M0O3 has the highly anisotropic layered structure illustrated in Fig. 2 [11], Because of the weak, secondary bonding between the sheets of corner and edge-sharing MoO^ octahedra perpendicular to [010], the crystals are micaceous. The most prominent facet on the growth form is {010}. The crystals are bounded laterally by {001} and {hkO} facets, where the h/k ratio is about 2. In some case, the {001} and {hkO} facets may be replaced by {101} and {100} facets, respectively. The Mo atoms on the (010) surface have their full complement of six oxygen nearest neighbors. However, termination of the bulk structure in other

487

[010]

Fig. 2. The structure of a-MoOs crystals, (a) The structure is built of double octahedral layers stacked along [010]. The spacing between adjacent double layers is 6.9 A. (b) A projection on the (010) plane shows two superimposed rectangular nets of comer-sharing octahedra. The shaded net is below the unshaded one. (c) A typical growth form of aM0O3. All indexing is with respect to space group Fbnm.

directions leaves the surface Mo atoms on lateral facets with fewer than six neighboring O. Investigations of molybdenum trioxide's catalytic anisotropics have been based on the hypothesis that this difference in coordination number will lead to differences in the catalytic properties of the different facets. This hypothesis has been experimentally tested in the following way. Chemically identical particulate samples were prepared in different ways so that they exhibited different characteristic particle habits and aspect ratios. Under the assumption that all of the particles in a given sample have approximately the same shapes, the fractional contribution that each type of facet makes to the total surface area of each specimen was determined by microscopic inspection. The role of each facet in a given reaction can then be inferred by comparing the properties of specimens with different distributions of external surfaces. Based on such studies, it has been concluded that the lateral facets are active for the oxidation of methane to formaldahyde [10] and the oxidative anmionalysis of toluene [7], while the (010) facet is active for the conversion of methanol to formalahyde [1]. Studies of the oxidation of propene to acrolein illustrate that it is not always easy to relate overall activities or selectivities to the presence of a single face [3, 5, 8, 9]. Since the overall reaction is composed of several elementary steps, it is possible that different steps occur on different facets. For example, it has been proposed that the mechanism for the oxidation of propene to acrolein begins with the activation to an allyl intermediate on a lateral facet and ends with the addition of O on a basal facet [5]. The (210) facet, which is thought to consist of terraces with (010) character and ledges with (100) character, should be able to perform both elementary steps. This explanation has been used to rationalize the observation that the (210) surface is especially active for the conversion of propene of acrolein [9]. Using similar

488

experimental techniques, the functionality of specific facets in other layered materials, including vanadium pentoxide [12] and vanadyl pyrophosphate [13], has been identified. While the knowledge gained from these studies is valuable, a number of limitations should be mentioned. First, the conclusions rest mainly on the observed macroscopic shape of the crystals and little is known about the morphology and structure of the individual facets on the particles. Furthermore, the surface morphology of the catalyst is expected to change during reactions, as stoichiometry compensating defects are continuously created and annihilated. Finally, it should be noted that conclusions from studies of particulate materials are not always consistent with observations from other experiments. For example, by comparing the properties of different particulate samples, Tatibouet and Germain [2] identified the (010) facet of M0O3 as being responsible for the conversion of methanol to formaldehyde while Farneth et al. [14] concluded, on the basis of temperature programmed desorption studies, that this facet was largely inactive for the partial oxidation of methanol. We shall see in section 5 that such discrepancies might be explained by morphological changes that occur on the surface of the oxide while in service. 2.2 Single crystal surfaces Perhaps the most convincing demonstrations of structure sensitivity come from highly controlled single crystal experiments in ultrahigh vacuum [15-18]. For example, Vohs and Barteau [15] compared the ability of the ZnO(OOOl) and ZnO(OOOT) to abstract protons from alkynes. While the alkynes (C2H2, CH3CCH, C6H5CCH) decomposed on the (0001) Zn-terminated surface, the (000T) O-terminated surface was observed to be inactive for dehydrogenation. Barteau's [16-18] group conducted parallel studies on Ti02 surfaces in UHV. Flat TiO2(001) orientated crystals can be induced to facet into nonplanar configurations terminated by {101} or {114} facets, depending on the thermal treatment. This phenomenon was exploited to show that when either methanol [16] or acetic acid [17] was reacted with Ti02, the distribution of products was governed by whether the crystal was terminated by {001}, {101}, and {114} surfaces. While the effect of orientation was convoluted with the level of surface oxidation, it was still possible to draw some conclusions about the property-character relationships for these surfaces. For example, the fourcoordinate Ti"^"^ cations found on the {114} faceted surface are capable of coupling carbon containing species to form higher order products (for example, dimethyl ether from methanol and acetone from acetic acid) while the fivecoordinate cations on the {101} faceted surface initiate only dehydration and dehydrogenation reactions. The rationale for this difference is that the lower coordinate cations on the {114} surface are able to adsorb both of the necessary

489

reactants to a single site while the cations on the {101} surface can adsorb only one [18]. 2.3 Thin films As noted briefly in the introductory section of this chapter, specifying only the orientation of a surface can sometimes lead to an ambiguity with respect to the chemical composition of the termination layer. Consider, for example, a crystal with the perovskite structure in the (100) orientation (see Fig. 3). The structure can be terminated either on the AO plane or on the BO2 plane. The composition of the termination layer has been studied most extensively for SrTiOg, which is commonly used as a substrate for the growth of superconducting cuprate phases such as YBa2Cu307. SrTiO3(100) substrates annealed at about 1000 °C exhibit two characteristic types of terraces. Terraces with smoothly curved edges are associated with SrO terminated surface; Ti02 terminated terraces have polygonized edges made of and step segments [19]. It has been demonstrated that it is possible to prepare surfaces that have a uniform termination layer. Ti02-terminated (100) surfaces can be prepared by chemical etching [20, 21] and SrO surfaces can be prepared by depositing a monolayer thick film of SrO [22, 23] or Sr2Ti04 [24]. It has been observed that the SrO and Ti02 surfaces stabilize different polytypes of SrCu02 during thin film deposition [24] and that YBa2Cu307 films are less susceptible to decomposition when deposited on SrO terminated substrates [23]. While SrTi03 is the most well-studied example, the influence of termination layer on film growth has been demonstrated to occur in other systems as well. For example, [0001] oriented sapphire can be terminated by three distinct planes. One is oxygen terminated and the other two are Al terminated. Bench et al. [25] have observed that sapphire (0001) surfaces heated

AO layer

BO2 layer (001)

B Fig. 3. Schematic illustration of the perovskite structure. Note that surfaces perpendicular to (001) can be terminated by and AO or a BO2 layer.

490

^ A

#

»

111 #

». •

*

*

Fig. 4. Scanning electron microscope image of Cu.O deposits on the Al.O.lOOOl) surface. The different particle sizes and morphologies correlate to different surface compositions [25].

for 6 h in air at 1400 °C exhibit a combination of these possible terminations on a sample with a constant average orientation. When CU2O islands were deposited on these surfaces, two distinct island sizes and morphologies were observed that correlated to differences in the terrace termination (see Fig. 4). When Ti02 was deposited on these surfaces, one of the terrace types promoted the growth of an unanticipated (and unidentified) phase that was suspected to be a reduced form of titania or an Al-Ti-0 ternary phase. The results cited above were selected to clearly illustrate the influence that the character of a surface has on its properties. To be more specific, the ability of an oxide surface to catalyze specific reactions or nucleate specific phases is known to depend on both the surface orientation and the termination layer. We now turn to the question of what determines the surface character and what characters are reasonably attainable. 3. FACTORS DETERMINING SURFACE MORPHOLOGY In equilibrium, the morphology of a surface is determined by the anisotropy of the surface energy. While the surface morphologies that we observe in most practical situations are not in global equilibrium, local equilibrium can be achieved at the intersection of two facets or at the point were a grain boundary intersects the surface. Observations of surfaces in local equilibrium give us the

491

opportunity to determine the anisotropy of the surface energy. Conversely, knowledge of the surface energy anisotropy allows us to specify possible morphologies. The two important conclusions from this section are that because of anisotropy, it is usually not possible to realize all surface orientations and that experimental observations can be used to determine the range of allowable orientations. The equilibrium shape of a free particle is determined by the surface energy, which for crystalline solids is usually anisotropic. The microscopic origin of this anisotropy can be understood if one considers that the work required to create a new surface should (ignoring relaxations or reconstruction) increase with the number of pair-wise interatomic bonds that must be broken. Those surfaces with a higher density of broken bonds will have a higher relative energy. Therefore, we can assume the surface energy is a function of the surface normal, Y(n), and for any element of the crystal surface, dA, the surface energy is y(h)dA, where fi is normal to the element dA. For a single crystal with fixed volume in thermodynamic equilibrium, the shape is that which minimizes the total surface energy, E [26]: E = JY(n)dA

(1)

For an isotropic material, the equilibrium shape is a sphere. For an anisotropic material, the equilibrium shape will depart from sphericity so that crystallographic orientations with relatively low values of y will be exposed. Based on Eq. 1, it has been demonstrated that for a crystal in equilibrium, the ratio of the energies of two orientations (y^ and Y2) is equal to the ratio of the perpendicular distances from the center of the crystal to the facet (Z^ and l^) [27]:

h

yi

In other words, the lower the surface energy of the facet, the closer the facet is to the center of the crystal and the greater the surface area of the facet. This is the basis of the Wulff [28] construction, which specifies the relationship between the surface energy of a crystal and its equilibrium shape. To apply the Wulff construction, we begin by imagining that a normal vector represents each distinct crystallographic orientation and the magnitude of each normal vector is equal to the surface energy of the orientation that it represents. Further, we assume that the vectors share a common origin. If a perpendicular plane is situated at the end of each vector, then the inner envelope of the planes defines the equilibrium crystal shape.

492

The Wulff construction is important because it allows the relative energies of different observed surfaces can be determined by direct microscopic measurement of crystals (or voids within crystals) with an equilibrium shape. By measuring the distances from the center of a crystal shape to the various facets (//), the relative surface energies (7/) can be computed using Eq. 2. While this method has been applied to many crystals, applications to oxides have been less frequent [29-31], Most recently, two groups reported results for AI2O3 (see Table 1) [32, 33]. Underpinning such studies is the assumption that the particle (or small cavity) is in equilibrium. To reach this state, it is necessary to have sufficient mass transport and surface attachment kinetics and, in the case of faceted particles, step producing defects that reduce the barrier for the nucleation of new layers [34]. Table 1. Experimentally measured surface energies for AI2O3 Plane (0001) (1011) (1012) (1120) (1123)

1600 °C Ref. 32

Ref. 33

1800 °C Ref. 33

1.0

1.0

1.0

1.07±0.02 1.05±0.02 1.22±0.05 1.06±0.02

0.947±0.016 0.855±0.017 0.947±0.026 0.957±0.026

1.042±0.019 0.950±0.030 1.080±0.017 1.029±0.016

That the data in Table 1 are not in agreement illustrates the difficulty of making such measurements. The constraint of global equilibrium is particularly challenging. An alternative method for experimentally assessing the anisotropy of the surface energy has been developed based on the assumption of local equilibrium between free surfaces and grain boundaries at thermal grooves formed by surface diffusion [35]. The results of this measurement are illustrated in Fig. 5. While the observed anisotropy is not large, the effects on surface morphology are significant. Furthermore, it should be pointed out that there are geometric limits to observable anisotropics for nonperpendicular surfaces in equilibrium experiments. For example, if the energy of the (110) surface of a cubic crystal were more than V2 times the energy of the (100) surface, then the surface would lower its energy by faceting into (100) and (010) surfaces that would meet along lines in the [001] direction and form ridges. Therefore, the energy of the (110) orientation has an upper limit of V2YJQQ. Similarly, if the energy of the (111) surface were more than V3 times the energy of the (100) surface, it would lower its energy by faceting into trigonal pyramids bound by (100), (010), and (001) facets. In fact, at lower temperatures where the anisotropy of the surface energy is expected to be higher, the MgO(lll) surface has been observed to break up into (100) facets [36]. Calculations for MgO

493

1.08 T 1.06 4o o

>=^ 1.04 +

(100)

3 ^

90

120

angle, Fig. 5. Plot of the relative surface energy of MgO in air at 1400 °C, around the perimeter of the unit triangle, from (100) to (111), then to (110), and back to (100).

suggest an anisotropy so high that only {100} facets should be observed [37-40]. While this is consistent with relatively low temperature vacuum experiments, it is not consistent with higher temperature observations in air [35, 41]. In the present context, where our interest is in controlling surface character, it is important to note that surface energy anisotropy frequently leads to missing crystallographic orientations. For example, consider the two-

i (a)

«2 (b)

Fig. 6. (a) A hypothetical two-dimensional equilibrium crystal shape. A continuous range of orientations exists between hi and n^, the complex facet. Between he and %, there is a range of missing orientations (n^ is an example), (b) The surface of a crystal whose orientation is macroscopically constrained to a missing orientation will break up into facets with orientations n^ and n2.

494

dimensional Wulff shape illustrated in Fig. 6. The surfaces of constrained crystals with orientations between fi^ and n^ will be flat, as this will minimize the total surface energy. However, surfaces with orientations between n^ and ^2 will be unstable with respect to faceting and in equilibrium, the surface will be covered with tent-shaped ridges bounded by surfaces with orientations fi^ and ^2- The relative proportion of the two surfaces will depend on how far the orientation is from either of the stable orientations. If it is very close to ^2' then the surface will be composed of large terraces with small steps whose orientations are n^. On the other hand, if the orientation is half way between the two stable surfaces, the surface will be composed of equal parts of the two orientations. The appearance of facets on surfaces with unstable orientations is one conmion manifestation of anisotropy. Surfaces with macroscopic orientations that are missing from the equilibrium shape can be prepared at low temperature by polishing. When heated, such surfaces form microscopic facets while maintainihg their global orientation. The ranges of stable and unstable orientations can be mapped out by examining the surfaces of individual grains in a polycrystalline ceramic that has been polished and then annealed at high temperature. Data for the orientation stability of MgO at 1400 °C in air are shown in Fig. 7. The orientations of more than 100 grain surfaces were determined by electron backscattered diffraction. AFM images of the same surfaces were used to discriminate flat from faceted surfaces (see Fig. 7).

(100)

(110)

Fig. 7. An orientation stability map for magnesia at 1400 °C. Each point corresponds to the orientation of an observed grain. Faceted orientations are marked with an x and smooth orientations are marked with an empty square. An AFM image of a typical faceted surface is shown in the upper left hand part of the figure. The facets are approximately 500 A high.

495

Because of the relatively high energy of the polar (111) surface, this orientation and orientations within about 20 ° decompose into lower energy complex facets [35]. The relative stability of the different orientations is conveniently represented on a stereographic projection, as proposed by Cahn and Handwerker [42]. These orientation stability diagrams are composed according to an analogy with isothermal sections of ternary phase diagrams. Ranges of orientations that are flat (or form part of a continuously curved surface on a particle) are represented as a shaded region and are analogous to a solid solution. Orientations that break up into two separate surfaces are placed on tie lines that connect the two stable facets, in analogy with two-phase regions; orientations that break up into three facets make up empty polygons whose vertices are stable

Fig. 8. (a) Orientation stability diagram for SrTiOg at 1200 °C, in air. There are a range of stable orientations near {100}, The {110} (black circles) and {111} (white triangles) orientations are also stable. Typical AFM images of surfaces from (b) a stable region, (c) a two facet region, and (d) a three facet region are illustrated below. Each shows a 5|xm x 5|im topographic image.

496

orientations. As an example, Fig. 8 illustrates an orientation stability diagram for SrTiOg annealed at 1200 °C in air for 6 h. The observations indicate that the {100}, {110}, and {111} surfaces are stable under these conditions. In addition, there is an extended range of stable surfaces vicinal to {100}. We can therefore say that the stable shape of SrTiOg (after heating for 6 h in air at 1200 °C) has continuously curved surfaces at each of the cube faces which meet flat {110} and {111} surfaces at sharp edges. In practice, it is much easier to derive information on the equilibrium crystal shape from observations of orientation stability than from small crystals or cavities. This is because orientation stability measurements require only local equilibrium at the points where facets meet, whereas global equilibrium is required for equilibrium crystal shape measurements. However, because only a portion of the equilibrium form is observed at any time, the relative energies of different continuously curved regions of the equilibrium crystal shape can not be determined. It is, however, possible to estimate the relative energies of any curved and flat regions of the equilibrium form that intersect [43]. Consider the example discussed above where and orientation h^ breaks up into to other orientations, fi2 and h^. Under the assumption of local equilibrium at the intersection between h2 and n^, then the Herring [44] equation gives the following relation: ^ = cos0-—^sin0 Yc

Yc de

(3)

If the anisotropy is small enough that the differential term can be ignored, then the ratio of the surface energies is simply related to the cosine of the angle at which they meet. Note that one criterion for the application of Eq. 3 is that the facets are in local equilibrium. For this to be true, there must be no net growth or evaporation. If this is true, then thermal grooves at surface-grain boundary intersections should maintain a quasistatic profile and increase in width with the one quarter power of time [45]. 4. FACTORS INFLUENCING SURFACE STOICHIOMETRY The factors influencing the surface composition of pure materials are described in this section. While the factors influencing surface morphology and surface composition are treated in different sections, it should be emphasized that they are closely related. We begin by considering the solid-vapor equilibrium and how this influences surface and bulk stoichiometry. The equilibrium defect structure, like the equilibrium morphological structure, might not always be obtained. Regardless, it remains useful to know the destination state of a surface whose composition is evolving at elevated temperature. The section concludes

497

with a brief consideration of the composition of oxide surfaces produced at low temperature by cleavage. 4.1 The solid-vapor equilibrium We begin by noting that the bulk stoichiometry of an oxide is a function of the temperature and composition of the surrounding atmosphere [46]. For a hypothetical metal oxide (MO), we can write the appropriate defect chemical reactions in Kroger-Vink notation and equilibrium constants for congruent evaporation (kô)? the production of vacancies by evaporation (k^ and kg), and Shottkey defect formation (kg): MM + Oo = I/2O2 ^,, + M f,,

PM PO2'^ = kMo

(4)

MM = M^^^ + VM*

PM[VM*] = kM

(5)

Oo = 1/20, ,,^ + Vo*

Po2'"[Vo*] = ko

(6)

0 = VM* + VO*

[VM*][Vo*] = ks

(7)

For the purpose of simplicity, the vacancies are taken to be neutral and we do not consider the Frenkel or anti-Frenkel processes. In both cases, extending the analysis does not alter the important parts of the results. From Eqs. 5-7, it can be shown that: [Vo*]=koPo2-^^'

(8)

[VM*] = {^,IK)Po2"

(9)

The equations are represented schematically in Fig. 9. There are several important conclusions to draw from this simple analysis. First, the defect concentration is a function both of temperature (through the equilibrium constants) and the partial pressure of oxygen. Second, there is a single composition of the gas phase where the compound is stoichiometric. Furthermore, the range of accessible compositions is limited by either the appearance of new phases or by experimental limitations. For example, the ranges of stoichiometry available to Fe304 are limited at high and low partial pressures of oxygen by the formation of Fe203 and FeO, respectively. For Ti02, the appearance of the Magneli phases at low partial pressures of O2 limit the maximum concentration of oxygen vacancies at 1000°C to x = 0.008 in Ti02.x [47]; the upper limit of oxygen pressures is determined only by the strength of the vessel that contains the experiment.

498

O

[Vnn*]

X''' [Vo*]

O

o o

o >

lnpo2 Fig. 9. Schematic graph illustrating how the composition of the surrounding atmosphere influences the stoichiometry of a metal oxide.

If we take the total pressure to be p^ + PQ2, we can use Eqs. 4-6 to show that the total pressure is: Piot ~ P02 "•" ^MO Po2

(10)

By minimizing Eq. 10 with respect to Po2' it can be shown that there is a minimum partial pressure of oxygen at which the compound can be in equilibrium. In this case: P02,

,= l/2(2kMo^)"^

(11)

As long as the partial pressure of O2 is greater than PQ2, min' the stoichiometry of the compound will adjust to equilibrate with its surroundings. Below this minimum, some of the crystal will have to sublime to establish the minimum vapor pressure. In a dynamic vacuum, where subliming material is constantly pumped away, this is clearly not possible. Under such conditions, the stoichiometry of the compound is determined by kinetic factors and it might not be possible to conduct experiments with reproducible stoichiometrics. While the preceding statements apply to the bulk of the crystal, the surface of the crystal is subject to the same constraints with some additional complexity. Specifically, because different point defects are likely to have different free energies of formation, there is an enhanced concentration of one of the two defects in the near surface region and a compensating layer of charge on the

499

surface [48, 49]. Suppose, for example, that it is easier to create O vacancies than M vacancies. The enhanced concentration of O vacancies that would result from this situation would lead to a net positive charge in the near surface region that would be have to be compensated by negative charge on the surface. It has been pointed out that the amount of charge that it is possible to store in the subsurface region will be limited by the capacity of the surface to accommodate the compensating charge [50]. As with the bulk situation, equilibrium surface defect concentrations will be established only at a fixed and sufficiently high partial pressure of oxygen. One of the more well-documented cases that illustrates the effect of the annealing atmosphere on the surface composition concerns the Al2O3(0001) surface [51]. After heating in air, 1 X 1 LEED patterns are observed. High temperature annealing (1350 °C) in UHV leads to the loss of O and the formation of an Al terminated surface that has a (VSl X V31)R±9° LEED pattern. The model for this reconstruction is two pure and nearly perfect Al planes on an O terminated Al2O3(0001) surface. The same structure can be produced by depositing Al metal from the vapor phase onto an Al2O3(0001) surface with the 1 X 1 structure [52, 53]. Therefore, and alternate way to view this reconstruction is that the Al2O3(0001) surface is decomposing during high temperature annealing in a low partial pressure of oxygen atmosphere to form Al, which adopts a low energy epitaxial structure. Similarly, Fe2O3(0001) surfaces are found to be enriched in metal and transform to Fe304(lll) and FeO(lll)-like reconstructions after being ion bombarded, annealed in UHV, and then heated in 1 x 10"^ mbar of O2 [54, 55]. As in the case of the Al2O3(0001) surface, the observed restructuring of the surface is consistent with O loss in a low PQ2 environment. One certain implication of the solid-vapor equilibrium is that when an oxide is ion bombarded or heated in UHV, the resulting bulk stoichiometry is neither well defined nor likely to be reproducible. This is a likely explanation for the diverse range of observations that have been reported for the TiO2(110) surface which include not only a (1 X 1) structure [56-62], but also structures related to Magneli phases [63,64], and (1x2) [56,57,59,61], (1x3) [61], (2x2) [59], c(2xl) [56], and (2x3) [65] structures. Despite the difficulty of producing oxides with reproducible bulk stoichiometrics in vacuum, it must be emphasized that kinetically stable surfaces with reproducible structures have been prepared in the UHV environment. The creation of the TiO2(110) (1x1) surface in UHV is an excellent example [62]. Not surprisingly, STM images of the surface illustrate that it has a remarkably high concentration of oxygen vacancies. Surfaces prepared by cleavage at ambient temperature, if not annealed or ion bombarded in vacuum, are likely to have relatively well defined and reproducible stoichiometrics. While not generally in equilibrium, they should be kinetically stable at room temperature. However, the composition of such a

500

surface is not always easy to specify. In very simple situations, cleavage can produce two identical surfaces and there is no ambiguity about the surface termination. In other cases, cleavage leads to two distinct surfaces. It is also possible that there are multiple possible cleavage planes. Consider, for example, the VÔjg structure shown in Fig. 10 [66]. Cleavage at the points labeled A, B, or C lead to distinct terminations. The analysis of STM and AFM images of the surface and comparison to calculated electron density maps are consistent with the conclusion that cleavage occurs at C. Based on studies of tungstates, molybdates, and vanadates, it is possible to make the general statement that oxides cleave at the longest, weakest set of bonds perpendicular to the cleavage plane [66-70].

(a)

(b)

Fig. 10. (a) Idealized polyhedral projection of the V6O13 structure, showing the three possible planes of cleavage, (b) An STM image of the plane, identified as C.

501

5. REACTION ANISOTROPIES ON M0O3 SURFACES In section 2 of this chapter, surface character-property relationships for the catalytic properties of M0O3 were described. These relationships were deduced largely from experiments on particulate materials. In this section, AFM experiments are described that were designed to discriminate the functions of the basal and lateral facets by direct observation. Experiments were conducted on single crystals surfaces with a clearly defined orientation and composition that were treated in a reactor system capable of reproducing the range of temperatures and environmental conditions relevant to the oxidation of alcohols [71-75]. Under these conditions, both the morphology and the stoichiometry of the surface evolved in a complex way. Room temperature AFM imaging at the conclusion of the treatment was used to determine how the surface structure changed. The AFM images in Fig. 11 show a variety MoOgCOlO) surfaces before and after reactions with methanol [73, 74]. The images in Fig. 11a and b show the surface structure before the reaction. Cleavage surfaces, such as those illustrated in Fig. 11a, are characterized by atomically flat terraces separated by steps that have heights that an integer multiple of about 7 A. This characteristic height is half the unit cell repeat length along [010] and corresponds to the distance between the van der Waals gaps which separate the adjacent layers of the structure (see Fig. 2). The steps on cleavage surfaces are most frequently aligned along [001]. The configuration of steps on a growth surface (see Fig. lib) is determined by the positions of the screw dislocations. Defects such as the one near the center of the image in Fig. l i b provide a continuous source of steps that facilitates crystal growth. In the reactivity studies described below, there was no detectable difference in properties of the growth surfaces and the cleavage surfaces. When MoO3(010) surfaces are reacted with air-N^-MeOH mixtures at 300 °C, loops of steps surrounding shallow pits nucleate on the previously flat terraces (see Fig. llc-e). The pits in these images are always bounded by half unit steps and are therefore only 7 A deep. While the pitting process was observed in all of the air-N2-MeOH mixtures tested (airiN^ ratios between 100:0 and 02:98), the shapes of the pits and their evolution depended on the composition of the mixture. Under oxygen rich conditions, the pits had a rectangular or ovular shape, and were always elongated along . As the air concentration is reduced, the pits elongate along and assume a triangular habit (see Fig. lie). The pits are only formed when some methanol is included in the feed. When air is completely eliminated from the feed, H^MoG^ precipitates (the white contrast in Fig. llf). These precipitates are formed when the hydrogen liberated during the chemisorption of the alcohol reacts with the M0O3 to form the molybenum bronze, rather than reacting with lattice O to form

502

water [73]. The same reaction occurs when MeOH is replaced by higher order alcohols.

Fig. 11. AFM images of the MoO3(010) surface. The arrows show the direction of the [001] axis. A cleavage (a) and growth (b) surface before the reaction, (c) After 2 h at 300 °C in air saturated with methanol, (d) After 2 h at 300 °C in a 50:50 air : N2 mixture saturated with methanol, (e) After 2 h at 300 °C in a 30:70 air : N2 mixture saturated with methanol, (f) After 6 min at 330 °C in N2 saturated with methanol.

503

The most plausible explanation for the formation of the pits during the reactions with the methanol-air mixtures is that the surface methoxide species that form during the initial dissociative chemisorption of methanol are capable of desorbing as molybdenum-oxide-methoxide molecules, which are then carried away in the flowing gas. The observation that crystalline MoÔ^COCH^)^ forms at the reactor exhaust supports this mechanism for evaportation [74]. These observations, made far from equilibrium, illustrate that the morphological structure of the M0O3 surface in service is continuously evolving and depends on the composition of the gas feed. Of particular interest is the generation of step edges on the surface. As discussed in section 2, the step edges provide undercoordinated Mo sites (similar to those found on the lateral facets of the

Fig. 12. AFM images of defect related pits that form on the MoO3(010) surface. In each case, the arrow indicates the [001] direction, (a) After heating for 3.5 h in air at 400 °C, shallow pits form at surface-dislocation intersections, (b) After 8 min at 400 °C in a 90:10 N2:H2 mixture containing 20 ppm water, (c) After 5 min at 400 °C in a 90:10 N2:H2 mixture containing 20 ppm water, (d) After 8 min at 400 °C in a 90:10 N2:H2 mixture containing 20 ppm water, followed by a 3.5 h anneal in air at the same temperature.

504

crystal) that are thought to have a different function than the basal sites. It has been observed that the elastic strain associated with extended crystal defects also promotes the formation of pits in the (010) surface when heated to approximately 400 °C in the presence of water vapor. It is therefore possible to generate controlled distributions of undercoordinated Mo to test their relative importance in the reaction with methanol. There are two primary crystal defects that create enough strain to nucleate a pit during reactions with water vapor: dislocations and crystallographic shear planes. Several examples of pits that can be formed at the sites where defects intersect the surface are illustrated in Fig. 12. Water vapor is known to enhance the sublimation of M0O3 through the formation of volatile molybdenum oxyhydroxides such as Mo02(OH)2 and is thought to play a similar role in the formation of pits near the points where defects intersect the surface [75]. For crystals heated to 400 °C in oxidizing atmospheres containing water vapor, pits form only at the points where screw dislocations intersect the surface. The image in Fig. 12a shows decending spiral steps into the etch pit formed around a dislcoation with a Burgers vector of approximately 14 A, which is equal to the lattice spacing along [010]. By examing dislocations with different Burgers vectors, it was concluded that the amount of material removed in a fixed time period was proportional to the stored elastic strain around the defect. Because the average dislocation density was only 1.5 x lOVcm^, such pits are rather rare. However, in reducing atmospheres, the wide spread formation of crystallographic shear planes (defects that acconmiodate oxygen loss) leads to pits densities up to four orders of magnitude greater [71]. A surface with a high density of small pits (black contrast) is shown in Fig. 12b. Higher resolution images show that 1.5 A high steps that are characteristic of surface/CS plane intersections invariably intesect the pits (see Fig. 12c). When a surface such as the one shown in Fig. 12b is reheated in an oxidizing atmosphere, the CS planes are eventually annihilated and and the remaining pits reorient to form acicular trenches along the direction, as illustrated in Fig. 12d. Because pit formation is activated by elastic strain from extended defects (which can be deliberately introduced), we can exert some degree of control over the size and density of the pits and, therefore, the available population of undercoordinated Mo. To test the influence that these sites have on the chemisorption of alcohols, a series of model surfaces were created with different densities of undercoordinated Mo associated with pits and then reacted with methanol saturated N2 at 330 °C. These conditions lead to the formation of the H^Mo03 phase, which was used as a quantitative and local indicator of the dissociative chemisorption reaction. The images in Fig. 13 illustrate selected results from these experiments. In Fig. 13a and b, it is clear that the acicular H^MoOg precipitates are growing from the lateral walls of the surface pits [75].

505

The images in Fig. 13c and d show that when the density of the pits increases and the distance between adjacent pits is small compared to the size of the precipitates, the preferential nucleation is not obvious. However, the total amount of the hydrogen bronze phase produced still correlates with the total area of the pits. Both crystals in Fig. 13 c and d were reduced for 5 min at 400 °C in forming gas to form the pits. One was then oxidized in ambient air at 410 °C for 1 h so that the pits could grow, while the other was not. As a results, the pits shown in Fig. 13c are all less than 200 A deep while those on the surface that received the oxidation treatment are all greater than 500 A deep (see Fig. 13d).

mi.% ~*l-% -*4,î

Fig. 13. AFM topographs of MoOsCOlO) surfaces, (a) A crystal that was reduced for 10 min at 400°C in forming gas and then oxidized for 2 hr at 400°C in the ambient prior to being reacted for 20 min at 330 °C in N2 saturated with MeOH. The black-to-white contrast is 100 A. (b) A surface that was reduced in forming gas at 400°C for 10 min, oxidized for 3 hr at 410°C in air, and then exposed to N2 saturated with MeOH for 15 min at 330 °C. The black-to-white contrast in the image is 100 A. (c & d) comparison of precipitation on surfaces with different pit sizes (see text). Each was reacted for 10 min at 330 °C with N2 saturated with MeOH. The black-to-white contrast in (c) is 200 A, while that in (d) is 400 A.

506

Both surfaces were then reacted for 10 min at 330 °C in N^ saturated with MeOH at room temperature; the precipitates on the surface with the larger pits (Fig. 13d) are larger and more numerous than those on the surface with more shallow pits (Fig. 13c). When the total precipitate volume on each surface is estimated, the results indicate that the surface with the larger pits produced four times as much of the H^MoOg phase. It should be pointed out that the stoichiometry of these two surfaces was not the same. The one that was heated in air was more oxidized than the other and, therefore, had fewer surface oxygen vacancies. Since surface oxygen vacancies are also sources of undercoordinated Mo, one must consider the possibility that they compete with the lateral facets as chemisorption sites. However, when the upper limit of the oxygen vacancy concentration (lO'"^) is compared to an estimate of the number of undercoordinated Mo associated with the lateral walls of pits, one finds that the surface oxygen vacancy concentration is not a significant factor [75]. This is consistent with the observations that the sample with fewer oxygen vacancies (but deeper pits) formed more of the hydrogen bronze phase. In summary, the results presented in this section indicate that at the temperatures and pressures relevant for partial oxidation reactions, it is the {hOk} surfaces that are active for the dissociative chemisorption of alcohols. Furthermore, the morphological structure of the catalyst particles is continuously evolving during the reactions, introducing a higher than anticipated density of undercoordinated Mo on {hOk} surfaces while the catalyst is in service. While proposed mechanisms for partial oxidation usually concentrate on the exchange of O between the catalyst and the gas phase, the current results suggest that the exchange of H and Mo may also be important. 6. THE ORIENTATION DEPENDENCE OF THE PHOTOCHEMICAL REACTIVITY OF TiOj The photochemical properties of titania surfaces are of interest for several reasons. They determine the stability of pigmented paint systems [76], the rate at which pollutants can be degraded in systems designed to purify air and water [77], and are the root cause of poorly understood phenomena such as water photolysis [78] and "super hydrophilicity" [79]. Using thin rutile epilayers with five low index orientations, it has been shown that the relative rates of photochemical reactions catalyzed by titania depend on the surface orientation [80]. In this chapter, experiments used to map the complete orientation dependence of the relative photochemical reactivity of TiO^ are described [8183]. In this case, the relevant reactions are carried out at room temperature and this gives us the opportunity to fix both the surface morphological structure and stoichiometry. For the reactions described here, all of the surfaces were

507

annealed in air for several hours at 1200 °C to define the stoichiometry and morphological structure. The challenge associated with measuring the photochemical properties of titania as a function of orientation lies in the number of required measurements. Each possible surface orientation is defined by two independent variables. If each orientation parameter can be measured to within ±3°, then for a tetragonal material such as rutile, there are 436 distinct surface orientations. It is clear that the preparation and analysis of oriented single crystal surfaces would be prohibitively tedious. This challenge can be overcome by using polycrystalline specimens. Assuming a 50 micron grain size, a 5 mm^ area of a polished section through a random microstructure reveals approximately 2000 grain surfaces. Single grain orientations were determined based on electron backscattered diffraction patterns (EBSPs) obtained in a scanning electron microscope (SEM). For unfaceted grains, the index of the surface plane can be determined directly from the EBSP. Faceted grains, however, are terminated by more than one surface plane as descibrd in section 3. In such cases, atomic force microscopy (AFM) was used to measure the inclination of each facet with respect to the surface normal; when these data are combined with knowledge of the grain orientation from the EBSPs, the facet planes can be indexed. To make a local measurement of the photochemical activity of each grain, we used a well established probe reaction (the reduction of aqueous Ag+ to Ag^) that deposits metallic silver on the surface as a reaction product [84-85]. The amount of silver deposited on each grain's surface during a given reaction, which is determined from atomic force microscopy (AFM) images, is taken to be a quantitative indicator of the grain's relative photochemical reactivity. The reactivity can then be correlated to surface orientation and/or the relative area of each facet on the surface. The AFM images in Fig. 14 illustrate how this technique is used to determine the anisotropy of the reactivity [81]. The rutile poly crystal used in this experiment was polished and then annealed in air at 1473 K for four hours to heal polishing damage. Note that grain boundaries and pores appear in the images as black contrast. Figure 14a shows several contiguous grains in the specimen. Fig. Ic shows a triple grain junction, and Fig. 14e shows the step structure on the surface of a single grain. The sample was immersed in aqueous AgNOs, exposed to ultraviolet (UV) light, and the same areas were located and imaged again. The results are shown in Fig. 14 (b,d,f). Note that after the reaction, heterogeneous white contrast (corresponding to elevated regions) is apparent in the images. On the basis of energy dispersive X-ray spectroscopy conducted in an SEM, there is a one-to-one correlation between the elevated regions of white contrast and high concentrations of Ag. Fig. lb illustrates that some surfaces appear to be almost completely covered with silver, while others have very few deposits. On the basis of these

508

observations, three types of grains (illustrated in Fig. 14d) are distinguished: very reactive (upper left), moderately reactive (lower left), and relatively unreactive (right). The very reactive grains are uniformly covered with silver so that the underlying surface structure can no longer be distinguished. This occurs when there are more than 4 Ag islands per jLim^. The moderately reactive grains

Fig. 14. AFM images of polycrystalline rutile surfaces before (a,c,e) and after (b,d,f) the photochemical deposition of silver (white contrast).

509

have easily distinguishable silver islands with a density between 2 and 4 per |im2, while the morphology of the TiO^ surface remains visible in the background. The relatively unreactive grains have an inhomogeneous distribution of Ag and there are less than 2 silver islands per îm^. While parts of unreactive grains in the vicinity of grain boundaries, residual polishing scratches, pores, and other defects do show deposited silver, the flat regions of the surface are indistinguishable before and after the reaction. It is the inhomogeneous distribution of silver that distinguishes the relatively unreactive grains from the moderately reactive grains. Based on knowledge of the grain orientation and surface inclinations measured by AFM, it is possible to index all of the grain surfaces. The dependence of the reactivity on the grain and surface orientation is illustrated simultaneously in Fig. 15, which shows an inverse pole figure that combines photochemical reactivity and surface orientation stability data. As discussed in section 3, the areas with the constant gray shading are regions where the surface orientation is stable and the grain has a flat surface. In such areas, the surface orientation and the grain orientation are the same. The areas filled by tie lines indicate orientations that break up into two facets. In this case, the two stable surface facets are found at the ends of the tie lines. The white areas show orientations that break up into three stable surfaces whose orientations are found at the vertices of the surrounding triangle. Based on the results in Fig. 15, we can make the following statements. First, the most reactive grains have orientations near {101}. Second, with the exception of a few outliers, the grains that have high reactivity have surface {001

•

100}

> 4 deposits/cm^

X > 2 deposits/cm^ O < 2 deposits/cm^ 110} Fig. 15 Inverse pole figure summarizing reactivity and orientation stability data for titania.

510

orientations that are unstable and break up into two or three separate facets. Furthermore, one of the stable facets is always {101}. Further analysis indicated that the relative amount of silver on surfaces containing {101} facets correlates with the fractional {101} area. Based on this result, we conclude that the {101} surface is the most reactive and this is the key microstructural feature for the design of high reactivity materials. This conclusion allows microstructures optimized for high reactivity to be specified. Planar structures should be oriented such that the surface normal breaks up into complementary {101} facets (the surface area of a faceted structure is greater than a single {101} type facet). Particulate catalysts should have a pseudo-octahedral habit with {101} planes bounding the crystal. The observation that the most photochemically reactive surface orientations all contain {101} facets suggests that the enhanced reactivity is not associated specificially with the bulk crystal orientation, but is a property of this particular surface plane. It is likely that there are special atomic configurations on this plane that either alter the efficiency with which photogenerated carriers are trapped at the surface or the rate at which they are transferred across the solid-liquid interface. Based on bulk geometry alone, there is nothing that sets the {101} plane apart from less reactive surfaces. On this surface, Ti cations are coordinated by five O, as they are on the (100) surface (which is inert). In the absence of higher resolution microscopy results, it is not possible to say if special molecular configurations are created by reconstruction or defect formation. 7. CONCLUSION The surface properties of metal oxides are anisotropic and because of this, there is a potential to control the properties of surface active materials by manipulating the distribution of exposed surface characters. Under equilibrium conditions, the distribution of surface characters is determined by the anisotropy of the surface energy and the solid-vapor equilibrium. While one is not necessarily limited to equilibrium configurations, materials at sufficiently high temperatures will evolve in this direction and significant morphological and stoichiometric changes can occur. By systematically evaluating the properties of surfaces with different characters, it is possible to define optimized morphological structures. ACKNOWLEDGEMENT The author acknowledges the support of the US National Science Foundation under grant DMR 0072151. The work was supported in part by the MRSEC program of the National Science Foundation under award number DMR0079996. Contributions to the research described in this chapter from Richard

511

L. Smith, Jennifer B. Lowekamp, David M. Saylor, Jennifer L. Giocondi, P.A. Morris Hotsenpiller, J.D. Bolt, and W.E. Farneth are also gratefully acknowledged. The author thanks Prof. C.B. Carter for the permission to use Fig. 4. REFERENCES [I] J.E. Germain in: M. Che, G.C. Bond, (Eds.), Adsorption and Catalysis on Oxide Surfaces, Elsevier Science Publishers, B.V., Ansterdam, 1985, p. 355. [2] J.M. Tatibouet, J.E. Germain, J. Catal., 72 (1981) 375. [3] J.C. Volta, J.M. Tatibouet, J. Catal., 93 (1985) 467. [4] J.M. Tatibouet, Ch. Phichitkul, J.E. Germain, J. Catal., 99 (1986) 231. [5] K. Briickman, R. Grabowski, J. Haber, A. Mazurkiewicz, J. Sloczynski, T. Wiltowski, J. Catal., 104(1987)71. [6] K. Bruckman, J. Haber, T. Wiltowski, J. Catal., 106 (1987) 188. [7] A. Anderssen, S. Hansen, J. Catal., 114 (1988) 332. [8] B. Mingot, N. Floquet, O. Bertrand, M. Treilleux, J.J. Heizmann, J. Massardier, M. Abon, J. Catal. 118(1989)424. [9] M. Abon, J. Massardier, B. Mingot, J.C. Volta, N. Floquet, O. Bertrand, J. Catal, 134 (1992) 542. [10] M.R. Smith, U.S. Ozkan, J. Catal., 141 (1993) 124. [II] L. Kihlborg, Arkiv Kemi, 21 (1963) 471. [12] M. Gasior, T. Machej, J. Catal., 83 (1983) 472. [13] E. Bordes, Catal. Today, 6 (1992) 21. [14] W.E. Farneth, F. Ohuchi, R.H. Staley, U. Chowdhry, A.W. Sleight, J. Phys. Chem., 89 (1985) 2493. [15] J.M. Vohs, M.A. Barteau, J. Phys. Chem., 91 (1987) 4766. [16] K.S. Kim, M.A. Barteau, Surf. Sci., 223 (1989) 13. [17] K.S. Kim, M.A. Barteau, J. Catal., 125 (1990) 353. [18] M.A. Barteau, J. Vac. Sci. Technol. A, 11 (1993) 2162. [19] J. Fompeyrine, R. Berger, H.P. Lang, J. Perret, E. Machler, Ch. Gerber, J.-P. Locquet, Appl. Phys. Lett., 72 (1998) 1697. [20] M. Kawasaki, K. Takahashi, T. Maeda, R. Tsuchiya, M. Shinohara, O. Ishiyama, T. Yonezawa, M. Yoshimoto, H. Koinuma, Science, 226 (1994) 1540. [21] G. Koster, B.L. Kropman, G.J.H.M. Rijnders, D.H.A. Blank, H. Rogalla, Appl. Phys. Lett., 73 (1998) 2920. [22] H. Tanaka, H. Tabata, T. Kawai, Thin Sohd Films, 342 (1999) 4. [23] M. Kawasaki, O. Ohtomo, R. Tsuchiya, J. Nishino, K. Koinuma, Mat. Res. Soc. Symp. Proc, 474 (1997) 303. [24] P.A. Salvador, B. Mercey, O. Perez, A.M. Haghiri-Gosnet, T.-D. Doan, B. Raveau, Mat. Res. Soc. Symp. Proc, 587 (2000). [25] M.W. Bench, P.G. Kotula. C.B. Carter, Surf. Sci., 391 (1997) 183.

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514

Chapter 13

Vibrational Spectroscopy at Oxide Surfaces Brian E. Hayden Department of Chemistry, University of Southampton, Highfield, Southampton, S017 IBJ, United Kingdom.

1.

INTRODUCTION

Vibrational spectroscopy of surfaces has proved itself a powerful tool for the study of both clean and adsorbate covered surfaces. The characterisation of the vibrational modes on clean surfaces provides information concerning interatomic potentials that can be used in the interpretation of structural and dynamical properties. Vibrational spectroscopy of adsorbates provides one of the most important sources of information concerning the chemical and physical nature of the adsorbed species. This chapter concerns the development and application of vibrational spectroscopy on well characterised metal oxide surfaces to the measurement of adsorbate and localised surface phonon modes of the substrate. The development of the techniques applied to metal oxides has been proceeded by their more extensive application to studies on metal and semi-conductor surfaces, for which there is a comprehensive review literature [1-8]. The most important and commonly applied techniques include various forms of infrared spectroscopy, such as Reflection Absorption InfraRed Spectroscopy (RAIRS), Surface Electromagnetic Wave Spectroscopy (SEWS), Emission Spectroscopy (ES) and Multiple Internal Reflection Spectroscopy (MIRS). High Resolution Electron Energy Loss Spectroscopy (HREELS) and Inelastic Atom Beam Scattering (lABS) have also found wide application to surface vibrational spectroscopy, and have the advantage of an intrinsic surface sensitivity in the scattering event. ES and SEWS are the least commonly applied of the infrared techniques, and no attempt has been made to apply them to metal oxide surfaces. MIRS, which has been extensively applied to semi-conductor surfaces, has not so far been applied to metal oxide surfaces. This is probably because absorption in the substrate resulting from more extensive and strong dipole active phonon modes in metal oxides would make internal reflection experiments difficult. The spectral limitations of lABS to ^

Ss Fig. 4. Reflection of IR radiation at a surface, showing incident vectors for P- and Spolarised radiation, and the components of the P-polarised radiation at the surface.

522

Consider an IR beam incident on the surface at an angle (p with respect to the surface normal (Fig. 4). The incident radiation can be resolved into components parallel (S-polarised) and normal (P-polarised) to the incident plane. The S-polarised radiation only has a component (S) parallel to the surface (in the y direction). However the p-polarised radiation has components parallel or tangential (Pt) to the surface, and perpendicular (?„) to the surface. Each layer (vacuum (e =1), adsorbate (e) and substrate (cs)) is characterised by an isotropic complex dielectric constant (e) which is defined as 8 = (n + ik/, where n in the refractive index, and k is the absorption coefficient. The change in reflectivity (AR) resulting from the adsorbate layer of thickness d for S- and Ppolarised radiation is usually expressed as a ratio to the reflectivity (AR/R): ARs =87id coscp Im fe - ER) Rs I (l-8s)

(4)

ARp =87cd C0S9 Im (e - eg) 0 - (\/ZER)(E + ER) sin^9) Rp X (l-8~s) (1 - (l/8s)(l + 8s) sin^G)

(5)

Given that, for a metal |8s| » |8| ~ 1, and d/X « 1 for IR frequencies of ~10"^m and d « 0.1 nm, it can be seen that the reflectivity change of S-polarised light is very small. In order to consider the effect of change in the optical response of the substrate between the two extremes of a metal (ks=30), and a non-absorbing (ks=0) or slightly absorbing (ks=0.05) semi-conductor or insulator. Fig. 5 shows the values of AR/R as a function of (p calculated using equations 1 and 2 for the various values of ks , with 8s = (3+ ksi)^. A moderately absorbing adsorbate layer has been assumed, with 8 = (1.3+O.li)^. AR is defined as R-Rads, where R and Rads are the reflectivities of the clean and adsorbate covered surfaces respectively: A positive value of AR corresponds to an adsorbate induced absorption band. Such a calculation shows why for a metal substrate (ks = 30), there is no coupling of the dipole to the parallel S component of the radiation, and the optimum incident angle for RAIRS using P polarised radiation is cp = 80-90° [36, 38, 39]. Under such conditions, one can expect an absorption band (AR/R is positive) at the adsorbate resonant frequency. The results in Fig. 5 show that for RAIRS on an oxide surface (ks » 0), a complex behaviour of AR/R is exhibited. The most significant influence of the absorbing overlayer on the P-polarised radiation is observed on either side of the Brewster angle, cpe . The absorbing overlayer will induce a reduction in reflectivity (absorption band) below cpe, and an increase in reflectivity (transmission band) above cps . S-polarised radiation also interacts with the absorbing overlayer, all be it weakly, with an increase in reflectivity induced at all (p. For P-polarised radiation, maximum sensitivity in RAIRS would therefore be expected just above or below cpB with ks=0. A small amount of absorption by

523

the substrate (ks = 0.05) results in a retention of the sensitivity for cp > cpB, but a reduction for (p > (ps. Typical values [40] of the optical constants for metal oxides in the range 1000-3000 cm'^ are ns=1.5-3.5, and ks=0-0.07, similar to silicon, with stronger absorption at energies < 1000cm"\ 0.003 1

0.002

ks=30

0.00000 -0.00002 H

f^ 900cm"^ on Si, it has been demonstrated that internal reflection methods produce higher sensitivity [34], and the majority of IR measurements on silicon, for example, are carried out using such an approach [33, 34, 43-45]. It is clear, however, that the presence of any absorption in the substrate will restrict measurements to external reflection. The extension of phonon absorption to higher energies on many metal oxide surfaces makes the application of internal reflection measurements less attractive on such surfaces. An alternative on semiconductor surfaces has been to use metal ion implanted substrates which exhibit dielectric propertied in the IR more similar to metal substrates [46-50]. Under such conditions, the conventional RAIRS geometry for metal surfaces of P-polarised radiation at grazing incidence is optimal, with strong screening of parallel components and absorption bands at the resonant frequency. An important consequence of carrying out RAIRS on a substrate exhibiting no absorption, or only weak absorption, is that coupling to both parallel and perpendicular components of the surface dynamic dipole can take place. This is evident clearly in the case of S-polarised radiation (Fig. 5) which predicts a small transmission band for all values of (p. For P-polarised radiation, the isotropic three layer model (Fig. 5) predicts a transmission band at the resonant frequency for (p > Brewster angle, and absorption bands for cp < (pB. This simple model, however, is insufficient to separate the coupling of the normal Pn and tangential Pt components of the IR field to the isotropic adsorbing overlayer (Fig.4). It is also ineffective in treating the anisotropic and microscopic characteristics of molecular adsorption, such as molecular alignment or dipole coupling.

525

The microscopic characteristics of a real adsorbate layer have been considered [33, 34] by separating the x, y and z components of the electric field at the interface (Fig.4), and applying the Lorentz oscillator model to microscopically represent the adsorbate in the three-layer model. For the case of external reflection at a vacuum/semi-conducting, where Ss is real (no absorption) and isotropic, w e can write: ARS = Rs

271V ly (8"ydy) coscp

(6)

ARp = .iTiv. {Ix (8"xdx) + Iz _ s s ' _ ( s " z d z ) } Rp

coscp

(7)

8'z^ + S"z^

Where Ix,y,z are the corresponding components of the field intensities: ly = 4cos^(p

(8)

1-8S

Iz = 4cos^(p l-8s

l/8s sin^cp {((l+8s)/8s) sin's-1}

(9)

Ix = 4cos'(p 1-Bs

WER sin'cp-n {((l+8s)/ss)sin'(p-l}

(10)

While the three layer "macroscopic" model (Eqn.4,5 and Fig.5) allows a calculation of AR/R for substrates in which 8s is complex, there are no simple explicit expressions corresponding to Eqn. 6,7 for complex 8s, and in practice normal reflection or transmission has been used for which the equations are greatly simplified [51, 52]. Nevertheless, Eqn. 6-10 can be considered a reasonable approximation for non-absorbing and weakly absorbing substrates. The separation of the field components, and the expression of anisotropic adsorbate dielectrics ( 8"x, 8"y , 8"z ) becomes usefiil when 8" is expressed in terms of meaningfiil microscopic quantities (eg. The dynamic charge of the adsorbate mode), and the Lorentzian oscillator [53] parameterisation of an adsorbate layer has proved usefiil [34] for this purpose. For any particular vibrational mode, the effective dielectric fimction 8(co) is given by: 8"(co) = 8oo + cop' coo^ - 0)' - iyco

(11)

526

(p = 45"

^ 4.00E-04

20|50

2070

2090 \

wavenumber/cm*

/2110 \[^

2130

2150

S

-8.00E-04 l.OOE-03

9 = 80" 20|50

2070

2090 1

2^0

wavenumber / cm'

Fig. 6 Model calculations (employing the Lorentzian ocillator approximation and a three layer optical model) of AR/R for RAIRS on semiconducting or insulating (isotropic and nonabsorbing) substrates (8s=3) for P- and S-polarised radiation. Calculations are shown for two values of incident angle ((|)), below and above cpe The adsorbate layer is assumed isotropic (8"x=8"y=8"z) with e*/e=0.5, v=2100cm"\ Y=5cm'\ The convention AR=R-Rads is used, so that positive resonances correspond to absorption bands, and negative values transmission bands.

Where Coo is the electronic part of the dielectric function, cop is the plasma frequency associated with the adsorbate layer, and y is the natural line width of the oscillator. Fig. 6 shows the resonance associated with the coupling of the Spolarised radiation to the y component of an oscillating dipole (S), and the coupling of P-polarised radiation to the x and z components through Pt and Pn

527

respectively, calculated using Eqn. 6-11. For the purpose of the calculation it has been assumed that the adsorbate layer is isotropic with a dynamic charge e*/e = 0.5 to represent a mode with a high extinction coefficient such as v(C-O). Calculations have been made at two incident angles, one below (cp = 45°) and one above (cp = 80°) (pB. The advantage of such an approach is that it starts to take account of the orientation of the adsorbate dynamic dipole in addition to the dynamical interactions resulting from dipole coupling in a thin layer or monolayer. A spectral shift from the singleton frequency can be expected in coupling to the z (normal) oscillator component as a result of the net dipole coupling of a two dimensional array of parallel oscillators [54]. This effect shows up as a thickness (d) dependence in the optical response in the z direction of the adsorbate layer (Eqn.7), within the modified three layer model. For a strongly adsorbing monolayer such as CO, this shift is calculated (Fig. 6) to be -18cm'\ in agreement with Chabal [33, 34]. The result is that for an isotropic adsorbed thin film on an oxide surface where the parallel mode may be observed, the parallel and normal components will be observed in RAIRS at different frequencies, and will be manifested as either absorption or transmission bands depending on (p (Fig.6). Measurement with S-polarised radiation at (p = 45° or (p = 80° (either side of (pe) will result in transmission bands (Fig. 6), as predicted also by the simpler model (Fig. 5). For P-polarised radiation at (p = 80°, the coupling of Pn and Pt to the adsorbate oscillator results in a transmission and an absorption band respectively. This situation is expected to be reversed at cp = 45°, below (p B. Owing to the relatively low reflectivity of semiconductor surfaces, and hence the low detected signal expected for the external reflection experiments, very few RAIRS studies of adsorbates on semiconductors have been performed, and most experiments have been carried out using internal reflection (MIRS). However, studies using P polarised radiation on the hydrided Si(lOO) surface at (p=85'' [35], and unpolarised IR of trimethylgallium deposited on GaAs(lOO) [55] confirm the existence of transmission bands predicted theoretically on semiconductors. For metal oxide surfaces, where the substrate is isotropic and nonabsorbing or weakly absorbing, the microscopic model predictions concerning the type of band and shifts for the coupling of the various polarised components (Fig. 6) will be valid. One may therefore expect (with cp > CPB) that for a thin film of an isotropic physisorbed layer, coupling of S or Pt to an adsorbate mode with a high extinction coefficient will result in transmission and absorption bands respectively at a frequency similar to the solid state value, as a result of the "infinite" dimensions of the film parallel to the surface. The coupling of the component Pn will result in a transmission band at a higher frequency (Fig. 6). Fig. 7 shows an FT-RAIRS spectrum taken [56] of a isotropic physisorbed multilayer of Rh2C04Cl2 on Ti02(l 10) with cp = 83°.

528

T 1 r""T 1—I I 2030 2120 2100 2080 2060 2040 2020 wavenumber / cm'^ Fig.7 FT-RAIRS spectrum of Rh2C04Cl2 on TiO2(110) [56]. Measurements are carried out at (p=83°, and can be compared directly with the predictions of the model used for an isotropic film on an isotropic, non-absorbing substrate (Fig.6). The lines indicate the expectation values for the modes in the condensed phase.

As predicted by the calculation, the absorption features of the P-polarised spectrum (coupling to Pt) correspond to the transmission features of the S polarised spectrum, and also correspond to the expectation values for the assigned modes of the bulk crystalline material [56]. The transmission features of the P-polarised spectrum (coupling to Pn) correspond to the blue shifted bands associated with a net dipole coupling of the normal components of the dynamic dipoles in a thin film. The microscopic model, however, cannot take into account net coupling of dynamic dipoles oriented parallel to the surface for the thin (microscopic) film. Such coupling in adsorbed monolayers has been shown [57] by probing an otherwise disallowed transition on a metal through a combination band, to result in a red shift from the singleton frequency. This effect of parallel and normal dipole components can be best exemplified by comparison of the RAIRS spectrum of an isotropic physisorbed molecule with a very strong dipole oscillator, v(C-O) in Mo(CO)6, with the gas phase value (singleton frequency) [58]

529

7 / / / /"Z

ttttftttttttt ttttttttttttt Fig. 8 FT-RAIRS spectrum of a physisorbed layer of Mo(CO)6 on TiO2(110) [58]. The net dipole coupling in the thin film shifts the singleton (gas phase) frequency coo either up or down in energy for normal and parallel components respectively.

This is shown in Fig. 8. The coupling of P polarised radiation to the normal and parallel components of the oscillator (the Tiu v(C-O) mode) results in the expected transmission and absorption bands respectively. The former is blue shifted, and the latter is red shifted. This can be understood simply by considering the dynamic dipole coupling lattice sum expected for aligned oscillators perpendicular and parallel to the surface [59], represented schematically in the inset of Fig. 8. The difficulty in carrying out RAIRS, with respect to metals, on metal oxides and semiconductors is that reflectivity and sensitivity are low. The advantages of RAIRS on metal oxides results from the fact that coupling of Ppolarised radiation to can take place to both normal and parallel components of the adsorbate dynamic dipole, not only in isotropic physisorbed layers, but more importantly in oriented chemisorbed monolayers. In addition the modes can be separated as a consequence of the net optical response (transmission or absorption bands), and the adsorbate geometry becomes accessible through the characteristic of both components. It will also become evident that since the parallel component of dynamic dipoles are IR active, the azimuthal dependence of the RAIRS spectrum can be used to obtain the alignment of adsorbed species. The number of RAIRS experiments carried out to date on single crystal oxide surfaces is relatively small. An alternative approach to studies on oxides is to use metal oxide substrates epitaxially grown on metal substrates. RAIRS on such systems is characterised by the dielectric response of the metal since the oxide films are thin with respect to IR wavelengths. The resuh is that only the

530

normal components of the dynamic dipole couple to the radiation, and all bands are absorption bands. The advantage of this approach is that sensitivity is concomitantly higher. In principle, the technique of enhancing RAIRS signals by metal ion implantation, a technique used for semiconductors [46-50], can be adapted to single crystal metal oxide systems. Indeed this approach has been applied to the study of SiOi supported CuO model catalysts on a metal implanted Si substrate [60]. The three-layer theory can, in principle, be extended to the more complex, but perhaps more interesting case with respect to modelling heterogeneous catalysis, of RAIRS on single crystal oxide supported metal particles or thin films. While no explicit calculations are available to date, it has been shown experimentally that RAIRS of adsorbates on small particles supported on oxide surfaces result in the dominance of the metal oxide dielectric response. As particle size (or film thickness) increases, the metallic response eventually dominates resulting in strong RAIRS absorption bands. The particle size or film thickness at which such a transition takes place is discussed below. RAIRS does provide a tool for differentiating between adsorption at the oxide itself, and on small and large particles supported on the oxide.

3.

SURFACE VIBRATIONAL MODES OF THE SUBSTRATE

The "surface" Fuchs-Kliewer modes, like the Rayleigh modes, should be regarded as "macroscopic" vibrations, and may be predicted from the bulk elastic or dielectric properties of the solid with the imposition of a surface boundary condition. Their projection deep into the bulk makes them insensitive to changes in local surface structure, or the adsorption of molecules at the surface. True localised surface modes are those which depend on details of the lattice dynamics of near surface ions which may be modified by surface reconstruction, relaxation or adsorbate bonding at the surface. Relatively little has been reported on the measurement of such phonon modes, although they have been the subject of lattice dynamical calculations [61-67]. Localised phonon excitations are in principle best studied by neutral-atom scattering, or off specular HREELS, in order to reduce the strong dipole excitation of Fuchs-Kliewer modes. Two off specular HREELS measurements on MgO(lOO) have been reported [25, 68], however there is some disagreement concerning the energy and assignment of the substrate derived loss peaks. Since the microscopic surface modes are expected to be sensitive to the surface structure, it has been suggested [9] that the differences may be associated with differences in surface preparation. Perhaps the most clear cut evidence for the observation of a localised surface phonon comes from a recent HREELS study a clean and adsorbate

531

covered CriOaCOOOl) surface [67]. Fig. 9 shows the specular HREELS spectrum of CriOsCOOOl) measured with Eo=7.5eV. The CriOsCOOOl) film was grown as a 40A film produced through the oxidation of Cr(l 10). Spectra are shown for the clean Crz^ÔsCOOOl) and Cr2^Ô3(0001) surface. The spectrum of the clean Cr2^Ô3(0001) surface is characterised by losses at 21.4, 51.7, 78.6, 85.0 and 88.5meV, and combinations of these losses at higher energies. The latter four are identified as Fuchs-Kliewer phonon modes, and the intensity and energy of these modes are found to be uninfluenced by the adsorption of CO or O2 at 90K.

84.6

88.5

Cr2 0 3

^lOcr/^03 I

0

"T" 50

^

I

100 150 energy / meV

r^ 200

Fig. 9. The HREELS spectrum of clean Cr2^Ô3(0001) and Cr2^Ô3(0001) [67]. The loss at 21.4meV in the spectrum of Cr2^Ô3(110) disappears on exposure to CO or O2, and is ascribed to a localised surface mode. Its isotopic shift of 2.4% is predicted theoretically, and is significantly smaller than that observed for the Fuchs-Kliewer modes of 4.5-6%.

532

Similarly the combination of two of these modes at 139.8meV (88.5+51.7meV) does not change on CO adsorption. In contrast, the loss at 21.4meV is attenuated with increasing CO and oxygen adsorption, and it is this attenuation which is given as evidence of a localised surface mode of the substrate. A weak loss peak observed at ca 40meV appearing at high CO coverages was associated with either the substrate-CO mode, or the CO modified Cr2O3(0001) counterpart of the clean surface localised surface phonon. The rather weak band at ca. llOmeV was attributed to a combination of the SS.SmeV Fuchs-Kliewer mode and the 21.4meV localised surface phonon, and consequently also disappears on CO adsorption. The small loss at llSmeV was assigned to oxygen atoms adsorbed (following diffusion from the bulk) at terminating chromium ions. Fig.9 shows that localised phonon mode at 21.4meV was also observed to undergo the smaller (theoretically predicted) isotopic shift than the Fuchs-Kliewer modes.

4.

VIBRATIONAL SPECTROSCOPY OF ADSORBATES

Both HREELS and RAIRS have been applied extensively of the study of adsorbates on metal surfaces. The extension of the techniques to semiconducting or insulating oxide surfaces is hampered by a number of problems. The result is that until even relatively recently [9] there were only a couple of examples of RAIRS studies on oxides, and these were confined to polycrystalline systems. Most early HREELS studies were concerned with the characterisation of the intrinsic phonon modes of metal oxide surfaces. This contrasts strongly with the extensive literature concerning the vibrational characterisation of adsorbates and intermediates on powdered oxide surfaces that have been obtained by transmission or diffuse reflection IR techniques. Both techniques will be influenced by the appearance of strong intrinsic phonon bands. This has been less of a problem for RAIRS measurements when confined to energies higher than the fundamental optical phonons. The intrinsic resolution of RAIRS also ensures that there is little contribution of absorption due to the phonon modes away from their oscillator frequency. In contrast the strong coupling of the scattered electron in HREELS to the phonons results in the excitation of not only the fundamental modes, but also their overtones, and these lie close in energy to even intramolecular adsorbate vibrational modes. This, coupled with lower spectral resolution, results a requirement to subtract the underlying phonon signal from the adsorbate spectrum in order to identify adsorbate bands. There are a number of methods that have been developed to deconvolute phonon overtones from HREELS spectra so that adsorbate losses can be observed (Section 2.1).

533

Fig. 10. FT-RAIRS spectrum of rhodium dicarbonyl Rh(C0)2 on TiO2(110) [56, 69, 70] following dissociative adsorption of Rh2C04Cl2 (a). The molecule is reformed after desorption of the CO at 450K following exposure to lOOL of CO (b). Reaction of Rh(C0)2 with hydrogen forms the monocarbonyl Rh(H)CO [71].

Because of the additional problem of sensitivity in RAIRS measurements made on semiconducting or insulating oxide surfaces, it is perhaps not surprising that in parallel to the history of RAIRS measurements on metal surfaces, the first measurements on a single crystal metal oxide surface involved a molecule with a carbonyl stretch. The dissociative adsorption of the MOCVD precursor Rh(CO)4Cl2 on TiO2(110) results in the chemisorption of the well known rhodium geminal dicarbonyl species Rh(C0)2 [56, 69, 70]. This species had been formed through the same route or through the reaction of CO with supported Rh^ particles, and characterised extensively, on high area powder oxides by IR because of its importance in heterogeneous catalysis. Two nondegenerate v(C-O) modes, the symmetric and anti-symmetric stretches, are expected for the oxide supported species at Vsym(C-O) » 2110cm"^ and Vasym(CO) « 2030 cm"^ The Rh(C0)2 species was formed exclusively on reaction of Rh(CO)4Cl2 with TiO2(110) at 300K to form 1/3 of a monolayer, and

534

characterised by a combination of XPS, TPD and FT-RAIRS [56, 69, 70]. The FT-RAIRS spectrum is shown in Fig. 10. The spectrum of Rh(C0)2 on TiO2(110) exhibits the two expected bands for the mono-dispersed geminal dicarbonyl, with Vsym(C-O) = 2112cm"^ observed as a transmission band, and Vasym(C-O) =2028cm"^ observed as an absorption band. Since the incident angle cp = 83° and the incident radiation was P polarised, the result is consistent Rh(C0)2 bound to TiO2(110) with C2v symmetry. The analysis of the theoretically expected changes in reflectivity due to adsorbates on a semiconducting or insulating substrate is given in Section 2.2. There is a coupling of Vsym(C-O) with Pn (AR is expected to be -ve, i.e. an increase in reflectivity following adsorption), and a coupling of Vasym(C-O) with Pt (AR is expected to be +ve, i.e. a decrease in reflectivity following adsorption). The advantages of RAIRS on the metal oxide surface are evident from this measurement in that modes polarised both normal and parallel to the surface can be observed, and can be differentiated by appearing as transmission or absorption bands in the RAIRS spectrum. A combination of FT-RAIRS, XPS and TPD allowed a study of the stability and reactivity of Rh(CO)2 on TiO2(110) [56]. Fig. 10 shows that following the thermal decomposition of Rh(CO)2 (a), the molecule can be reformed through carbonylation at 300K (b). The reaction with hydrogen produces a monocarbonyl Rh(H)CO in which v(CO) = 2065cm"^ is polarised normal to the surface, and couples to Pn resulting in a transmission band [71]. Since the parallel components of the dynamic dipole are active in RAIRS, it is possible to use the azimuthal dependence to obtain the orientation of the adsorbate at the surface. A similar technique has been applied to adsorbates on metals in HREELS measurements made off specular in order to observe parallel modes through impact or resonant scattering processes. This was first demonstrated for the Rh(CO)2 molecule on anisotropic TiO2(110) surface [72]. The results of this study also allow a test of the three layer model theory (Fig.5,6) as applied to S-polarised radiation. Fig. 11 shows the FT-RAIRS spectrum for 1/3 monolayer of Rh(C0)2 on TiO2(110) measured with P and S polarised radiation. Measurements have been made as a function of azimuthal angle, where d=0°corresponds to the incident radiation being incident along the azimuth. Note that this differs from the definition used in [72], but is chosen for consistency with the convention used elsewhere for FT-RAIRS measurements of formate on TiO2(110) [73] shown below. The angle of incidence cp = 83°. The mode polarised normal to the surface Vsym(C-O) only couples to the Ppolarised radiation through Pn (transmission band). The mode polarised parallel to the surface Vasym(C-O) couples to the P polarised radiation through Pt (absorption band), and also to the S-polarised radiation. The latter gives rise to a transmission band, as predicted by the three layer model (Fig. 5,6).

535

titanium

—I

1

2160

1

1

2100

1

1

2800

1

r

2400

2000

1960

wavenumber / cm

Fig. 11. Azimuthal dependence of FT-RAIRS spectra for TiO2(110)-Rh(CO)2 [72]. The azimuthal angle (j) is defined as 0° when the incident radiation is aligned in a plane parallel to the direction. The Vsym(C-O) dynamic dipole is aligned normal to the surface and couples to Pn (transmission band), and Vasym(C-O) is aligned parallel to the surface in the direction, and couples to Pt (absorption band). Two possible adsorption geometries consistent with the FT-RAIRS azimuthal dependence are shown for the gem-dicarbonyl.

The azimuthal dependence of the intensity of Vasym(C-O) in the Ppolarised radiation shows a maximum at d=90° indicating an ahgnment of the Rh(CO)2 in the direction. Since the S and Pt fields are orthogonal, using S-polarised radiation at S = 0° Vasym(C-O) is not observed, but is observed at S = 90°. The two most likely adsorption geometries of the adsorbed gem-dicarbonyl are shown in Fig. 11, both with the C-O bonds in a plane aligned in the direction. The Rh(C0)2 molecule has also been observed in FT-RAIRS studies of AI2O3 supported Rh particles [74]. This provides an interesting comparison with the single crystal TiOiCllO) measurements, since the alumina substrate was grown a thin film on NiAl(l 10). The film was sufficiently thin to ensure that the optical response in the IR at the oxide surface was dominated by the metallic substrate. Fig. 12 shows the FT-RAIRS spectra taken for CO adsorbed on metal particles of increasing size supported on the alumina film. For the smallest particle sizes, a sharp band at about 2120cm'^ is observed, and is associated with Vsym(C-O) of dispersed Rh(CO)2. The corresponding Vasym(C-O) band is not observed because of the screening of the parallel modes.

536

v^™(C-0) 9 13 20 30 55

100 200 450 720

I I

2000

I

I I I I I I j I I I I I I

1900 1800 1700 wavenumber/ cm"*

Fig. 12. A series of RAIRS spectra taken following deposition of increasing amounts of Rh on an alumina film grown on NiAl(llO) and subsequent CO saturation at 90K[74]. The average number of Rh atoms per particle is indicated. Vsym(C-O) of the gem-dicarbonyl is observed (indicated) as an absorption band at low Rh coverages, while Vasym(C-O) is screened by the underlying metal substrate. The large broader band observed at higher Rh coverages (particle sizes) results from v(C-O) of CO adsorbed on Rh° particles.

The broad band observed around 2080cm"\ predominating at higher particle sizes, is due to CO adsorbed on Rh^ particles. The RAIRS spectra obtained on oxide supported metal particles (single crystal oxides and thin film oxides) is discussed in detail in Section 5. A number of HREELS measurements of adsorbates have also been carried out on TiO2(110) single crystal surfaces [19-22]. One of these, the study of adsorbed formic acid [20], provides a useful comparison of HREELS and FTRAIRS on single crystal oxides. Fig. 13 shows the HREELS spectra at increasing exposure for formic acid adsorbed at llOK on TiO2(110). The spectra presented (including those of the clean surface) have been Fourier deconvoluted [23] in order to remove the contributions of multiple phonon losses. The positive and negative features in the background spectrum signify a less that ideal deconvolution of the data, probably a result of less than ideal spectrometer tuning and/or sample positioning [20]. It is such features, and the fundamental phonon features, that distort the adsorbate spectra, particularly in the spectral range -^

(3p'3d"^') (3p^3d"-') + e

555

The interference gives rise to a characteristic Fano-type lineshape in the ionisation cross section profile. Resonant photoemission, in which the photon energy is tuned to coincide with the maximum in the cross section profile, is of particular value in the study defect-induced metal d states. Measurement of constant initial state spectra (CIS), in which the photon energy and electron energy are varied in synchrony to keep fixed on the same initial state also provides a way of fingerprinting the atomic character of features in valence band photoemission in situations where the issue may be in doubt. For example, CIS spectra allowed identification of distinct Ti 3d and Fe 3d states when Fe was deposited on Ti02 [10]. Elsewhere it has been shown that bandgap states in Nbdoped TiOi are associated with electrons localised in Ti 3d orbitals rather than Nb 4d orbitals, whereas in V-doped Ti02 the analogous states are of V 3d character [11,12]. Angle resolved photoemission enables characterisation of the energy versus wavevector dispersion of electronic states [13,14]. The simplest experiment exploits momentum conservation parallel to the surface. For materials with layer-like structures in which there is little dispersion in electronic states normal to the surface, variations in binding energies with angle of off-take relative to the surface normal allows bands to be mapped out along directions parallel to the surface. Studies of this sort have proved to be particularly popular in the study of the dispersion of bulk states in layered oxides such as Lio.9Mo60i7 [15,16] and Nao.9Mo60i7 [17]. The technique would also be appropriate to measurement of the dispersion of 2D surface states on 3D materials. However for bulk states in 3D materials themselves interpretation of off-normal angle resolved spectra requires full knowledge of the bandstructure along a circular arc in k space. Direct band mapping in 3D materials can be achieved by the study of the dispersion of spectral features in normal photoemission spectra as a function of the energy of incident photons. This approach requires assumptions about the magnitude of the inner potential and the dispersion in the freeelectron-like final states. Variable photon energy normal photoemission experiments have now been carried out on a range of oxides. A general problem in these experiments is that the widths of the individual bands are quite large in relation to the separation between the individual bands within the bandstructure, so that individual spectral features are rarely well-resolved. In addition it is probable that the spectra contain a sizeable incoherent component arising fi-om breakdown of k conservation due to electron-phonon coupling [1]. Nonetheless the dispersion of O 2p states in materials such as ZnO [18,19] and SrTiOs [20] is in reasonable agreement with bandstructure calculations. Dispersion of 3d states in oxides such as CoO [21,22] and NiO [23] has also been observed, but it is questionable whether one-electron bandstructure calculations provide an appropriate framework for interpretation of these spectra.

556

Inverse photoemission involves a time reversed photoemission process in which electrons of varying energy incident upon a sample decay radiatively into empty electronic states. Photons of fixed frequency are counted and the energy dispersion of empty states may be investigated by varying the incident electron angle. Both photoemission and inverse photoemission require reasonable sample conductivity and their application to "hard" insulators such as MgO and AI2O3 is problematic. Both techniques also involve the complication that inelastic electron energy loss processes become convoluted with electron emission or decay. This may give rise to spectral features in regions where none are expected from the density of states [24,25] and care must always be taken to exclude these features before considering assignment to surface states. 3.2 Electron energy loss spectroscopy Two variants of the technique of electron energy loss spectroscopy (EELS) operated in the reflection mode have played a role in the characterisation of surface excitations on oxides. Moderate energy resolution of at best around 0.5 eV is achieved using non-monochromated electron beams whose energy width is determined by the thermal spread of energies in the cathode of the electron gun. The resolution may be further compromised if it necessary to apply a modulation in the analyser system in order to measure spectra in a differential dN(E)/dE mode. Nonetheless interband and other higher energy electronic excitations can be observed in N(E) or dN(E)/dE modes. The high resolution variant of the technique (HREELS) involves a monochromatised electron beam, allowing energy halfsvidths of 5 meV or better. Both phonon excitations and very low energy electronic excitations - such as the d to d excitations discussed in section 4.2 and low energy conduction plasmons in degenerately doped oxide semiconductors - can be observed in HREELS. At low beam energies inelastic scattering is usually dominated by a dipolar mechanism in which the electric fields associated with surface excitations couple over a long length range with the exciting electron in its trajectory toward and away from the solid [26]. The lobe of inelastically scattered electrons peaks strongly close to the specular direction with dominant k transfer of the order k=o}/v, where co is the frequency of the excitation and v is the electron beam velocity normal to the surface. The wavevector k in turn determines the decay length X of excitations into the bulk through X=l/k=v/co. Although the excitations do not propagate into the bulk, the decay lengths may be of the order of hundreds of Angstroms under typical experimental conditions, especially when dealing with surface phonons. Thus one is dealing with surface excitations in the sense that they arise from a dielectric discontinuity at the surface, but no information is conveyed about details of top layer structure. The

557

small angular spread of the inelastic dipole lobe when ^co«E allows analyser systems to collect all of the k integrated dipole loss intensity. Under these conditions the loss intensity I(co) as a function of frequency is related to the bulk dielectric function £(a>): I(co) =

2e^ ^ -1 Im hojvj^ £(co) -h 1

With increasing beam energy above about 100 eV, the mean free path of electrons into solids increases and the inelastic scattering is increasingly dominated by penetration of electrons into the bulk of the solid. The loss function is then dependent on Im(-l/£(co)), In the literature EELS data are often compared with bulk optical excitation data. However it should be remembered that bulk absorption spectroscopy measures Im(8(cD)) and that in general there are well-defined shifts between the transverse resonance frequencies found optically and the longitudinal frequencies for bulk losses in EELS. Surface loss energies lie between bulk transverse and longitudinal frequencies. The shifts are particularly obvious when dealing with vibrational losses, but the same principles apply to electronic excitations. 3.3 Scanning tunnelling spectroscopy Scanning tunnelling microscopy (STM) is a proximal probe technique in which an atomically sharp tip is held sufficiently close to a surface to allow significant overlap between the tails of the wavefunctions of electronic states associated with the tip and the surface respectively. Application of a small bias between the tip and the surface therefore allows electron tunnelling between the two. In the simplest models the magnitude of the tunnelling current varies as e'^, where d is the tip surface separation. The strong dependence on distance has allowed development of STM as a surface microscopy: as a tip is scanned across a surface the tunnel current will vary in a way that reflects the surface topography. However it must also be realised that STM is strongly influenced by electronic structure. Thus under conditions of positive sample bias, electrons tunnel into empty electronic states in the solid, whilst at negative bias electrons tunnel out of filled states in the solid into the tip. For an oxide with filled O 2p bands and empty metal-based bands it therefore follows that images taken at positive sample bias from a "flat" surface containing both O ions and metal ions in the same plane will be dominated by empty metal based states. At negative sample bias filled oxygen-localised states will be probed. In the TersoffHamman model [27] contours of constant tunnel current above a surface are simply contours that mirror the filled or empty state electron density in the relevant energy window. The variation of tunnelling current with sample bias is

558

intimately related to the density of states at the surface. Again assuming the simplest model, the density of states dN(E)/dE may be shown to be proportional to the normalised differential conductance i.e. dN{E) dI{E)ldV{E) dE °^ I(E)/V{E)

The measurement of tunnelling spectra in a scanning tunnelling microscope offers the potential of measuring the local density of states at spatially defined sites whose topography can be established at an atomic scale by STM. This information is only available however at the price of losing the information about the k-dependence of electronic states that is available in photoemission and inverse photoemission. In particular STS offers the prospect of measuring local densities of states at defect sites whose real space atomic structure can be established by STM. Given the potential of the technique there have been remarkably few studies of the electronic structure of defects by STS. One of the most satisfying early studies was of bandgap states at the (V5xV5)R26.6° surface of SrTiOsClOO) [28,29]. This surface was originally described as (2x2) [30], but is now believed to terminate in a TiOi surface plane with an ordered arrangement of in-plane oxygen vacancies. The vacancies are imaged as grey-scale maxima at negative sample bias in STM i.e. the image is dominated by filled electronic states just below the Fermi level. STS shows the vacancies to be associated with filled electronic state 1.35 eV below the Fermi energy. A similar state is observed in photoemission from oxygen deficient SrTiOsClOO) surfaces, but photoemission provides no guidance as to the spatial location of the state. The application of STS to defective TiOi and WO3 surfaces is discussed in section 4. 4. SELECTED CASE HISTORIES 4.1 The rocksalt MgO(lOO) surface*. MgO is a transparent insulator with a gap of about 7.8 eV separating a filled band of O 2p states from an empty band of Mg 3s states: Mg 3p states lie at higher energy. The lowest excitation in bulk MgO is not a simple band to band transition, but involves an excitonic process. "Conventional" excitons in oxides such as Li20 and SiOi are associated with localisation of the hole on an oxygen anions and the electron on an adjacent cation. The cation and anion involved undergo large asymmetric relaxations and there is therefore a substantial Stokes shift between the energies associated with the vertical excitonic absorption and * See the chapter by Freund et al for diagrams of the MgO(lOO) surface.

559

the excitonic luminescence from the self-trapped, relaxed configuration. In MgO by contrast it appears that the electron and the hole both reside on the same O site: the excitation may be viewed as an on-site O 2p->0 3s transition. The relaxation around the exciton is symmetrical and involves smaller energies than for the two-centre exciton. The Stokes shift is therefore small and the lower self trapping energy gives rise to much higher exciton mobility than is found for say Si02[31]. The (100) surface of rocksalt oxides typified by MgO represent perhaps the simplest of oxide surfaces. The ionic layer which terminates this surface is electrostatically neutral and contains equal number of Mg^"^ and O^" ions in interpenetrating square arrays. The (100) surface energy is lower than for any other surface of this material and bulk MgO crystals cleave easily to expose high quality (100) surfaces. LEED [32,33] and medium energy ion scattering demonstrate [34] that the rumpling of the Mg(lOO) surface is rather small, with inward relaxation of Mg ions and outward relaxation of O. In detail the surface structure as determined by surface X-ray diffraction involves [35] inward relaxation by 0.56% of the interlayer separation between MgO planes and 1.07% rumpling. These observations are reproduced by atomistic simulations [36], as well as periodic ab initio Hartree-Fock [37,38] and density functional calculations [39,40]. There is no particular indication of enhanced covalency at the surface as has been claimed on the basis of the low corrugation at the surface observed in He atom diffraction [41]. Experimental difficulties in the study of MgO surfaces by electron spectroscopic techniques may be anticipated because of the highly insulating nature of the material. Photoemission is indeed somewhat problematic, but techniques such as EELS and LEED, in which an electron beam impinges on the surface, prove to be viable even on bulk crystals. This is because the secondary electron yield coefficient (that is the number of secondary electrons emitted from the surface per incident electron) is greater than one above a threshold energy of around 20 eV and may reach values in excess of 10 at higher beam energy [42]. The sample therefore tends to charge positively and does not repel the incoming electron beam. HREELS experiments with beam energies below the 20 eV threshold are possible using a subsidiary high energy electron gun to promote an outward electron flux [43]. The alternative approach to dealing with sample charging is to prepare thin MgO overlayers on a metallic substrate. Goodman and co-workers [44-47] in particular have perfected recipes for growth of (100) oriented layers of MgO by evaporation of Mg onto an Mo(lOO) substrate in an oxygen ambient with pressures of the order 10'^ mbar. Two issues have been of interest particular in relation to the electronic structure of Mg(lOO) surfaces. The first is whether the electronic structure at the terrace surface differs in any significant and experimentally observable way

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from the bulk. The second concerns the nature of electronic excitations associated with defects at (100) surfaces. Most features in angle resolved He(II) photoemission measurements on cleaved MgO(lOO) surfaces can be adequately rationalised in terms of a direct transition model based on the bulk bandstructure, although some features fall outside the range of bulk dispersion around the M point [48]. Owing to sample charging it is difficult to fix the position of the valence band edge relative to the vacuum level: the value of over 9 eV that was proposed by Sawatzky and coworkers [48] is probably too high. A more surface sensitive way of exciting valence band electron spectra is provided by metastable impact electron spectroscopy (MIES) [49-51]. In this technique a beam of He atoms in a ^S excited state associated with the configuration ls^2s^ impinges on the surface. At an insulating wide band gap surface such as that of MgO, the excited atoms decay via an intra-atomic Auger mechanism (sometimes called Penning deexcitation) in which an electron from the surface fills the Is hole and the 2s electron is emitted with the excess energy available. The MIES spectra from MgO(lOO) are much narrower than the photoemission spectra and essentially reflect the density of states associated the O 2pz orbitals which protrude out of the surface [49]. A more sophisticated approach to calculation of the MIES spectra involves projection of the surface density of states onto the Is orbital of the He* ion in its trajectory toward the surface [51]. These calculations provide closer agreement with experimental spectra, but reinforce the conclusion that the band of O 2pz surface states is narrower than the bulk valence band. The most direct evidence for differences between bulk and surface electronic structure comes from EELS. Using the technique in the differential mode Henrich et al found a loss peak at just over 6 eV on non-defective MgO(lOO) surfaces, an energy well below the onset of 7.5 eV for bulk excitations [52,53]. The relative intensity of this loss feature decreased dramatically on increasing the energy of the exciting beam from 100 eV to 2000 eV, as expected for a surface excitation [53]. Mg 2p and 2s core excitation spectra show a similarly shifted surface peak to low energy of the bulk core excitation threshold. HREEL spectra excited using a monochromatised electron beam from both bulk and thin film single crystal surfaces (fig. 2) illustrate the lowering of the excitation threshold even more dramatically [54,55]. The EELS data were originally interpreted in terms of simple bandgap narrowing at the surface, as expected from the reduced Madelung potential at surface sites [52,53,55,56]. However, Williams argued that the reduction in Madelung constant from 1.74765 in bulk sites to 1.68155 in surface sites is too small to produce the observed reduction in the excitation energy: in any case reduced valence and conduction bandwidths at the surface will oppose the Madelung term and the net effect might actually be to increase the gap [57]. It is informative therefore to consider quantitative

561 350

HREELS MgO(100)/Mo(100) 262

Ep=:48eV

N

X

t 175 (0

z

lU

87

1 2

3

4

5

6

LOSS ENERGY (eV)

7

8

9

^.2

Fig. 2. HREELS of MgO(lOO) thin film on Mo(lOO) excited at 48 eV beam energy. In moving upward the initially stoichiometric film is subject to annealing to the progressively higher temperatures indicated. Adapted from ref. 72.

attempts to model the electronic structure of MgO surfaces using bandstructure and cluster models. The first periodic calculation of energy levels for MgO(lOO) surfaces used a semiempirical LCAO model applied to an unrelaxed surface. Surfaces states were found w^hich fell outside the range of energies of the projected bulk band away from the T point. However the lowest energy direct surface excitations were found at 7.6 eV, very little below the calculated bulk bandgap of 7.8 eV [58]. Similarly small differences were found when relaxation was allowed using a tight binding total energy method [59]. On the other hand discrete variational cluster calculations using MgOe^^" to represent bulk MgO and MgOs^" to represent the (100) surface gave a clearcut indication that the bandgap should be reduced at the surface [60,61]. Essentially the bandgap narrowing arises from mixing between Mg 3s and 3p levels under the influence of the electric field gradient at the surface. This splits a hybrid 3s-3p state away from the bottom of the conduction, thus narrowing the bandgap. There is no significant lifting of

562

states above the top of the O 2p valence band. More recently ab initio calculations have been performed based on density functional theory applied to a fully relaxed MgO(lOO) surface represented by a periodic arrangement of slabs separated by vacuum gaps. As in the cluster calculations, a hybrid state localised mainly above the surface Mg ions in an 3s-3p hybrid state splits off from the bottom of the conduction band [40]. This result is independent of slab thickness. Despite these theoretical results, it is probably too naive to interpret the EELS results simply in terms of surface interband transitions, given that the threshold bulk excitation are excitonic in nature. In a survey of HREEL spectra of a wide range of compounds with the rocksalt structure it was found that a difference between surface and bulk excitation thresholds was only found for a limited range of materials including MgO, CaO, LiF but not NaCl, NaBr, KCl and KBr [54,57]. For compounds in the first group it was argued that the conduction band edge lies above the vacuum level giving a negative electron affinity - the high secondary electron yield coefficient for MgO discussed above has been attributed to this unusual arrangement of energy levels [62,63]. In this situation the lowest surface excitation is from the valence band edge to the vacuum level: binding of the electron at the vacuum level to the surface by its image potential gives rise a novel form of surface exciton where the electron sits outside the surface. This qualitative picture for the surface exciton receives support from calculations of Shluger et aL, who used a semiempirical intermediate neglect of differential overlap (INDO) model in which a Mg24024 cluster is embedded in a 5 layer periodic slab. The basis set included 12 floating Slater functions [31]. The surface exciton involved on site excitation of O 2p electrons, with outward displacement of the five near neighbour magnesium cations. The electron population of the floating orbitals located outside the surface plane was larger than those inside in the excited state. This model predicted a surface excitation energy of 6.4 eV, in excellent agreement with the experimental value. The surface exciton model of Cox and Williams was criticised by Saiki et al [64] who argued that the model does not account for the low energy metal core excitations observed for MgO and fluorides such as NaF. However there seems to be no reason why there should not be two distinct excitonic processes involved, with metal core level excitation into a surface metal sp hybrid and valence excitation into an on-site anion 3s state. Very recent calculations using an embedded cluster model give a value of 6.7 eV for the energy of the top of the O 2p valence band relative to the vacuum level [65]. This is in agreement with MEIS measurements on thin film MgO(lOO). Given the bulk bandgap of 7.8 eV, these calculations further reinforce the view that the conduction band minimum does indeed sit above the vacuum level. We turn next to consider defective MgO(lOO) surfaces. In contrast to the transition metal oxides to be discussed subsequently, argon ion bombardment

563

has little effect on UV photoemission spectra of MgO [48]. However Henrich et al [52,53] observed a loss feature in dN/dE energy loss spectra of UHV cleaved MgO(lOO) at about 2 eV that was obviously a surface excitation, as gauged by the decreasing intensity with increasing beam energy. These observations were confirmed by subsequent work using higher resolution [56,57,66]. The intensity of the feature increased upon electron irradiation of the crystal surface, but decreased on oxygen exposure. In most oxides electron beams induce oxygen deficiency, so the empirical observations support assignment of the loss feature to a surface F centre, that is a an oxygen vacancy with two associated electrons trapped in the vacancy site [52,53]. However, the excitation energy for a surface F centre was calculated to be 5 eV [67,68], leading to an alternative assignment [56,66] to a surface cation vacancy or V-centre whose excitation energy should be close to the value of 2.3 eV observed in non-differentiated EEL spectra [56]. This assignment receives some support from the observation that the loss feature is quenched by deposition of Cu [69], which would be expected to refill cation vacancy sites. On the other hand Tegenkamp et al found that loss intensity around 2 eV is enhanced by partial desorption of excess Mg from an MgO film surface: this experiment is hardly compatible with formation oi metal vacancies [70]. In a more detailed experimental study, Wu et al. [71] studied the effects of high temperature annealing in UHV on HREEL spectra of thin film singlecrystal MgO(lOO). The loss onset was just above 6 eV for as-prepared films, but UHV annealing at 1400 K led to growth of loss features at 1.15 eV, 3.58 eV and 5.33 eV. The first of these loss peaks is at lower energy than in other energy loss studies. However, it is difficult to compare HREEL data with data taken at lower energy resolution using non-monochromatised electron beams, especially when data are acquired in the differential mode. The 1.15 eV feature was assigned to surface F centre excitation, the loss at 3.58 eV to F centre aggregates and the loss at 5.33 eV to F centres in the "bulk" of the thin film. The third assignment here was based on two considerations. First the bulk F centre excitation in electron irradiated MgO crystals is at about 5 eV. Second it was observed that the relative intensity of the 5.33 eV loss decreased as the beam energy increased from 15 eV to 48 eV. The electron pathlength in the solid is undoubtedly greater at 15 eV than 48 eV, but the argument that the intensity variations establish the 5.33 eV loss as a bulk excitation is not valid if the dominant mode of inelastic scattering is by the surface dipole mechanism. The empirical observations relating to the low energy loss peak all seem to favour assignment to oxygen rather than magnesium vacancies. On the other hand more recent theoretical work has not led to a complete understanding of the experimental results. Periodic LDA calculations due to of Kantorvich et al establish that surface F centres have a formation energy of about 9.8 eV, which is significantly lower than the formation energy of 10.5 eV in the bulk. The electronic level of the relaxed surface F centre was only 2.3 eV above the

564

valence band maximum and therefore below the corresponding bulk F centre, which is 2.7 eV above the valence band [40]. This is in fair agreement with results from an embedded cluster calculation which places the surface F centre 3.25 eV above the valence band maximum and 3.43 eV below the vacuum level [65]. However these were both essentially ground state calculation that did not seek to calculate excitation energies explicitly. By contrast Pachioni and coworkers used cluster models to evaluate energies of excited states directly. Both neutral F centres and singly ionised F^ centres were considered. These calculations allowed for configuration interaction to take due account of correlation in both ground and excited states. The bulk singlet to singlet excitation energies for F and F^ centres were found to be at 6.00 eV and 5.75 eV [72], as compared with experimental values of 5.03 eV and 4.96 eV [73]. The bulk transitions correspond to dipole allowed Aig-^ Tiu and Aig-> T lu excitations. At the surface the threefold degeneracy of the Tiu states is removed. Surface F centre excitations corresponding to Âi->Âi and Âi^Ê processes are found at 3.24 eV and 4.22 eV, with corresponding doublet to doublet transitions for the surface F"^ centre at 3.47 eV and 4.52 eV. Application of a scaling factor of 0.84, which accounts for the error in the calculated bulk excitation energy, leads to a best estimate of 2.72 eV for the Âi^Âi transition. This is of course higher than the experimental values of 2.3 eV and 1.15 eV discussed earlier. We note in passing that the Âi^Âi transition is surface dipole allowed but the Âi—>Ê excitations is forbidden and would not be observed in HREELS if a surface dipole scattering mechanism were predominant. There remain at least two further possibilities to account for the low energy surface excitations. The first is that the oxygen vacancies are found at 4- or 3-coordinate step or comer sites rather than 5-coordinate surface terrace sites. The lowest singlet to singlet excitation energies for step and comer F centres are calculated as 2.92 eV and 2.60 eV respectively [74], which scale down to 2.45 eV and 2.18 eV using the correction factor of 0.84. These energies are quite close to the values of 2.0-2.3 eV found in the earlier experimental work. Thermodynamically it is to be expected that oxygen vacancies will congregate at step and comer sites because the oxygen removal energies of 9.00 eV and 8.06 eV respectively are both lower than the value of 9.77 eV for terrace sites [40]. The second possibility is that low energy excitations are associated with F centre aggregates at the surface. Using a semiempirical Hartree-Fock approach Castanier and Noguera found that the energy required to create surface F centres decreased with increasing levels of oxygen deficiency, which can be interpreted in terms of an attractive interaction between oxygen vacancies [75,76]. However, later ab initio density functional calculations demonstrated that the interaction is repulsive [77]. The Hartree-Fock result is probably an artefact of the method, which requires the electron density associated with the F-

565

centre to be accommodated in neighbouring magnesium orbitals. This leads to a large quadrupole moment and attractive interactions between the quadrupoles. The more sophisticated density functional calculations suggest that the F centre electron density is trapped mainly in the oxygen vacancy site and that the localised electron densities repel each other. 4.2 Transition metal rocksalt monoxide (100) surfaces As discussed in section 2.1, the oxides MnO through to NiO are all insulators with localised d" configurations. In the simplest picture the bulk octahedral crystal field of the oxide ions splits the d levels into a lower tag set and an upper Cg set. This ordering is actually reversed in Hartree-Fock calculations on NiO, although the ground state configuration remains tigQg [78,79]. By contrast cations at the (100) surfaces sit in a site of C4V symmetry and the d levels split to give eSn 5p excitation threshold [14]. Despite earlier claims to the contrary [137], it also appears that surface reduction produces a chemically shifted Sn^^ core level peak [131]. The reason why the Sn^^ defect levels lie close to the valence band edge has proved to be controversial. Based on parameters from the tight binding calculation of Robertson [138], Munnix and Schmeits [139-141] used a scattering method to calculate the surface band structure for the bare (i.e. no bridging oxygen) (1x1) reconstruction of Sn02(l 10). Projecting the surface

576

bands onto the bulk bandstructure, two surface states were found in the so-called stomach gap toward the bottom of the main O 2p valence band. In addition a number of weak resonances were found in the upper part of the valence band. In the conduction band region, three surface states were found, the lowest of which was only 0.1 eV below the conduction band minimum at the T point. Thus the bandgap region was found to be essentially free of surface states. More recent tight binding calculations by Godin and LaFemina [142] reached a similar conclusion. A way out of this dilemma is provided by the suggestion of P.A. Cox et al [130] that the surface state arises from mixing between 5s and 5p atomic orbitals at surface sites where the essentially centrosymmetric coordination of bulk sites is lost, as shown schematically in fig. 9. In general mixing between s and p levels under the influence of a potential V is described by a matrix element and will be non-zero only if V has a component belonging to the same irreducible representation as one of the 5p orbitals. This is never possible at a site with a centre of symmetry because the 5s and 5p orbitals are of opposite parity and cannot be mixed by centrosymmetric crystal fields. The condition for mixing is however satisfied by tin ions in the 4- and 5coordinate surface sites of the bare (1x1) reconstruction and will generally pertain at any surface cation site where the coordination is reduced below its bulk value of 6. The essential physics of this qualitative model has recently been confirmed [143,144] by first principles density functional calculations on SnOaCllO). In particular Mannasidis et al, [143] found a broad band of states in the bulk bandgap for the fully reduced (1x1) surface: on the half-reduced surface the band of gap states contracts to have a width of 1.5 eV and has maximum intensity about 1.1 eV above the valence band maximum. In the calculations the electron density associated with the gap states is mainly localised above bridging tin sites, with some weight on in-plane oxygens and elsewhere. This accords closely with the chemical picture of a directional 5s5p hybrid lone pair localised on Sn(II) ions. Gradient corrections within the local density approximation had little effect on the relaxed surface structure, although the surface energies are reduced by about 30% [145] The failure of tight-binding calculations to reproduce the observed gap states arises simply from the fact that the tight-binding Hamiltonian has no term which accounts for the 5s-5p mixing and includes no effects due to surface fields [143]. Aside from giving rise to the occupied 5s5p surface state, the hybridisation depicted in fig. 9 should also produce an empty surface state above the main band of empty Sn 5p states. Growth of intensity in this region has been observed in inverse photoemission spectra of ion bombarded Sn02 surfaces [146].

577

4.4Ti02*

4.4.1 Comparison ofTi02 andSn02 Ti02 is another oxide adopting the tetragonal rutile structure with a=4.594lA and c=2.9589A. Despite the fact that the metal oxygen bondlength is smaller in Ti02 than Sn02, metal-oxygen covalency is less pronounced in the former and the dynamic charge derived from analysis of IR reflection spectra is greater in TiOi (0.61 of the formal ionic charge) than in Sn02 (0.49 of the formal ionic charge). This has the consequence that the O 2p valence bandwidth and bandgap in Ti02 (6.3 eV and 3.062 eV respectively) are less than in Sn02 (where the corresponding values are 10.0 eV and 3.596 eV) [147,148]. These differences appear to derive from the fact that the conduction band states in Ti02 are based on relatively contracted Ti 3d atomic orbitals, whereas in Sn02 the conduction bands involve more diffuse 5 s and 5p orbitals. It should be emphasised that the covalency in Ti02 is not insignificant. In particular strong admixture of Ti 3d character into O 2p states toward the bottom of the valence band provokes strong resonance effects in photoemission at the Ti 3p->Ti 3d excitation threshold [149]. Resonance behaviour is also observed at the Ti 3p^'Ti 4s threshold [150]. Of course UV photoemission is intrinsically surface sensitive so that these experiments are in effect probing covalency at surfaces: using on and resonance data the total density of states can be decomposed into atomic components. Courths and co-workers [151,152] have raised the intriguing possibility of exploring differences between partial densities of states derived on the one hand from resonant photoemission and on the other from core and valence level photoelectron diffraction, which is not so strongly surface sensitive. However these ideas remain to be fully exploited. The bulk oxygen vacancy defect chemistry of Ti02 is very different to that of Sn02. In the tin-oxygen phase diagram, Sn02 tolerates only a narrow range of non-stoichiometry and there are no ordered defect structures connecting Sn02 and the lower oxide SnO. By contrast in the titanium-oxygen system, an essentially continuous range of compositions between Ti02 and TiOi.75 (i.e. Ti407) is possible. For very small values of x in Ti02-x, non-stoichiometry is accommodated by point defects, although there is some ambiguity as to whether these are oxygen vacancies or titanium interstitials. In the simplest ionic picture each oxygen vacancy is associated with reduction of two Ti"^^ ions to Ti^^. For x values greater than about 10"^ the defects cluster to give first Ti^^-Ti^^ pairs in face sharing octahedra, which then order into shear planes running along {132} crystallographic directions. The {132} shear planes order at yet higher deviation from ideal stoichiometry to produce members of the homologous series of Magneli phases Tin02n-i (15:-o.9 TiO.(llO) ^ , ^ „, hv"=22eV

II101

A

-PL

PC

Z-0.8i -0.9 -\

binding energy / eV

-1.0 -I

Fig. 11. Left hand panel: bandgap feature in photoemission (hv = 22eV) of thermally annealed TiOaCllO) as aftinctionof exposure to oxygen (a) OL (b) 10 L (c) 100 L (d) 2500L. Oxygen is seen to quench the gap state. Right hand panel: dispersion of the bandgap feature along rX and FY directions. Adapted from ref. 211.

eV above the Fermi level for perfect surfaces [165,214], new features appear at 3 eV for defective surfaces [215]. These are predicted by the calculations of Munnix and Schmeits discussed above. We turn next to consider shear planes at TiO2(110) surfaces. The earliest STM experiments on Ti02(l 10) surfaces that had been subject to very prolonged high temperature annealing in UHV revealed structures that wêre attributed to shear planes [216,217], although this interpretation was later called into question [155]. More recently Norenberg et al [218] observed periodic steps on Ti02(l 10) that were consistent with formation of {112} shear planes. However, shear planes of this sort are not a feature of the bulk defect chemistry and the interpretation of these experiments was complicated by the simultaneous segregation of Ca to the (110) surface [219]. Most recently Bennet et al [220] have provided clearcut LEED and STM evidence for formation of the expected {132} series of shear planes at TiO2(110) surfaces following UHV annealing at 1223K for several hours. It was predicted many years ago that the 3d electron states associated with trapping of electrons at shear plane sites are lower in energy by about 0.3 eV as compared with simple self trapped electron states [221]. It would therefore be of interest to compare photoemission spectra from surfaces where shear planes provide the dominant mode of accommodation of non-stoichiometry with spectra from surfaces where point defects predominate. This experiment remains to be performed. Similarly there is considerable scope

583

for careful STS experiments which could in principle establish differences (if any) between the Ti 3d electronic levels associated with bridging and subsurface oxygen point vacancies and Ti 3d levels at shear plane sites. Early STS experiments on the effects of oxygen deficiency on the electronic structure of TiO2(110) surfaces were mostly performed under conditions where the topographic resolution was rather poor [222,223]. Nonetheless, oxygen deficiency was seen to produce clear signatures in lA'^ spectra. At stoichiometric surfaces, the spectra showed typical n-type semiconductor behaviour with significant empty state tunnelling currents at low positive sample bias, but essentially zero tunnelling current at negative sample bias. However upon reduction, filled state tunnelling was observed, the spectra from ion bombarded surface being almost metallic in appearance. Following these experiments little effort seems to have been devoted to STS under conditions where good topographic resolution is obtained on Ti02(l 10) and further work in this area is urgently needed. Finally we consider reconstructions induced by oxygen deficiency on TiO2(110). Moller and coworkers first characterised a (1x2) reconstruction on the (110) surface [224]. The initial model for TiO2(110) (1x2) was based on missing rows of bridging O ions, as discussed above in the context of Sn02(l 10) (1x2). Electron book-keeping dictates that if there are bridging oxygen vacancies, some surface or near surface Ti"*^ ions must be reduced to Ti^^. In agreement with these ideas strong bandgap intensity was found in photoemission from the (1x2) surface. Atomically resolved STM images appeared initially to provide support for the missing row model [175] but also gave evidence for significant lateral relaxations away from bulk terminated positions. However several problems were encountered with the model. Firstly, LDA calculations predicted that a (2x1) arrangement of bridging oxygen vacancies would be more stable than a (1x2) arrangement [157*, 225]. Secondly and most surprisingly Moller and coworkers in a second paper contradicted their own initial observations and found that a (1x2) surface could be prepared which essentially gave zero bandgap intensity in resonance photoemission [150]. Thirdly it was observed that the (1x2) reconstruction grew as an overlayer on top of the (1x1) surface when a reduced subsurface was annealed at elevated temperatures either in UHV or oxygen [170,173,178,226]. On the basis of dynamic STM experiments at high temperatures Onishi et ah suggested a radically different model for TiO2(110) (2x1) involving added Ti203 rows [170,173]. This model was shown to be consistent with measurements of O^ electron stimulated desorption ion angular distributions (ESDIAD), which revealed two parallel stripes rather than the single stripe expected for the missing row model [227]. Note however that in reference [157] the (1x2) reconstruction discussed here is called (2x1).

584

The added rows of the Onishi model are oxygen deficient with respect to a stoichiometric termination of the surface and simple considerations therefore suggest that 3d^ states associated with Ti^^ must be produced at such a surface. It is not therefore possible to reconcile this model with MoUer's observations [150]. The assertion that the (1x2) reconstruction supports no bandgap states has however itself been called into question by Patel et al [228], who found pronounced Ti 3d intensity in the gap in photoemission spectra excited with 47 eV photons. But this study in turn threw up the puzzling result that the bandgap intensity was not associated with chemically shifted Ti 3p structure in shallow core photoemission. More recently Thornton and coworkers have suggested a fiirther alternative model for the (1x2) reconstruction [229]. This involves added rows in which the Ti and O positions are close to those for a bulk terminated surface, but the added rows were presumed to terminate without bridging O ions. Thus we have another "reduced" termination, but unlike in the Onishi model it is possible to envisage a stoichiometric version of this reconstruction with bridging oxygen left in place. Thornton's model was later criticised by Tanner et al. [185], who carried out atomically resolved STM measurements that achieved better resolution of structure within the unit cell than in earlier work. They claimed that the Thornton model gave the wrong registry between (1x1) and (1x2) reconstructions. However, this claim has in itself provoked fiirther controversy [230,231]. The most recent dynamic elevated temperature STM measurements suggest that the two different added row (1x2) reconstructions may co-exist [232,233]. Further complexities in the surface chemistry of TiO2(110) have been uncovered by Diebold and coworkers who have made extensive and systematic studies of the "restructuring" of TiO2(110) surfaces by annealing bulk reduced (and therefore blue!) crystals in oxygen or UHV [234-238]. Restructuring arises by elimination of excess Ti present as either interstitials or in shear planes in the sub-surface region. Lightly reduced crystals only display (1x1) reconstructions after UHV annealing, whereas more heavily reduced crystals display areas of a (1x2) reconstruction that probably correspond to Onishi Ti203 added rows. The most heavily reduced crystals give rise to pseudo-hexagonal rosette structures upon annealing in oxygen. The rosette structure involves an incomplete Ti02 layer with both Ti and O vacancies. However the remaining ions occupy essentially bulk-like positions. The rosette overlayer is essentially stoichiometric and supports no occupied bandgap states. However, an LDA electronic structure calculation of the rosette overlayer predicted a significant widening of the bandgap as compared to the bulk value due to enhanced Ti-0 covalency within the rosettes [236]. In summary we conclude that despite the enormous effort devoted to the study of Ti02(l 10), the relationships between geometric and electronic structure

585

are not entirely clearcut. In particular there is still scope for definitive spatially resolved scanning tunnelling spectroscopic measurements that could in principle define the distinctive features of electronic structure associated with missing bridging oxygen atoms and subsurface oxygen vacancies on TiO2(110) (1x1); the differing (1x2) reconstructions; oxygen deficient shear planes in their intersection with the (110) surface; and the stoichiometric rosettes. It is probably impossible to prepare a surface where these different features do not co-exist to some extent, so that photoemission is too crude a tool to explore these subtleties of electronic structure. 4.43 TiO2(100) The TiO2(100) surface provides another example of the way in which understanding of the relationship between the geometric and electronic structure associated with defects at surfaces has evolved in the past 10 years. The earliest experiments on cleaved (1x1) TiO2(100) surfaces found very weak photoemission intensity in the bulk bandgap at hv=21.2 eV [153], presumably associated with defects generated during cleavage. In previous work it had been found that bombardment and annealing of TiO2(100) surfaces led to formation of superstructures, originally described as (1x3), (1x5) and (1x7) [239,240], the latter two associated with 3d electron states in the bulk bandgap. However, subsequent work showed that there are probably only two reconstructions, both of which can be produced by ion bombardment followed by well-defined annealing procedures. A stoichiometric (1x1) surface was obtained by annealing a bombarded surface in 10"^ mbar O2 at 870 K for 30 minutes, whereas a (1x3) surface was produced by annealing at the higher temperature of 1020-1070 K in UHV. Most surprisingly it was found that a small miscut which would have been expected to produce steps along the [001] direction led to LEED spot splitting in the [010] azimuth: the spot splitting may be misinterpreted in terms of higher order (Ixn) reconstructions [241]. Photoemission spectra measured at the resonance energy for the Ti 3p core threshold hv=47 eV (fig. 12) showed that the (1x1) surface supported no bandgap states, whereas for the (1x3) surface there was pronounced emission intensity in the gap [241]. The gap state also showed resonance enhancement at the Ti 2p core threshold [242]. Similar bandgap intensity is produced by Kdeposition [243,244]. A simple model for the (1x3) reconstruction was put forward, involving removal of every third row of on-top oxygen ions from an essentially bulk terminated (100) surface. The intensity of the bandgap emission was consistent with this model. However, grazing incidence X-ray diffi*action [245] led to development of a radically different model for the (1x3) surface

586

1 8

1

1

r

6 4 2 Binding Ener©^ (eV)

Fig. 12. Normal photoemission spectra of the (1x1) and (1x3) reconstructions of TiO2(100) excited at photon energies indicated. At 47 eV there is resonant enhancement of structure associated with Ti 3d states. The angle integrated spectrum at 47 eV is also shown. Adapted fromref 241.

based on {110} nanofacets*: the driving force for this reconstruction is simply the lower surface energy of rutile (110) as compared with rutile (100). Atomically resolved STM images of TiO2(100) (1x3) are consistent with the nanofacet model [247,248], which also provides a simple rational for the updown steps along [010] in terms of maximisation of the area of {110} facets [249]. The nanofacet model suggested on the basis of the XRD data involved an essentially stoichiometric termination and was not therefore consistent with the photoemission data. Thornton and co-workers therefore suggested that oxygen ions on the ridges of the nanofacets must be removed to give "bare" Ti ions at the top of the ridges [247,248]. Charge book-keeping then requires that some Ti ions in the near-surface region must be reduced. In agreement with the valence photoemission data, Ti 3p shallow core level photoemission spectra showed a chemically shifted component associated with Ti^^ for TiO2(100) (1x3). The * In the original literature the facets are referred to as microfacets. However, the length scales involved are in the nanometre range and we prefer the term nanofacet [246].

587

bare Ti ions on top of the ridge are the obvious candidates for reduction to Ti^"^. To explore this possibility, the azimuthal angular dependence of the intensity of the chemically shifted Ti 3p emission was measured and compared with photoelectron diffraction (PED) curves calculated using a multiple scattering formalism [250]. Qualitatively, computed PED curves with O vacancies and Ti^^ ions on the tops of the ridges gave better agreement with the experimental data than could be achieved with models assuming O vacancies on other possible sites on the nanofacets. R factor analysis confirmed this conclusion. More recent STM work has investigated surfaces where the transformation between the stoichiometric (1x1) reconstruction and the facetted (1x3) reconstruction is not complete. Two intermediate reconstructions with (1x3) periodicity have been identified and a plausible mechanism for their interconversion has been suggested [251]. These ideas have been further refined in the light of very recent non-contact atomic force microscopy (AFM) [252,253] and further surface XRD data [254] which has explored the magnitude of relaxations within the nanaofacet structure. On the basis of the discussion thus far, it appears that there is an excellent understanding of the relationship between the structure of TiO2(100) (1x3) surfaces and the nature of defect states. However two issues do remain to be resolved. First, in the light of the STS data for oxygen deficient SrTiOs alluded to in section 3.3 it might be anticipated that spatially resolved scanning tunnelling spectroscopy would provide evidence for occupied electronic states just below the Fermi energy on top of the ridges of the nanofacets. In fact the tunnelling spectra show that on the rows there is no filled state tunnelling and the I-V spectra are completely dominated by tunnelling into empty states. By contrast filled state tunnelling is found for sample biases negative of -1 V between the rows [247,248]. Second, a completely new model for the structure of Ti02 has been proposed on the basis of the most recent analysis of X-ray diffraction experiments [255]. An ab initio direct methods approach was used to re-analyse the original surface X-ray diffraction data set. The direct methods solution to the phase problem is widely used in determination of bulk structures from XRD data and leads to electron density maps without presupposition of a trial structure. The alternative model for the (100) surface derived in this way involves up-down topography with 3a periodicity. However the structure is quite distinct to that of the nanofacet model, which retains essentially the same linkage between TiOe octahedra as is found in the bulk structure. The new model involves extensive O deficiency in the near surface region and fusion of TiOe octahedra to eliminate O point defects in a fashion similar to that found in bulk shear planes. Clearly the new structure will be associated with bandgap states in photoemission but it remains to be seen if the model can be reconciled with the observed photoemission intensity, with the shallow core photoelectron diffraction data discussed above and with the extensive body of STM and AFM observations.

588

4.5W03andNaxW03 4,5.1 WOsiOOl) Tungsten trioxide (WO3) is a prototype d^ oxide based on a framework of comer sharing W06 octahedra. The formal oxidation state of tungsten is 6^, although due to strong covalency, the W ions do not carry the full 6^ charge. The neighbouring oxide ReOs has a simple cubic structure with an Re-O-Re angle of 180°, but in WO3 rotation and distortion of the WOe octahedra gives rise to a series of lower symmetry phases: there are at least five bulk phase transitions between 100 K and lOOOK. A monoclinic phase Y-WO3 is stable between 290K and 603K. It has cell parameters a = 7.297A, b = 7.539A, c = 7.688A and P = 90.91° and belongs to the space group P2i/n. The y-WOs structure represents a (2x2x2) superstructure based on an idealised cubic unit cell of dimension about 3.7A [256]. Crystals of WO3 cleave most easily to expose [001] surfaces, which because of the monoclinic symmetry are distinct from [100] and [010] surfaces. Stoichiometric monoclinic tungsten trioxide is an insulator at 300K with a band gap of about 2.6eV. However, WO3 readily becomes oxygen deficient to form WO3-X with variable oxygen composition parameter x. This oxygen deficiency greatly influences the bulk electronic transport properties by introducing donor electronic states into the upper half of the bulk bandgap: charge bookkeeping requires that each oxygen vacancy converts two W(VI) ions to W(V). For defect concentrations of x >10''^, point defects are eliminated by the formation of crystallographic shear planes belonging to the {ImO} series. These involve edge-sharing W06 octahedra [257]. In general the higher the degree of oxygen deficiency, the smaller m. Thus for very lightly reduced WO3, the shear planes are {120}, swinging round toward {010} for very highly reduced material. Electron spin resonance measurements on WO3.X (x>10"'^) in the low temperature 8 phase indicate that the 5d^ electrons associated with reduced W^^ centres are paired to give bipolarons [258]. The lattice distortion within the polaronic centres probably involves motion of the two W^^ ions toward each other. A simple rationale for the driving force for shear plane formation therefore involves bond formation between W^^ ions in edge-sharing pairs. A number of early studies by X-ray photoelectron spectroscopy (XPS) demonstrated that surface reduction of WO3 by thermal treatment in UHV or by exposure to electron and ion beams or UV photons could be monitored by changes in W 4f core level spectra [259-265]. Subsequently Bringans and coworkers used UV photoemission to follow changes in valence band spectra resulting from electron or ion bombardment of WO3(001) [266-267]. It was found that reduced surfaces were characterised by a broad and generally illdefined band of W 5d states extending from the valence band edge to the Fermi level: The overall spread of these defect induced states is much greater than for

589

TiOi. Annealing the highly oxygen deficient surfaces in a significant partial pressure of oxygen quenched the W 5d intensity, but annealing in UHV alone led eventually to a surface which was characterised by a sharp W 5d peak close to the Fermi energy. There was no attempt in any of this earlier work to relate the features observed in XPS or valence photoemission to specific atomic arrangements at the surface and indeed there was no real prospect of pursuing such ideas at that time. However it has recently been possible to combine STM and STS on WOsCOOl) surfaces with resonant photoemission [268]. The idealised cubic structure of WO3 may be described alternatively in terms of alternating layers of stoichiometry {0}-{W02}-{0}-{W02}- stacked along the [001] direction. These carry formal ionic charges {2-}-{2+}-{2-}-{2+}. To avoid termination of the [001] surface with a repeating normal dipole it is necessary to remove half the oxygen ions in the outermost oxygen ionic layer to give a sequence {Oo.5}-{W02}-{0}-{W02}. The charges can then be grouped into repeating quadrupolar units (fig. 13). Atomistic simulation show that the favoured configuration for the half monolayer of on-top oxygen involves an ordered c(2x2) (that is (V2xV2)R45°) arrangement [269]. Thus the "stoichiometric" [001] surface displays a c(2x2) reconstruction, which has been imaged by STM [269-271]. As with TiO2(110) (1x1) it was initially unclear whether the uppermost O ions or the bare metal cations in the WO2 plane appear as greyscale maxima at positive sample, although a simple consideration

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m •a ,-0

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Fig. 13. Schematic representation of the structure of WO3 in terms of layers of ionic planes. A stoichiometric (001) surface cannot terminate in either a complete O or WO2 plane as this gives rise to a repeating dipole normal to the surface. The preferred termination with a half monolayer of on-top oxygen to give a c(2x2) reconstruction is shown in the left hand panel.

590

of the extent of W-O found in bandstructure calculations [272] suggested that there should be sufficient W 5d character in the states associated with the on-top oxygen for them to predominate. Very recently Altmann and co-workers have observed that alkoxide ions adsorb on the sites which appear as greyscale minima in STM images [273]. Chemically it only makes sense to assume that alkoxide species adsorb on bare cation sites: this in turn implies that the greyscale maxima must correspond to on-top oxygen ions. In contrast to TiO2(110) (1x1) it therefore appears that STM images are dominated by geometric rather than electronic structure effects. This difference must be attributed in part to the fact that the geometric height difference between the ontop O and the W ions is greater than for TiOaCllO). Additionally WO3 is more covalent than Ti02 so there is greater admixture of O 2p character into conduction band states. STM reveals that the oxygen deficiency produced by ion bombardment and/or UHV annealing of WO3(001) surfaces is initially associated with striking "trough defects" that dissect the c(2x2) terraces (fig. 14). The dominant photoemission intensity from these surfaces lies deep within the bulk bandgap with maximum intensity at about 1.5 eV binding energy [268]. This binding energy is similar to that associated with W-W a- bonding states in the oxide WO2, whose rutile-like structure is based on edge sharing WOe octahedra [274].

Fig. 14. 280 A X 280 A STM images of WOsCOOl) acquired at +2.0V sample bias and InA timnel current. Left hand panel is of a typical surface showing a c(2x2) reconstruction dissected by dark "defect troughs". Right hand panel shows a surface after prolonged annealing in UHV. The bright "raft" running diagonally across the image supports a p(lxl) reconstruction. Adapted from ref. 268.

591

Topographically resolved scanning tunnelling spectroscopy confirms that the occupied electronic states are associated with the trough like defects, which are therefore believed to correspond to shear-plane structures intersecting the surface (fig. 15). Metal-metal bonding between adjacent W ions within the shear planes pushes the states deep down into the gap [268,270,271]. Prolonged annealing of WOsCOOl) in UHV leads eventually to the growth of rafts on the c(2x2) terraces which support a new (1x1) reconstruction as shown in fig. 14 [275]. The simplest model for this reconstruction involves a termination in a bare WO2 layer which can be accommodated on top of the stoichiometric c(2x2) reconstruction provided that half a monolayer of oxygen ions are added to the O0.5 layer. Thus the overall stoichiometry of the raft structure relative to that of the stoichiometric surface is WO2.5 and the W ions are all formally 5d^ species in a +5 oxidation state. The raft structure appears to grow on the c(2x2) terraces by elimination of W-rich shear planes from the nearsurface region: this growth by net transport of cations from the bulk is analogous to the model proposed for growth of added Ti203 rows on Ti02(l 10) (1x2).

>

-n

1\

« 1

1 2 - 2 - 1 Sample bias / V

Fig. 15. Scanning tunnelling spectra (dl/dV versus sample bias) for WO3(001). Left hand panel: solid line, c(2x2) terraces; dashed line, defect troughs. Right hand panel: p(lxl) rafts. Partly adapted from ref 268.

592

Fig.l6. STM images of WO3(001) (1x1). Left hand panel: 30 A x 30 A empty state image taken at +2.0V, InA. Right hand panel: 25 A x 25 A filled state image taken at -0.4V, InA. Filled state imaging of this reconstruction is possible because of the metallic nature of the surface. Adapted from ref 275

0 - 1

3

binding energy / eV

Fig. 17. Resonant photoemission at 57.5 eV photon energy from WOsCOOl). Left hand panel: surface with defect troughs, giving defect states predominantly deep in the bandgap at 2 eV. Right hand panel: surface supporting large areas of the p(lxl) reconstruction. There is now a stronger peak at the Fermi energy with a metallic density of states. Adapted from ref 275.

593

Tunnelling spectra from the (1x1) rafts are characteristic of an essentially metallic material (fig. 15). In agreement with this observation it is possible to image the (1x1) reconstruction at both positive and negative sample bias, as shown in fig. 16. Photoemission spectra such as shown in fig. 17 from surfaces demonstrated by STM to support large areas of the (1x1) reconstruction are different to those where shear plane defects are dominant: the deep gap state observed for the latter is much attenuated and there is higher intensity near the Fermi energy. This suggests that the (1x1) surface is indeed metallic, in agreement with the conclusion from scanning tunnelling spectra, from (1x1) [275]. 4.5,2 Na^WOsiOOl) Introduction of sodium into empty 12-coordinate sites of the WO3 lattice results in a series of oxide bronzes. These have general formula NaxWOs, where 0.0<x0.43 the bronzes adopt an essentially cubic perovskite structure [278] closely related to that of Re03 (fig. 18). The Na 3s levels lie about 10 eV above the bottom of the W 5d bands and each added Na atom is therefore ionised to Na^, with donation of one electron into the W 5d band of local t2g symmetry [279]. For x values of less than 0.26, the 5d electrons are localised, probably by an interplay between polaronic effects, disorder =3.84A x=0.665

r*i O-

Fig. 18. Schematic representation of the structure of NaxWOs. WO2 and NaxO surface planes derived by bulk truncation are shown on the right.

594

induced (Anderson) localisation and on-site Coulomb repulsion [280,281]. However the cubic bronzes have metallic conductivities comparable with that of Cu. In this regime, the conduction electrons occupy a band which has an almost free-electron-like density of states, although the width of the band does not show an ^'^ variation with carrier concentration because Na intercalation leads to simultaneous narrowing of the W 5d conduction bands [279,282]. In addition the disorder and on-site Coulomb repulsion induce a minimum in the density of states or pseudogap at the Fermi energy [283,284] which may be seen in tunnelling spectra [285]. The gap becomes deeper as the metal to non-metal transition is approached and becomes apparent even at much lower resolution in photoemission for x values of around 0.4 [280,281]. In the metallic cubic phase the gap between the bottom of the occupied conduction band and the top of the valence band is much smaller than in the distorted monoclinic phase of WO3. Bandstructure calculations suggest that the bandgap narrowing could be attributed in part to structural effects [272, 286]. It should also be realised that bandgap shrinkage is a general phenomenon in passing from a non-metallic to a metallic state even when there is no structural change: the bandgap renormalisations arises from a Coulombic attraction associated with the donor ions and self-screening of the repulsion between valence and conduction electrons by the conduction electron gas [287]. Thus shrinkage of the intrinsic gap between valence and conduction bands has been observed in photoemission from systems such as CdO [288,289] and Sn02 [290] which may be doped to become metallic without structural change. In the present context, the sodium tungsten bronzes are of particular interest in view of the claim by Hochst et al that pronounced intrinsic surface states intensity is found in photoemission between the O 2p band and the metal conduction band at vacuum cleaved NaxWO3(001) surfaces [291]. It was suggested that the mutual Coulomb repulsion that pushes states out of the gap for insulating oxides [292] is screened out in a metallic oxide, thus allowing the emergence of gap states which appeared as strong features in He(I) photoemission spectra. The gap states showed pronounced dispersion with wavevector in the surface plane, as expected for a delocalised surface state. However it emerged that that most of the gap intensity in this work could be ascribed to a simple experimental artefact, namely satellite structure associated with He(I)p emission in the He lamp used to excite the photoemission spectra. On carefully vacuum annealed (001) surfaces the gap intensity in He(I) photoemission was found to very weak [293]. As with Na2/3WO3(110) [294] most of this intensity is attributable to a plasmon loss satellite of the conduction band peak. However, the experiments on the (001) surface [293] suffer from the ambiguity that the surface reconstruction was described as p(2x2). In the light of

595

the STM data discussed above this must correspond to the co-existence of c(2x2) and (2xl)+(lx2) reconstructions. STM and LEED demonstrates that NaxWOsCOOl) (with x values around 2/3) surfaces annealed to high temperatures above 600°C in UHV support a c(2x2) reconstruction analagous in most ways to the c(2x2) reconstruction of WO3(001). However, because the tungsten bronzes are metallic it is possible to image the reconstruction at both positive and negative sample bias, as shown in fig. 19 [295]. Very weak greyscale maxima appear in the centres of the squares that define the (V2xV2) R45° cell at positive sample bias: these are strongly enhanced at negative sample bias. The interpretation of these observations is that on-top oxygen atoms are imaged as the strong grey scale maxima and the bare W ions in the underlying WO2 plane appear as subsidiary maxima: the position of the Na ions is unclear from the STM images. The enhancement of the W ions at positive sample bias arises from an interesting aspect of electronic structure. The extent of mixing of O 2p states with W 5d states in the conduction band

-i

10

1

1

20

1

1-

30

Distance along [100] /Angstrom

Fig. 19. Right hand panels: STM images of the (V2xV2)R45° reconstruction on Nao.67W03(001) acquired at -0.6 V sample bias (upper) and +0.6 V sample bias (lower). Corresponding corrugation profiles are shown in the left hand panels, emphasising the fact that W ions in the centres of the squares of O ions which define the reconstruction are more pronounced in the filled state image. Adapted from ref 295.

596

increases in moving up through the t2g band: in the bulk inversion symmetry means that there is no mixing at all at T point where the bottom of the t2g band is found. Thus the W 5d states below the Fermi level that are interrogated at negative sample bias have less O 2p admixture than the empty states above the Fermi level that are involved in images taken with positive sample bias. This effect increases the contribution the W ions relative to on-top O in filled state images. Annealing (001) tungsten bronzes at temperatures below 600°C in UHV leads to a (1x2) reconstruction that has been studied by two groups by STM [296-298]. Jones et al. obtained well resolved images which were interpreted in terms of termination of the surface of a crystal with composition close to Nai/sWOs with a plane of stoichiometry NaxO [297,298]. The x2 periodicity apparently arises from dimerisation of oxygen ions in the surface plane to give "peroxide-like" surface species with an O-O separation less than twice the ionic radius of O^" (although the distance was still somewhat greater than in a true peroxide). The development of this model structure is shown alongside an STM image in fig. 20. A speculative model was proposed which involves reduction of

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0@ 0 9 9® ^ ^ Fig. 20. Left hand panel: 23 A x 23 A STM image of the (2x1) reconstruction on Nao.67W03(100) taken at +0.4 V sample bias and 1 nA tunnel current. Top right hand panel: unrelaxed Nao.sO surface plane. The oxygen ions (large spheres) and sodium ions (small spheres) are assigned their conventional Shannon-Prewitt radii. Bottom right panel: schematic of the relaxed (2x1) reconstruction with "peroxide-like" oxygen ion dimers. Adapted from ref 298.

597

10

8

6

4

-2

binding energy / eV

Fig. 21. Photoemission spectra of Nao.67W03(001) (2x1) excited at the different photon energies indicated. The dashed line and arrow in the spectrum excited at 100 eV emphasise the feature attributed to the a bonding level in peroxide-like surface dimers. Adapted from ref 298.

charge in the surface layer by a mechanism different to that prevailing in the c(2x2) reconstruction. As two O^' ions move together, the bonding Gg and itu levels within the full-shell configuration fj^TiîZg^^ of the 02^^' unit will move down in energy. But the antibonding 7Cg and GU levels will move upward. If the Gu level crosses the Fermi energy, electron density will flow out of the localised orbital reducing the surface charge. A feature of this model is that it should be possible to observe the molecular-like orbitals associated with the dimer in a photoemission experiment. Using synchrotron radiation at around hv=100 eV in order to maximise the surface sensitivity and to be well off the W 5d resonance associated with the W 6p core threshold, a feature was observed (fig. 21) below the bottom of the valence in photoemission spectra at about 11 eV binding energy [298]. This was attributed to the ag bonding states. Structure is observed at a similar energy when K deposited on metals [299] or Ti02 [300] is exposed to oxygen and in this earlier work the feature is also ascribed to peroxide-like

598

surface species. However a problem with the model is that there was no firm indication of the GU antibonding levels in the vicinity of the Fermi level.

5. CONCLUDING REMARKS There has been very significant progress in the past decade in clarifying the nature of electronic states at oxides surfaces. Particular highlights include the clear-cut demonstration of the influence of reduced symmetry at surface sites in producing first order changes in the pattern of ligand field excited states in rocksalt transition metal monoxides (section 4.2); and verification of the importance of second order effects in inducing mixing between s and p states (which have opposite parity) at surface cation sites in MgO (section 4.1) and Sn02 (section 4.3). Second order effects do not appear to have a major influence on surface electronic structure for d^ oxides such as Ti02 and WO3, but for these materials the lowest empty nd and (n+l)s states have the same parity. The recent progress has depended in part on further refined application of established techniques such as photoemission and HREELS, but it is clear that STM and STS are now beginning to make a major impact. Work which attempts to correlate spatially resolved STS spectroscopic measurements with results from photoemission is in its infancy and there is undoubtedly scope for further work using STS on topographically well defined surfaces, especially when defects are present. The increased ability of theoretical models to capture the essential physics of surface electronic structure in a satisfactory fashion is a further significant development. However it is sobering to reflect on the fact that it is only in the past 5 years that first principles calculations have been able to account, for example, for surface states on SnO2(110). The chemical model of 5s-5p hybridisation at surface sites was first introduced in the early 1980s [130] based on ideas that had been prevalent in the chemical literature since the 1940s [301].

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Chapter 15

Desorption Induced by Electronic Transitions, DIET, at Oxide Surfaces J. L. de Segovia^'* and E. M. Williams^ Înstitute de Ciencia de Materiales de Madrid. CSIC, E-28049 Cantoblanco Madrid, Spain ^ Surface Science Research Centre, The University of Liverpool, Liverpool L69 3BX, UK 1. HISTORICAL INTRODUCTION TO ESD AND PSD The electron stimulated desorption of ions, positive or negative, and neutrals induced by electronic transitions, ESD, has been widely studied for many years. In 1965 Lichtman and McQuistan reviewed several aspects of ESD [1], Madey and Yates in 1971 [2] and in the mid to late 80's the present authors reviewed the field of ESD and its application for surface adsorbate analysis [3,4]. An extensive review was published in 1991 by Ramsier and Yates [5] and in 1994 Madey reviewed the history of ESD [6]. The same phenomena can be observed when photons are the incident agents at the surface. In this case it is recognised as photon stimulated desorption, PSD. In addition to ions and neutrals, excitation leading to desorption as metastable species can also occur [7], as well as surface induced reactions. Thus, from 1982 the phenomenon was recognised in a more general form as Desorption Induced by Electronic Transitions, DIET [8, 9]. Apart from the general bibliography, a significant contribution to this field can be followed in the Proceedings of the Conference series on Desorption Induced by Electronic Transitions, DIET. In this introduction, and ahead of a discussion of specific oxide surfaces, it is of interest to consider some of the most relevant discoveries with the development of ESD. We shall also restrict the study to the desorption of positive ions from oxidised and oxide surfaces. The desorption of ions from adsorbed gases on surfaces was observed as early as 1918 by Dempster [10] using the then recently discovered mass spec* Author for correspondence:

[email protected]/[email protected]

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trometer. The same approach was pursued later in 1950 by Plumlee and Smith [11] in 1961 using a magnetic deflection mass spectrometer. Most of the earlier studies of ESD were performed by detecting pressure changes within a vacuum system, such as in the pioneering work of Ishikawa in 1943 [12]. The first systematic studied of the role of impinging electrons in causing desorption of gases adsorbed on surfaces was performed by G. E. Moore in 1961 [13] using a magnetic deflection mass spectrometer. He observed desorption of O^ and CO"*" ions from CO adsorbed on molybdenum and tungsten surfaces. One of his most important findings was to show that ESD is a first order phenomenon, i.e. the number of desorbed ions is proportional to the number of impinging electrons. He also determined the dependence of the desorbed ion yield on the incident electron energy and found a maximum at about 70 eV. The minimum energy necessary to produce an O"^ was between 17 and 18.3 eV. The last years of the 50's and beginning of the 60's saw an increasing interest in the behaviour of ion currents in ionisation gauges when exposed to actives gases such us CO and O2. An anomaly arose in that gauges indicated pressures much higher than the true pressures in ultra high vacuum systems. The groups of Alpert, at the University of Illinois [14], Redhead [15] at National Research Council, Otawa, and Robins [16], independently discovered the role of the surface ions desorbed from the grids of gauges and mass spectrometers by impinging electrons. These ions contribute significantly to the apparent ion current, thus indicating values much higher than the true system pressure, even with the measuring instrument operated at low electron emission current. Fig. 1 shows a spectrum recorded by

Total pressure = 5 x lO'^torr

28 20 22119 44

Fig. 1. Mass spectrum recorded with a 90° magnetic deflexion mass spectrometer of the Davis Vanderslice type. Note the ESD surface ions at mass 19 (F"^), and the shoulder at mass 16 (O"^) with 5 eV kinetic energy. From Ref. [14]

610

Alpert's group with a 90° magnetic deflexion mass spectrometer of the Davis VandersHce type [17]. Note the peak at mass 19 assigned to ESD of adsorbed fluorine and the double peak at mass 16 corresponding to O^ from the gas phase and surface ions from ESD. The calculated kinetic energy of this surface ion at 5 eV was considerably in excess of the thermal range of energies normally associated with gas phase ionisation. The discovery of the threshold and energy distribution of ejected ions attracted much excitement amongst investigators since their form should reflected the character of the bond between the adsorbate and the surface. Redhead [18] measured the ion energy distribution of CO"^ ions from carbon monoxide at tungsten and found a distribution with two most probable kinetic energies at 1 and 7 eV. However, O"^ ions from O2 adsorbed on the same surface showed only a single peak at 7 eV. These findings led in 1964 to the presentation of the first theoretical interpretation of the ESD process independently by Redhead [19] and by Menzel and Gomer [20], later known as the MGR model. The systematic study of DIET demanded special instruments to study the phenomena, specifically to determine the desorption yield, the ion kinetic energy distribution and the energy threshold for the ejected ions. The focus upon ion detection, desorbed with a typical efficiency of around 10'^ ions/electron, naturally served as the most accessible monitor of the desorption process. Most early investigations of ion desorption parameters made use of the so-called electron desorption cell, with the exception of the contribution of Moore mentioned earlier. In its simplest form this consists of a sample situated in front of a filament, with a collector of either planar or hemispherical form, to the rear of the filament. In the hemispherical form the sample face will be entirely surrounded by the collector, with the sample at the centre of curvature. The arrangement can also incorporate further grids to prevent secondary electrons from the sample reaching the collector and to suppress photoelectron production at the collector due to the action of soft X-rays produced at the sample by the impinging electrons. This was the technique used by Redhead [21] to study ion energy distributions, ion yields and thresholds for ESD with oxygen on molybdenum and with carbon monoxide at tungsten and molybdenum. Degras and Lecante [22] used a similar device to study the adsorption of CO on tungsten, while Madey and Yates developed a similar electron desorption cell at NBS (now NIST) [23] to perform ESD studies of the reactivity of CO, NO and O2 with tungsten [24]. In an approach taken by Lopez-Sancho and de Segovia [25] an ionization gauge of the suppressor type [26] was adapted to study the reactivity of O2 on tungsten surfaces. The threshold for O^ desorption from O2 at W was determined at 24.2 eV with a value of most probable kinetic energy of 6.7 eV. Several devices based on the similar concept of a desorption cell have been used by other labo-

611

ratories; however, a common limitation was that the identity of desorbed ions could not be explicitly determined. Thus, to circumvent this limitation the next logical development was to couple an electron desorption cell to a mass spectrometer. Magnetic sector mass spectrometers had been utilised to investigate phenomenological aspects of ESD in the early 1960's as in the work of Young [27] and the systematic study of G. Moore [13] related with the adsorption of CO on W and Mo. One of the undesirable features, however, of magnetic deflection mass spectrometers is that mass separation is energy dependent. This was overcome with the advent of the quadruple mass spectrometer, QMS, which in addition offered a far more compact structure free from magnetic fields. Two of the first instruments to combine an electron desorption cell (of hemispherical form) with a quadrupole mass spectrometer were developed at NIST by Madey and Yates [28] and at Liverpool by Leek and Williams [29]. In both instruments the hemispherical collector contained an aperture just in front of the entrance to the QMS, so as to sample a fraction of the desorbed ion flux. While the hemispherical collector could detect the entire angular integrated ion yield, the QMS detected only those ions emitted in the solid angle subtended by the aperture. An instrument based on this principle was later developed in Madrid with the objective also to identify neutral desorbed species [30-32]. There has also been in addition a significant contribution to ESD using experimental methods based on changes in surface parameters as opposed to the direct detection of desorbed products. Menzel and Gomer [33] performed studies on the reaction of H2, CO, O2 and Ba adsorbed at W in a field emission electron microscope, monitoring changes in the work function of the tip. Dylla, King and Cardillo [34] using Auger electron spectroscopy, AES, and mass spectrometry techniques to study the adsorption of CO on plasma and oxygen-etched silicon surfaces, obtained data on total cross sections with both techniques which showed good mutual agreement. A combination of a double pass cylindrical mirror analyser with time of flight of ions was used for Traum and Woodruff [35] The ESD of CO adsorbed on Ir (111) has been studied by LEED using the intensity of the CO (VsxVs)/? 30°pattern (Shek, Withrow and Weinberg [36]); they obtained data on cross sections and rates of induced dissociation at the surface. Auger electron spectroscopy and mass spectrometer techniques have been applied to ESD of O2 on tungsten carbide for a broad range of incident electron energies, 80 eV to 2.5 keV, by Stori, Braun and Gomer [37], who found a maximum in ion yield at 300 eV. Using a similar instrumental approach, Priegge, Niehus and Baueri [38] investigated the ESD of O2 on W (111) and determinated a cross section of 2x10'^"^ cm^, much lower than the value 7x10'^^ cm^ reported by Lopez-Sancho and de Segovia at Madrid [25]. However, diffe-

612

rent results using different methods at the same sample can be found; Williams at Liverpool found cross sections of 1.3x10"^^ cm^ and 1.4x10'^^ cm^ for O2 on W using respectively a time of flight ion detection technique and a cylindrical mirror analyser, CMA, for Auger analysis [39]. The authors proposed that the explanation for the differences lay with the specificity of the ion response in ESD, dependent on the adsorption states contributing to the response. A review to 1989 on the use of ESD as tool for surface adsorption studies can be seen in Refs. [3,40,41]. In 1972 Madey at NIST [42] discovered that O^ ions desorbed from oxygen on tungsten had angular distributions with sharp maxima in the direction normal to the surface. Later, Madey, Czyzewski and Yates developed a sophisticated instrumentation incorporating microchannel plates [43] to improve the study and visualisation of the phenomenon. Patterns were observed with sharp structures, and so emerged the technique of Electron Stimulated Desorption Ion Angular Distribution, ESDI AD. The principle of the technique as identified from the work at NIST is that ions are initially emitted along the bond direction at the surface. This discovery stimulated theorists to refine the dynamics of the process taking into account the contribution of re-neutralisation processes of ions and of image and other short range forces. To a first approximation under typical desorption conditions, for angles up to around 40"" these opposing influences leave the distribution largely preserved [44-53]. Thus, by the early 80's ESDIAD was confirmed as a useful technique for characterisation of the bonding geometry of adsorbed molecules or surface atoms. In 1988 the ESDIAD technique was further improved by the addition of digital video data acquisition to obtain excellent patterns of surface-adsorbate structures [54]. As mentioned earlier desorption of ions, neutrals and metastables can also be caused by photons incident at the surface, in the process known as Photon Stimulated Desorption, PSD. In principle, there is an important difference between ESD and PSD. While in the first case, the energy of the primary electron interacting with the electronic structure of either the substrate or adsorbate atom can be partially transferred, in PSD the whole energy of the primary photon is transferred to the interacting electron. This raised the question as to whether this difference in primary excitation mechanism gives a different response with both processes. The first studies using photons of well-defined energies were performed by the Knotek group [55] at the Stanford storage ring. They found that ion yield curves are sharper in PSD than in ESD, although the onset of ion desorption is the same in both cases. Several laboratories were motivated to perform further experiments to compare ESD and PSD. Madey's group at NIST and Eastman's group at Wisconsin [56] in a joint experiment compared the anisotropy of ion emission with ESD and PSD, but found the angular distribution for O^ from W (HI) to be identical with both methods. PSD experiments by

613

Stockbauer, Madey and Hanson [57] using the SURF-II facility at NIST with C from Ti02 again showed similarity with BSD experiments. A review of ejection mechanism related to these different studies was presented by Feibelman in 1983 [58]. He analysed the areas where BSD can be useful as an analytical technique, and discussed experiments to elucidate the desorption pathways at the surface. At the end of this introduction, we can conclude that by the middle of the 1980's it was well established that electrons or photons of certain energy, generally higher than 5 eV, impinging at surface atoms or adsorbate produce electronic transitions, which yield ions, neutrals and metastables to the gas phase. So, once the principles of the BSD and PSD were established and that both primary excitation of the atoms yield similar results, the process was recognised by the general title of Desorption Induced by Blectronic Transitions, DIET. The main findings can be summarised as follows: (i) Exposure of surface atoms or adsorbate atoms to electron or photon fluxes above certain energy, gives rise to the desorption of ions, neutrals, fragments and metastables. (ii) The basic mechanism of ejection involves electronic excitation of the substrate atoms or adsorbed entities, (iii) Ions as well as neutrals are produced at specific threshold energies of the impinging electrons or photons, (iv) Ions and neutrals are produced with kinetic energies of a few electron volts; neutrals can also be produced in their ground or metastable states. (v) Cross sections for ion production are lower than those corresponding to gas phase ionisation, dependent on the adsorbate. Ion desorption cross sections are typically around 10"'' ions/electron with neutral desorption around 10"* neutrals/electron. (vi)

Ions are ejected in specific directions, which correspond to the metal atom-adatom bonding direction. Studies at this time related to several metal substrates with mainly small molecular gases (i) Hydrogen [30, 31, 59-65]. (ii) CO [2, 10, 24, 45,46, 66-73]. (iii) Oxygen [16, 23, 45, 74-79], (iv) H2O [45, 60, 61]. The latter two gases reflect a first step of enquiry into oxidation and corrosion problems at surfaces. Progress with alkaline atoms on several surfaces was pioneered by the Ageev group at St. Petersburg [80- 83]. Interest in DIET studies at oxides surfaces underwent an explosion in 1978, when Knotek and Feibelman at Sandia Laboratories [84] reported a new mechanism for BSD, the so-called "Auger decay of a core hole". The model explained well the thresholds and ion kinetic energy distributions for the ejection of O"*",

614

OH"^ and F"^ from transition metal oxides of maximal valence such as Ti02, V2O5 and WO3, for which, the MGR theory could offer no satisfactorily explanation. The authors subsequently published a reinterpretation of previous results from chemisorption systems to conclude that with ionic bonding the Auger mechanisms for desorption explained well the observed thresholds for ion production [85]. The new mechanism has become known familiarly as the KF model. In their survey Feibelman and Knotek entered into an in-depth review of thresholds in which they compared findings of energy loss spectroscopy, ELS, with thresholds as identified with changes in the slope of the ion yield as a function of electron energy. The significant finding was that singularities, inflexions and slope changes in the ion yield curve, could be unambiguously identified with core level excitation of the surface atom and/or electron levels of the adsorbate. For Ti02 and other full valence oxides they found that the O^ threshold began at about 22 eV, which corresponds to the ionisation of the O 2s level. A new increase of the O"*" ion yield at about 34 eV, could be correlated with ionisation of the Ti 3p level at 33.6 eV. Developments around the KF model gave rise to a wide interest in stimulated desorption processes and their role at oxide surfaces. Covalent systems were also found to reflect some features of the Auger decay model, which will be discussed in the next paragraph. 2. EJECTION MECHANISMS: THE FEIBELMAN-KNOTEK MODEL The basic principle of the Feibelman-Knotek model, KF model, as first published [84] for full valence oxides is based on an inter-atomic Auger neutralisation of a hole produced in the core level of the surface atom by the primary electron or photon. The model as related to Ti02 is sketched in Fig.2. The impinging electron (e) or photon (hv) creates a hole in the Ti 3p level (33.6 eV). Since in a full valence oxide the valence band is empty, one electron from the O 2p (13.5 eV) level decays to the 3p core level. In the situation two O 2p electrons take the excess energy and leave the atom; the initial Ti'^'^0^' configuration is thus transformed to a final state Ti"^"^ O"*". Here the O"^ is subject to a very high Coulomb repulsion force and, as a consequence, is ejected into vacuum. In a later communication Feibelman and Knotek [85] reviewed the general concepts of ESD over a range of chemisorbed systems and showed that many features can be easily understand using the Auger decay model instead of the MGR theory. For the Ti02 system the work by Knotek, Jones and Rehn [86] showed that the desorption by ESD and also PSD of H^, GH"^ and F"^ following water exposure (with residual fluorine) has major thresholds at the G 2s and F 2s core levels, consistent with the applicability of Auger decay model also to covalent che-

615 Auger electrons

Vacuum level

n

'(3d) Valence

r n Ti(3p)

rl r 1

'\^^^dy/)^/////Â

1

0^0 Ti02 bonding Final state yielding an ion by coulorabian repulsion

Fig. 2. Schematic of the Feibelman-Knotek model for the TiOa case. [3]

misorption systems. The Menzel group at Munchen [87, 88] studied the case of O2 and CO on W(IOO) and Ru(OOl), where the ion desorption yield shows thresholds associated with the ionisation of the C Is and O Is core levels (respectively to 270 eV and 510 eV). This was in addition to the usual valence band excitation (MGR). Water and fluorine adsorbed on several oxides have also shown ion desorption by Auger decay [89, 90]. Threshold for H"" ions were seen to be dominant with ESD of water at Al and W surfaces as found in joint experiments at Liverpool and Madrid [ 91,92] with thresholds associated with excitation within OH radicals. Thresholds of H"" with water at Nb were shown to be linked to the core Nb levels [93]. These findings, generally, supported the applicability of the KF model for desorption within some covalent systems. Knotek and Feibelman [94] examined the modification to a surface when exposed to ionising radiation and assesed the damage that can be produced. They addressed the stability of ionically bonded surfaces, where the KF mechanism applies, and concluded that Auger induced decomposition only occurs when the cation in the solid is ionised to relatively deep core levels. In the case of non-maximal oxides as with NiO, Freund's group [95] showed that whilst desorption of neutral NO and CO from NiO(lOO) and (111) surfaces has thresholds at the C Is, N Is and O Is core levels, it proceeds mainly on the basis of the MGR model, involving an excited state of the adsorbate. An overview of electronic desorption presented by Feibelman in 1983 [96] examined particularly the stability of the multiple-hole final state configuration leading to desorption. The presence of multiple holes , and associated hole-hole correlation

616

effects, is seen to lead to inherently longer iSnal state times life and so can facilitate desorption in covalent systems. These many body effects in desorption are also addressed in the work of Ramaker [97] and have their origin in the CiniSawatzky Configuration-Interaction theory of hole interaction. ESD induced by electrons extracted at the tip of a scanning tunnelling microscope (STM) has recently applied to the study of the H-Si and Cl-GaAs(l 10) systems by Shen and Avouris [98]. A feature of the STM is the capability to provide an intense electron density (10^^ A/cm^) of atomic dimensions directly at the surface atoms. Under these conditions, desorption can also proceed with multiple-vibrational excitation. A combination of ESD and STM holds promise as a tool for atomistic DIET, also as a means to asses the influence of electron irradiation at surfaces [99]. In ESDIAD an interesting aspect is the determination of the trajectory of the ion on leaving the surface, which will be discussed later in relation to the TiOi surface. Laser stimulated desorption has been also studied as in the work of Freund's group [100]. Low energy excitation and desorption dynamics from NO adsorbed on NiO (100) and (111) oxide surfaces is stimulated by a Ar-Fexcime laser with 6.4 eV energy. They found total cross sections of 6x10"^^ cm^ and explained their results on the basis of the MGR model. Induced ion desorption by resonant photon-stimulated desorption has been also observed. The yield of HT desorbed ions as a function of the incident photon energy from CeOi has been studied by the group at Sandia Laboratory. [101] They observed a sharp resonance in PSD of H^ using photon energies near the Ce 4d edge at 110 eV. The resonance is due to autoionisation from excited states from the interaction of the Ce 4d and 4f levels. The contribution of secondary electrons to DIET has also been studied. Jaeger and Stoehr proposed that secondary electrons are the predominant contribution to the H^ ion yield in the NHs/Ni system [102]. In a study of PSD ion yield of YC from the OH/YbO-Sm system, Schmidt-May, Senf and Voss explain the structure on the basis of the secondary electron contribution to the ion yield [103]. Hutson, Ramaker, Bermudez and Hoffbauer [104] found a significant contribution of secondary electrons to the ESD ion yield of OH^ from TiOi at incident electron energies higher than 50 eV. An approach taken by Madey's group was to study medium-energy backscattered electron diffraction at TiO2(110) [105]. A maximum in the backscattered electrons from the second layer of atoms was found at an incident electron angle of 45°, which corresponds to a close packed direction for the (110) surface. Under this conditions the incident electrons are focused onto the second layer atoms by the atoms in the first layer. In their study of ESD ion yield from a well characterised TiOi (110) surface, the Madrid group [106, 107] found that in addition to the threshold at the

617 200

2500

175 2000

150
3d excitations. One of the most active groups in ESD studies of the adsorption of Ba, Li, Na, K and Ca at oxidised metals has been the group of Ageev at the Joffe Institute, St. Petersburg [112-116]. They made use of a surface ionisation detector to detect directly the desorbed alkali atoms. The atom kinetic energy distribution typically has multiple peaks, with energies over the range 0.2 Hi^Ô+^Ô is seen as being incorporated into the surface oxide. 3.2. Stoichiometric Oxides 3.2.1. The TiO2 Model Surface. Most of the work on stoichiometric oxides has been carried out on TiOi, especially samples with the (110) orientation which offers an interesting and contrasting character of possible adsorption sites. In addition this surface is the most stable and does not facet with thermal treatments. Fig.4 is the ball model of this surface [119]. Two kind of O"^ ions can be well differentiated at the surface: (i) ions located in bridging sites, which protrude out of the surface plane, and (ii) ions located at the in-plane sites. Ti atoms are six or five co-ordinated corresponding to positions associated with O"^ ions at type (i) or type (ii) sites, respectively. One of the first contributions of DIET to oxides was to correlate

619

THE STOICHIOMETRIC TiO, (110) SURFACE [110]

0 BRIDGING

[001]

0 IN-PLANE

Ti ATOMS

Fig. 4. The TiO2(110) ball model [158].

ESD witli the local geometry. In 1985 Kurtz, Stoc]

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