Conjugated Polymer & Molecular Interfaces

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ISBN: 0-8247-0588-2 This book is printed on acid-free paper. Headquarters Marcel Dekker, Inc. 270 Madison Avenue, New York, NY 10016 tel: 212-696-9000; fax: 212-685-4540 Eastern Hemisphere Distribution Marcel Dekker AG Hutgasse 4, Postfach 812, CH-4001 Basel, Switzerland tel: 41-61-261-8482; fax: 41-61-261-8896 World Wide Web http://www.dekker.com The publisher offers discounts on this book when ordered in bulk quantities. For more information, write to Special Sales/Professional Marketing at the headquarters address above. Copyright 䉷 2002 by Marcel Dekker, Inc. All Rights Reserved. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming, and recording, or by any information storage and retrieval system, without permission in writing from the publisher. Current printing (last digit): 10 9 8 7 6 5 4 3

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Foreword

Although it would be difficult to describe in detail the evolution of the science of condensed organic molecular solids over the past century, recognizable periods of activity can be identified. An abbreviated outline of the lineage might be as follows: Evolving from a diverse array of work from many different backgrounds and disciplines, there was a surge in interest in the 1970s in the metallic one-dimensional charge transfer salts, as represented by the molecular crystal TTF-TCNQ, and in parallel, in the intercalated graphites. Based initially on these two classes of materials, the terminology synthetic metals was coined. The discovery of electrically conducting organic polymers in the late 1970s provided a new focus for the field and resulted in a natural scientific reorganization. During the 1980s, conducting polymers matured as a unique class of materials: stable electronic materials that combined the electrical and optical properties of metals and semiconductors with the processing advantage and mechanical properties of polymers. Organic (small-molecule) light-emitting devices (OLEDs) emerged in the 1980s and became a new focus for research and development in the 1990s. Light-emitting devices based on luminescent, semiconducting polymers (PLEDs) evolved in the 1990s along with the more general development of ‘‘plastic electronics,’’ including photodiodes, photovoltaic cells, thin-film transistors, and lasers. Although we have only just entered the first decade of the new century, it seems that a principal focus will be on the use of advanced printing technologies to fabricate such plastic electronic devices. These developments in electronic applications of organic molecular and polymeric materials are discussed in this book. Novel electronic devices iii

Preface

Electronic devices based on organic materials have recently made their debut in the marketplace. Many companies, especially new and small ones, are entering the field. Novel materials in pure, well-characterized form have been generated, especially over the last decade. The organic materials can be classified broadly as falling into one of two categories: (1) conjugated, or semiconducting, polymers and (2) so-called ‘‘small molecules.’’ The latter term is somewhat misleading, since ‘‘small molecules’’ refers mainly to molecules smaller than polymeric macromolecules but otherwise still relatively large. Although some fundamental differences exist between these materials, there are also many underlying similarities. In terms of applications, these materials have many common target markets. It may safely be stated that, should progress in the development of new materials suddenly come to a halt, device performance could continue to improve through engineering of the interfaces, in particular the cathode and anode interfaces. Many scientists believe that the interface issues related to organic molecular-based devices are, at this particular moment, some of the most important with regard to the overall device concept and performance. Engineering more efficient interfaces may be determinant in keeping up with the rate of progress in the development of polymer- and smallmolecule-based electronic devices. This book provides an up-to-date fundamental understanding of the properties of organic interfaces and a practical approach to the formation and improvement of these interfaces. We initially intended to divide the chapters into groups reflecting chemistry, physics, engineering, and other subjects. In the process of compiling the manuscript, however, we realized that even with a directive to the v

vi

Preface

authors it would be impossible to categorize all the chapters in such a manner. This issue in itself indicates the interdisciplinary nature of the field of polymeric and molecular organic semiconductors in electronics applications. Similarly, we had hoped to assemble a chapter for beignners. This also proved almost impossible, as it appeared that each group of authors had their own special set of procedures. However, the sample preparation sections in each of the device-oriented chapters will serve as a guideline for beginners. A considerable amount of work on organic interfaces has been accomplished over the past few years and the time is right to present some of the principal work in the field in this book. We have prevailed on some of the eminent leaders involved in work on interfaces in electronic devices based on organic materials. These scientists, from around the world, have contributed chapters to this book by gathering together their specialist expertise in interface science. The content of these chapters spans a considerable range of issues from basic or fundamental materials science through device applications, pointing out the role of interfaces in each case. We have endeavored to represent all subfactions in the field as fairly as possible. The result establishes a snapshot in time, representing the state of the art in interface science in the area of electronic applications of organic materials. Essentially all the problems met in practicing the art of interface science with these organic materials are represented. We hope that the readers of this book will come away with the same general opinions. Finally, we thank all those who have contributed to this book, for their time and efforts, and for the scientific dexterity it took to carry out the works reported. There are numerous financial sources that need to be acknowledged, which is done at the end of each chapter. William R. Salaneck Kazuhiko Seki Antoine Kahn Jean-Jacques Pireaux

Contents

Foreword Alan J. Heeger Preface Contributors 1.

Functional Dendritic Thin Films and Monolayers Albertus P. H. J. Schenning, J. W. Weener, and E. W. Meijer

2.

A Theoretical Study of Electrical Contacts to Self-Assembled Molecular Wires on Conducting Substrates Sven Stafstro¨m, A˚sa Johansson, and Ma˚rten Bjo¨rketun

3.

4.

5.

A Theoretical Insight into Organic Interfaces in Electro-Optic Devices Je´roˆme Cornil, Donizetti dos Santos, Roberto Lazzaroni, Massimo Malagoli, and Jean-Luc Bre´das Photoelectron Spectroscopy of Interfaces for Polymer-Based Electronic Devices Michael Lo¨gdlund, Grzegorz Greczynski, Annica Crispin, Mats Fahlman, William R. Salaneck, and Thomas Kugler Role of Interfaces in Semiconducting Polymer Optoelectronic Devices Paul F. van Hutten, Victor V. Krasnikov, and Georges Hadziioannou

iii v xi 1

23

39

73

113

vii

viii

6.

7.

8.

9.

Contents

Studies of Organic Thin Films and Interfaces by Various Electron Spectroscopies: Molecular Orientation, Electronic States, and Reactions Nobuo Ueno, Koji Kamiya Okudaira, Yasushi Azuma, Satoshi Kera, and Shinji Hasegawa Organic–Metal Interfaces: From Physisorption to Covalent Bonding Norbert Koch, Gu¨nther Leising, Aparna Rajagopal, and Jean-Jacques Pireaux

241

Indium-Tin-Oxide Thin Films for Characterization of Electrochemical Processes in Molecular Assemblies: Surface Characterization and Surface Modification Carrie L. Donley, Darren Dunphy, Rebecca A. P. Zangmeister, Kenneth W. Nebesny, and Neal R. Armstrong

269

Energy Level Alignment at Organic–Metal Interfaces Hisao Ishii and Kazuhiko Seki

11.

Organic Molecular Interfaces: Investigations of Electronic Structure, Chemistry, and Carrier Injection Properties Chongfei Shen, Antoine Kahn, and Ian Hill

13.

14.

205

Electronic Structure of Interfaces Between Organic Molecules and van der Waals Surfaces Toshihiro Shimada and Atsushi Koma

10.

12.

153

Metal–Polyfluorene Interface and Surface: Structures and Stability Liang-Sheng Liao, C. S. Lee, Shuit Tong Lee, M. Inbasekaran, and W. W. Wu Femtosecond Photoemission Study of Relaxation and Interface Charge Transfer Dynamics in Organic Photoreceptors Antti J. Ma¨kinen, Stephan Schoemann, Yongli Gao, Max G. Mason, Annabel A. Muenter, and Andrew R. Melnyk STM-Excited Electroluminescence and Spectroscopy of Conjugated Polymers Santos F. Alvarado

293

351

401

443

473

Contents

15.

Contact-Limited Hole Current in Poly(p-phenylenevinylene) P. W. M. Blom, D. W. Black, and M. C. J. M. Vissenberg

16.

Organic Light-Emitting Diodes Using Akaline-Earth Fluorides as an Electron Injection Layer Shizuo Tokito, Hisayoshi Fujikawa, Tomohiko Mori, and Yasunori Taga

ix

505

525

17.

Characterization of Metal Interfaces to Molecular Media from Analysis of Transient and Steady-State Electrical Measurements 545 Martin A. Abkowitz

18.

Molecular-Sized Structural Trap at Organic–Metal Interface and Photocurrent Multiplication Phenomenon Masahiro Hiramoto, Ken-ichi Nakayama, Masaaki Yokoyama, and Tadashi Katsume

585

19.

Charge Injection into Disordered Organic Semiconductors Vladimir I. Arkhipov, Ulrich Wolf, Siegfried Barth, and Heinz Ba¨ssler

20.

Energy Structures of Molecular Thin-Film/Metal Interfaces in Air Yutaka Harima and Kazuo Yamashita

651

Metal–Organic Schottky Energy Barriers: Measurements and Device Implications I. H. Campbell and Darryl L. Smith

693

21.

613

22.

Nanometric Electrostatic Phenomena at Molecular Interfaces Mitsumasa Iwamoto

23.

Modifying Interfaces to Semiconducting Polymers: PEDOT in Polymer Microelectronics Lucimara Stolz Roman and Olle Ingana¨s

777

Electron and Hole Injection Processes in OLEDs Consisting of Low Molecular Weight Materials Michio Matsumura

793

24.

733

x

25.

26.

Index

Contents

Preparation of Single Crystalline van der Waals Surfaces: Tips for Beginners Toshihiro Shimada and Atsushi Koma Molecular Control of Electron and Hole Injection at Electrodes and at Organic Layer Interfaces in Organic Electroluminescent Devices Masamichi Fujihira and Chimed Ganzorig

811

817

859

Contributors

Martin A. Abkowitz, Ph.D. Department of Chemistry, University of Rochester, Rochester, New York Santos F. Alvarado, Ph.D. Ru¨schlikon, Switzerland

Zurich Research Laboratory, IBM Research,

Vladimir I. Arkhipov, Ph.D., D.Sc. Institute of Physical, Nuclear and Macromolecular Chemistry, Philipps University Marburg, Marburg, Germany Neal R. Armstrong, Ph.D. Optical Sciences, Department of Chemistry, University of Arizona, Tucson, Arizona Yasushi Azuma, D.Eng. Materials Characterization Division, National Metrology Institute of Japan, Tsukuba, Japan Siegfried Barth, Ph.D. benhausen, Germany

Display Technology, Optrex Europe GmbH, Ba-

Heinz Ba¨ssler, Ph.D. Institute of Physical, Nuclear and Macromolecular Chemistry, Philipps University Marburg, Marburg, Germany Ma˚rten Bjo¨rketun, M.Sci. Department of Physics and Measurement Technology, IFM, Linko¨ping University, Linko¨ping, Sweden xi

xii

Contributors

D. W. Black Philips Research Laboratories, Eindhoven, The Netherlands, and University of Glasgow, Glasgow, United Kingdom P. W. M. Blom, Ph.D. Materials Science Centre, University of Groningen, Groningen, The Netherlands Jean-Luc Bre´das, Ph.D. Tucson, Arizona

Department of Chemistry, University of Arizona,

I. H. Campbell, Ph.D. New Mexico

Los Alamos National Laboratory, Los Alamos,

Je´roˆme Cornil, Ph.D. Laboratory for Chemistry of Novel Materials, University of Mons-Hainaut, Mons, Belgium Annica Crispin Department of Physics and Measurement Technology, IFM, Linko¨ping University, Linko¨ping, Sweden Carrie L. Donley, B.S. Tucson, Arizona

Department of Chemistry, University of Arizona,

Donizetti dos Santos, Ph.D. Laboratory for Chemistry of Novel Materials, University of Mons-Hainaut, Mons, Belgium Darren Dunphy, Ph.D. Mexico

Sandia National Laboratories, Albuquerque, New

Mats Fahlman, Ph.D. Department of Science and Technology, Linko¨ping University, Norrko¨ping, Sweden Masamichi Fujihira, D.Eng. Department of Biomolecular Engineering, Tokyo Institute of Technology, Yokohama, Japan Hisayoshi Fujikawa, Ph.D. Toyota Central Research and Development Laboratories, Incorporated, Nagakute, Japan Chimed Ganzorig, D.Eng. Department of Biomolecular Engineering, Tokyo Institute of Technology, Yokohama, Japan Yongli Gao, Ph.D. Department of Physics and Astronomy, University of Rochester, Rochester, New York

Contributors

xiii

Grzegorz Greczynski, M.Sc. Department of Physics and Measurement Technology, IFM, Linko¨ping University, Linko¨ping, Sweden Georges Hadziioannou, Ph.D. Department of Polymer Chemistry and Materials Science Centre, University of Groningen, Groningen, The Netherlands Yutaka Harima, Ph.D. Faculty of Integrated Arts and Sciences, Hiroshima University, Higashi-Hiroshima, Japan Shinji Hasegawa, D.Sc. Molecular Assemblies, Institute for Molecular Science, Okazaki National Research Institute, Okazaki, Japan Ian Hill, Ph.D.

Sarnoff Corporation, Princeton, New Jersey

Masahiro Hiramoto, Ph.D. Material and Life Science, Graduate School of Engineering, Osaka University, Osaka, Japan M. Inbasekaran Central and New Businesses Research and Development, The Dow Chemical Company, Midland, Michigan Olle Ingana¨s, Ph.D. Biomolecular and Organic Electronics, IFM, Linko¨ping University, Linko¨ping, Sweden Hisao Ishii, Ph.D. Department of Chemistry, Graduate School of Science, Nagoya University, Nagoya, Japan Mitsumasa Iwamoto, D.Eng. Department of Physical Electronics, Tokyo Institute of Technology, Tokyo, Japan ˚ sa Johansson, M.Sc. Department of Physics and Measurement TechnolA ogy, IFM, Linko¨ping University, Linko¨ping, Sweden Antoine Kahn, Ph.D. Department of Electrical Engineering, Princeton University, Princeton, New Jersey Tadashi Katsume, Ph.D. Material and Life Science, Graduate School of Engineering, Osaka University, Osaka, Japan Satoshi Kera, Ph.D. Department of Materials Technology, Faculty of Engineering, Chiba University, Chiba, Japan

xiv

Contributors

Norbert Koch, Ph.D. Institute of Nanostructured Materials and Photonics, Joanneum Research Forschungs GmbH, Weiz, Austria Atsushi Koma, D.Eng. Tokyo, Japan

Department of Chemistry, University of Tokyo,

Victor V. Krasnikov, Ph.D. Department of Polymer Chemistry and Materials Science Centre, University of Groningen, Groningen, The Netherlands Thomas Kugler, Ph.D.

Acreo AB, Norrko¨ping, Sweden

Roberto Lazzaroni, Ph.D. Laboratory for Chemistry of Novel Materials, University of Mons-Hainaut, Mons, Belgium C. S. Lee, Ph.D. Department of Physics and Materials Science, City University of Hong Kong, Hong Kong, China Shuit Tong Lee, Ph.D. Department of Physics and Materials Science, City University of Hong Kong, Hong Kong, China Gu¨nther Leising Institute of Nanostructured Materials and Photonics, Joanneum Research Forschungs GmbH, Weiz, Austria Liang-Sheng Liao, Ph.D. Department of Physics and Materials Science, City University of Hong Kong, Hong Kong, China Michael Lo¨gdlund, Ph.D. ko¨ping, Sweden

Interconnect and Packaging, Acreo AB, Norr-

Antti J. Ma¨kinen, Ph.D. Department of Physics and Astronomy, and Center for Photoinduced Charge Transfer, University of Rochester, Rochester, New York Massimo Malagoli, Ph.D. zona, Tucson, Arizona Max G. Mason, Ph.D.† Rochester, New York †

Deceased

Department of Chemistry, University of Ari-

Eastman Kodak Imaging Materials and Media,

Contributors

xv

Michio Matsumura, Ph.D. Research Center for Photoenergetics of Organic Materials, Osaka University, Osaka, Japan E. W. Meijer Laboratory of Macromolecular and Organic Chemistry, Eindhoven University of Technology, Eindhoven, The Netherlands Andrew R. Melnyk, Ph.D. Webster, New York

Webster Research Center, Xerox Corporation,

Tomohiko Mori, M.D. Toyota Central Research and Development Laboratories, Incorporated, Nagakute, Japan Annabel A. Muenter, Ph.D. dia, Rochester, New York

Eastman Kodak Imaging Materials and Me-

Ken-ichi Nakayama, Ph.D. Material and Life Science, Graduate School of Engineering, Osaka University, Osaka, Japan Kenneth W. Nebesny, Ph.D. izona, Tucson, Arizona

Department of Chemistry, University of Ar-

Koji Kamiya Okudaira, D.Sc. Graduate School of Science and Technology, Chiba University, Chiba, Japan Jean-Jacques Pireaux, Ph.D. Department of Physics, Laboratoire Lise, Faculte´s Universitaires Notre-Dame de la Paix, Namur, Belgium Aparna Rajagopal, Ph.D. Department of Physics and Astronomy, Rutgers University, Piscataway, New Jersey Lucimara Stolz Roman, Ph.D. of Parana´, Parana´, Brazil

Department of Physics, Federal University

William R. Salaneck, Ph.D. Department of Physics and Measurement Technology, IFM, Linko¨ping University, Linko¨ping, Sweden Albertus P. H. J. Schenning, Ph.D. Laboratory of Macromolecular and Organic Chemistry, Eindhoven University of Technology, Eindhoven, The Netherlands Stephan Schoemann, Ph.D.

Infineon Technology, Munich, Germany

Kazuhiko Seki, Ph.D. Research Center for Materials Science, Nagoya University, Nagoya, Japan

xvi

Contributors

Chongfei Shen, Ph.D. Department of Electrical Engineering, Princeton University, Princeton, New Jersey Toshihiro Shimada, D.Sc. Tokyo, Japan

Department of Chemistry, University of Tokyo,

Darryl L. Smith, Ph.D. Materials Science and Technology Division, Los Alamos National Laboratory, Los Alamos, New Mexico Sven Stafstro¨m, Ph.D. Department of Physics and Measurement Technology, IFM, Linko¨ping University, Linko¨ping, Sweden Yasunori Taga, Ph.D. Toyota Centeral Research and Development Laboratories, Incorporated, Nagakute, Japan Shizuo Tokito, Ph.D. Science and Technical Research Laboratories, Japan Broadcasting Corporation, Tokyo, Japan Nobuo Ueno, D.Eng. Department of Materials Technology, Faculty of Engineering, Chiba University, Chiba, Japan Paul F. van Hutten, Ph.D. Department of Polymer Chemistry and Materials Science Centre, University of Groningen, Groningen, The Netherlands M. C. J. M. Vissenberg, Ph.D. The Netherlands

Philips Research Laboratories, Eindhoven,

J. W. Weener, Ph.D. Laboratory of Macromolecular and Organic Chemistry, Eindhoven University of Technology, Eindhoven, The Netherlands Ulrich Wolf, Ph.D. Germany

Elkamet Kunststofftechnik GmbH, Biedenkopf,

W. W. Wu Central and New Businesses Research and Development, The Dow Chemical Company, Midland, Michigan Kazuo Yamashita, D.Sc. Faculty of Integrated Arts and Sciences, Hiroshima University, Higashi-Hiroshima, Japan Masaaki Yokoyama, Ph.D. Material and Life Science, Graduate School of Engineering, Osaka University, Osaka, Japan Rebecca A. P. Zangmeister, Ph.D. of Arizona, Tucson, Arizona

Department of Chemistry, University

1 Functional Dendritic Thin Films and Monolayers Albertus P. H. J. Schenning, J. W. Weener, and E. W. Meijer Eindhoven University of Technology, Eindhoven, The Netherlands

I.

INTRODUCTION

Over the last 150 years, organic chemists have primarily been concerned with the synthesis of relatively small molecules with precise dimensions and of exact composition. Extraordinary achievements have been obtained in natural product synthesis primarily due to advances made in synthetic methodologies and developments achieved in novel characterization techniques. The polymer chemist, on the other hand, has relied on the synthesis of large but polydisperse linear and cross-linked polymer chains. Despite the molecular weight distributions present in high molecular weight materials, polymers fulfill a multitude of important roles in society. It is impossible to imagine life today without the contributions of either of the two fields mentioned above. Despite these accomplishments, no synthetic strategy has yet approached the diversity and precise molecular morphogenesis that nature has mastered in the evolutionary development from simple organic molecules to complex macromolecules and assemblies. Both synthetic organic and polymer chemists have long sought ways to obtain control over synthetic polymer structure and composition at a level comparable to that observed in biological systems. In this quest a new category of polymers has emerged during the past decade, termed ‘‘dendrimers,’’ which are beginning to display control over features such as chemical composition, size, and shape. Dendrimers (dendron = tree, branch; meros = part) are well-defined, highly branched, globular macromolecules that emanate from a central core. 1

2

Schenning et al.

They are produced in an iterative reaction sequence and, at least in theory, lack polydispersity (i.e., D = 1). The synthesis and characterization of dendrimers, also known as ‘‘arborols’’ [1], ‘‘starburst molecules’’ [2] or ‘‘cascade polymers’’ [3], has been reviewed extensively [4]. Two major approaches may lead to these novel macromolecules: convergent and divergent. Divergent synthesis was introduced independently by Newkome et al. and Tomalia et al. and involves an approach in which a dendritic structure is grown by repetitive reaction steps starting from a multifunctional core [1,2]. With every sequence of (generally two) reactions, one generation is added. In the divergent methodology, the number of functional groups at the periphery in successive generations increases exponentially, and a defect in growth in an early generation has a major and predictable impact on the final structure of the dendrimer. Therefore, high-yield reactions have to be used in combination with a large excess of reagents in each dendrimer growth step. Imperfections are nevertheless often encountered in higher generations, while purification becomes exceedingly difficult because the perfect dendrimers differ from the defective ones by only one or a few monomer units. The first dendritic structures prepared according to the divergent strategy that received widespread attention were Tomalia’s poly(amidoamine) (PAMAM) dendrimers [2] and Newkome’s arborols [1]. On the basis of the original work by Vo¨gtle [3], divergently prepared poly(propylene imine) (PPI) dendrimers were reported by Wo¨rner [5] and De Brabander [6]. The synthesis and structure of these dendrimers is given in Scheme 1.

Scheme 1

Functional Dendritic Thin Films and Monolayers

3

Today both PAMAM and PPI dendrimers are commercially available. The second step-by-step strategy leading to dendritic macromolecules, the convergent approach, starts at the periphery of the dendrimer and proceeds inward. This approach was introduced by Hawker and Fre´chet through the synthesis of poly(benzyl ether) dendrimers [7]. The strength of the convergent synthesis lies in the separate preparation of polymeric wedges (or monodendrons) that can be purified prior to their connection to a multifunctional core. These purifications are inconceivable for divergently prepared dendrimers of higher generations. Therefore, the convergent strategy can be referred to as the organic approach to dendrimers: products can be purified and will, in general, be defect-free. Following the above-described synthetic methodologies, a large variety of dendrimers have been prepared and characterized. The initial research on these novel macromolecules was mainly driven by their aesthetically pleasing appearance in combination with curiosity and speculation. In the meantime, a number of these speculations have become reality and dendrimers are now used in fields as diverse as host–guest chemistry, medicinal chemistry, and catalysis. The emphasis in this chapter will be on the use of dendrimers as nanoscopic building blocks in functional monolayers. Back in 1993, Dvornic and Tomalia presented the idea that dendrimers are ideal, well-defined macromolecules for creating self-assembled nanostructures [8]. Since then, dendrimers have been shown to be useful in the construction of mono- and multilayers, liquid crystals, and supramolecular architectures in solution. To successfully line up molecules in an ordered array one generally needs welldefined but flexible building blocks that control their shape by switching between different conformational states in the process of self-assembly. Examples in this chapter show that the successful application of dendrimers in the areas mentioned above relies not only on their uniformity but on their tunable architectures. The emphasis will be on the use of dendrimers as multifunctional building blocks in the construction of mono- and multilayers, sensors, liquid crystals, and light-emitting diodes. Functional materials that utilize specific dendritic properties will be discussed. The often unique conformational behavior of different dendrimer types within these supramolecular assemblies will be explained.

II.

DENDRIMERS ON SURFACES: CONFORMATIONAL BEHAVIOR

The preparation of thin films and layers from dendrimers by self-assembly is a topic of great current interest because it allows the construction of

4

Schenning et al.

functional interfaces that use specific dendritic properties such as size, shape, porosity, end-group density, and multifunctionality. It is also an area of research that greatly benefits from the flexible character intrinsic to most dendrimers. Mansfield et al. was the first to predict deformations to occur for dendrimers deposited on a surface using Monte Carlo simulations [9]. More importantly, these calculations predicted an increase in flattening for dendrimers of higher generation and/or with increasing interaction strength, as indicated in the phase diagram in Figure 1. Tsukruk et al. confirmed the predictions made by Mansfield et al. by using PAMAM dendrimers (generations G3.5–G10) in the construction of self-assembled monolayers (SAMs) [10,11]. PAMAMs of two adjacent generations with surface amine and carboxylic acid end-groups, respectively, were used in an electrostatic layer-by-layer deposition. The resulting films were studied using scanning probe microscopy and X-ray reflectivity. The dendrimers were found to possess an oblate shape with axial ratios in the range from 1:3 to 1:6. Similar results were obtained upon deposition of carboxylated PAMAMs on positively charged surfaces [12,13]. Next to electrostatic interactions, various other types of secondary interactions have been used to adhere dendrimers to surfaces, ultimately leading to flattened dendrimer structures. Sheiko et al. observed strong deformations of hydroxylterminated carbosilane dendrimers on mica due to multiple hydrogen bonding interactions [14]. These interactions were absent in the case of substrates that were first coated with a semifluorinated polymer resulting in a surface

Figure 1 ‘‘Phase diagram’’ displaying the influence of the interaction strength and generation number on the shape of a dendrimer at a surface [9].

Functional Dendritic Thin Films and Monolayers

7

other hand, the poly(benzyl ether) dendrimers are hydrophobic and want to minimize contact with the water surface. This property results in a conformational shape change from ellipsoidal for the lower generations to oblate for the higher generations [22]. Studies on dendrimers at interfaces have shed new light on the conformational behavior of these macromolecules. Conventionally, dendrimers were viewed as spherical macromolecules with a hollow-core, dense-shell character, according to the theoretical model of de Gennes and Hervet [25]. Later, Mansfield and others argued against this model because their calculations indicated that the end-groups were also back-folded into the dendritic interior to a nonnegligible extend [26–29]. When we take recent literature into account, we realize that next to the dendrimer structure itself, the medium around the dendrimer influences its conformational behavior significantly. Dendrimers can swell or shrink in response to a change in solvent type [30,31] or pH [32]—a property that obviously affects the end-group localization as well. In this context, the behavior of dendrimers at surfaces provides us with the most extreme examples of conformational changes in response to external stimuli. Depending on the strength and type of interaction between the dendrimer and the surface, they can flatten out almost completely. Although the unique spherical shape of the dendrimer is very often lost in self-assembled films, other specific properties, such as multifunctionality, porosity, and so forth, are retained. This important feature has led to the birth of a new and exiting area of research—functional dendritic surfaces—and will be discussed in the next section.

III.

PHOTO- AND CHEMICALLY RESPONSIVE DENDRITIC THIN FILMS

Dendrimers are highly potential materials for chemical sensing because of their morphological characteristics, high degree of synthetic flexibility, and multiply reactive sites [33]. The self-assembled monolayers on Au developed by Crooks et al. are very effective chemical sensors for volatile organic compounds (VOCs) [18]. The PAMAM dendrimer shown in Figure 2B, for instance, displays ideal chemical sensor features towards butanol: rapid response, low signal-to-noise ratio, and full reversibility. The extent of penetration of the VOC in the dendrimer framework is governed by the size and surface density of the dendrimer and hence by the generation number. Gorman et al. used the same permeability strategy in the construction of monolayers on gold from organothiol dendrons of the first through third generation [34]. The ability to trap and hold small molecules within these layers was found to be generation dependent. Miller et al. modified polyamido dendri-

Functional Dendritic Thin Films and Monolayers

9

[43,44]. The ferrocenyl units are located at the end of long, flexible, siliconcontaining branches and serve to electrically connect the enzyme to the electrode. The flexibility of the dendrimer is proposed to play an important role in the interaction with the redox center of glucose oxidase. Fre´chet et al. showed that both ionically and covalently bound SAMs of poly(benzyl ether) dendrimers [45,46] can serve as resists for scanning probe lithography [47,48]. The tip of a scanning probe microscope was used to write a pattern in a dendritic monolayer by oxidation of the underlying silicon substrate. The oxide pattern thus obtained could be removed selectively under aqueous hydrofluoric acid etching conditions. The dendrimer monolayers proved resistant to the HF(aq) etchant and in this way a positive tone image of the written pattern could be obtained with dimensions below 60 nm. The increased stability of the dendritic monolayers toward the etchant as compared to monolayers prepared from low molecular weight molecules is ascribed to the dense nature of the dendrimers used, resulting in a better protection of the anchoring group. Future applications of these dendritic resists are foreseen in the area of high-density data storage systems. Weener et al. prepared photo-responsive monolayers from azobenzenemodified PPI dendrimers that also hold promise in the area of optical data storage [49]. A fifth-generation PPI dendrimer was functionalized with equal amounts of palmitoyl and azobenzene containing alkyl chains, resulting in the formation of an amphiphilic copolymer with a random shell structure (Figure 5). Atomic force microscopy (AFM) and grazing X-ray reflectivity measurements revealed an orientation of the dendrimers within the monolayers identical to the one previously established by Schenning et al. (Figure 3) [20]. The azobenzene groups displayed facile and reversible isomerization in Langmuir and Langmuir–Blodgett monolayers as shown in Figure 5. The dendrimer used can be viewed as a well-defined, head-group-polymerized, amphiphilic system. In this case the use of the dendritic scaffold resulted in the formation of stable monolayers and, in addition, prevented microphase separation of the azobenzene units within the monolayers, resulting in reversible switching behavior. Vo¨gtle and Balzani et al. recently reported on PPI dendrimers up to generation 4 functionalized with azobenzene groups in the periphery with potential use as materials for holographic data storage [50]. Repetitive cistrans isomerizations of the azobenzene groups in solvent cast films resulted in a preferred orientation of the azobenzene moieties, perpendicular to the polarization direction of the incident light. AFM measurements performed on films after irradiation revealed the presence of relief gratings with heights up to 1500 nm.

10 Schenning et al.

Figure 5 Reversible switching of amphiphilic azobenzene–modified PPI dendrimer on quartz and at the air–water interface (inset).

Functional Dendritic Thin Films and Monolayers

IV.

11

DENDRIMERS IN LIGHT-EMITTING DIODES

Dendrimers are promising candidates in the area of organic light-emitting diodes (LEDs); they can enhance durability in these devices due to prevention of crystallization of the active component as a result of branching that in turn leads to amorphous films. This is important because crystallization or melting, caused by heat or short-circuit current, often results in device damage. Shirota et al. have synthesized so-called starburst molecules having a triphenylamine core [51]. These molecules, which can be regarded as first-generation dendrimers, have been successfully applied as hole transport layers in organic LEDs, and increased thermal stability of the devices was found. Based on the concept of Shirota et al., Strohriegl et al. constructed starburst molecules based on oxadiazoles [52,53]. These materials were tested as electron-transporting layers in LED together with poly( p-phenylenevinylene) (PPV). The two-layer LEDs showed an improved quantum yield and brightness compared to LEDs based on PPV only. A second reason that conjugated dendrimers are of interest as materials for application in organic LED is that they can be designed in such a way that the core defines the color of the light emission, while the surface groups control the processing properties and the branching units act as transporting units [51]. Moore et al. have reported for the first time this concept in the fabrication of a single-layer LED using dendrimers consisting of a highly fluorescent 9,10-bis(phenylethynyl)anthracene core, phenylacetylene branches as electron capture, and a triphenylamine periphery to promote efficient hole transport and recombination [54]. High-quality amorphous films could be obtained by spin coating on indium-tin-oxide (ITO) wafers with thickness ˚ . After deposition of the electron-injecting electrode between 1200 and 600 A (Mg/Ag), emission of relatively low intensity of the anthracene core was observed. Burn et al. have synthesized conjugated dendrimers for LEDs based on the phenylenevinylene structure and have investigated the effect of generation on the photophysical and electroluminescence properties of the materials [55,56]. The dendrimers consisted of a distyryl benzene blue emissive core, stilbene branches as charge-transporting units, and t-butyl surface groups for good processability. All three dendritic generations can be spincoated to give good-quality thin films. The devices show an evolution of the color of emission from blue to white on a time scale of seconds to minutes. The authors also observed an effect of the generation number on the device efficiency. The external quantum efficiency of the single-layer devices was found to be 0.01% for generation 1, 0.07–0.09% for generation 2, and 0.03% for generation 3. An important feature of incorporating chromophores as cores in dendrimers is that it provides the chromophores in a

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solution-processible form, while aggregation is reduced. Burn et al. have demonstrated this concept in controlling the color in LEDs using stilbene dendrimers having three different types of cores [57,58]. A distyrylbenzene core was chosen for the blue emission, an anthrachene derivative for the green emission, and a free-base porphyrin for the red emission. In all cases, good-quality thin films could be spin-coated and LEDs were successfully constructed with external quantum efficiencies between 0.1% and 0.01%. Freeman and Fre´chet have synthesized dendrimers labeled peripherically with hole transmitting amines and with emissive dyes as cores [59]. Single-layer devices on ITO wafer could be prepared with an average thick˚ by spin casting. The devices that used an Mg-Ag cathode ness of 800–900 A had moderately high turn-on voltages between 12 and 15 V and could be driven at voltages exceeding 30 V. Lupton et al. have used a first generation distyrylbenzene cored dendrimer, which utilizes triazine rings to the link the surface groups [60]. Triazine units, which can be regarded as electron acceptors, were used as branching points, whereas the core unit, a disytrylbenzene moiety, acted as a blue-emitting material. Blue-emitting single-layer devices with an external quantum efficiency of 0.003% were successfully constructed. However, Lupton et al. found that the intensity decreased in a few minutes and the emitted light was observed to change irreversibly from blue–white to white. Recently, Schenning et al. presented a very universal approach to adjusting the emission wavelength of LEDs using dendrimer/dye assemblies [61]. PPI dendrimers up to generation 5 were functionalized with ␲-conjugated oligo(p-phenylene vinylene)s (OPVs) (Figure 6), yielding amphiphilic macromolecules that self-assemble at the air–water interface forming homogeneous thin films. Moreover, these dendrimers proved to be good hosts for anionic dye molecules. Thin films prepared from these host–guest complexes showed efficient (>90%) energy transfer from the OPV units to the encapsulated dye molecules (Figure 6). By choosing the appropriate dye, the emission wavelength could be tuned. Furthermore, these supramolecular systems compatibilize ionic dyes with organic semiconducting polymers, and energy transfer is observed from the organic polymer to the dyes. Preliminary experiments showed that it was possible to construct LED’s from the dye–dendrimer system. The red color of the dye was seen, whereas when only dendrimers were used, the green emission of the OPV units was visible. V.

LIQUID CRYSTALLINE DENDRIMERS

Thermotropic liquid crystalline phases or mesophases are usually formed by rod-like (calamitic) or disk-like (discotic) molecules. Spheroidal dendrimers

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13

Figure 6 Schematic representation of the energy transfer process between OPVPPI dendrimer and a dye guest molecule.

are therefore incapable of forming mesophases unless they are flexible, because this would allow them to deform and subsequently line up in a common orientation. Percec et al. were the first to report on flexible, nonspherical dendrimers with AB2 mesogens in the branches that exhibited nematic and smectic liquid crystalline behavior [62,63]. Since then several reports have been made concerning the decoration of dendrimers with mesogenic units on the periphery. Frey et al. attached several mesogenic units to carbosilane dendrimers, such as cyanobiphenyl [64] and cholesteryl [65]. Besides these clas-

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sical calamitic mesogens, carbosilane dendrimers carrying perfluoroalkyl groups have also been reported [66]. In all cases, smectic mesophases were observed in which the dendrimers were deformed (stretched) in a cooperative way to adjust to the superstructure imposed by the mesogens [67]. However, reduction of the spacer length and/or increasing the generation number complicated the formation of well-developed smectic phases [64]. Latterman and coworkers observed a comparable generation effect by using functionalized PPI dendrimers with mesogenic 3,4-bis(decyloxy)benzoyl groups [68,69]. Generations 1–4 exhibited a hexagonal columnar mesophase in which the dendrimers had a cylindrical conformation (Figure 7a), whereas generation 5 was not liquid crystalline. The lack of mesomorphism for the highest generation dendrimer was attributed to the incapability of this dendrimer to reorganize into a cylindrical shape. Baars et al. functionalized PPI dendrimers (generations 1, 3, 5) with pentaoxycyanobiphenyl and decyloxycyanobiphenyl mesogens [70]. All dendrimers were found to exhibit smectic A mesophases. The smectic A layer spacings observed were independent of the dendrimer generation for both spacer lengths, indicating that the dendritic part was completely distorted, even for the higher generations (Fig. 7b). These findings were later confirmed by Yonetake et al., who also used cyanobiphenyl-substituted PPI dendrimers [71]. PAMAM dendrimers (generations 1–4) derivatized with mesogenic ester units [4-(4⬘-decyloxybenzyloxy)salicylaldehyde] also displayed a smectic A mesophase. The molecular picture to explain this behavior is identical to the flattened dendrimer model previously described by Baars et al. [70]. For PPI dendrimers functionalized with palmitoyl alkyl chains ( — C15H31), lamellae structures were present in the bulk according to X-ray

Figure 7 Conformational changes of PPI dendrimers in liquid crystalline materials from spherical to (a) cylindrical, in hexagonal columnar mesophase, and (b) ellipsoid, in smectic A mesophase.

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diffraction, transmission electron microscopy [72], and small-angle neutron scattering (SANS) [73] measurements. The interlamellae distance found was 5 nm and independent of dendrimer generation, a feature that again could be explained by assuming a flattened dendrimer shape. Stark et al. recently described a similar observation in a SANS/small angle x-ray scattering (SAXS) study of perfluorinated carbosilane dendrimers. In this case, the strong tendency of the perfluorinated end-groups to form layered structures led to strong deformations of the dendritic core [74]. The above-mentioned examples show that in general two opposing forces compete in the formation of liquid crystalline phases from dendrimers terminated with mesogenic units. The functionalization of PAMAM, PPI, and carbosilane dendrimers with mesogenic groups results in an initial spherical arrangement, whereas the mesogenic units tend to interact with each other to form a parallel arrangement at the same time. Therefore, the flexibility of dendrimers is a prerequisite for the successful preparation of liquid crystalline materials from these macromolecules. Recent exciting developments in the field of liquid crystalline dendrimers include control over polymer shape by rational design [75–79] and applications of dendrimers in liquid crystalline displays [80].

VI.

CONCLUSIONS AND PROSPECTS

The successful application of dendrimers as functional building blocks in the construction of mono- and multilayers and liquid crystals is mainly due to the flexible nature intrinsic to most dendrimers. The numerous examples involving assemblies of dendrimers described in this chapter clearly illustrate that the conventional viewpoint of dendrimers, as spherical macromolecules with a dense shell, often has to be adjusted when dendrimers are used as building blocks in the construction of higher order aggregates. The shape of most dendrimers changes drastically from spherical to oblate or prolate under the influence of secondary interactions when they are incorporated in monolayers and liquid crystalline materials. In this context it is important to realize that next to the dendrimer structure itself, the medium around the dendrimer influences its conformational behavior significantly. External stimuli, such as hydrophobic effects, phase separation, electrostatic interactions, and hydrogen bonding interactions, are driving forces that determine the localization of the end-groups, and hence the dendrimer shape, to a large extent. The fact that unique dendritic properties such as multifunctionality, porosity, and chemical tunability are often retained within assemblies prepared from these macromolecules explains their successful application in the

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fields of chemical and biological sensors, lithography, photoresponsive materials, organic light-emitting diodes, and liquid crystalline materials.

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2 A Theoretical Study of Electrical Contacts to Self-Assembled Molecular Wires on Conducting Substrates ˚ sa Johansson, and Ma˚rten Bjo¨rketun Sven Stafstro¨m, A Linko¨ping University, Linko¨ping, Sweden

I.

INTRODUCTION

The ability to control structural and electronic properties at a nanometer scale has during the last decade opened up a new field of research, nanotechnology. Carbon-based nanostructured materials have had an important role in this development, primarily because of carbon nanotubes [1,2], but also other types of molecular wires have been widely studied [3–6]. One common problem to all such molecular-based nanoelectronics is how to make the contact to the molecules. The contact resistance is determined by the interactions between the molecules and the contact surface. In order to control these interactions it is necessary to have a detailed understanding of the interfacial interactions and how these interactions relate to the contact resistance. This type of study is also relevant in the context of other soft material electronics, such as polymer/molecular based light-emitting diodes [7] and molecular field effect transistors [8]. The description of the electronic structure of the organo/metallic contact is somewhat of a mixture between the band structure picture, used to describe the contact (or lead), and the discrete orbital picture of the molecular wire. Transport across the interface can be viewed as a scattering process in which the transmission (forward scattering) is determined by the interactions between the wavefunctions of the contacts (leads) and the mo23

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lecular orbitals as well as the localization (delocalization) of the molecular orbitals over the molecule. In this chapter we do not discuss the actual transport process; instead we focus on the interactions that take place at the interface. Therefore, it is not necessary to describe the electronic structure of the metallic contact using the band structure picture. Rather, we model the metallic contacts with a cluster that is large enough to cover most of the interaction energies at the interface. The attractive interaction at the interface between the contact and the molecular wire is usually classified as chemisorption or physisorption. Chemisorption denotes the process in which a chemical bond is formed between the adsorbed molecule and the surface. The bond could be covalent, i.e., exchange of electrons between the interacting substances, or ionic for which an electron transfer takes place at the interface. In the case of physisorption the molecule is held onto the surface by attractive so-called dispersive forces that arise due to a correlated motion of the electrons on the two subsystems [9]. Since this correlation increases in strength the closer the subsystems come to each other, this force is always attractive. The equilibrium distance is determined by a balance between this attraction and repulsive short-range interactions. In general, the dispersive forces are weaker than the covalent or ionic forces, which results in larger equilibrium separation between the subsystems in the case of physisorption compared to chemisorption. As will be shown below, this does not imply that the electronic coupling between the conducting states of the molecular wire and the metal contact differ in the same way. In the case of chemisorption, the formation of stronger covalent or ionic bonds between the surface and the molecule brings the molecule in closer contact with the outermost layer of the surface. Charge transport across such a bond is strongly energy dependent and reflects the local electronic structure of the system. Orbital wavefunctions that are localized to the bond are not involved in the transport process, whereas orbitals that are delocalized over both the contact and the molecule give rise to a conducting channel. In this chapter we will discuss the interaction between two types of systems, one describing chemisorption and one describing physisorption. The standard approach to chemisorption of molecular wires to a metal contact has been to use gold as the contact material and to attach a thiol group to the molecule forming the conducting wire, either in the form of a selfassembled monolayer (SAM) [3,4] or in the gap of a mechanically controlled break junction [5]. The thiol group forms a strong chemical bond to the gold surface, which determines the orientation of the molecule on top of the surface. The symmetry of the gold orbitals forming the bond to sulfur is

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important, as is the coupling of sulfur to the rest of the molecular wire, which could be a saturated hydrocarbon chain or a conjugated oligomer. The case of physisorption applies for instance to the studies of conductance through C60 [10] but also to all cases in which a scanning tunneling microscope (STM) tip is used as one of the leads to contact the molecular wire and to the case of depositing carbon nanotubes (NTs) on top of metallic contacts [2]. Molecular–metal interaction in the case of physisorption are usually described as a weak hopping between neighboring atoms on the metal and the molecule. In this chapter we have restricted the description of the possible contact metals to graphite. Graphite has been used in contacts to overlayers of polycyclic aromatic hydrocarbons (PAHs) [11]. These molecules can self-organize in so-called ␲ stacks, i.e., stacks of PAH that interact via the ␲ systems of the molecules [12–15]. These ␲ stacks are highly conducting along the stacking direction and form, in such a way, a molecular wire. Here, we will discuss the interaction energies of the molecules within the ␲ stacks as well as of the contact to the graphite surface.

II.

METHODOLOGY

Detailed theoretical investigations of molecules adsorbed on the surface of a metal contact have to be performed at an ab initio quantum chemical level of modeling [16]. In principle, both density functional theory (DFT) and Hartree–Fock (HF) theory can be used. Both these theories are based on an iterative approach to determine the electronic structure (or charge density) that minimizes the total energy of the system. This so-called self-consistent field (SCF) approach gives a very good description of chemisorbed species, i.e., subsystems bond together by covalent bonds. In the calculations of the gold-thiol system we have used a double-␨ basis set of contracted Gaussian functions including a relativistic effective core potential for gold (LanL2DZ). The (111) gold surface was modeled by a 10-atom cluster (Figure 1). The geometry of the cluster was fixed to that ˚ ). It has been shown [17] that of a gold lattice (Au-Au bond length is 2.88 A an underlying layer of atoms is of great importance for the electronic structure of thiolates on gold. Therefore, the 10-atom cluster consists of a twolayer structure, one upper layer of 7 atoms and one lower layer of 3 atoms. As molecular wire we use a single phenyl monomer with a sulfur attached to one of the carbon atoms in the phenyl ring. The formation of a chemical bond with the thiol group is a fairly local type of interaction. Thus, even though the gold cluster is very small, we believe that it is large enough to reveal the correct symmetry of the chemical bond and thus to give some

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Figure 1 Sulfur-bonded phenyl on a hollow site of the gold cluster (top) and pyrene (C16H10) on a model system of graphite (bottom).

insight into what type of coupling there is between the ␲ system of the phenyl molecule and the gold (contact). For practical reasons it is not possible to allow for a complete geometry relaxation of all the degrees of freedom related to the orientation of the molecule relative to the gold cluster in one single calculation. Since the gold cluster is quite small, there are large edge effects that can lead to energy minima that reflect the properties of the cluster rather than of the gold surface. Since we are interested in the properties of the surface, we have to restrict the positioning of the molecule to the interior of the cluster. Figure 1 shows the total system configuration. The sulfur atom is placed in its equilibrium position in the middle of a three-atom gold triangle and right above a gold atom in the second layer. The rest of the orientation of the molecule relative to the gold cluster is described in terms of the direction of the S-C bond both relative to the surface normal angle ␪ in Figure 1 and

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the BSSE is quite dominating and has to be calculated. The standard procedure is to use the so-called counterpoise corrections [20] in which the interaction energy of the molecular complex is defined as the difference in total energy of the complex and the two individual subsystems using the basis set of the molecular complex also for calculations of the individual molecules.

III.

RESULTS AND DISCUSSION

A.

Thiols on Gold

Geometry optimization of the free C6H5SH molecule results in an S-C dis˚ . All C-C bond lengths are almost identical in the phenyl tance of 1.83 A ˚ . As an intermediate step we have reoptimized the moring, around 1.39 A lecular geometry by replacing the hydrogen with a single gold atom. This ˚ . The ring results in a very small change of the S-C bond length, to 1.84 A structure also remains unchanged with an aromatic structure in which the ˚ . The Au-S bond length is 2.41 A ˚ , which is C-C bond lengths are 1.40 A very close to the value obtained for the 10-atom cluster (see below). Using this geometry we checked first the lateral position of the molecule on the gold surface. Experimentally it is known that the preferred position of the sulfur is in the hollow site, on top of a gold atom in the second layer of the gold substrate [21] (Figure 1); it has also been calculated by, e.g., Sellers et al. [17]. We have calculated the total energy of the complex for different position of the sulfur relative to the surface and observed a clear minimum for the experimentally predicted position. This result indicates that the Au10 cluster is a relevant model at least for the more qualitative features of the gold–thiol interaction. We have also performed similar studies with the phenyl ring replaced by an alkyl chain and obtained the same equilibrium position of the sulfur atom relative to the atoms of the gold surface [22]. Thus, the type of interactions that determine the position of the thiol-substituted molecule to the surface are independent on the detailed structure of the molecule attached to sulfur. To obtain the ground state distance between the gold surface and the sulfur atom (placed above a hollow site) we performed a series of total˚ and 5.0 A ˚ (Figure energy calculations, varying this distance between 2.2 A 2). When going from small distances to higher (curve 1), making use of the density matrix from the next lower distance, we see a clear energy minimum ˚ and a continuous increase in total energy as the for the distance 2.50 A distance increases. However, when going from large to small distances (curve 2), using the density matrix from the next higher distance, for large distances the energy is much lower than that calculated when going upward

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Figure 2 Interaction energy of sulfur-bonded phenyl as a function of the distance to the gold surface. The reference energy is the total energy of the two systems fully separated.

in distance. Obviously, this discrepancy is a computational artifact, i.e., the SCF calculation converges to two different local minima depending on the input density matrix. This has to be due to the fact that there is a substantial energy barrier in redistributing the charge density from that corresponding to the case when a chemical bond is formed (curve 1) to that of no chemical bond. In order to put this in more quantitative terms, Table 1 shows the Mulliken charges of the sulfur atom and the S-C bond as well as the Au/S bond order matrix element (related to the bond strength) for different distances between the sulfur atom and the surface. The upper six rows are values from curve 1 and the two last rows are values from curve 2. The values of the bond order give a clear indication that curve 2 corresponds to a state for which the molecule and the surface are not bonded to each other (low bond order), whereas curve 1 corresponds to a state for which they are bonded to each other (high bond order). The difference in total energy between the bonded ground state at 2.50 ˚ and the sum of the total energies for the gold cluster and the free C6H5S A

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Table 1 Mulliken charges of the sulfur, the S-C group, and bond order matrix element for the S-Au bond for different sulfur–surface distancesa ˚) Distance (A

Charge S

Charge S-C

Bond order S-Au

2.20 2.50 2.60 3.10 3.40 3.48

⫺0.1202 ⫺0.1712 ⫺0.1871 ⫺0.2431 ⫺0.2708 ⫺0.2729

⫺0.3120 ⫺0.4183 ⫺0.4480 ⫺0.5544 ⫺0.6062 ⫺0.6109

0.2433 0.2686 0.2778 0.2952 0.2878 0.2802

3.50 5.00

⫹0.1627 ⫹0.0926

⫹0.0465 ⫺0.0326

0.0821 0.0372

a

The six upper rows correspond to curve 1 and the two lowest to curve 2 in Figure 2.

molecule (geometry optimized) corresponds to the total dissociation energy of the thiol-phenyl system from the gold surface. Due to the computational artifact discussed above, it is difficult to estimate the dissociation barrier, ˚ i.e., the barrier when moving the molecule from the ground state at 2.50 A to infinite separation (see curve 1). It might be the case that the system remains in its local energy minimum, i.e., the bonded state, even though the total energy is lower for the nonbonded electronic configuration. However, with thermal fluctuations included we can expect that the energy barrier will be overcome fairly close to the crossing of the two potential curves. A lower limit of the barrier would therefore be approximately 0.3 eV. With this barrier it is likely that the molecule will remain bonded to the gold substrate even in the case of local heating, which can be expected when a current is passing through the molecular wire. The calculated value for the interaction energy is 0.1 eV. We also stress that, based on our calculations, conclusions can only be drawn concerning the dissociation process, not regarding a selfassembling reaction when the approaching molecule is C6H5S and a hydrogen molecule might be created in the surface reaction [21]. ˚ above Given the position of the sulfur atom in a hollow site, 2.50 A the gold surface, we reoptimized the orientation and the structure of the molecule (Figure 1). The angle between C-S and the surface normal is ␪ = 37⬚. Sellers et al. [16] have calculated this angle to 0⬚ and 76⬚ corresponding to an sp2 and sp3 hybridization of sulfur, respectively. The rotation angle, ␾, defined as 0⬚ when the molecule is directed to the center atom of the cluster, is 50⬚, and the tilt angle, ␣, defined as 0⬚ when the plane of the phenyl ring is perpendicular to the surface, is 80⬚. The overall result is that shown in Figure 1. It must be stressed, however, that the energy surface for changes

Electrical Contacts to Molecular Wires

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in all these angles is very flat and that the molecule is likely to rotate quite freely at ambient temperatures. Since the gold cluster is very small, the charge distribution in this cluster could be slightly different from the image type of charge distribution that is expected for a semiinfinite gold substrate. Thus, the ionic contribution to the gold-sulfur bond is less well described in the present model and the value of the dissociation energy discussed above is therefore not very precise. The Mulliken charges shown in Table 1 are also more of a qualitative description of the actual charge transfer at the interface. The charge transfer process is more related to the thiol group than to the molecular wire attached to the thiol and is therefore a general result for this type of connection site. In the context of molecular wires, the dipole layer at the interface will certainly affect the transport of electrons through the thiol group and into the ␲ system. B.

Polycyclic Hydrocarbons on Graphite

As discussed above, a ␲ stack of PAHs is one candidate for the formation of a conducting path at the molecular scale. The intermolecular interactions within the ␲ stack are of interest for the electronic delocalization of the electronic wavefunction over the stack. Furthermore, the stack is usually contacted by means of physisorption on a graphite surface [12]. The interactions involved in the contact problem as well as in the ␲ stack are closely related, as will be apparent from the discussion below. In Figure 3 is shown the interaction energies for pyrene (C16H10) dimers as a function of the separation between the two molecules. We find that, as in graphite, the A-B stacking is the most stable orientation. Without taking ˚ in the A-B BSSE into account, the calculated equilibrium distance is 3.3 A ˚ . This case. The corresponding equilibrium distance for A-A stacking is 3.7 A difference simply reflects the fact that binding in the case of A-B stacking is stronger than in the case of A-A stacking: at the equilibrium distance the interaction energy for A-B stacking is 0.47 eV. The corresponding number for A-A stacking is 0.29 eV. As stated above, there are a number of sources of errors in these estimates of the interaction energies and equilibrium distances. The MP2 result certainly overestimates the bond strength [19]. Since the systems in which we are interested are quite large, we do not have the possibility of going beyond the MP2 level of the perturbation treatment. Calculations on smaller systems have shown that the MP3 and MP4 indeed give lower interaction energies and longer equilibrium distances [19,23]. The trend is oscillatory: HF largely underestimates the interaction energy, MP2 gives an overestimate, MP3 gives an interaction energy somewhere between the MP2 and

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Figure 3 Interaction energy of a pyrene dimer as a function of the distance between the parallel planes of the two molecules. The reference energy is the total energy of the two systems fully separated.

HF results, and the MP4 result, finally, is somewhere between the MP2 and MP3 results in terms of both the interaction energy and the equilibrium ˚ between the distance. The difference in separation distance is about 0.3 A MP2 and MP4 results on benzene-like systems [23]. The effect of using an incomplete basis set is another source of error, as concerns the total energy of a given calculation but also due to the BSSE, which provide an effectively more extended basis set the closer the two subsystems are. Again, due to the size of the system we are unable to perform calculations for larger basis sets than the typical 3-21G** basis set (on the pyrene dimer). Using a minimal STO-3G basis set instead of the 3-21G ˚ to 3.5 A ˚ for the gives an increase in the equilibrium separation from 3.3 A A-B orientation, in very close agreement with the interlayer separation in graphite. Using the 3-21G** basis gives instead an equilibrium distance of ˚ . With the counterpoise [20] correction of the BSSE the result is dra3.2 A matically altered. Calculation with the STO-3G basis set results in a negative interaction energy. Using the 3-21G basis set the equilibrium distance is

Electrical Contacts to Molecular Wires

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˚ with an interaction energy as low as 0.12 eV (Figure 3). shifted to 3.9 A ˚ . Naturally, With the 3-21G** basis set the equilibrium distance is also 3.9 A the BSSE is reduced when going to very large basis sets and will be zero in the limit of a complete basis. Note that, for a finite basis, the ‘‘correction’’ is always too large since the basis functions in the counterpoise correction are used exclusively to improve the result for the single molecule, whereas in the dimer complex the coefficients are optimized to describe both molecules. Thus, an estimate of the truly corrected BSSE should give an equilibrium distance somewhere between the uncorrected and the counterpoise-corrected values. At the MP2 level and with the 3-21G basis set, this would end up with a value some˚. where between 3.3 and 3.9 A It is clear from the relatively large variations in the results described above that it is difficult to present a result that is fully adequate from the point of view of both correlation treatment and size of the basis set. The better the basis set is, the shorter the equilibrium distance obtained. For the correlation treatment, any improvement of the MP2 result increases the equilibrium distance. Obviously, these two errors cancel to some extent. Together with experimental estimates of, for instance, the equilibrium distance in the stacks [24], we find that an MP2 calculation with a 3-21G basis set slightly overestimates the binding strength but that nonetheless this method describes the interaction of the dimer both structurally and electronically. The interaction energies at the equilibrium separation distance have also been calculated for coronene (C24H12) dimers. As in the case of pyrene, the A-B stacking is preferred. THe MP2/3-21G interaction energy to form the dimer is 0.93 eV. Comparison with the results on pyrene shows that an increase in the number of atoms by 1.5, when going from pyrene to coronene, results in a near-doubling of the interaction energy. In other terms, by dividing the total interaction energy by the number of carbon atoms in the monomer the resulting interaction energies per carbon atom are 29 and 39 meV in the pyrene and coronene dimers, respectively. This nonlinear increase in the interaction energy with number of atoms in the system is a clear indication that the stacking becomes increasingly stable the larger the PAH molecule is. Experimentally, ␲ stacking has not been observed for smaller molecules such as pyrene. Perylene (C20H12) has a stable dimer phase [24] but no extended stacking has been observed. However, hexaperihexabenzocoronene (C42H18) forms stacks on graphite. Also the bulk crystal structure of the corresponding molecular solid contains extended ␲ stacks. Now we turn to the issue of physisorption of PAH on a graphite surface. The interaction energy of pyrene (C16H10) on a model system of a graphite surface (C42H18; see Figure 1) was calculated at the MP2 level, using the 3-21G basis set. The interaction energies are shown in Figure 4 for the

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Figure 4 Interaction energy of pyrene on a model system of graphite (C42H18) as a function of the distance between the parallel planes of the two molecules. The reference energy is the total energy of the two systems fully separated.

A-B and A-A stacking orientations of pyrene on top of graphite. As in the systems discussed above, the A-B stacking is the most stable orientation. At the equilibrium distance this orientation is 0.22 eV lower in energy than the A-A stacking. Without taking BSSE into account the calculated equilibrium ˚ in the A-B case, about 0.4 A ˚ shorter than the interlayer distance is 3.1 A separation in graphite. The corresponding equilibrium distance for A-A ˚. stacking is 3.3 A The interaction energy at the equilibrium separation distance is 1.01 eV, which corresponds to 63 meV per carbon atom in pyrene. This can be compared with pyrene-on-pyrene for which the interaction energy per carbon is 29 meV. This shows again that the intermolecular ␲-␲ interaction is quite long range; that is, the carbon atoms of the model system for the graphite surface that have no nearest-neighbor carbon atom from the pyrene system actually contribute to the binding. Of course, the present system also suffers from both the incompleteness of the correlation treatment and the basis set. As discussed above, these effects cause errors of different sign that partially

Electrical Contacts to Molecular Wires

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cancel. Most probably, however, using the MP2 method with the 3-21G basis set we slightly overestimate the interaction energy as a result of the tendency of MP2 to include too much binding contribution to the correlation treat˚ should therefore be adment. The calculated equilibrium distance of 3.1 A justed to a somewhat larger value.

IV.

SUMMARY AND CONCLUSIONS

It is clear that there are fundamental differences in the behavior of the interaction between the two cases presented above. Dissociation of a thiol group from a gold surface has a barrier of about 0.3 eV followed by a repulsive potential. Note that the adsorption process in this case is not the reverse of dissociation since the adsorption also involves the release of a hydrogen atom. The energy barrier for this process is thus more complex and involves breaking of an S-H bond together with the formation of the S-Au bond. PAHs physisorption on a graphite surface involves no chemical reactions, and adsorption and dissociation can to a first approximation be considered as reversible processes. There is no barrier for adsorption, and the resulting interaction energy per carbon atom of the adsorbed molecule is as large as 0.06 eV for pyrene, which results in a total interaction energy of about 1 eV for pyrene. This number increases with increasing size of the PAH molecules. Since the PAHs lie flat on the surface, with essentially all carbon atoms contributing to the bonding, the total interaction energy of large PAH is very large. These molecules will therefore stick to the surface. The difference in interaction energy between A-B and A-A stacking also indicates that the lateral motion along the graphite surface is quite restricted. The relatively (per atom) strong chemisorption bond manifests itself ˚ as compared to the physisorbed in a shorter equilibrium distance, i.e., 2.5 A substances. For charge transport, this would indicate a better contact between the contact surface and the molecular wire. However, as compared to the PAH physisorbed on graphite, the thiol group is a very limited in terms of the number of channels (eigen states) that can transport charge across the interface [25]. In the case of larger PAHs, not only the total interaction energy is larger but also the contact resistance is smaller since the number of active channels that allow for charge transport at a certain energy is larger. We can consider the system as having N parallel channels, where N is the number of atoms in the PAH molecule [15]. Thus, in constructing lowresistance contacts to molecular wires, it is important to consider both the aspect of the actual binding of the wire to the contact and the number of conducting channels that will be accessible for charge transport. From the results presented here, stacks of PAHs physisorbed on a graphite surface is

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an ideal system for such a contact. The PAHs also introduce an additional flexibility because it would be possible to vary the contact resistance by choosing PAH molecules of different sizes.

ACKNOWLEDGMENTS Computational resources were provided by the Swedish Council for High Performance Computing (NSC). Financial support from the Swedish Research Council for Engineering Science (TFR) and the Swedish Natural Science Research Council (NFR) is gratefully acknowledged.

REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.

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S. Iijima. Nature 354:56, 1991. S. J. Tans, M. H. Devoret, H. Dai, A. Thess, R. E. Smalley, L. J. Geerligs, C. Dekker. Nature 386:474, 1997. M. P. Samanta. Phys Rev B 53:R 7626, 1996. L. A. Bumm, J. J. Arnold, M. T. Cygan, T. D. Dunbar, T. P. Burgin, L. Jones II, D. L. Allara, J. M. Tour, P. S. Weiss. Science 271:1705, 1996. M. A. Reed, C. Zhou, C. J. Muller, T. P. Burgin, J. M. Tour. Science 278:252, 1997. V. Mujica, M. Kemo, M. A. Ratner. J Chem Phys 101:6856, 1994. J. H. Burroghs, D. D. C. Bradley, A. R. Brown, R. N. Marks, K. Mackay, R. H. Friend, P. L. Burns, A. B. Holmes. Nature 347:539, 1990. J. H. Scho¨n, Ch. Kloc, B. Batlogg. Nature 406:702, 2000. A. J. Stone, The Theory of Intermolecular Forces. Oxford: Clarendon Press, 1996. C. Joachim, J. K. Gimzewski. Chem Phys Lett 265:353, 1997. P. Herwig, C. W. Kayser, K. Mu¨llen, H. W. Spiess. Adv Mater 8:510, 1996. K. Mullen, J. P. Rabe. Ann NY Acad Sci 33:205, 1998. A. M. van de Craats, J. M. Warman, K. Mu¨llen, Y. Geerts, J. D. Band. Adv Mater 10:36, 1996. M. Kiel, P. Samorı´, D. A. dos Santos, T. Kugler, S. Stafstro¨m, J. D. Brand, K. Mu¨llen, J. L. Bre´das, J. P. Rabe, W. R. Salaneck. J Phys Chem B 104:3967, 2000. M. Paulsson, S. Strafstro¨m. J Phys C 12:9433, 2000. M. J. Frisch et al. GAUSSIAN 98, Gaussian Inc., Pittsburg, PA, 1998. H. Sellers, A. Ulman, Y. Shnidman, J. E. Eilers. J Am Chem Soc 115:9389, 1993. C. A. Hunter, J. K. Sanders. J Am Chem Soc 112:5525, 1990. S. Tsuzuki, T. Uchimaru, K. Matsumura, M. Mikami, K. Tanabe. Chem Phys Lett 319:547, 2000. S. F. Boys, F. Bernardi. Mol Phys 19:553, 1970.

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P. Fenter, A. Eberhardt, P. Eisenberger. Science 266:1216, 1994. J. Krohn, Master’s thesis, Linko¨ping University, Sweden, 1998. M. Pickholtz, S. Stafstro¨m. Chemical Physics (in press). J Tanaka. Bull Chem Soc Jpn 36:1237, 1963. ˚ . Johansson, S. Stafstro¨m. Chem Phys Lett 322:301, 2000. A

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3 A Theoretical Insight into Organic Interfaces in Electro-Optic Devices Je´roˆme Cornil, Donizetti dos Santos, and Roberto Lazzaroni University of Mons-Hainaut, Mons, Belgium

Massimo Malagoli and Jean-Luc Bre´das University of Arizona, Tucson, Arizona

I.

INTRODUCTION

The reports of efficient electroluminescence from organic conjugated molecules [1] and polymers [2] have triggered over the recent years spectacular developments in the field of organic-based electro-optic devices, such as light-emitting diodes (LEDs), photodiodes, and solid-state lasers. The typical architecture of a single-layer LED consists of an organic film sandwiched between two metallic contacts [3,4]. A first process in the operation of an LED is the injection of electrons and holes in the organic layer under the application of a voltage in forward bias between the two electrodes. In order to reduce the energy barriers at the metal–organic interfaces and hence the driving voltage required for producing light out of the plastic layer, the metal used at the cathode (typically Al, Ca, Mg, or alloys) thereof has a low work function with the Fermi energy matching at best the energy of the lowest unoccupied level of the organic compound; in contrast, the metal at the anode (indium-tin-oxide, or ITO, in most devices due to its metallic and transparent character) has a larger work function with a Fermi energy lying close to the energy of the highest occupied level of the molecule. The nature of the chemical interactions taking place at the metal–organic interface plays a critical role in determining the efficiency of emitting devices and has been extensively described in previous theoretical and experimental works [3,5]. 39

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It has been shown that the metal atoms of the electrodes can diffuse into the near-surface region of the organic layer and give rise to (a) formation of covalent bonds between the metal atoms and the molecular backbones, thus leading to a disruption of conjugation along the chains; or (b) charge transfer processes (redox chemistry) between the metallic centers and the conjugated chains. It is worth stressing that the creation of such interfacial regions between the bulk of the metal and that of the organic semiconductor is not necessarily detrimental to the device performance. We also note that nowadays the injection process is helped by the incorporation of an extremely thin interfacial layer between the electrode and the organic material (such as a conducting polymer on top of ITO [6] or a lithium halide on top of an aluminum electrode [7]). Once injected, electrons and holes start drifting in opposite directions within the organic layer under the influence of the static electric field until there is a possibility to recombine under the form of bound electron-hole pairs (also referred to as excitons). According to simple quantum mechanical considerations, a statistical spin recombination would be expected to generate singlet and triplet excitons in a 1:3 ratio; this would set the upper limit of electroluminescence quantum yield to 25% since only singlet excitons decay radiatively to the ground state. However, recent experimental studies have pointed to spin-dependent recombination rates leading to ratio between the electroluminescence and photoluminescence quantum yield as large as 50% [3,8,9]; this is further supported by recent theoretical estimates of the cross sections for the recombination of an electron and a hole on two separate chains into the lowest intramolecular singlet vs. triplet excited state [10]. Over the last decade, joint efforts in the fields of synthetic chemistry, devices fabrication and characterization, and theory have allowed for a good understanding of the mechanisms governing LED performances [11]. This contributed to an increase in the external quantum efficiency reported for the first diodes by several orders of magnitude up to state-of-the-art values around 8–9% [12]. Organic LEDs can now emit over the whole visible range and have recently reached the stage of industrial production for the development of backlights and low-resolution matrix displays [13,14]. High electroluminescence quantum yield in LEDs requires balanced injection of electrons and holes within the organic layer [15] and restriction of the nonradiative decay routes of the singlet excitons, e.g., following captures by quenching centers; these goals can be fulfilled through the design of multilayer devices [16,17] where an adequate choice of the constituting layers may help: (a) to tune the barrier between the Fermi energy of the bulk at the anode metal (when no conducting layer is deposited on its surface) and the energy of the highest occupied level of the molecules in the

Organic Interfaces in Electro-Optic Devices

41

adjacent layer and independently, the barrier at the cathode between the Fermi energy of the bulk metal and the energy of the lowest unoccupied level of the neighboring organic compound. The organic layers used to facilitate the injection of holes and electrons in the plastic matrix are often referred to as hole- and electron-transporting layers, respectively, although this terminology can be misleading because it suggests that the charges necessarily have a high mobility in these layers; (b) to create energy barriers at the interface between two organic phases by inducing an offset between the energy of the highest occupied and/or lowest unoccupied levels of two different organic constituents (Figure 1). These barriers tend to confine electrons and holes at a molecular interface or in a central emitting layer within the organic matrix, thus generating a preferential zone for the radiative recombination of the electron-hole pairs; the barriers can also prevent the recombination to occur too close to the organic–metal interface, which generally acts as a quenching center [16,18]. The previous considerations highlight the fact that the design of efficient electro-optic devices requires fine control of molecular interfaces, which can be ultimately achieved only through understanding of the electronic structure of the constituent molecules. In this context, ultraviolet photoelectron spectroscopy (UPS) represents a method of choice [5]. In short,

Figure 1 (a) Schematic diagram of the various processes taking place in a doublelayer organic light-emitting diode under forward bias; electrons and holes are initially injected at the metallic electrodes, then drift in the organic layer and get trapped at the organic interface where the radiative recombination takes place to produce light escaping from the diode by the ITO surface. (b) Corresponding diagram for a doublelayer photodiode operating in the short-circuit regime; the incident light entering in the device through the ITO surface generates electron-hole pairs, which then migrate to the organic interface where their dissociation occurs under the form of charges of opposite sign that are ultimately collected at the two electrodes.

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this experimental technique allows for the determination of the binding energy of the valence electrons by measuring the difference between the energy of monochromatic incident photons used to ionize the molecules and the kinetic energy of the ejected electrons. The use of synchroton irradiation and high-resolution spectrometry has opened the way for a detailed characterization of the density of valence electronic states, in particular the ␲ levels responsible for the unique electronic and optical properties of organic conjugated chains [19]. Thus, UPS spectra give access to the binding energy of the highest occupied levels of organic molecules and the way the relative positions of the electronic levels can be modulated by changes in the molecular topology, chain size, or upon attachment of electroactive substituents on conjugated backbones; this is illustrated in Section II of this chapter where we described a simple way to simulate UPS spectra of organic molecules with the help of semiempirical Hartree–Fock methods; we assess the reliability of this approach by comparing our theoretical results to available experimental UPS spectra. The study is further extended to the unoccupied levels of organic molecules in relation to experimental inverse photoemission (IUPS) spectra; in the latter case, the binding energy of the unoccupied levels is measured as the difference between the kinetic energy of the electrons in the incident beam and the energy of the released photons following population of the unoccupied levels [20]. The theoretical and experimental UPS and IUPS data obtained for single molecules can in turn be used to provide quantitative estimates of energy barriers expected at molecular interfaces in electro-optic devices. However, we stress that the predicted offset of the energy levels might be modulated in devices by the appearance of a dipole layer at the interface between the two organic phases, as evidenced by recent experimental data [21,22]. The energy level alignment at a molecular interface also plays a crucial role in the design of organic photodiodes and solar cells [23–26] where typically a homogeneous blend of two organic components or a bilayer structure is sandwiched between two metallic contacts. In such devices, the conversion of incident light into an electrical current results from three consecutive processes: (a) the macromolecules are promoted into an excited state by the incident light and decay rapidly by internal conversion processes to the lowest excited state; (b) the photogenerated excitons migrate to an interface between the two organic components where their dissociation into charge carriers occurs following a photoinduced charge transfer process; charges of opposite sign thus accumulate in the two components of the blend (Figure 1); (c) the charges follow pathways within a single organic phase to be collected at the metallic electrodes, and their amount accounts for the efficiency of charge carrier generation. However, other mechanisms can take place at a molecular interface: the exciton located on a conjugated chain

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can, for instance, be transferred to the second material (i.e., in a process referred to as energy transfer) or remain on the first material to subsequently decay radiatively or nonradiatively. The relative stability of the lowest intramolecular excited state of each organic component with respect to the lowest charge transfer excited state between the two macromolecules thus rules the occurrence of charge vs. energy transfer processes at molecular interfaces; this will be exemplified in Section III of this chapter with the help of a very simple theoretical model that relies on the relative positions of the frontier levels of two interacting materials. Note that the present considerations are relevant not only for the dissociation of excitons into charge carriers in photodiodes or solar cells but also for the inverse process in light-emitting devices based on the recombination of charge carriers into intramolecular excited states. This chapter is constructed as follows: in Section II, we describe different theoretical approaches to simulate the UPS spectra of organic conjugated molecules and dwell on the strategy adopted in the present work. Section III describes the one-electron structure of important conjugated compounds and assesses the reliability of our theoretical approach by confronting UPS and IUPS simulations to corresponding experimental spectra. Sections IV and V illustrate, respectively, that energy barriers at molecular interfaces can be modulated by the introduction of electroactive substituents on a conjugated backbone and that the relative positions of the frontier levels of two materials can be exploited to predict the occurrence of charge vs. energy transfer processes at their interface. The last section is devoted to our conclusions and perspectives. The chemical structure of all the organic compounds involved in this chapter are sketched in Figure 2.

II.

THEORETICAL METHODOLOGY

A rigorous simulation of UPS spectra would give rise to a UPS feature for each molecular level at a binding energy corresponding to the difference between the total energy of the molecule in the neutral state and that in the singly ionized state obtained by extracting an electron from the considered level; we should also take into account the occurrence of shake-up phenomena for which the ionization process is accompanied by the photogeneration of an electron-hole pair. In principle, ab initio Hartree–Fock and post– Hartree–Fock approaches using large basis sets are best adapted for this task. However, even today, such calculations performed on large conjugated molecules represent a time-consuming enterprise, thus calling for simpler alternatives. One approach is to adopt Koopman’s approximation in conjunction with Hartree–Fock calculations to estimate the binding energy of

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Organic Interfaces in Electro-Optic Devices 45

Figure 2 Chemical structures of the various organic compounds under study: (a) oligomers of polyparaphenylenevinylene, PPV; (b) PPV oligomers substituted by tertiobutyl groups; (c) N,N ⬘-diphenyl-N,N ⬘-bis(1-naphthyl)-1,1⬘-biphenyl-4,4⬘-diamine, ␣-NPD; (d) tris(8-hydroxyquinoline)aluminum, Alq3; (e) 4,4⬘-bis(m-tolyphenylamino)biphenyl, TPD (R1, R2, R3 = H), and substituted derivatives; (f) poly(2,5-dimethyloctylsilylparaphenylenevinylene), DMOS-PPV; (g) poly(2-methoxy-5-(2⬘ethylhexyloxy)paraphenylenevinylene), MEH-PPV; and (h) poly(2,5,2⬘,5⬘-tetrahexyloxy-7,8⬘-dicyanodiparaphenylenevinylene), CN-PPV.

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a level as the inverse of the level energy calculated for the neutral molecule. However, this should be restricted to the low binding energy region of the spectra where compensation usually occurs between the electronic relaxation and electronic correlation effects that are both neglected in such calculations. Koopman’s approximation has been extensively used and tested in our earlier theoretical works based on the valence effective Hamiltonian (VEH) [27,28] and proved very useful in terms of giving a detailed characterization of the valence electronic structure of conjugated polymers [29–33]. This nonempirical technique has been developed to reproduce the valence oneelectron energies obtained at the ab initio double-␰ level but suffers from the fact that parameters are available only for a limited set of atoms (H, C, N, S, O, and Si). The CNDO/S (complete neglect of differential overlap technique with a spectroscopic parameterization) [34] was the first semiempirical Hartree–Fock method to be successfully applied to simulate and interpret the UPS spectra of organic conjugated molecules [35–37]; however, a major drawback of the CNDO Hamiltonian is its inability to distinguish singlet and triplet excited states of the molecules when addressing their optical properties. This motivated us to validate the use of the semiempirical Hartree–Fock INDO (intermediate neglect of differential overlap) Hamiltonian [38] to describe the UPS spectra of conjugated systems, as will be reviewed in the next section. This choice is mainly driven by the fact that the spectroscopic version of INDO has been parameterized by Zerner and coworkers [39] to reproduce the optical absorption spectrum of molecules in the neutral state when coupled to a single configuration interaction (SCI) scheme; the reliability of INDO/SCI approaches to describe the nature of the lowest singlet and triplet excited states of organic conjugated systems has been demonstrated in our recent work [40–42]. Furthermore, this method is currently available in a large number of quantum chemistry packages running on PC workstations. For the sake of completeness, we emphasize that recent theoretical studies have also addressed the applicability of the AM1 (Austin Model 1) [43], MNDO (modified neglect of differential overlap) [44], and PM3 (Parameterized Method 3) [45] techniques to simulate the UPS spectra of conjugated systems. Another sophisticated alternative to compute UPS spectra is to apply techniques related to density functional theory (DFT) [46]. In that case, the DFT eigenenergies cannot be readily used within Koopman’s approximation since the error compensation occurring with Hartree–Fock techniques is lost by the fact that the correlation effects are explicitly included while the electronic relaxation effects are neglected [47]. This problem can be overcome with the help of the generalized transition state (GTS) [48,49] method or derivations thereof. We also stress that the DFT methods have been primarily developed to characterize ground-state properties and have so far a limited

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access to the description of excited states, which has to be tackled so as to depict optical processes at molecular interfaces. However, recent developments in this field are very promising and might give access to a reliable description of excited states from DFT-based techniques in the near future [50,51]. The UPS spectra of the selected set of organic conjugated systems shown in the next section have been obtained within Koopman’s approximation on the basis of the one-electron levels calculated with the INDO Hamiltonian. We have applied systematic corrections to the bare results in order to match the theoretical simulations to the corresponding experimental spectra [43], as sketched in Figure 3. The spectra are initially obtained from a peak convolution based on Gaussian functions with a full width at half maximum (FWHM) chosen to match the experimental spectral resolution. The intensities of the peaks do not account for the photoionization cross sections; this is a source of discrepancy between the intensities of the experimental peaks and those observed in the simulations. We then compress the energy scale of the calculated spectrum (starting from the HOMO level) to compensate for the neglect of electronic relaxation and correlation effects. When using the INDO Hamiltonian, we make use of a compression factor of 1.3 to simulate the UPS spectra (if not stated otherwise), which is very similar to the values best adapted to deal with VEH data [29–33]. In contrast, the compression of the energy scale is found not to be required to simulate the inverse UPS spectra of conjugated systems with the help of the INDO Hamiltonian; even though this procedure cannot be easily understood and justified, it is consistent with the fact that the INDO parameterization systematically yields a much larger width for the band made of all the occupied valence levels than for the corresponding unoccupied band. We stress that the high-energy part of the spectra should be addressed in all cases with much caution due to the drastic approximations used in our approach. Finally, when the measurements are obtained for solid samples, an additional correction is to shift rigidly the whole energy scale to account for (a) the difference in the energy reference (vacuum level in gas phase vs. Fermi energy in the solid state) and (b) the stabilization of the photogenerated holes by the surrounding medium (i.e., solid-state polarization effects), which are neglected in our gas phase calculations. Note that the amplitude of the shift to be applied for a given compound depends on the choice of the semiempirical approach; this is illustrated, for instance, by the fact that the ionization potential (calculated as the inverse of the HOMO level energy) is typically increased by 1.0 eV when going from INDO to AM1.

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Figure 4 INDO-simulated UPS spectrum of the three-ring unsubstituted PPV oligomer (FWHM = 0.6 eV) together with the corresponding experimental solid-state spectrum. The energy scale is referred to the vacuum level. We also show the LCAO patterns of the molecular orbitals giving rise to the three UPS peaks with the lowest binding energies.

grows in intensity as the chain size is increased and always appears at almost the same energy; this is consistent with the localized nature of the molecular orbitals contributing to the peak intensity, which have no significant electronic coupling between the monomer units and hence make the energy of these levels insensitive to the extent of the conjugated backbone. In contrast, the number and positions of the peaks at lower binding energy vary with chain size; this points to their delocalized character, which is affected by the extent of conjugation. We also observe that the energy of the peak associated to the HOMO (highest occupied molecular orbital) level is progressively shifted to lower binding energy as the chain grows; this evolution is consistent with experimental data showing a linear relationship between the

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Figure 5 INDO-simulated UPS spectra of unsubstituted PPV oligomers containing three to seven phenylene rings (FWHM = 0.4 eV in all cases). The inset shows the experimental solid-state UPS spectra obtained for the same oligomers substituted by two tertiobutyl groups on the external rings, as extracted from Ref. 59. The energy scale is referenced to the vacuum level.

ionization potential and inverse chain length in conjugated oligomers of increasing size [59–61]. In spite of the overall good agreement between the global shape of the experimental spectra and the corresponding simulations, we observe significant deviations between the lineshapes in the energy range between ⫺7.0 and ⫺7.5 eV. Since the experimental data are obtained from tertiobutylsubstituted derivatives, we have investigated the possible influence of these groups on the relative positions of the highest occupied levels, as illustrated in Figure 6 for the three-ring oligomer. The main changes in the shape of the UPS spectrum upon substitution result from a lift of degeneracy of the three localized levels; the two localized orbitals having significant weight on the external phenylene rings are indeed destabilized upon substitution while the third level localized on the central ring is weakly affected. This

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Figure 6 INDO-simulated UPS spectrum (FWHM = 0.6 eV) of the unsubstituted three-ring PPV oligomer (bottom) compared to that obtained for the same oligomer substituted by two tertiobutyl groups on the terminal rings (top). We also indicate the localized (l) vs. delocalized (d) nature of the levels contributing to the UPS features at the lowest binding energies (H refers to the HOMO level). The energy scale is referred to the vacuum level.

rationalizes the significant increase in intensity around ⫺7.0 eV when going from the unsubstituted to the substituted oligomer; this now leads to a remarkable agreement between the theoretical and experimental data. In the higher binding energy range, the substitution also gives rise to energy shifts of the ␲-delocalized levels having a significant electron density on the substituted rings. These results emphasize that an explicit account of the full chemical structure is required to provide the best match between experimental and calculated UPS spectra. We report in Figures 7 and 8 the solid-state experimental UPS and IUPS spectra of the ␣-NPD [N,N⬘-diphenyl-N,N⬘-bis(1-napthyl)-1,1⬘ biphenyl-4,4⬘ diamine] and Alq3 [tris(8-hydroxyquinoline)aluminum] molecules, respectively, together with the corresponding theoretical simulations;

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Figure 7 Comparison between the UPS and IUPS spectra measured from ␣-NPD thin film (top) and the corresponding INDO-simulated spectra (bottom). The energy scale is referenced to the Fermi level and the FWHM is set to 0.5 eV prior to compression. Vertical bars at the bottom of the graph refer to the calculated energies of the molecular orbitals.

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Figure 8 Comparison between UPS and IUPS spectra measured from an Alq3 thin film (top) and the corresponding INDO-simulated spectra (bottom). The energy scale is referenced to the Fermi level. A compression factor of 1.2 is here used to simulate the UPS spectrum and the FWHM is set to 0.5 eV prior to compression. Vertical bars at the bottom of the graph refer to the calculated energies of the molecular orbitals.

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separation between the HOMO and LUMO levels of the two isolated constituents; these values are, however, directly comparable to experimental data only in situations where the chemical interactions at the interface can be neglected. If chemical interactions do occur, the theoretical results can then be exploited to detect and quantify energy level shifts linked to the appearance of dipole layers at molecular interfaces; the origin of such interfacial effects are far from being understood, which points to the need for additional theoretical and experimental investigation [21]. In addition to these aspects, it is also highly desirable to find strategies to fine-tune energy barriers in order to enhance device performances. The most frequently proposed solution is to rely on experimental determinations of the HOMO and LUMO energies of a large number of synthesized compounds with varying molecular topology to drive the search toward the best matching partners. A different alternative is to graft electroactive substituents on a given conjugated backbone in order to modulate the positions of the frontier levels over a large range of energies, as recently exemplified for TPD [74]. In a previous theoretical study performed on PPV oligomers, we showed that the introduction of ␲-donor and ␲-acceptor substituents can lead to substantial shifts in the energy of the HOMO and LUMO levels, which can be modulated by the nature, number, and relative positions of the attached groups [75]; this was further exemplified by similar calculations achieved on other PPV derivatives [76]. In all cases, the stabilization vs. destabilization of the frontier levels results from a subtle interplay between inductive effects in the ␴-skeleton and ␲-mesomer effects. Unfortunately, gas phase calculations cannot be readily compared to solid-state UPS or electrochemical data due to the fact that the conformational behavior of the molecules and the solid-state polarization effects are generally not well defined. However, this major drawback can be overcome by confronting the theoretical results to corresponding gas phase experimental measurements when available. This is illustrated hereafter through a joint experimental and theoretical analysis of the electronic properties of a series of hole-transporting materials based on the TPD molecule [64]; these are obtained by substituting the TPD backbone by methoxy groups or fluorine atoms, as shown in Figure 2. Here we first characterized the geometrical and electronic structure of TPD and its constituent moieties, triphenylamine (TPA) and biphenyl, at the DFT level, using the B3LYP functional by Becke [65], Lee et al. [66], and the 6-31G** basis set. Since TPD is widely used as hole-transporting material in multilayer LEDs, we have also compared the geometry of the neutral molecule to that of the radical cation so as to account for the geometrical relaxation taking place when a hole is injected into the system. This should contribute to future developments aimed at clarifying the relationships be-

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tween the geometrical and electronic structure and the transport properties of these materials [67]. The main results of the geometry optimizations are presented in Figure 9, where the geometries of the neutral and cation systems are also compared. For biphenyl, the calculations in the neutral state give an interring bond ˚ ; the torsion angle between the two aromatic rings is estilength of 1.49 A mated to be 38.4⬚, in agreement with X-ray diffraction experiments on gaseous biphenyl [77]. The geometry of the radical cation is shifted toward planarity; the value of the torsion angle is reduced to 19.3⬚ and the bond length alternation acquires a quinoid-like character (the bond connecting the ˚ and the bond length alternation in the rings is rings is reduced to 1.44 A increased with respect to that in the neutral molecule). For triphenylamine, the optimized geometry of the neutral molecule yields an sp2 hybridization (planar configuration) for the nitrogen atom, in agreement with earlier theoretical works [78]; the three phenyl groups assume a propeller-like structure, with a torsion angle of 41.7⬚ with respect to the plane defined by the three C-N bonds. The geometrical structure of the radical cation differs only slightly from that of the neutral one: the amino group remains planar, the torsion angles of the phenyl groups decrease to ˚ in the neutral molecule 38.9⬚, and the C-N bond lengths change from 1.42 A ˚ to 1.41 A in the cation. In the case of TPD, the neutral molecule has a geometry very similar to that of each of its separated components: the central part has a structure nearly coincident with the one found for biphenyl, with an interring bond ˚ and a twist angle of 33.8⬚, both very close to the values obtained of 1.48 A for the isolated biphenyl molecule. The same applies to the triphenylamine moiety of neutral TPD, for which a planar configuration of the C-N bonds ˚ and torsion angles of 42.3⬚ and is found, with C-N bond lengths of 1.42 A 40.6⬚ for the terminal tolyl and phenyl units, respectively, and 42.0⬚ for the central ring. The picture changes markedly when going from neutral TPD to the positive ion; in this case, while the similarity between the central part of the molecule and biphenyl is maintained, the structure of the triphenylamine moieties of TPD⫹• looks quite different from that of the triphenylamine cation. The group of C-N bonds remains planar, but TPD⫹• shows a marked asymmetry in the C-N bond lengths, the C-N bond with the central biphenyl ˚ vs. 1.43 A ˚ ). The part being significantly shorter than the other two (1.39 A same asymmetry is encountered for the torsion angles between the phenyl groups and the plane of C-N bonds, with the central phenyl being shifted toward planarity (25.8⬚) and the terminal rings toward orthogonality (49.3⬚ and 47.6⬚).

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Figure 9 DFT-optimized geometrical parameters for the neutral and cation struc˚ ) are written in plain tures of biphenyl, triphenylamine, and TPD. Bond lengths (in A characters; bond angles and dihedral angles (in degrees) are in italics.

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Figure 9

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Continued

These results indicate that the TPD molecule, when changing from the neutral to the radical cation state, undergoes a geometry relaxation similar to the one occurring in biphenyl, and sensibly different from what can be expected from triphenylamine alone. The analysis of the frontier orbitals is consistent with this picture, showing that the HOMO and LUMO levels of

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TPD are mainly localized over the central biphenyl core and the two nitrogen atoms. It is worth stressing that a similar description of the geometrical structure of TPD in its neutral state is provided by calculations performed at the AM1 level; accordingly, this approach has been adopted to optimize the full ground-state geometry of the other TPD derivatives. In all cases, the results point to highly nonplanar conformations characterized by (a) an almost pure sp2 character for the nitrogen atoms; (b) a torsion angle around 40⬚ between the two rings of the central biphenyl core; and (c) a propeller shape for the three rings connected to a nitrogen atom, with torsion angles (with respect to the plane defined by the nitrogen atom and the three attached carbon atoms) ranging from 30⬚ to 40⬚ as a function of the substitution pattern. On the basis of the structural data, we have estimated at the INDO level the evolution in the energy of the frontier levels when going from TPD to the substituted derivatives. The experimental gas phase UPS spectra of all the derivatives are characterized in the low binding energy range by a doublet feature assigned to the HOMO and HOMO-1 levels [64]; this is illustrated in Figure 10 where we report the gas phase spectrum of TPD together with the corresponding INDO simulation, which again reveals an excellent agreement between the theoretical and experimental lineshapes. We collect in Table 1 the calculated energy of the HOMO and HOMO-1 levels together with the binding energy of the corresponding peaks in the gas phase UPS spectra; this points to a remarkable parallelism between the evolution of the experimental and theoretical data. It is interesting to note that a large number of conformers having very similar electronic properties are expected for all of the derivatives in the gas phase. The HOMO and HOMO-1 levels of the TPD derivatives are characterized by very small LCAO coefficients on the ␲-electroactive substituents (i.e., on the methoxy groups or the fluorine atoms), indicating that the shifts of the highest two occupied levels are driven by purely inductive effects. This is supported by the results in Figure 11 showing the evolution of the HOMO energy in the various compounds as a function of the total charge (calculated at the AM1 level by means of a Mulliken population analysis) on the biphenyl core and the nitrogen atoms where the wavefunction is mostly confined. This also demonstrates that the amount of charge transferred from the external substituted rings to the central biphenyl core is the dominant parameter controlling the energy of the HOMO level. A very similar evolution is calculated for the energy of the LUMO level (see Table 1), as expected from its localization on the central biphenyl core. The changes in the energy of the HOMO-1 level are also governed by inductive effects; the evolution results in that case from a subtle interplay between the localization of the molecular orbital and the total charge distribution over the molecule. We emphasize that the calculated separation between the highest

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Figure 10 Comparison between the gas phase UPS spectrum of the TPD molecule and the corresponding INDO simulation (FWHM = 0.3 eV). The energy scale is referenced to the vacuum level.

two occupied levels is overestimated with respect to the measured energy spacing between the corresponding UPS peaks; this is fully consistent with the fact that the INDO simulations of UPS spectra require a compression of the energy scale, as described in the previous section. It is also of interest to note that the INDO/SCI-calculated lowest optically allowed transition energy of TPD is weakly affected by the various substitution patterns [64], in spite of the significant modulation of the energy of the frontier levels. This can be understood by the fact that this transition is mostly described by an one-electron excitation between the HOMO and LUMO level and thus is not affected by the rigid shift of these two levels upon substitution. This does not hold true when ␲-mesomer effects start playing a role, leading in most cases to an asymmetrical stabilization or destabilization of the HOMO and LUMO levels, and hence to a shift of the lowest optical transition [75,76].

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Figure 11 AM1-calculated evolution of the HOMO energy of TPD and derivatives as a function of the net charge on the biphenyl core and nitrogen atoms; the labels I–VII used for the molecules are defined in Table 1.

V.

CHARGE VS. ENERGY TRANSFER AT MOLECULAR INTERFACES

When two chains of different nature interact, it is of prime importance to determine whether the lowest excited state has a dominant intrachain or charge transfer character in order to understand the nature of the electronic processes taking place at the molecular interface. The intrachain character is highly desirable in blends to be used in light-emitting devices, whereas the charge transfer character is required for the operation of solar cells and photodiodes. These two competing processes are nicely illustrated by a recent joint experimental and theoretical work investigating the properties of homogeneous blends formed, on the one hand, by MEH-PPV [poly(2methoxy-5-(2⬘-ethylhexyloxy)paraphenylenevinylene] and CN-PPV [poly(2,5,2⬘,5⬘-tetrahexyloxy)-7,8⬘-dicyanodiparaphenylenevinylene] and, on the other hand, by CN-PPV and DMOS-PPV [poly(2,5-dimethyloctylsilylparaphenylenevinylene] [79]. The chemical structures of the three polymers are sketched in Figure 2. We have modeled these chains by first optimizing the

64 Table 1 Energy (eV) of the HOMO (H) and HOMO-1 (H-1) levels in TPD and substituted derivatives, as extracted from gas phase UPS spectra and calculated at the INDO levela UPS Molecule I II III IV V VI VII

— R2 — R1 —H; —H; R3 — —H R1 — —F; R2 — —H; R3 — —H R1 — —F; R2 — —H; R3 — —F R1 — R —F; 2 — —F; R3 — —F R1 — —H; R2 — —F; R3 — —H R1 — —OCH3; R2 — —H; R3 — —H R1 — —H; R2 — —OCH3; R3 — —H

INDO

H

H-1

H

H-1

L

⫺6.69 ⫺6.89 (⫺0.20) ⫺7.08 (⫺0.39) ⫺7.14 (⫺0.45) ⫺6.74 (⫺0.05) ⫺6.75 (⫺0.06) ⫺6.59 (⫹0.10)

⫺6.98 ⫺7.18 (⫺0.20) ⫺7.39 (⫺0.41) ⫺7.44 (⫺0.46) ⫺7.10 (⫺0.12) ⫺7.04 (⫺0.06) ⫺6.95 (⫹0.03)

⫺6.74 ⫺6.92 (⫺0.18) ⫺7.08 (⫺0.34) ⫺7.24 (⫺0.50) ⫺6.89 (⫺0.15) ⫺6.77 (⫺0.03) ⫺6.68 (⫹0.06)

⫺7.13 ⫺7.32 (⫺0.19) ⫺7.49 (⫺0.36) ⫺7.64 (⫺0.51) ⫺7.29 (⫺0.16) ⫺7.17 (⫺0.04) ⫺7.05 (⫹0.08)

0.21 0.04 (⫺0.17) ⫺0.08 (⫺0.29) ⫺0.29 (⫺0.50) 0.04 (⫺0.17) 0.18 (⫺0.03) 0.21 (0.00)

a

In each case, we report in parentheses the energy shifts calculated with respect to the TPD molecule. We also collect for each compound the calculated energy of the LUMO level.

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would suggest that the excitation of DMOS-PPV or MEH-PPV in the blends will give rise in both cases to a photoinduced electron transfer toward CNPPV, leaving behind the hole, as a result of the large energy difference between the LUMO levels of the two chains. However, this is contradicted by photovoltaic and photoluminescence measurements, giving firm evidence that the mixing of MEH-PPV and CN-PPV leads to photoinduced charge transfer processes at the interface between the two polymers whereas blends of DMOS-PPV and CN-PPV show energy transfer toward the CN-PPV chains. We show below that these results can be interpreted with the help of quantum chemical calculations from estimates of the relative energies of the lowest intrachain vs. interchain excited states in the blends. In order to provide a reliable description of energy vs. charge transfer at the polymer–polymer interface, the amount of energy required to dissociate an intrachain exciton into a charge transfer exciton is a very important parameter to assess. Correlated calculations performed on cofacial dimers formed by two five-ring PPV oligomers provide a reasonable estimate of this parameter on the order of 0.75 eV [79]. The results also indicate that this value is weakly affected by changes in the interchain separation, chain size, and number of chains in interaction. A second factor to be accounted for is the influence of the polarization effects induced by the surrounding medium [80]. These effects are expected to give rise to a significant stabilization of the lowest charge transfer excited state with respect to the lowest intramolecular excited states. Unfortunately, the amplitude of the solid-state polarization effects cannot be readily quantified from quantum chemical calculations but can be inferred indirectly from the present experimental measurements, as discussed below. For the sake of illustration, we have adopted in the following a value of 0.4 eV, which makes the penalty to go from an intrachain exciton to an interchain exciton on the order of 0.75 eV ⫺ 0.40 eV = 0.35 eV. We stress that this value is reasonable because it typically corresponds to estimates of the exciton binding energy in PPV chains and substituted derivatives [81,82]. This also implies that it does not cost that much energy to transform a charge transfer exciton, where two polarons of opposite sign located on adjacent chains are bound by Coulomb attraction, into free carriers; this is consistent with the high efficiency of short-circuit charge carrier generation in photodiodes and can be understood by the fact that the energy required to separate the charges is partly compensated by the solid-state polarization effects. Thus we have all the required parameters to determine whether the lowest excitation is localized on one component or corresponds to charge transfer from one component to the other in the DMOS-PPV/CN-PPV and MEH-PPV/CN-PPV blends. Figure 12 provides the estimated ordering of the lowest excited states in the two blends from a one-electron picture; this is relevant here since the

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lowest intrachain transition of each chain is dominantly described by an excitation between the HOMO and LUMO levels [75], whereas the lowest interchain transition is essentially depicted by an excitation between the highest occupied level lying on one chain and the lowest unoccupied level located on the other chain. In going from PPV to CN-PPV, we calculate the lowest intramolecular excited state to be red-shifted by 0.38 eV due to the asymmetrical stabilization of the HOMO and LUMO levels. The stabilization of the LUMO level by 0.55 eV induced by the cyano groups significantly lowers the energy of the lowest charge transfer excited state; however, this stabilization is partly offset by the 0.35 eV required to transform the intrachain exciton into an interchain exciton. Therefore, we estimate the lowest charge transfer excited state in the PPV/CN-PPV blend to be located 0.20 eV below the lowest excited state of PPV, and hence 0.18 eV above the lowest intrachain transition of CN-PPV. Thus we expect an energy transfer to the CN-PPV chains to take place in the DMOS-PPV/CN-PPV blend; this is clearly supported by the experimental data [79]. Since a polarization energy on the order of or larger than 0.6 eV is required to make the charge transfer excited state the lowest in energy in this blend, we also conclude on the basis of this model that the amplitude of the polarization effects is lower than 0.6 eV when going from an intrachain exciton to a charge transfer exciton. In the case of the CN-PPV/MEH-PPV pair, analysis of the one-electron structure reveals that the lowest optical transition of CN-PPV is red-shifted by only 0.19 eV with respect to that of MEH-PPV. The LUMO level of CNPPV is 0.63 eV lower than that of MEH-PPV due to the strong acceptor character of the cyano groups; however, stabilization of the lowest charge transfer excited state by 0.63 eV with respect to the intrachain transition of MEH-PPV has to be compensated by the energy required to separate the electron and the hole. As a result, the lowest charge transfer excited state in the MEH-PPV/CN-PPV blend is estimated to be about 0.28 eV below the lowest intrachain transition of MEH-PPV and, hence, some 0.10 eV below the lowest excited state of CN-PPV. Note that we are led to the same conclusion when considering that the CN-PPV chain is initially photoexcited and that the charge transfer state is reached upon photoinduced hole transfer. In this blend, we thus predict the occurrence of a charge transfer process at the polymer–polymer interface, in full agreement with the experimental observations [79]. It is worth stressing that the charge transfer excited state could not be the lowest in energy in the MEH-PPV/CN-PPV pair without an explicit account of the increased polarization energy in going from intrachain to interchain exciton; in the framework of our model, our calculations indicate that a choice of amplitude for the polarization energies between 0.3 and 0.6 eV brings agreement with the experimental data.

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In this section we have presented a simple model, and one that awaits further experimental and theoretical developments, to show that the use of conjugated backbones with different substitution patterns can make blends acting as materials of choice for use in photodiodes where charge transfer processes need to prevail or in light-emitting devices where energy transfer processes must take place. The calculations have also underlined the key role played by solid-state polarization effects to predict the correct ordering of the lowest intramolecular vs. charge transfer excited states at molecular interfaces. When charge transfer takes place, the energy separation between the initially photoexcited state and the final charge transfer state also plays a crucial role in determining the rate of the photoinduced charge transfer process, as exemplified in the framework of Marcus’s theory [83,84].

VI.

CONCLUSIONS AND PERSPECTIVES

We have illustrated throughout this chapter that simple theoretical approaches can prove very useful in our effort to understand the electronic properties of organic conjugated compounds in relation to their use in electro-optic devices. We have shown in partcular that the characterization of the relative positions and nature of the frontier levels of individual molecules allows us to shed light into the nature of the processes taking place at molecular interfaces in such devices. Further theoretical developments will require an explicit consideration of the possible chemical interactions taking place at the interface between two materials in order to understand and control such interfacial effects in the search for devices with enhanced performances. It goes without saying that this will be ultimately achieved only through joint theoretical and experimental efforts, as evidenced by the strong complementarity between the theoretical and experimental results presented here.

ACKNOWLEDGMENTS We are grateful to N. R. Armstrong, R. H. Friend, J. J. M. Halls, I. G. Hill, A. B. Holmes, D. H. Hwang, A. Kahn, P. A. Lee, D. L. Lichtenberger, M. Lo¨gdlund, W. R. Salaneck, R. Silbey, S. Vanderdonckt, and L. M. Yu for stimulating discussions and the work done in collaboration. We are especially indebted to N. E. Gruhn for the communication of gas phase UPS data prior to publication. This work has been partly supported by the Belgian Federal Government (Poˆle d’Attraction Interuniversiatire en Chimie Supramole´culaire et Catalyse, PAI 4/11); the European Commission under Brite-

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Euram Contract No. BRPR-CT97-0469 (OSCA); the Belgian National Fund for Scientific Research (FNRS); ONR in the framework of the MURI Center on Advanced Multifunctional Polymers at the University of Arizona; and the Francqui Foundation, Belgium, through award of a postdoctoral fellowship to M.M. J. C. and R. L. are Chercheur Qualifie’ and Maıˆtre de Recherches of FNRS.

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4 Photoelectron Spectroscopy of Interfaces for Polymer-Based Electronic Devices Michael Lo¨gdlund and Thomas Kugler Acreo AB, Norrko¨ping, Sweden

Grzegorz Greczynski, Annica Crispin, and William R. Salaneck Linko¨ping University, Linko¨ping, Sweden

Mats Fahlman Linko¨ping University, Norrko¨ping, Sweden

I.

INTRODUCTION

In electronic devices based on organic materials, and for light-emitting diodes (LEDs) in particular, the characteristics of carrier injection from a metal contact to a polymer layer is a crucial issue in determining performance. The details of energy level alignment at interfaces between organic layers and metal electrodes, a fundamental issue for designing better devices, is still the subject of controversy [1–4]. Most detailed spectroscopic work in this area has been performed on the interfaces prepared under the ultrahighvacuum conditions (UHV), by means of the vapor deposition of an organic overlayer onto ultraclean metallic substrates, thereby minimizing potential contamination. As has been shown for these ‘‘ideal’’ contacts, the systems studied are not necessarily in the Schottky–Mott limit [5], and alignment of the vacuum levels of the organic film on the metal substrate does not occur in most cases. The reported results of recent studies using photoelectron spectroscopy [1,2,4,5] emphasize the importance of the influence of several factors that may lead to the formation of interfacial dipole layers at inter73

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faces. These factors include charge transfer through the interface, interfacial electronic states, chemical interactions, and image forces, among others. The nature of energy band edge alignment at polymer-on-metal interfaces prepared under the ambient (air) conditions, often by spin coating techniques, are very different from the ‘‘ideal’’ interfaces formed by vapordeposition of organic molecules on ultra-clean metal surfaces in UHV. The nature of the (spin coating) preparation process itself (presence of, for example, solvents) leads to an unavoidable passivation of the polymer-substrate interface. Although this makes the situation rather complex from a point of view of fundamental studies, the information obtained on such reallife interfaces is directly applicable to the analysis of real device performance. In this chapter, first some fundamental principles of photoelectron spectroscopy are reviewed, as introduction to the technique used in many of the chapters to follow. Then a portion of the interface work carried out in Linko¨ping over the past several years is reviewed. In this review, some examples of organic polymer interface formations studied by photoelectron spectroscopy are discussed. The emphasis is on band-level alignment between organic materials and electrode materials. The outline of the chapter is as follows: In section II, the experimental photoelectron spectroscopy is described; in section III, band level alignment is discussed. In sections IV and V some examples of conjugated polymers interacting with ITO and other surfaces in the polymer-on-metal situation are given. Finally, in sections VI and VII, the interaction between poly(9,9-dioctylfluorene) and various overlayers, such as Li, LiF, Al, and Na, are discussed. The examples are chosen to illustrate various interface aspects such as the effect of substrate work– function variation, ideal vs. ‘‘real’’ surfaces, and ‘‘buried interfaces.’’

II.

PHOTOELECTRON SPECTROSCOPY

Because photoelectron spectroscopy (PES) is an important tool in studies of polymer surfaces and interfaces, especially involving conjugated polymers [6], it is appropriate to include an introduction to this method at the beginning of this chapter. The use of soft X-ray sources led to the development of X-ray photoelectron spectroscopy (XPS), originally known as electron spectroscopy for chemical application (ESCA) [7], indicating the applicability of the method to surface chemical analysis. More or less simultaneous with the development of XPS, ultraviolet photoelectron spectroscopy (UPS) [8], i.e., the use of ultraviolet photon sources in PES, was applied to studies of gases and bulk as well as surface electronic structures of solids. However, the

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There are several possible single-electron molecular configurations, as shown schematically in Figure 1. In the neutral molecule in the ground state, represented as M0, all of the electrons in the molecule are occupying only the lowest allowed energy levels (Vi), while the V* i levels are empty. The other panels illustrate the generalized excited state of the (photoionized) molecular cation, M * ⫹, where an electron has been removed from the core level, Ci, but the remaining electrons do not occupy the lowest possible single-electron states (the special case of ionization from the HOMO level is illustrated as M⫹); and the optically excited neutral molecules, M * 0 (where one electron has been excited from the HOMO to the LUMO in the illus-

Figure 1 Schematic illustrations of some single-electron molecular configurations are displayed. The different configurations are discussed in the text.

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During the photoelectron emission event there are electronic relaxation effects occurring, which usually are divided into intra- and intermolecular relaxation effects. These effects can be rationalized, in a classical picture, as follows. An electron in a photoelectron emission event leaves a molecule typically within about 10⫺15 s. The nuclear geometrical relaxation time is around 10⫺13 s (i.e., which corresponds to an optical phonon frequency) while the corresponding electronic relaxation time is about 10⫺16 s [17]. Thus, the main line corresponds to the binding energy of an electron in a ionized molecule where the electrons have had time to relax, i.e., the hole is fully screened, but the nuclei are frozen during the process; this is referred to as the adiabatic peak. The intramolecular electronic relaxation effects occur in response to the creation of a hole state in PES of an isolated molecule in the gas phase. In the solid state, there are also intermolecular relaxation effects occurring due to electronic and atomic polarization of the molecules surrounding the particular molecule on which the hole state is created. The intermolecular relaxation energy is given to the escaping electron and thus the kinetic energy is increased, i.e., the corresponding peak in the photoelectron spectra appears at lower binding energy compared to photoelectron emission from the same level in the gas phase. This energy difference, i.e., the polarization energy, is in the order of 1–3 eV [18] throughout the valence region. However, the relaxation energy for core levels, can be much larger due to a higher degree of localization of the core holes [16,17]. A.

X-Ray Photoelectron Spectroscopy

With XPS, as commonly carried out with Al(k␣) and Mg(k␣) laboratory sources, it is possible to perform qualitative and even semiquantitative analysis of chemical composition in the near-surface region of a solid sample. Changes in the valence electron density will be reflected as small but significant shifts in the core level binding energies, so-called chemical shifts, although the core electrons are not involved in the chemical bonds [7]. Of course, XPS can be (and has been extensively) used for studies of the valence band region, but the photoionization cross section for the valence photoelectrons in XPS are approximately one order of magnitude lower than that for the core photoelectrons, which leads to more time-consuming experiments. Also, the photoionization cross sections are such that it is not very convenient to study conjugated polymers, where interest is in the states near the valence band edge, which are very weak in XPS valence band spectra. Another feature in PES spectra is the so-called shake-up structures, appearing as weak satellites on the low binding energy side of the main line.

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The shake-up structure reflects the spectrum of the one-electron two-hole states generated in connection with photoionization, and can give useful information about the valence ␲-electronic structure of a molecular ion. The PES measurements are performed, with reference to the Fermi level of the photoelectron spectrometer, in solid specimens, as dealt with here, by how the spectroscopy works. Thus, in cases when the Fermi level shifts due to some chemical modifications of the sample, e.g., in the intercalation of graphite or other layered compounds [19] or in the doping of conjugated polymers [20], it is necessary to account for the change in the Fermi energy level before interpreting spectra. This can be done, for example, by combining the XPS core level spectra with UPS valence band spectra, as described in more detail in the next section, or by referencing to model compounds [16].

B.

Ultraviolet Photoelectron Spectroscopy

The most commonly used laboratory photon source for UPS is the helium resonance lamp, although the use of synchrotron radiation as a photon source is in extensive use today. For valence band spectroscopy, UPS has two advantages over XPS. First, the usual in-house photon sources have highenergy resolution, a full width at half maximum (FWHM) of about 30 meV for the He lines, and second, the photoionization cross section for electrons is higher in the valence region for UPS. However, the natural linewidths in condensed molecular solids, from both homogeneous and inhomogeneous broadening effects, can approach 1 eV at room temperature [12,21,22], and a large fraction of the photon energy resolution advantage is not fully realized. As mentioned above, PES measurements of solid specimens are performed relative to the Fermi level. However, in some cases, it can be desirable to derive the binding energies relative to the vacuum level from a photoelectron spectrum. This is particularly important before interpreting different core level chemical shifts induced upon, for example, doping or interaction with metals. In fact, in some cases, the apparent binding energy shift can appear to be in the wrong direction before the Fermi energy shift is taken into account [20,23]. The Fermi energy shift can be deduced by measuring the position of the zero kinetic energy cutoff of the secondary electron distribution, as illustrated in Figure 3. The vacuum level can be located simply by adding the photon energy to the secondary electron cutoff energy; measuring down from the vacuum level energy to the known position of the Fermi energy (EF) determines the work function (␾s) of the substrate [24].

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Figure 3 A UPS spectrum of Au illustrating the procedure to reference spectra to the vacuum level.

III.

BAND LEVEL ALIGNMENT BETWEEN CONJUGATED SYSTEMS AND INDIUM-TIN-OXIDE

Although indium-tin-oxide (ITO) is not a well-defined material, it is by far the most common transparent electrode used in the area of polymer LEDs. Also, the work function of ITO can be varied by different surface treatments. In this section, the interaction between bis-DMOS-PPV and PEDOT-PSS and ITO surfaces with different work functions is discussed. The chemical structure of bis-DMOS-PPV and PEDOT-PSS are displayed in Figure 4. The ITO surfaces with different work functions were prepared using several chemical techniques: (a) washing with acetone, followed by isopropanol, 5 min each, in an ultrasonic cleaner; (b) treatment with a diluted aqua regia solution after (a); or (c) treatment with a solution of 30 wt% H2O2, 25 wt% NH4OH, and deionized water at 80⬚C for 20 min, followed by rinsing with deionized water [25] after (a). The work functions of these surfaces were reproducible, with the following values: (a) 4.4, (b) 4.6, and (c) 4.8 eV (⫾0.1 eV). Also, the effect of an ultrathin interfacial layer, consisting of either poly(4-styrenesulfonate), denoted PSSH, or of an electrically conducting polymer blend, namely poly(3,4-ethylenedioxythiophene) doped with poly(4-styrenesulfonate), denoted PEDOT-PSS [26], have been studied. Thin polymer films were prepared by spin-coating onto freshly prepared ITO sur˚ , as estimated from XPS faces. The resulting thickness was about 100 A

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Figure 4 The chemical structure of the polymer materials discussed in the text are shown: poly[bis-(2-dimethyloctylsilyl)-1,4-phenylenevinylene], or bis-DMOSPPV; poly(3,4-ethylenedioxythiophene), or PEDOT; poly(styrenesulfonic acid), or PSSH.

measurements. Films of bis-DMOS-PPV on PEDOT-PSS-on-ITO, and bisDMOS-PPV on PSSH-on-ITO were prepared successively. When ITO substrates are irradiated with unfiltered Mg(k␣) radiation, a reproducible change in work function is observed [27], presumably as a result of X-ray-induced changes in the oxygen defect distribution, which is the basis for the electrical conductivity in ITO. Therefore, by using prolonged exposure to Mg(k␣) radiation, the work function of the ITO substrates could be modified in situ, with a polymer overlayer in place, and the corresponding changes in the UPS observed, as discussed below. In Figure 5 is shown the energy level scheme and the corresponding experimental UPS intensity curves for bis-DMOS-PPV, coated on a gold substrate, which establishes the parameters to be discussed below. The nomenclature and structure of the diagrams are now standard and follow that of Seki et al. and Kahn et al. [2,4,5,28–30]. It is worth pointing out the following: First, the difference in kinetic energy between the fastest photoelectrons from the gold substrate and from the polymer overlayer corresponds to the offset between the valence band edge in the polymer overlayer and the Fermi energy in the conducting gold substrate, E VB F . It determines the barrier for hole injection from the substrate into the valence band of the polymer. Second, the energy difference between the vacuum level of the polymer overlayer and the Fermi energy in the conducting substrate, E VAC , F max VB is given by E VAC = h ␯ ⫺ (BDMOS ⫺ PPV) ⫹ When the width (E E ). F k F of the band-bending region at the substrate–polymer interface is well below the polymer film thickness, E VAC , as probed by UPS, is equal to the work F function of the polymer (which is a material constant). However, this requires a high concentration of free charge carriers, which is only the case

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Figure 5 The band parameters that may be derived from the UPS spectra, as presented in previously reported studies [Ishii, 1997 #34; Hill, 1998 #1], are diagrammed. The UPS spectra shown are for a gold substrate cleaned in UV-generated ozone for 5 min followed by heating to 120⬚C in UHV for 2 h (left) and bis-DMOSPPV on an identical substrate (right).

in electrically conducting polymers (due to doping). Third, the offset between the vacuum levels of the polymer coated and the uncoated gold substrate is ⌬ = E VAC ⫺ ⌽Au (where in Figure 5, ⌽Au is the work function of F the gold substrate). Fourth, the ionization potential IP for the polymer, as probed by UPS, is IP = E VB ⫹ E VAC [6]. It is a specific property of the F F polymer and therefore independent of the energy level alignment at the interface. In UPS measurements it is important to clearly define the significance of the term ‘‘vacuum level’’: The vacuum level of a metal substrate as

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measured by UPS [denoted as E VAC(Au) in Figure 5] corresponds to the energy of an electron at rest just outside the sample [3] and is sometimes referred to in literature as ‘‘surface vacuum level’’ in contrast to vacuum level defined for the electron at rest in the infinite distance from the sample. In Figure 6 are a series of values of E VAC , ⌬, E FVB, and IP displayed. The F ITO work function ranges from 4.4 eV [method (a)] to 4.8 eV [method (c)] [25]. Even higher values were obtained for ultrathin layers of PEDOT-PSS and PSSH on ITO: 5.0 and 5.2 eV. For spin-coated bis-DMOS-PPV, the position of the vacuum level, E VAC , follows the work function of the substrate, ranging from 4.4 eV [for F ITO (a)] to 5.0 eV (for PSSH-coated ITO). The offsets, ⌬, for bis-DMOSPPV are very small (⫺0.2 eV average) and independent of the substrate work function. An important consequence of this vacuum level alignment is that an increase of the substrate work function results in a corresponding decrease

Figure 6 The band parameters E VAC , ⌬, E VB F F , and IP (see text) of bis-DMOS-PPV are shown as a function of the substrate work function. The substrates are as follows: full circles, ITO after surface preparation type (a); full triangles, ITO following treatment of type (b); full squares, ITO following treatment of type (c); open circles, PEDOT-PSS on ITO; and open triangles, PSSH on ITO [Kugler, 1999 #68]. Two decimal places are given because the data are averaged over a large number of measurements.

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of the difference between the valence band edge in the bis-DMOS-PPV layer and the Fermi level of the ITO substrate: E VB ranges from 1.6 eV for bisF DMOS-PPV on ITO (a) to 0.9 eV for bis-DMOS-PPV on PSSH on ITO. This result is highly relevant in the context of polymer LEDs, since E VB F determines the barrier for hole injection (⌽ pB) from the ITO electrode into the electroluminescent bis-DMOS-PPV layer. An increase of the ITO work function reduces the barrier for hole injection, thereby following the Schottky–Mott rule for the case of a zero offset ‘‘⌬’’ [31]. The ionization potential IP, as probed by UPS, corresponds to the sum of the distances of VAC the valence band edge, E VB , from the Fermi F , and of the vacuum level, E F level of the ITO substrate. As E VB and E VAC evolve strictly antiparallel with F F changing substrate work function, their sum is constant (approximately 5.9 eV) and independent of the energy level alignment at the interface (see Figure 6). In Figure 7 is the position of the vacuum level, E VAC , of the semiconF ductor overlayer (bis-DMOS-PPV) displayed, following the changes in the work function of the ITO substrate from in situ X-ray exposure. The changes in the work function in each case were determined by exposing two samples to the X-rays: one sample with the overlayer under study and one control sample without any overlayer. There are three different situations: 1. For bis-DMOS-PPV on ITO, and bis-DMOS-PPV on PSSH on ITO, the slope dE VAC /d⌽ is approximately 1.0 and the data points lie close F to the diagonal in the E VAC vs. ⌽ diagram. This indicates an almost perfect F vacuum level alignment (‘‘⌬’’ ⬃ ⫺0.2 eV). The (weak) interfacial dipole layer is independent of the change of the substrate work function induced by irradiation with X-rays. 2. For PSSH on ITO, the slope dE VAC /d⌽ is still close to 1.0. HowF ever, there is a large-offset ‘‘⌬’’ (⫹0.4 eV) from the diagonal in the E VAC F vs. ⌽ diagram. This indicates a modification of the respective vacuum level alignment by the presence of a substantial interfacial dipole layer. A likely source for the formation of this dipole layer is the protonation of the ITO surface (PSSH is a strong acid). The abrupt change of the potential across the dipole layer [32] leads to an increase of the vacuum level of PSSH on ITO relative to the vacuum level of the bare ITO substrates (positive sign of ⌬ in Figure 7). The presence of an interfacial dipole layer modifies the barrier for hole injection from the ITO substrate into the electroluminescent bis-DMOS-PPV layer, ⌽ pB(⌬) (see Figure 5): ⌽ Bp (⌬) = IP ⫺ (⌽ITO ⫹ ⌬) = ⌽ pB ⫺ ⌬

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Figure 7 The changes in E VAC of bis-DMOS-PPV as the work functions of the F different ITO substrates are changed by irradiation with X-rays. The samples are coded as in the text of Figure 6. The gray diagonal line of unit slope corresponds to an ideal vacuum level alignment.

3. For PEDOT-PSS on ITO, and bis-DMOS-PPV on PEDOT-PSS on ITO, a completely different behavior is observed. In both cases, the vacuum level E VAC remains constant (within experimental accuracy) upon changing F the work function ⌽ of the ITO substrate by X-ray irradiation. As a consequence, the offset, ⌬, for PEDOT-PSS on ITO changes from ⬃0.0 to ⫹0.5 eV, as seen in Figure 7. In the case of PSSH on ITO, the offset, ⌬, remains unchanged upon irradiation with X-rays. Given the fact that PEDOT-PSS films are made up of nanometer-sized grains covered with surface-segregated PSSH [33], a modification of ⌬ due to X-ray-induced changes of the chemical structure at the PEDOT-PSS/ITO interface can be excluded. In contrast to PSSH, PEDOT-PSS is an electrically conducting polymer with a high concentration of free charge carriers. Therefore, the charge in the PEDOTPSS layer arising from the Fermi level alignment with the ITO substrate is

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contained within a distance from the ITO interface much smaller than the thickness of the PEDOT-PSS film, thus resulting in the constancy of the vacuum level E VAC observed by UPS. F To illustrate this point further, consider bis-DMOS-PPV spin-coated onto ITO substrates with strongly different work functions, both with and without a thin, interfacial PEDOT-PSS layer: ITOs (a) and (c). The respective substrate work functions are 4.4 and 4.8 eV (⫾0.1 eV). The results are displayed in Figure 8. Without interfacial PEDOT-PSS layer, E VAC for F the bis-DMOS-PPV layer is strongly dependent on the substrate work function. The values are 1.6 and 1.2 eV, respectively. With an interfacial PEDOTPSS layer, however, E VB is 1.0 eV for bis-DMOS-PPV on both ITOs with F

Figure 8 The band diagrams corresponding to the data of Figure 7 are shown. On the left side is the case of bis-DMOS-PPV on different substrates: ITO after surface preparations (a) and (c). On the right is the case of bis-DMOS-PPV on the same substrates coated with an interfacial layer of PEDOT-PSS.

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a PEDOT-PSS interfacial layer. This confirms the establishment of the flatband situation within a distance from the ITO interface smaller than the PEDOT-PSS film thickness. In a recent study of device performance as a function of the nature of the ITO substrate, Kim et al. observed a dependence of injection-onset parameters on the treatment conditions of the ITO substrate, despite the presence of a PEDOT-PSS intermediate layer [34]. This result is in contradiction to the results reported here. However, the ITO substrates by Kim et al. had different chemical compositions, different surface morphologies, different resistances, and different work functions. It has already been shown [35] that ITO surfaces with similar work functions but different morphologies lead to notable differences in LED performance (e.g., turn-on voltages, etc.). The present results are confined to identical starting ITO samples, with a fixed surface morphology, whereas the work functions were varied independently. The work function changes reported here do not appear to lead to any major changes in morphology or surface chemistry.

IV.

INTERACTION OF POLY(p-PHENYLENEVINYLENE) WITH ITO

PPV, i.e., unsubstituted PPV, is insoluble. In order to prepare a thin film of PPV on an ITO substrate, a soluble, so-called precursor polymer is used. The most common precursor polymer is a polyelectrolyte in the form of a polysulfonium chloride or bromide salt. This particular precursor polymer is stable at low temperatures, in the dark, for years [27]. Thin films of the precursor polymer, spin-coated on ITO surfaces, are thermally converted, either under inert atmosphere or under vacuum. During the thermal conversion process, tetrahydrothiophene and HCl (or HBr) are eliminated from the precursor polymer, yielding fully conjugated PPV, as illustrated in Figure 9. These chemical species may then migrate to the surface of the PPV and be pumped away. Moreover, the strong acidity of HCl or HBr may result in chemical interactions with the surface of the ITO substrate [36], leading to the formation of undesirable chemical species at the polymer–ITO interface. For example, indium chloride (InCl3) may be formed, which serves as a ptype dopant in PPV [37]. The presence of such dopant species not only influences the electrical characteristics of the device but may also cause considerable photoluminescence quenching [38]. Previously related published works comprising Rutherford back scattering (RBS) and depth profiling using X-ray photoelectron spectroscopy (XPS) were used to demonstrate the correlation between the presence of indium and the degradation of the electroluminescence spectra in polymer-

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Figure 9

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Reaction scheme for the thermal conversion of precursor-PPV to PPV.

based LEDs [39]. Furthermore, capacitance-voltage measurements have indicated doping of PPV films from species derived from chemical reactions of the ‘‘leaving groups’’ of precursor PPV with the ITO substrate [38,40,41]. In addition, thermally stimulated current (TSC) techniques have been used to establish the existence of trap distributions in PPV LEDs, with ITO as the hole-injecting electrode [40]. Finally, depth profiling of ITO/PPV/Al devices with secondary ion mass spectroscopy (SIMS) has shown enhanced chlorine and indium count rates throughout the entire PPV layer [42]. In this context, the initial steps in the conversion process have been studied in ultrathin films of precursor-PPV. The studies have been performed using photoelectron spectroscopy on thin precursor films spin-coated onto both ITO and naturally oxidized surfaces of optically flat Si-100 wafers. Samples for study were made by spin-coating thin films (<20 nm) of the precursor-PPV onto either ITO (which has first been cleaned by a UV-ozone treatment) or the naturally oxidized surface of optically flat Si-100 wafer substrates. The coating procedure was calibrated to yield a precursor polymer thickness of about 20 nm. Since thermal conversion results in a loss of about 50% of the molecular weight of the precursor polymer, the thickness of a converted PPV film was about 10 nm. The use of high signal-to-noise electron detection systems in modern ESCA spectrometers, and the careful choice of polymer–substrate combinations, allows the study of interfaces ‘‘buried’’ under 10 nm (up to about 20 nm) of polymer layer. The chemical species involved in the thermal conversion of precursorPPV may be identified from the XPS core electron spectra of precursor-PPV, as shown in Figure 10, for a thin film coated on an oxidized Si-100 substrate

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Figure 11 Photoelectron spectra of as-prepared films of precursor-PPV on silicon. A narrow energy spectrum of the Cl(2p) components, where the experimental data were fitted with four components. The two peaks toward low binding energy derive from the precursor sulfonium chloride/bromide salt, whereas the two peaks toward high binding energy (again, spin-split components) derive from the eliminated HCl. The formation of HCl on the time scale of three successive scans (20 min/scan) indicates a partial conversion of precursor-PPV in UHV at room temperature.

in the spectra may be identified (from left to right) as follows. The two highbinding energy components of the Cl(2p) spectra correspond to eliminated HCl, and the two low-binding energy components of the Cl(2p) spectra correspond to the precursor sulfonium chloride salt. In the figure it is clear that the content of the chloride anion (Cl⫺) decreases over the time scale of the figure, whereas the content of hydrochloric acid (HCl) increases. This is an indication of the time scale of the conversion process. In other words, after less than 1 h at room temperature a partial elimination of HCl from the precursor sulfonium chloride salt occurs. Obviously, the four-peak spectrum may be fitted well with only two spin-split components of given energy

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splitting and intensity ratio. However, for precursor-PPV on ITO substrates the situation is more complex. For precursor-PPV spin-coated on ITO, the detailed features in the Cl(2p) spectra, shown in Figure 12, are clearly different from those for a silicon substrate (Figure 11). The individual components are labeled. Clearly here, the intensities of the features may not be fitted with only two spinsplit components. A third set of spin-split components, open circles in the figure, is necessary to account for the observed intensity distribution. By comparison with our own reference spectra, the third and lowest intensity doublet corresponds to Cl(2p) in InCl3. Within an our at room temperature, InCl3 is present at the polymer–ITO interface as a result of the chemical reaction of the HCl, released in the thermal elimination reaction, with the ITO substrate. By studying the various spectra of Figures 11 and 12 more closely, it may be seen that the rate of conversion, i.e., the elimination of HCl, is faster for precursor-PPV on ITO than on the natural oxide of silicon. Already in the first scan in each figure, a significant difference may be seen. Differences persist even at later stages of conversion. It appears that ITO or, more likely, the indium chloride generated at the polymer–ITO interface acts as a catalyst in the thermal elimination reaction. Moreover, from the angular dependence of the Cl(2p) spectrum, it is clear that InCl3 lies closer to the ITO substrate than the polymer surface; probably because the InCl3 is actually formed at the ITO–polymer interface, from where it diffuses into the polymer film. This result is in agreement with both a measurement, using temperaturedependent impedance spectroscopy, and theoretical modeling, performed by Scherbel et al. [43]. They found that a PPV film on ITO is composed of a highly doped region at the ITO interface and a region with lower doping at a higher distance to the ITO. Moreover, the boundary between these two regions is not sharp, but there is a gradual change in dopant concentration. Complementary information confirming the above model may be obtained from an analysis of the changes in the S(2p) spectra (not shown here), corresponding to the elimination of neutral tetrahydrothiophene molecules from the polysulfonium precursor. UPS spectra of the samples discussed above (not shown here) provide additional information. PPV films prepared on ITO substrates, by ‘‘complete thermal conversion’’ according to published recommendations, are p-type, due to residual InCl3, which is removable only under rather stringent conditions of temperature and time in UHV (e.g., 200⬚C for 8 h). The presence of doping may be seen as a shift in the Fermi energy and as a shift of the C(1s) peak to lower binding energy, in converted, but doped, films on ITO. On the other hand, a second layer applied on top of the first, converted, PPV layer, and then thermally converted, does not appear to be doped, at least

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in the surface region accessed by UPS. Since the second layer is not in contact with the ITO substrate, the interfacial interactions described above do no occur. The residual InCl3 in the first PPV layer does not appear to diffuse into the second layer during the (second) thermal treatment.

V.

POLY(9,9-DIOCTYLFLUORENE) ON METALS

Poly(9,9-dioctylfluorene) (PFO), the chemical structure of which is shown in Figure 13, was supplied by Cambridge Display Technology (Cambridge, UK). Thin films were prepared by means of spin-coating from a 5 mg/ml xylene solution which had been filtered with 0.2 ␮m polytetrafluoroethylene (PTFE) syringe filter. Resulting film thicknesses were determined by ellipsometry, and checked by observing the attenuation of the substrate XPS ˚ thick), signal in the presence of the thin-film overlayer (for films up to 100 A ˚ ). Very and by a DEKTAK instrument for thicker films (starting from 500 A good overall agreement was found in the regions of overlap. Three kinds of substrate, with considerably different work functions, were used: aluminum ˚ -thick native oxide layer (AlxOy); optically flat Si-100 with a 10- to 20-A ˚ layer of native oxide (SiO2); and UV-ozonewafers, with a nominally 10-A etched gold (Au). The first two substrate types were cleaned in standard organic solvents prior to the spin-coating of the polymer films, whereas the gold substrates were first cleaned in organic solvents and then cleaned by the UV-ozone treatment. The energy band parameters have been obtained from the UPS data as shown in Figure 14, where experimental spectra for a gold substrate (black dots) and a PFO polymer film (white dots) are shown. The nomenclature used is that employed in earlier studies by Seki et al. [2,5,29] and Kahn et al. [4,30], and was previously introduced in Section III. Two related issues should be emphasized. First, the sum of E VAC and E FVB is equal to the IP of F the polymer and constitutes a specific property of the material, independent of the energy level alignment at the interface. Second, the position of the Fermi level in the polymer film cannot be measured directly in UPS.

Figure 13

Chemical structure of poly(9,9-dioctylfluorene), or PFO.

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Figure 14 Illustration of the UPS method, and the energy band parameters, as discussed in the text.

In Figure 15 are shown the main results, changes in E VAC as a function F of the work function of the substrate, on the pristine PFO films. The work functions of the substrates, denoted AlxOy, SiO2, and Au, were found to be 3.9, 4.4, and 5.4 eV, respectively. PFO films of different thicknesses, in the ˚ , were studied on each of the three substrates. In each case, range 50–1100 A the position of the vacuum level of PFO as measured from the Fermi level, E VAC , follows the Schottky–Mott limit (long dashed line), although a small F deviation in slope is apparent. As the Fermi energy changes from 3.9 eV for aluminum oxide to 5.4 eV for ozone-treated gold, E VAC increases from 3.9 F eV to 5.2 eV, resulting in a negligible interfacial shift ⌬, which implies alignment of the vacuum levels at the interface between PFO and conducting substrates. As a consequence, the entire UPS spectrum experiences a rigid shift in binding energy, up to 1.3 eV, when the work function (substrate) is varied. In particular, E VB changes by the same amount as E VAC because they F F are closely related through the IP. This issue is very important for polymer LEDs because such significant shifts cause large changes in the barrier

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Figure 15 Dependence of E VAC on the work function of the substrate for PFO F ˚. films of different thickness: 50, 100, 220, 500, and 1100 A

heights for carrier injection at the interfaces between the polymer and the electrode contacts. The above results indicate that PFO films on the substrates indicated follow rather closely the Schottky–Mott limit independently of the film thick˚ , dependence of E VAC ness. Even for the thickest film, 1100 A on the substrate F work function is evident. This means that the flat-band regime extends at least to this distance from the interface, i.e., the length of the depletion region is larger than the film thickness. Such large depletion lengths are the consequence of the very low concentrations of mobile charge carriers in the PFO polymer layer. Since in wide-band gap semiconductors at room temperature the charge carriers are mostly extrinsic [3], the upper limit for the carrier density may be estimated from the relationship of the depletion length, ␭, to the density of impurity centers N [44]: N = 2␧␧0VD /e␭2

(4)

In an organic medium with a static dielectric constant, ␧, of about 3, and the diffusion potential, VD, close to 1 eV, Eq. (4) gives N on the order of 1014 impurity sites/cm3. This number indicates the high purity of the PFO samples, as the typical values reported are on the order of 1016 –1017 sites/

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tributed within the polymer film, as determined by angle-dependent XPS data. The main results are plotted in the three diagrams of Figure 16. Starting VAC , and from the bottom, successive panels show the variation of E VB F , EF their sum, plotted against the dopant concentration for both cases: PFO/AlxOy (white dots) and PFO/Au (black dots). Note that the average concentration of the dopant atoms, as determined by XPS, are given. The most important points are as follows: (a) In each case, as the doping level increases, E VB F also increases while E VAC decreases, such that their sum remains constant; F VAC (b) both parameters E VB tend to reach the same limits (long dashes) F and E F independent of the substrate; and (c) for PFO on AlxOy the saturation values (corresponding to horizontal lines) for E VB and E FVAC are reached at considF erably lower concentration of dopant atoms than for PFO on Au, as discussed below. Note that changes in the valence band (UPS) spectra of PFO are accompanied by expected shifts in the carbon C(1s) core level spectra (not shown). There are two main reasons for the observed changes in E VB and F VAC E F . First, doping occurs via charge transfer from the sodium atoms to the ␲ system of the polymer chains [48]. The added electrons lead to the generation of self-localized states, known as solitons, polarons, or bipolarons depending on the symmetry of the ground state of the system under investigation [45,49]. In PFO, two new features, corresponding to bipolaron states, appear in the previously forbidden band gap, which in turn leads to a shift in the position of the Fermi level. The Fermi level in the fully doped polymer is then determined by the electronic structure of the polymer itself rather than the substrate (as in the case of the undoped polymer). Also, the Fermi level moves gradually to a position halfway between the lower binding energy bipolaron state and the conduction band edge [50], i.e., midway between a filled band (the upper bipolaron band) and an empty band (the conduction band). On the other hand, in principle, the generation of free charges in the polymer allows band bending to occur. As significant amounts of charge carriers become available on the polymer side of the interface, the equilibrium between the conducting substrate and the organic layer becomes determined by charge transfer across the interface. When bipolaron states generated in the polymer (by doping) lie above (at lower binding energy than) the Fermi energy of the substrate, electron transfer from PFO take places, until equilibrium has been reached. As a consequence, band bending may occur in the interface region of PFO film, with the depletion length given by Eq. (1). It must be pointed out that for low concentrations of the dopant, the depletion length might still be much longer than the film thickness. Hence, a shortening of the depletion region caused by an increase in density of extra charge carriers in the sample results in a continuous shift

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Figure 16 Changes in the E VAC , E VB F F , and their sum upon increasing the concentration of a dopant in PFO film. The horizontal axes, common for all three panels, are scaled either in average concentration of sodium atoms (bottom) or in a percentage doping level (top).

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of the energy levels of the PFO toward higher binding energy. This process is observed until the depletion length becomes shorter than the film thickness, beyond which it is not observable by surface-sensitive UPS, even if it should continue with further doping. The overall shifts are very significant: about 1.2 eV for PFO/AlxOy and 2.4 eV in the case of PFO/Au. The fact that the sum E VAC ⫹ E VB remains F F constant for all doping levels plays an important role in the interpretation of the experimental data. For instance, possible modification of the dipole moment contribution to the work function of PFO by means of formation of the metallic islands on the surface, and so forth, would lead to the changes only in E VAC which, in turn, should result in a variation of the sum. The F fact that this does not happen (within the experimental error) is in agreement with the uniform diffusion of the sodium atoms/ions into the film. As indicated above, E VAC tends to reach the same saturation value (2.7 F eV) at higher doping levels, independent of the substrate type used. E VB F tends to reach 3.2 eV in the same way. Increasing the sodium concentration further does not lead to changes in the energy level positions of the PFO. Once the conductivity of a PFO film is high enough, the energy band parameters become completely independent of the work function of the substrate. This behavior is in full agreement with the results of Kugler et al. [27], for electrically conducting films of PEDOT-PSS on ITO, discussed in Section III. In the present case, however, there is a continuous transition from the case of a strong dependence of the energy level alignment on the substrate work function (pristine PFO film), to the situation where the specific values of E VAC and E VB depend only on the intrinsic properties of the F F polymer overlayer itself (sodium-doped PFO above the saturation threshold). It is also worth pointing out that the saturation values of E VAC and F VB E F are reached at quite different doping levels, depending on the kind of substrate used. For PFO on AlxOy, roughly 0.2 ⫻ 1021 atoms/cm3 is needed, whereas for PFO-on-Au, this value increases up to 1021 atoms/cm3. This difference is a consequence of the fact that in the latter case more charge is necessary to bring the entire electronic structure of the PFO film to its final position. In other words, the larger the energy difference between the work function of the polymer and that of the substrate, the more charge has to be transferred before saturation is reached. B.

Lithium on Poly(9,9-dioctylfluorene)

Lithium was deposited on thin films of PFO in UHV. The results are depicted in Figure 17. The overall changes in the valence band region are displayed in the left panel. The vertical bars indicate the position of the Fermi level. Starting from the bottom-most spectrum (corresponding to the pristine poly-

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Figure 17 Evolution of the HeI spectrum of PFO upon increased Li doping. The Fermi region is magnified in the right panel. The binding energy is given with the vacuum level as reference.

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mer), lithium deposition increases upward. The estimated average doping level range from 0.05 lithium atom per PFO repeat unit (second spectrum from bottom of figure) up to 1.5 (top-most spectrum). As the doping level increases, two new features evolve in the previously forbidden band gap at approximately 3.4 eV and 5.8 eV binding energy (arrows). The higher binding energy one gains its intensity from the old valence band edge as it moves deeper into the band gap. Since the latter shift in binding energy is not very pronounced, the newly formed feature is hardly resolved from the original valence band edge, which still possesses some remaining intensity due to undoped sites. This is not the case for the second new state appearing, as the continuous growth of the peak at 3.4 eV binding energy upon successive deposition steps is clearly visible. The position of the Fermi energy coincides with the newly formed feature at 3.4 eV for low doping levels, up to 0.1 Li/repeat unit (right panel, Figure 17). However, as the doping level increases, the Fermi energy moves away from the center of the newly formed peak toward the vacuum level. This observation suggests that at low doping levels the new states correspond to polarons, whereas for higher doping levels bipolarons are formed. Polaron-to-bipolaron transition upon increasing doping level was reported previously for poly(p-phenylenevinylene) [51]. Thus, if LiF undergoes dissociation at the interface, resulting in a full or partial charge donation to the polymer, significant changes in the vertical bars region of the photoelectron spectra corresponding to either polaron or bipolaron formation would be expected. However, if the dissociation leads instead to the formation of Li⫹ and F⫺ ions and no charge is donated to the neighboring polymer chains, then change in the relative binding energy of the F(1s) and Li(1s) core levels with respect to the value typical for bulk LiF (629.2 eV) remains the only reliable way to detect this process.

VII.

INSULATING OVERLAYERS ON POLY(9,9-DIOCTYLFLUORENE)

Recently, the use of insulating thin layers between the metal contacts and the organic/polymer layer(s) has been shown to dramatically improve lifetimes and quantum efficiency in organic-based LEDs. These insulating layers, typically LiF [52–54], are formed between the cathode and the organic material. Other insulating layers, such as Al2O3 [55–57], or insulating polymers [58] at the interface can also improve performance. Below is a short summary of the various theories that have been presented to explain the improvements seen. ˚ ) aluminum oxide layer between the Al The addition of a thin (5–20 A cathode and Alq3 is thought to create a barrier through which electrons can

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tunnel [56]. Considerable voltage drop would occur over the Al oxide layer when a bias is applied, resulting in an improved alignment of the Fermi level of the Al cathode and the conduction level of Alq3, thus enhancing the injection of electrons [56]. The luminescence efficiency also may be enhanced by a reduction of quenching sites at the cathode–Alq3 interface, since the Al oxide layer will act as a buffer, reducing the Al bonding–induced degradation of Alq3. Low work function metals interacting with organic (light-emitting) materials are known to significantly modify the electronic structure around the Fermi level, often introducing gap states [51,59–61]. Luminescence efficiency is enhanced even at very low Al oxide coverage ˚ ) for devices based on amino-oxadiazole-fluorenes. It is unlikely that a (8 A continuous oxide film is formed at the interface or that a significant voltage drop can occur over such a thin layer; hence, the reduction of quenching sites at the cathode interface seems to be the more likely scenario [57]. Deposition of LiF at the cathode interface has been shown to enhance charge injection in Alq-based [52,62] and 4,4⬘-bis(2,2-diphenylvinyl)-1,1⬘˚ was found to biphenyl-based [54] LEDs. Even LiF coverage as low as 1 A increase charge injection [62]. Here too the lack of significant LiF coverage contradicts the concept of a large voltage drop over an insulating barrier with resulting improvement of the alignment of the metal Fermi level and the emissive material’s conduction band or LUMO level, although this has been proposed for doped Alq3 devices [53]. Instead, band bending (>1 eV in Alq devices) at the LiF–organic interface is presented as an alternative explanation [52]. The LiF did not dissociate upon deposition on Alq, so a mechanism for the large band bending was not presented. However, dissociation can take place when aluminum is deposited on top of the LiF layer in Alq3 devices [63], producing F⫺ and Li⫹ ions that can interact with the organic layer and/or the metal cathode. Other explanations for the effectiveness of an LiF interfacial layer include the proposed reduction of the emissive material’s surface potential due to the dipole moment of LiF and reduction of quenching sites at the interface similar to that proposed for Al oxide [54]. Based on XPS studies of the LiF-Al interface, dissociation of LiF when in contact with Al forming a Li-Al alloy has been proposed [54]. Since it is known that alloying aluminum with lithium enhances electron injection in organic LEDs [64], dissociation of LiF when in contact with Al also could explain the enhanced performance of the devices with LiF interfacial layers [54]. Ultrathin films of PFO were spin-coated onto various substrates from mixed xylenes 5 mg/ml solution previously filtered with 0.2-␮m PTFE sy˚ as determined from ringe filter resulting in film thickness of about 100 A ellipsometry and XPS measurements. The Li, LiF, and Al deposition was done in situ at pressures in the low 10⫺9 –high 10⫺10 mbar region. The Li

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was deposited from a SAES getter source. The LiF and Al were deposited using an Omicron UHV evaporator EFM 3.

A.

LiF-PFO Interface

In Table 1 are shown XPS data on the core level features of pristine PFO and the early stages of interface formation between LiF and PFO. Going from pristine polymer to increased LiF coverage, the position of the C(1s) core level peak does not shift more than about 0.1 eV to higher binding energy. The F(1s) core level peak undergoes a similar evolution with the binding energy increasing by about 0.1 eV as the interface is built up, and no significant shifts are seen for the Li(1s) peak. The difference in binding energy between Li(1s) and F(1s) signals remains roughly constant, and the shape and width of the core level peaks are not significantly affected by the LiF deposition. Thus, there is no evidence from the core level spectroscopy that LiF dissociates at the interface to PFO, nor is there evidence of a significant dipole formation (or band bending) given the lack of binding energy shifts of the core level features. Figure 18 depicts a comparison between HeI spectra of pristine PFO ˚ up to 21 A ˚ ). Continuous and LiF/PFO at different LiF coverages (from 1 A smearing out of the characteristic features in the vertical bars region of PFO takes place along with increasing LiF coverage. However, even at high coverage of LiF, the essential features of PFO remain visible. There is no significant change in the work function when going from pristine PFO to 21

Table 1 XPS data on core level features of pristine PFO and early stages of LiF–PFO interface formation C(1s) LiF ˚) (A 0 1 3 4 5 6 9

F(1s)

Li(1s)

BE (eV)

FWHM (eV)

BE (eV)

FWHM (eV)

BE (eV)

FWHM (eV)

284.96 285.00 285.02 285.04 285.06 285.06 285.08

1.02 1.03 1.05 1.05 1.04 1.07 1.07

N.A. 685.54 685.62 685.58 685.62 685.66 685.67

N.A. 1.66 1.72 1.67 1.63 1.63 1.67

N.A. — — 56.45 56.32 56.46 56.45

N.A. — — 1.52 1.54 1.48 1.54

XPS, X-ray photoelectron spectroscopy; PFO, poly(9,9-dioctylfluorene); BE binding energy; FWHM, full width at half maximum; NA, not applicable.

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Figure 18 Comparison between HeI spectra of pristine PFO and LiF/PFO for different LiF coverage. The position of the Fermi level is marked with vertical bars.

˚ coverage of LiF. This is in contrast to the case of Li deposition on PFO, A where there was an immediate shift in work function and two new states appeared in the band gap. Hence, there is no evidence from UPS that LiF dissociates at the interface to PFO donating the charge to the polymer, in agreement with the XPS data. B.

LiF-Al Interface

LiF was deposited on in situ sputter-cleaned Al substrates. In Figure 19 are shown the three most representative steps of the evolution of the F(1s) core level spectrum upon increasing deposition of LiF. There are two distinct components in the F(1s) signal present at all the LiF deposition stages shown. For each step the F(1s) core level was fitted with two Voigt functions ˚ ), the low with the Shirley background subtracted. At low coverage (1.3 A

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Figure 19 F(1s) core level spectra for three successive steps of the LiF–Al interface formation. The numbers in the corner give the effective thickness of the LiF layer. Solid lines represent the two fitted components (Voigt functions).

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binding energy feature (686.2 eV) dominates over the high binding energy feature (687.0 eV). The relative intensity of the two peaks changes as the coverage increases, with the high binding energy feature becoming the dom˚ . The FWHM decreases as inant one for LiF overlayers thicker than 3.9 A the high binding energy feature becomes dominant, which would not be the case if the shifts were due to sample charging. The Li(1s) core level exhibits similar behavior, a broadening to the low binding energy side at low coverage that disappears at high coverage. The F(1s) interface component present at the lower binding energy could in principle indicate the dissociation of LiF into Li⫹ and F⫺ ions, as in this case the corresponding fluorine atoms would have accepted negative charge from the Li ions. However, this would require that the Li(1s) interface component be shifted toward slightly higher binding energies, having donated charge to the fluorine. That is in contrast to our results, where similar shift of the Li(1s) peak toward lower binding energy is seen at the interface as for F(1s)! Hence, our results would preclude dissociation of LiF at the Al interface. Instead, the shift in both F(1s) and Li(1s) toward lower binding energy at the interface is likely the result of a combination of surface dipole formation and screening from the metal. C.

Al-PFO Interface

Aluminum also was deposited in situ onto as-prepared PFO films. Figure 20 shows the changes in the valence band region of PFO upon (increased) Al coverage. A significant smearing out of the ‘‘␲-band edge’’ and the peak positioned at 8 eV binding energy occurs upon Al deposition, leading to almost complete disappearance at high Al coverage. Note that a relatively small number of Al atoms is needed (effective thickness of the deposited Al ˚ ) in order to smear out the valence band features of PFO layer equals 3 A significantly, suggesting chemical interactions between the aluminum atoms and the carbon atoms in the polymer backbone. For the sake of comparison, the result from the section on LiF–PFO interface can be recalled, where the corresponding changes in the valence band region caused by the presence of the LiF layer are less pronounced even for the thickness of the LiF layer ˚ . The interaction between the Al atoms and polymer backlarger than 10 A bone likely leads to the formation of Al–carbon complexes at the interface, disrupting ␲ conjugation along the backbone. This would cause the enhanced smearing out of the low binding energy features seen in the UPS spectra, as compared to that caused only by an overlayer effect (LiF case). The described changes are accompanied by an abrupt decrease in the work function of PFO, from 4.3 eV to 4.1 eV, when going from spectrum of pristine material to the one corresponding to the lowest Al coverage. Further

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Figure 20 Evolution of the valence band region of PFO during Al deposition. Vertical bars indicate the position of the Fermi level.

deposition of the Al does not affect the work function of the system. The small decrease of the work function affects the position of the carbon C(1s) core level, which then shifts by roughly 0.2 eV toward higher binding energy and thus remains constant with respect to the vacuum level. The FWHM of C(1s) increases from 0.85 eV to 1.00 eV upon Al deposition, which again indicates formation of new (C-Al) complexes. The fact that the increase in the FWHM of C(1s) is not very pronounced can be rationalized by means

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of the strong contribution to this signal coming from the alkali side chains, which are not going to be significantly affected by the Al deposition. D.

Al–LiF–PFO Interface

Since results have been presented whereby LiF dissociates when in contact with Al and Alq3 but not when in contact with Al or Alq3 separately [20], we carried out Al deposition onto LiF/PFO films. Two types of LiF/PFO ˚ or 10 A ˚. films were studied, the difference being in LiF thickness: 5 A However, the experimental data did not differ between these samples, so ˚ LiF/PFO samples are presented here. only the results from the 5 A A series of HeI UPS spectra for increasing coverage of Al on a LiF/ PFO film are depicted in Figure 21. Upon deposition of Al, the ␲-band edge features of the spectra are somewhat smeared out, but there is no significant shift in the work function (less than 0.1 eV), nor are there any new features appearing in the band gap, in contrast to the report on Al/LiF/Alq3 [63]. The

Figure 21 film.

A series of HeI spectra for increasing coverage of Al on an LiF/PFO

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FWHM of the Al(2p), C(1s), F(1s), and Li(1s) core level peaks do not change by more than 0.05 eV as compared to the ‘‘prinstine’’ LiF/PFO spectra. No significant binding energy shifts are seen for any of the core levels upon deposition or as the coverage increases. Hence, the Al–LiF– PFO interfaces are dramatically different from those reported for Alq3. In PFO, no dissociation of LiF takes place and no new gap states are formed. There is no significant surface dipole formation at the LiF–PFO interface, though there is an evidence of a weak dipole at the interface between LiF and (clean) Al. It should be pointed out that the smearing out of the valence band region features for Al–LiF–PFO does not extend to the limits observed for the Al–PFO interface, even though the effective thickness of the Al layer ˚ ) is larger than in the later case (3 A ˚ ). This, deposited onto LiF (5.7 A together with the fact that there is no increase in the FWHM of the C(1s) core level for Al–LiF–PFO interface, clearly indicates that the LiF layer protects PFO from Al-C complex–forming interactions.

VIII.

SUMMARY

In the context of polymer interfaces in device applications, some examples of organic polymer interface formations studied by photoelectron spectroscopy have been discussed. The band level alignment between organic materials and electrode materials has been emphasized. The examples were chosen to illustrate various interface aspects, such as the effect of substrate work–function variation, ideal vs. ‘‘real’’ surfaces, and ‘‘buried interfaces.’’ Engineering of such interfaces is expected to be central to the development of polymer-based electronics.

ACKNOWLEDGMENTS This work was supported by the European Commission within a Brite/ EuRam project (OSCA, project number 4438), and the Center for Advanced Molecular Materials, CAMM, supported by the Swedish Foundation for Strategic Research (SSF). In addition, work on polymer surfaces and interfaces in Linko¨ping is supported by a Research Training Network (LAMINATE, project number 00135), a Training and Mobility of Researchers (TMR) network (SISITOMAS, project number 0261), the Swedish Natural Sciences Research Council (NFR), the Swedish Research Council for Engineering Sciences (TFR), and the Carl Tryggers Foundation.

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5 Role of Interfaces in Semiconducting Polymer Optoelectronic Devices Paul F. van Hutten, Victor V. Krasnikov, and Georges Hadziioannou University of Groningen, Groningen, The Netherlands

I.

INTRODUCTION

Recent advances in the performance of organic semiconductor materials in device configurations have been dramatic. In particular for light-emitting diodes (LEDs) [1–3], stability and durability have improved to levels that warrant their application in everyday life. The performance of organic optoelectronic devices has been found to depend crucially on the properties of the interfaces. This concerns electronic as well as structural aspects of interfaces. This feature can be exploited to improve the characteristics of a device. A well-known example is the incorporation of charge-transporting and charge-blocking layers in LEDs. Other functions of a layer between the active film and the electrodes may be to change the interface roughness or to act as a barrier for diffusion of oxygen or metal. In this chapter, the emphasis will be on the geometry of the interface between active components or, in other words, on the morphology of the active layer. The role of morphology is especially significant when there is more than one active organic component, i.e., in a composite active layer. (Interfaces can be a major factor within a one-component system too, e.g., the influence of grain boundaries on charge transport or on the propagation of light.) Currently, the prime example of a composite system is the photovoltaic material based on the donor–acceptor concept. The major issues here 113

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are obviously the interface area and the continuity of the component phases, but details are expected to be important in the inherently asymmetrical configuration of a device. The microstructure of an organic-based optoelectronic device is commonly realized through a series of processing steps, and each of these can be engineered to optimize performance. If a polymer is involved, it will be found that the microstructure of the polymer layer depends strongly on the processing history. This is due to the nonequilibrium, metastable character of polymer solids. This aspect needs careful consideration but can be exploited to one’s advantage. However, one may use an entirely different approach and apply the self-organization properties of cleverly designed molecules to obtain a desired microstructure for a specific purpose (optoelectronic function). This is the approach that we pursue, and the first steps toward this goal will be presented in Section III of this chapter. We focus on the use of block copolymers for microstructure formation. A block copolymer molecule consists of two (or more) chemically distinct macromolecular chains (the blocks) that are permanently linked, usually end to end. In most cases, the blocks tend to segregate into nanometer scale domains (microphases) of their own kind, for thermodynamic reasons. As a result, the block copolymer molecule is not found on one side of the interface or on the other, but exactly across it! These molecules have their part of the interface built into them, and their architecture is the key factor in shaping that interface as a whole. Our ability to vary the chemical structure and the length of the blocks gives us a powerful tool for controlling the nature and geometry of the interface formed during the microphase separation. Copolymers have been synthesized for applications in optoelectronics, in most cases merely to effect a ‘‘chemical tuning’’ of their optoelectronic properties [4–7]. The striking phase behavior of block copolymers has not been invoked for the realization of morphologies that may be particularly effective. For this strategy to work, the physicochemical characteristics that drive the microphase separation must be combined with the appropriate optoelectronic functionalities. Further ‘‘goodies’’ of the interfaces in block copolymers are a direct consequence of the chemical connectivity between the blocks. Since the interface is not created by bringing two materials in contact, the problem of wetting between dissimilar materials is avoided, resulting in optimum contact. For the same reason, the interface will be pure, uncontaminated. Since the covalent bond is a strong bond, ‘‘adhesion’’ is also optimal. Hence, we expect to create a robust interface by using block copolymers. For a basic understanding, the interface must be characterized in terms of geometry and nature: area, curvature, local roughness, and degree of molecular mixing.

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Clearly, in this order, these features represent a decreasing scale of observation, and different research tools are required at each level. The tunability of block copolymers extends to all of these features. The purpose of this chapter is to give the reader the flavor of the concepts presented above. After some simple considerations on interfaces in polymer-based systems, the issue of interfaces and morphology in relation to functionality is illustrated by studies of donor–acceptor systems for photovoltaic devices. These include small-molecule composites as well as examples of polymer systems designed for this purpose. The final section deals with electronic phenomena accompanying the formation of interfaces.

II.

INTERFACES IN POLYMER-BASED OPTOELECTRONIC SYSTEMS

A.

Polymer Blends

For combinations of macromolecular chains, immiscibility is the rule rather than the exception. It is a consequence of the long-chain nature of the molecules. Since the number of molecules per unit volume is very low, the entropy gained upon mixing is very low and is easily offset by enthalpic factors originating in segment–segment interactions. In many cases, like contacts (A–A and B–B) are preferred over mixed contacts (A–B), and phase separation leads to a lowering of the free energy. The criterion for miscibility in a 1:1 blend of two homopolymers (flexible chains) was derived by Flory and Huggins [8]: (␹ N)blend < 2

(1)

Here N is the degree of polymerization (taken identical for each of the polymers), and ␹ (T ) is the Flory–Huggins interaction parameter:

␹ = (kBT)⫺1[␧AB ⫺ (␧AA ⫹ ␧ BB)/2]

(2)

where ␧ij are the interaction energies for contacts between i and j (i, j = A, B) segments. (These ␧ij are not constants, however, but may depend on several state parameters.) For polymers that have dispersion forces as the primary interactions, ␹ ⱖ 0. If there are no special entropy effects involved in the association of unlike polymer segments, ␹ varies more or less inversely proportional to T, and as a result, mixing is favored by a temperature increase. Higher ␹ N means stronger segregation (purer phases) and a sharper interface. Based on the above, one might expect interfaces to be formed spontaneously in a polymer blend. However, the situation is more complicated. Due to the kinetics of diffusion, the phases encountered at the service tem-

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perature are not those corresponding to thermodynamic equilibrium. Changes in the sizes or composition of domains in phase-separated morphologies require mass transport, even over macroscopic distances. Diffusive transport is very much impeded by the presence of entanglements, which exist in the relaxed polymer bulk for N exceeding a certain value, Ne. For vinyl polymers, of which polystyrene and poly(methyl methacrylate) (PMMA) are well-known examples, the value of Ne is about 200, corresponding to a molecular weight (molar mass, in g/mol) of about 20,000. For chains longer than Ne, the ‘‘disentanglement time’’ is proportional to N 3 and the diffusion coefficient of a chain varies as N ⫺2. Above the glass transition temperature, Tg, the time it takes for a chain with N = 200 to travel a distance of 1 ␮m would be of the order of 1 h. Some tens of degrees below Tg, mass transport is completely stopped. In semicrystalline samples, the long-range diffusion of a chain will also be impeded when it is partially incorporated in a crystallite. Because of the lack of sp3-configured atoms in the backbone, the chains of conjugated polymers are somewhat inflexible, and entanglements do not occur except at much larger chain lengths, probably much larger than the length of many actual chains. The chain stiffness, expressed as its persistence length, is a ‘‘classical’’ aspect of macromolecules that has not been explored for conjugated polymers, except in a few cases [9–11]. Therefore, it is very difficult to estimate the importance of diffusion in multicomponent films of conjugated polymers at operating conditions. It is important to note that diffusion in a polymer solid is very much impeded by the presence of side chains. Side chains slow down phase separation, but by the same token they stabilize sharp interfaces that have been artificially created by processing, such as those in double-layer films. Side-chain-substituted derivatives of conjugated polymers are much more soluble than their parent polymers and are therefore often employed. A difference in the flexibility of the chains is another thermodynamic driving factor for the demixing of polymer blends. Clearly, this could be a relevant factor for composite materials based on conjugated polymers. The basic science of the thermodynamics and phase behavior of such systems is still in its infancy; the phenomena are especially relevant for block copolymers (see Section III.B). Since the diffusion processes in the polymer ‘‘solid’’ are many orders of magnitude slower than the transport modes operating under the conditions of processing, the latter conditions will primarily determine the morphology. Nonequilibrium processes, such as flow and heat transport, have a major influence. Even when polymer layers are deposited one atop the other, the final polymer–polymer interface(s) will depend strongly on the methods of processing of the various polymer layers. In particular, the solvent(s) used to cast the layers will determine the degree of mixing at the interface. Blur-

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ring of the interface is expected if a common solvent is used to cast the second layer on top of the first, up to the point where the first layer is totally disrupted. (But note that a ‘‘common’’ solvent does not imply that both polymers can be dissolved together to any desired concentration; the phase behavior of such a three-component system generally shows polymer demixing as well.) Spin-casting, the technique of choice for obtaining thin polymer films, introduces additional nonequilibrium features. The solidification process is affected by the centrifugal flow of the solution. The rapid evaporation of the solvent may also promote a nonequilibrium conformation of the chains. In particular, and most essential, the extent of chain entanglement in the solution, which is essentially nil, may be carried over to the film. This leaves the solid film in a nonequilibrium state. Obviously, due to the various kinetic factors coming into play during processing, the microstructure found in phase-separated polymer systems may not conform to the equilibrium phase diagram. Depending on the extent of diffusion required to transform the initial structure into a more stable one, the morphology may be found to change over time. The entire composition profile will adjust itself according to temperature and molecular weights of the polymers (i.e., as dictated by the product ␹ N). This process may be so slow as to be unnoticeable at room temperature, but it will accelerate with increasing temperature. This implies that polymer–polymer interfaces are often metastable. The fact that each polymer has a distribution of molecular weights complicates the situation; the shorter molecules tend to mix more readily and their presence will therefore promote a diffuse interface. An important consequence of the metastable character of phase-separated polymers is that their microstructure may change dramatically, and irreversibly, upon thermal annealing. Hence, the properties of such materials may be highly temperature-sensitive. Especially during the 1960s and 1970s, much scientific effort was devoted to relating the mechanical properties of solid polymers (and their dependence on temperature) to their detailed microstructure, down to the macromolecular level [12]. Chain ends, branches, chemical imperfections, and more specific features of particular chains were invoked to describe the mechanisms behind the tensile behavior of polymers. This is not to say that the issue has been fully resolved. In the authors’ opinion, the community that is currently involved in research on (semi)conducting polymers now faces a similar challenge, this time concerning the relationship between microstructure and electronic transport properties. B.

Interfaces Between Polymers and Other Device Materials; Model Oligomers

The fact that the materials that constitute an optoelectronic device—the electrode materials on the one hand and the organic layer(s) on the other—

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are very different in many respects has important consequences for the properties of their interfaces. When two materials are brought into contact, the properties of each material’s surface not only govern wetting and adhesion, but may also affect the structure of a layer as a whole, in particular the organic one. Preferential adsorption of one component from a blend at the interface, or one block of a copolymer, will influence the microstructure. Order and alignment in the organic phase may be promoted through a (grapho)epitaxy mechanism. In low molecular weight organic compounds, the growth of one crystal polymorph may be favored over another, and so forth. (Of course, to some extent, this consideration is also valid for organic–organic interfaces.) Electron diffraction methods form the most obvious tool to study the phenomena of epitaxy and polymorphism. What is the nature of the surface that a conjugated polymer exposes to another material, such as a metal electrode? Are there specific interactions between the two materials? While many questions concerning the bulk structure of semiconducting polymers remain, characterization of surface and interface structure is even more challenging. It is not obvious how the orientation of conjugated segments at an interface could be assessed. This is one issue on which the study of model oligomers may contribute to our understanding at the molecular level. These model oligomers are low molecular weight analogs of a polymer but of well-defined length. Being a single chemical species, they often form highly crystalline thin films. In some cases, single crystals suitable for crystallographic studies can be obtained, from which the molecular packing may be resolved. The crystallographic data further enable one to find the molecular orientation in thin films. For example, one of the issues is to find the most favorable alignment and mode of packing of conjugated molecules for a given direction of charge transport. As a first step, one may study the deposition and growth of crystals of conjugated oligomers on a variety of substrates that are relevant to optoelectronic devices. This opens the way to an investigation of the charge carrier mobility in relation to solid-state structure and morphology. An additional advantage of oligomers over the corresponding polymers is that they can be processed under ultraclean conditions. Films can be prepared by vapor deposition under ultrahigh-vacuum (UHV) conditions, thus avoiding the uncontrolled influences of a solvent. Deposits can be as thin as a monolayer, which opens the way to spectroscopic studies that may reveal how the substrate and overlayer interact. Section IV reports on an investigation of the electronic energy level structure of such interfaces by means of ultraviolet photoelectron spectroscopy. This is probably as close as one can get to the actual interface. It may well be that an even more local probe of the structure and its defects is indispensable for the understanding of, for example, the mechanisms of charge trapping.

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Interface Geometry in Two-Phase Thin Films

One can look at the geometry of interfaces at many different levels or scales. At the end of the (chemical) scale is the roughness of the interface at the molecular level. How important is molecular scale roughness? It increases the interfacial area between two dissimilar species. One could imagine that this makes a difference for transport phenomena that depend on the proximity of molecules, e.g., tunneling. For many other transport mechanisms, e.g., variable-range hopping, exciton diffusion, or dipole–dipole interactions, a range of a few nanometers is probably appropriate. It therefore seems relevant to be able to characterize and control the microstructure of a material at the nanometer scale, even if the molecular scale remains poorly accessible. As stated in the Introduction, control of the microstructure at the nanometer level (and up) can be designed directly into the molecule in the case of polymers. What techniques allow one to probe the interface geometry? If one has access to a surface that will become an interface in the final material or device, scanning probe microscopy appears to offer the closest approach to it. On conducting surfaces, scanning tunneling microscopy (STM) furnishes the required detail but for organic semiconductors, atomic force microscopy (AFM)-derived techniques may be the sole option. The softness of organic surfaces makes it hard to attain the desired resolution. For observing the actual interface between materials, electron microscopy of cross sections is a valuable option. It has the added advantage that diffraction can be studied at the same time. However, during the necessary procedures of embedding and ultramicrotomy, artifacts may be easily introduced. Another possibility is measuring the density profile by means of reflectivity techniques (X-ray, neutron) [13]. The resolution is somewhere be˚ ngstro¨ms and a few nanometers, and the information is limited tween a few A to the direction normal to the surface probed. For interfaces parallel to that surface, the profile reflects the extent of mixing at the interface. Reflectivity methods are by far not as common and accessible as the scanning probe techniques. A number of criteria can be employed to characterize the microstructure of a two-component film between electrodes (Figure 1). First, one has to find whether there is molecular mixing or whether there is segregation (phase separation) into domains of the respective components. A primary criterion for a two-phase structure is its scale (the ‘‘characteristic size’’ of the domains or channels). Additional criteria are whether it is disperse (one phase continuous) or bicontinuous; whether it is (macroscopically) isotropic or somehow oriented; whether it is symmetrical or asymmetrical across its

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Figure 1 Examples of two-phase film structure between electrodes (hatched). A) double layer; B) interdigitating double layer; C) island layer of black phase; D) columnar bridging black phase; E) fully interpenetrating phases; F) disperse black phase. A, B, and C are asymmetrical films; D, E, and F are symmetrical, where E and F are isotropic but D is uniaxially oriented. C and F have a continuous white phase from electrode to electrode, whereas D and E have both phases continuous.

thickness. It may be useful to go into more detail: whether the interface runs in preferential directions (which is relevant in connection with electric fields and drift of charge carriers), whether the interface has a well-defined curvature, etc. Of course, intermediate cases will be found in practice, such as island domains coexisting with channels across the film thickness. In the next section it will be demonstrated that such microstructural details play an important role to the efficiency of optoelectronic functionality.

III.

NOVEL MATERIALS FOR PHOTOVOLTAIC APPLICATIONS

A.

Small-Molecule Composites: Oligo(phenylenevinylene):C60 Blends

1.

Introduction

A novel concept for photovoltaic devices (phototdetectors [14–16] and solar cells [17–20]), based on interpenetrating blends of donors and acceptors sandwiched between two asymmetrical contacts (two metals with different work functions), has been proposed [17,18]. When the donor molecule is optically excited, a valence electron is promoted to the lowest unoccupied molecular orbital (LUMO) and, accompanied by a slight relaxation of the molecular structure, a bound electron–hole pair or ‘‘exciton’’ is created (Figure 2A). The presence of the acceptor facilitates dissociation of the exciton into free charge carriers by allowing the promoted electron to lower its energy by transfer to the LUMO of the acceptor (Figure 2B). Under the influence of the built-in electric field caused by the asymmetrical contacts,

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Figure 2 A) Creation of the bound electron–hole pair (exciton, designated with *) in an organic molecular semiconductor upon excitation by light and subsequent relaxation. B) Electron transfer between donor* and acceptor: a hole carrier (donor⫹) and an electron carrier (acceptor⫺) are created.

separation of the opposite charges takes place, with the holes being transported in the donor phase and the electrons in the acceptor phase. In this way, an interpenetrating blend can be considered as a network of donor– acceptor heterojunctions that allows efficient exciton dissociation and balanced bipolar transport throughout its whole volume. The performance of this type of device is very sensitive to the morphology of the blend since exciton dissociation as described above occurs only at the donor–acceptor interface. To ensure efficient exciton dissociation, an acceptor species should be within the exciton diffusion range from any donor species and vice versa. To put it differently, the characteristic size of the channels of the phases should be matched to this diffusion range. Moreover, both the donor and the acceptor phases should form a continuous network to allow bipolar charge transport. Since the exciton diffusion range is typically about 10 nm and hence shorter than the light absorption depth, a properly dimensioned network of donor–acceptor heterojunctions should be more efficient in terms of the exciton dissociation than a double-layer structure with a single, planar donor–acceptor heterojunction (Figure 3). However, such a system of spatially distributed donor–acceptor heterojunc-

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Figure 3 Thin-film configurations for photovoltaic devices (donor phase white, acceptor phase black). A) Double-layer film; excitons within hatched layer may reach the interface by diffusion. B) Film having interpenetrating phases.

tions may contain imperfections that degrade performance, such as fully dispersed domains and cul-de-sac-type discontinuities of the donor and/or acceptor phases. Furthermore, the extended interface area is accompanied by an enlarged phase boundary volume in which donor and acceptor species are molecularly mixed. This is likely to result in increased energy level disorder, causing an increase of the charge trap density and a reduction of the electron and/or hole mobility. Hence, though exciton dissociation may be enhanced, transport is impeded by the circuitous geometry of the interface. For these and other reasons, the overall device performance could be lower than expected. Obviously, a comparative study of both configurations requires the use of the same materials and identical preparation procedures. This requirement can be met by using materials suitable for vacuum deposition. 2.

Film Composition and Device Configurations

In this study, we investigate the morphology of blends and double layers of C60 and (E,E,E,E)-1,4-bis[(4-styryl)styryl]-2-methoxy-5-(2⬘-ethylhexyloxy)benzene (MEH-OPV5, Figure 4) and report on their photovoltaic characteristics [21]. Oligo(phenylenevinylene)s have been extensively studied as

Figure 4

The photovoltaic cell and its components.

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model compounds for the understanding of the optoelectronic properties of poly(phenylenevinylene)s and, more generally, conjugated polymers [22]. The efficient photoinduced electron transfer from conjugated polymers to C60 has been demonstrated [23]. C60 has a higher electron affinity and ionization potential than MEH-OPV5. The efficient quenching of photoluminescence in blends of MEH-OPV5 with C60 provides an indication for photoinduced charge transfer, with C60 acting as the electron acceptor. MEH-OPV5 was synthesized in a Wittig–Horner reaction using stilbenealdehyde and tetraethyl-2-methoxy-5-(2⬘-ethylhexyloxy)benzene-1,4bis(methylene)bisphosphonate according to literature procedures [24]. It was purified by column chromatography and subsequent recrystallization. 1HNMR and mass spectrometry showed no cis components or other impurities. C60 of purity greater than 99% was obtained from BuckyUSA (Houston, TX). The photovoltaic device (Figure 4) consists of a metal (Al) contact on the surface of an organic film on a partially coated indium–tin-oxide (ITO) glass. MEH-OPV5 and C60 were vacuum deposited on the ITO/glass substrate at a dynamic pressure below 10⫺6 mbar. The organic layer, with a typical total thickness of 80–100 nm, was prepared in two different ways: (a) codeposition, i.e., MEH-OPV5 and C60 are deposited simultaneously on an ITO substrate, and (b) consecutive deposition, i.e., the deposition of MEH-OPV5 on an ITO substrate is followed by deposition of C60. 3.

Photovoltaic Characteristics

For the photovoltaic characterization, an argon-ion laser (Spectra-Physics, BeamLok 2060) together with I–V source–measure unit (Keithley, Model 236) was employed. All measurements were done under dynamic vacuum better than 10⫺6 mbar. The current–voltage characteristics of ITO/codeposited MEH-OPV5: C60 (1:1)/Al devices in dark and under illumination, at ␭ = 458 nm with an intensity of 1 mW/cm2, are presented in Figure 5 (top). Forward bias is defined as positive voltage applied to the ITO electrode. It is known [25] that C60 can form ‘‘quasi-ohmic’’ contacts with Al and ITO electrodes. Therefore, the relatively high absolute value of the dark current, especially in reverse bias (approximately 10 ␮A/cm2 at ⫺1 V), is an indication of a continuous phase of C60 between the electrodes. We attribute the strong increase of the dark current in forward bias above 0.5 V to a contribution from a current through the MEH-OPV5 phase. This suggests that the MEHOPV5 phase also connects to both electrodes and hence that the blend is cocontinuous. Under illumination, the short-circuit current and the opencircuit voltage are about 0.1 mA/cm2 and 0.46 V, respectively. The photo-

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Figure 5 Characteristics of the ITO/codeposited MEH-OPV5:C60/Al device (top) and the ITO/consecutively deposited MEH-OPV5:C60/Al device (bottom). Left: I–V curves in the dark (solid circles) and under illumination at ␭ = 485 nm with an intensity of 1 mW/cm2 (open circles). Right: Dependence of photosensitivity (open squares) and open-circuit voltage (solid squares) on incident light intensity.

sensitivity of this device at zero bias is about 100 mA/W. The fill factor, a measure of the squareness of the I–V characteristic [26], is 0.3. On the righthand side in Figure 5, the variation of the photovoltaic sensitivity and the open-circuit voltage as a function of light intensity are shown. We observe saturation of the open-circuit voltage at about 0.5 V. This value is quite typical for the interpenetrating blends sandwiched between Al and ITO and corresponds to the difference of their work functions. The sensitivity shows only a slight decrease with increasing light intensity. These values are comparable to, or even slightly better than, those reported for organic donor– acceptor devices employing Al as the metal electrode [15–20].

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The current–voltage characteristics of ITO/consecutively deposited MEH-OPV5:C60/Al devices in the dark and under illumination, at ␭ = 458 nm with an intensity of 1 mW/cm2, are presented in Figure 5. Unlike the codeposited device, the consecutively evaporated one shows almost symmetrical I–V characteristics in the dark. In the latter device, MEH-OPV5 is not expected to form a continuous phase between the electrodes, and indeed, there is no sign of any significant contribution from it to the overall current in forward bias. The reason for the relatively high absolute magnitude of the dark current, which is only a few times smaller than that of the codeposited device, is not obvious and will be discussed later. Surprisingly, the photovoltaic characteristics of the consecutively deposited device at zero bias are similar to those of the codeposited one. It is all the more remarkable since, whereas luminescence from the latter device is completely quenched, the consecutively deposited device still shows fairly bright luminescence. For both devices the monochromatic power conversion efficiency is around 2%. The efficiency depends linearly on open-circuit voltage, short-circuit current, and fill factor. Therefore, a slight decrease in the photosensitivity of the consecutively deposited device is compensated by an increase in the open-circuit voltage and an improvement of the fill factor from 0.3 to 0.4. Note that we use Al as a metal electrode. Typically, the open-circuit voltage and the short-circuit current of organic photovoltaic cells are significantly better when Ca is used instead of Al [18]. Although a monochromatic power conversion efficiency of the order of a few percent has been reached, the spectral response of organic photodiodes must be matched to the solar spectrum to make them truly attractive for solar cell applications. 4.

Morphology and Its Role in Performance

To relate the device performances to the microstructure of the devices, the surface morphologies of vacuum-deposited C60, MEH-OPV5, and codeposited MEH-OPV5:C60 on a solid substrate (mica) have been studied by AFM. A Digital Nanoscope III with silicon cantilevers was used in tapping mode under ambient conditions to acquire the topography images shown in Figure 6. Deposited C60 forms continuous films of small aggregates. The deposition of MEH-OPV5 results in the formation of large islands, with an average height of around 30 nm. The dark domains correspond to the substrate surface. The codeposited film shows the best homogeneity and smallest roughness. The consecutively deposited film has a topography similar to that of deposited C60. The somewhat surprising properties of the consecutively deposited device can be explained on the basis of its morphology. According to Figure

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Figure 6 AFM images of (a) deposited MEH-OPV5, (c) deposited C60, (e) codeposited MEH-OPV5:C60, each on mica. Schematic illustrations of (b) deposited MEH-OPV5, (d) consecutively deposited MEH-OPV5:C60, and (f ) codeposited MEH-OPV5:C60.

6, the layer of MEH-OPV5, deposited first, forms islands on the substrate. By subsequently depositing C60 until it comprises the entire top surface, an asymmetrical device is created in which C60 makes contact with both electrodes but MEH-OPV5 does not. The continuous phase of C60 between the electrodes gives rise to the relatively high dark current. Because MEH-OPV5 is present as the initially formed islands, the interfacial donor–acceptor area is much larger than that of a planar double layer and is probably approaching (in terms of availability for dissociation of all light-created excitons) that of the codeposited structure. The interdigitated morphology of the consecutively deposited device can be viewed as intermediate between an interpenetrating blend and a truly planar double-layer film. This new morphology obviates some of the limitations to transport in interpenetrating networks that were discussed above, such as dispersed domains and dead ends. The increase in the fill factor of the consecutively deposited device can be attributed to a lower trap density and to the semi-double-layer morphology, which leads to a lower dark current. One may expect that a further reduction in the dark current will be obtained by disconnecting C60 from the ITO anode. Moreover, to ensure an efficient exciton dissociation, the characteristic radius of the islands should be of the order of the exciton diffusion length. The nonoptimized (too large) island radius of the consecutively deposited device, reflected in incomplete quenching of luminescence, may explain why its photosensitivity does not

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surpass that of the codeposited device. Hence, there are ample possibilities for the optimization of the interdigitated morphology, and we expect that the performance of an optimized structure of this type will be superior to that of an interpenetrating network. We have demonstrated that the active layer of a photovoltaic cell is a prime example of a multicomponent material that could benefit from morphological structuring. Our study suggests that an efficient photovoltaic composite should have the following feature: a bicontinuous structure with characteristic dimensions on the scale of the exciton diffusion length, ⬃10 nm, preferably of the interdigitated type. The required degree of control over the morphology is not easily achieved in blending of small molecules. As will be explained below, improved control over the microstructure may be exerted by employing the phenomenon of microphase separation in block copolymers. B.

Block Copolymers

1.

Phase Separation and Interfaces in Block Copolymers

In block copolymers, two (or more) dissimilar polymer chains (now referred to as ‘‘blocks’’) are chemically connected to form a single chain (see Figure 7 for a diblock copolymer). This connectivity has consequences for the thermodynamics of mixing and for the range of morphologies that may be formed. As in the case of blends, dissimilar blocks have a tendency to segregate into (nearly) pure domains of their own kind, unless favorable interactions between the different repeat units occur. However, segregation limits the number of available chain conformations since the junction point(s) on each polymer chain must reside in a relatively narrow region of space between the domains; it implies a distortion away from the random (Gaussian) conformation. Because of the entropy loss associated with this localization, demixing of copolymer blocks is less favorable than demixing of separate chains of the same length and type. For symmetrical A–B diblock copolymers as compared to blends, therefore, the critical point of the phase diagram is found at a higher value of ␹ N (Figure 8) [27]: (␹ N)diblock ⬇ 10

(3)

When crossing this point from below (increasing N, lowering temperature), a disorder–order transition takes place through spinodal decomposition. The higher the value of ␹ N, the better defined is the interface between domains of A and B. (The term ‘‘interphase’’ is often used to describe the layer between A and B domains where the majority of the A–B junction points on the chains is found.)

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Figure 7 Schematic representations of the molecular architecture and chain conformation of polyA-block-polyB, referred to as A–B diblock copolymer.

Figure 8 Phase diagram of A–B diblock copolymer [ f is the composition expressed as the fraction of B: f = NB/(NA ⫹ NB)], and associated two-phase morphologies. S, spheres; C, cylinders; L, lamellae; G, gyroid (bicontinuous) phase.

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Side-chain functional diblock copolymer poly(sPBD-b-sOPV3).

casting and annealed to remove solvent. Whereas a thin-film device of the blend shows efficient electroluminescence, the diblock copolymer film shows photovoltaic activity, its luminescence efficiency being two orders of magnitude lower than that of the blend. The architecture of the block copolymer forces the OPV3 substituents to be in close proximity to the PBD moieties, facilitating electron transfer. The polymer blend, on the other hand, will tend to exhibit macrophase separation, which reduces both the extent and the accessibility of the interface for exciton dissociation. The efficient electroluminescence of these particular blend films was attributed to a double-layer morphology, in which the PBD-containing phase functions as the electron-transporting layer. This layer formation was suggested by optical microscopy observations and is most likely induced by the different interfacial properties of the two electrodes and the spatial confinement of the film between these. In the block copolymer, on the other hand, microphase separation may provide separate pathways between the electrodes for each type of charge carrier. Charge separation rather than their recombination is promoted in such a morphology. The optoelectronic properties, including the I–V characteristics with a higher turn-on voltage for the blend as a result of the interface barrier (Figure 10), can be understood on the basis of the morphologies proposed here. Importantly, by using a diblock copolymer instead of a blend, we impose constraints on the phase separation and alter

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Figure 10 Top: Current-voltage (I–V ) characteristics of devices (ITO/polymer/Al) made with the diblock copolymer poly(sPBD-b-sOPV3) (continuous line), and the 2:1 blend of homopolymers poly(sPBD) and poly(sOPV3) (dashed line); polymer film thickness 100 nm. Bottom: Schematic morphologies of thin film devices of diblock copolymer and blend.

the morphology. As a result, we change the nature of the optoelectronic behavior from luminescent to photovoltaic. These results demonstrate the crucial role of the spatial distribution of the interface between the two active organic components in photonic behavior, and moreover they illustrate one of the ways in which the metal– organic interface comes into play. The photonic properties observed in our system resemble those reported by Friend and coworkers for composites of derivatives of poly(p-phenylenevinylene) (PPV). These authors studied a double-layer film formed by sequential spin-casting of PPV and cyano-PPV for LEDs [31] as well as films cast from a mixed solution of MEH-PPV and cyano-PPV for their photovoltaic behavior [17]. Phase separation was found to be in the 10- to 100-nm range in the latter system. The preliminary findings indicate the feasibility of our objective of designing and synthesizing improved materials with specific optoeletronic properties and to do so by controlling their microstructure through the phe-

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nomenon of (micro-)phase separation in block copolymers. The specific goal of the study described in this section is to enhance the photovoltaic properties of the PPV:C60 system by incorporating both active components in a self-structuring system 3.

Donor–Acceptor Diblock Copolymer for Photovoltaics

(a) Chain Architecture. Based on the earlier results and the general principles discussed above, we have designed and synthesized a self-structuring composite especially for photovoltaics. It is a diblock copolymer consisting of a PPV block and a block densely functionalized with C60 (Figure 11). Since PPV is a rigid-rod type of molecule, the second block should be a flexible one to enhance solubility. The synthesis of such a rod–coil polymer poses several challenges. To obtain a well-defined material that is not susceptible to large-scale demixing, good control over the length of both blocks is desirable. Furthermore, a sufficient amount of C60 should be incorporated to ensure the formation of a continuous path for electron transport throughout the material. The most prominent method that allows control of the degree of polymerization for block copolymers is living anionic polymerization [32]. This method yields products with very low polydispersities, but it has its drawbacks. One is the necessity of working under very strict conditions to avoid impurities such as water and oxygen. Moreover, the highly reactive carbanion prohibits the use of various functional groups [33]. Recently, this has been overcome to some extent by the development of the controlled/ ‘‘living’’ radical polymerization techniques [34] using either stable nitroxide

Figure 11

Molecular structure of the target compounds.

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counterradicals [35,36] (nitroxide-mediated living radical polymerization, NMRP) or atom transfer radical polymerization (ATRP) [37]. Both of these techniques employ the principle of an equilibrium between a low concentration of active radicals and a rather large number of dormant species. This suppresses bimolecular side reactions such as recombination or disproportionation to such a degree that the overall polymerization process shows living characteristics, e.g., the increase of the molecular weight with conversion and the possibility to synthesize block copolymers. We have employed the NMRP method for the synthesis of rod–coil block copolymers in which PPV forms the rigid block and the flexible block carries C60 as side groups [38]. NMRP does not interfere with the olefinic double bonds of the PPV part or the substituents on the rings. (b) Functionalization of Copolymers with C60. To be able to incorporate high amounts of C60, we decided to attach it as side groups and to employ radical addition of C60 [39] to poly(4-chloromethylstyrene) (PCMS). Radicals are generated at the chlorine sites in PCMS through an atom transfer mechanism [40]. However, the presence of two multifunctional components, PCMS and C60, makes the occurrence of cross-linking very likely. This was confirmed by experiments using PCMS and different amounts of C60; only insoluble products were obtained. We found that soluble C60-functionalized materials could be obtained by starting from statistical copolymers of styrene and 4-chloromethylstyrene (CMS), thus diluting the reactive groups with unreactive styrene moieties that act as spacers. To find the proper composition, random copolymers with different ratios of styrene to CMS (n = 1, 3, 5, with m = 1, x being in the range 10–20) were synthesized using NMRP with 2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO) according to the scheme in Figure 12. These polymers were subsequently functionalized with C60 and characterized by means of thermogravimetric analysis (TGA) and UV absorption spectroscopy to find the C60 content. Our results show that while the amount of C60 incorporated into the copolymer can be predetermined to a certain extent via the feed ratio of styrene to CMS, the length of the copolymer itself is controlled by the use of nitroxide-mediated free-radical polymerization. Other methods to incorporate fullerenes into copolymers shown so far in the literature (azide, cycloaddition, esterification) include multiple synthetic steps to obtain the functional monomers/prepolymers, whereas our approach involves only two steps (polymerization, radical addition) with commercially available compounds. (c) Synthesis of Rod–Coil Block Copolymers Based on PPV as the Rigid Rod. Most of the rod–coil block copolymers with conjugated polymers or oligomers as the rigid rod synthesized so far [41–47] were obtained

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by anionic polymerization of the coil polymer, followed by quenching, either with a reactive end-group on the rigid conjugated block or with a functional molecule to obtain an end-functionalized coil polymer that is subsequently coupled to the conjugated block. Our approach to the synthesis of rod–coil block copolymers is based on the use of an initiator for living polymerization that is attached to the rigid block. The synthetic routes employed are shown in Figure 13. For our block copolymers, poly(2,5-dioctyloxy-1,4-phenylenevinylene) was selected as the rigid part. Its synthesis was carried out via the Siegrist polycondensation as originally described by Kretzschmann et al. [48,49]. By this method, oligo(phenylenevinylene)s of narrow chain length distribution [49] that carry exactly one terminal aldehyde group per molecule are obtained. This is essential for the subsequent functionalization. Control of chain length can be exercised by varying the reaction temperature. At room temperature,

Figure 13 Synthesis of the alkoxyamine initiator, the aldehyde-functionalized PPV block, and their combination into the macroinitiator.

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the oligomers precipitate after an average of four addition steps. Temperatures around 80⬚C, on the other hand, may result in oligomer mixtures that reach or exceed the effective conjugation length of 12 repeating units [50]. An additional advantage is that this method yields only oligomers with transconfigured vinylic double bonds [51], rendering the conjugated block virtually defect-free. Attachment of the initiator to the rigid PPV block occurs via the nucleophilic attack of a Grignard reagent to the aldehyde group. An excess of the Grignard reagent guarantees complete functionalization. The alkoxyamine initiator containing a bromide group for the formation of the Grignard reagent can be easily obtained by the reaction of TEMPO with Br2, followed by the addition of styrene, as described by Kobatake et al. [52]. An average degree of polymerization of 10 repeat units was calculated for the macroinitiator PPV-TEMPO by means of 1H-NMR end-group analysis using the peaks of the methyl end-group at 2.2 ppm and the signals for the OCH2 groups at 4.0 ppm. On the other hand, the UV–Vis spectrum shows a maximum absorption at 467 nm, which corresponds to seven phenylenevinylene repeat units [50]. The comparison with the absorption spectrum of the seven-ring oligomer of Ref. 50 indicates the presence of additional higher oligomers in our PPV material. The calculations of molecular weights presented below are based on a value of seven repeat units (Mn = 2.5 ⫻ 103 g/mol) for the PPV part. The PPV-TEMPO macroinitiator has been successfully used in the polymerization of styrene, yielding the rod–coil polymer PPV-b-PS (Figure 14, top). The molecular weight increases with reaction time, as expected for living polymerizations [38]. Our objective being to bring PPV and C60 together in one molecule, we have synthesized a random copolymer of styrene and CMS from our PPV-based macroinitiator (Figure 14, bottom). Based on results from the series of statistical copolymers of styrene and chloromethylstyrene, a styrene/CMS feed ratio of 2:1 was chosen to ensure solubility of the C60-containing material. 1H-NMR analysis of the block copolymer indicated that the actual ratio of styrene to CMS is 1.5 and that the molecular weight of the styrenic block is 9 ⫻ 103 g/mol. Subsequently, this block copolymer was functionalized with C60. Figure 15 shows UV spectra and TGA of the starting PPV-b-PSCl and the C60containing block copolymer (PPV-b-PSFu). Whereas C60 is stable up to 550⬚C, the polymers start to decompose at much lower temperatures. Since this decomposition is almost complete at 550⬚C, it is safe to assume that the residue at this temperature corresponds to the C60 content. The increase in residue at 550⬚C from 14 wt % to 60 wt % upon functionalization of the rod–coil block copolymer with C60 is clearly visible in the TGA graph of Figure 15 (top). This indicates an average of around 19 fullerene molecules

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Figure 14 NMRP.

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Block copolymer synthesis using the PPV-TEMPO macroinitiator for

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Figure 15 Comparison of the starting diblock copolymer PPV-b-PSCl (---) and the C60-functionalized diblock copolymer PPV-b-PSFu (—). Top: TGA, heating rate 10⬚C/min. Bottom: Absorption spectra in CHCl3; the spectrum of the PPV-TEMPO macroinitiator is shown for comparison (⭈ ⭈ ⭈).

per chain, and means that two out of every three initial CMS groups has been functionalized with C60. UV–Vis spectrometry might be another convenient method to determine the amount of C60 in a copolymer. For this method to be quantitative, values of the molar extinction coefficient are required. However, the fine structure of the absorption spectra of well-defined derivatives of C60 (and hence their absorption coefficients at a given wavelength) were shown to be dependent on the substitution pattern as well as on the number of substituents [53]. The randomness in our copolymer will therefore render this method somewhat inaccurate. Still, the absorption spectra can be used to give a qualitative picture of the degree of functionalization. At the bottom of Figure 15, weight-normalized absorption spectra are shown. The absorption band of the PPV block is centered around 465

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nm. Comparison with the spectrum of the PPV macroinitiator suggests the weight ratio of the PS block to the PPV block to be around 3. C60 functionalization by means of radical addition little affects the absorption band of the PPV block, indicating that the conjugated block remains intact. When the spectrum of a substituted fullerene derivative of known C60 content (C122H88O15, [C60] = 40 wt %, Ref. 54) is used as the reference for a spectroscopic determination of the C60 content of PPV-b-PSFu, we find that the weight-normalized value of the absorbance at 330 nm corresponds to 46 wt %, or 14 C60 substituents per chain. (d) Properties of PPV–(C 60-Functionalized Polystyrene) Diblock Copolymer Films. The suitability of the type of block copolymer described above for application in photovoltaic devices relies on, among other things, the abilities of the blocks to function as electron donor and acceptor, respectively, and as charge transport media. As outlined earlier, the microstructure is a crucial factor to both these functions. In an AFM investigation of spin-cast films of our diblock copolymer, we have observed nanoscale features but no order (Fig. 16). Extended annealing at 160⬚C does not change the morphology. The driving force for microphase separation is apparently too small to give rise to sharp phase boundaries. This may be due to the chemical similarity of the blocks, which both have a high aromatic content. The charge transport properties of this material were found to be poor, which we attribute to insufficient continuity of at least one phase. Stronger segregation and higher order may be enforced by increasing the length of the

Figure 16 Left: TEM image of the PPV-b-PS diblock copolymer film. Right: AFM image (tapping mode) of the PPV-b-PSFu diblock copolymer film.

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flexible block or by replacing it by a nonaromatic polymer; these routes are the subjects of further investigations. Even though the phase separation into a bicontinuous morphology may be incomplete, the presence of donor and acceptor in one molecule ensures that they are in close proximity. The efficiency of the electron transfer between the excited donor (PPV) and the acceptor (C60) may be assessed from the extent to which luminescence from the donor block is suppressed (‘‘quenched’’). Figure 17 shows luminescence decay curves of spin-cast thin films of our block copolymers, both the rod–coil diblock (PPV-b-PSCl) and its C60-functionalized form (PPV-b-PSFu). While the former has a dominant photoluminescence decay time of around 250 ps, the decay time found for the donor–acceptor block copolymer is equal to or less than the instrumental response time of about 50 ps. The quenching indicates that electron transfer from an excited PPV block to a C60 substituent of the PS block occurs efficiently. This result is very preliminary since no control has been exerted over the microstructure. It is a promising starting point for further studies dealing with its optimization. (e) Micrometer-Scale Pattern Formation in Diblock Copolymers. Upon casting films from CS2 solution (a selective solvent for polystyrene), the formation of honeycomb patterns on a micrometer scale was observed for the PPV-b-PS copolymer as well as for PPV-b-PSFu (Figure 18) [55]. The diameter of the holes and their spacing is a few micrometers, three orders of magnitude higher than the characteristic size of the macromolecules and that of the mesophases in common microphase-separated block

Figure 17 Photoluminescence decay curves of spin-cast films of the rod–coil diblock copolymer PPV-b-PSCl and its functionalized derivative, the donor–acceptor diblock PPV-b-PSFu.

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Figure 18 Left: Scanning electron micrograph of a honeycomb structure in a solution-cast film of PPV-b-PS. Right: Luminescence micrographs of honeycomb structures in solution-cast films (scale bars: 20 ␮m). (A) PPV-b-PS; (B) PPV-b-PSFu; due to strong quenching of the luminescence from PPV by C60 in this copolymer, the intensity was found to be reduced by three orders of magnitude relative to that from a PPV-b-PS honeycomb.

copolymers! Moreover, hole formation is not observed in block copolymer microphase separation. Clearly, other physical phenomena are at play here. Similar hexagonal patterns have been previously observed and studied in detail in thin liquid layers that have a free upper surface and are heated from below [56]. In the polymer field, this phenomenon has been reported by Franc¸ois and coworkers [57,58] for poly(p-phenylene)–polystyrene block copolymers and by Jenekhe and Chen [59] for poly(phenylquinoline)– polystyrene, among others. Whereas honeycomb material obtained from PPV-b-PS shows a strong yellow–orange luminescence, the luminescence intensity is very low in PPV-b-PSFu. Luminescence is quenched to such an extent by the C60 side groups that to obtain images of similar brightness an exposure time approximately 1000 times longer than that for PPV-b-PS is necessary for PPV-b-PSFu. The honeycomb material discussed here may be useful for photovoltaic devices, since the porous nature and the increased surface area reduce the reflection of light at the face of a film and hence increase energy harvesting. The regularity of the structure obtained may be exploited to fabricate optical band gap materials or patterned LEDs, for instance. The use of the material as a mask for the deposition of a hexagonal array of aluminum dots has been demonstrated [55]. A fundamental understanding of the physicochemical principles underlying the pattern formation is still lacking and urgently needed for engineering and controlled fabrication.

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IV.

ELECTRONIC ENERGY LEVEL ALIGNMENT AT THE OLIGO(PHENYLENEVINYLENE)–METAL INTERFACE

A.

Introduction

In the previous sections we paid attention to two main issues: (a) finding an optimum microstructure for a given photonic functionality; (b) designing polymer molecules that combine functionalities and are designed for selforganization into a favorable microstructure. In addition to these efforts toward obtaining control over the creation of interfaces, a more detailed study of each type of interface encountered in a device configuration is required for the understanding and, ultimately, control of charge transport. To meet these objectives, the structural and electronic signatures of each interface must be explored. The study described here demonstrates that the formation of an interface between conjugated organic materials (oligomers) and metals affects the electronic properties of the materials involved. The charge transfer process at the interface between the electrode contact and the active layer is a function of the barrier height between the Fermi energy level of the electrode and the frontier orbital levels of the organic layer. In order to estimate the injection barrier one usually assumes a rigid band model and an alignment of the vacuum levels at the interface. However, these assumptions ignore the possible formation of dipole layers or covalent bonds at the interface and doping of the interfacial region. Although measurements of the vacuum level alignment are important for understanding device behavior and optimizing device efficiency, they have seldom been reported for the interface between metals and phenylenevinylene oligomers [60–62]. Here we report ultraviolet photoelectron spectroscopy (UPS) measurements on the interfaces between metals (Au and Ag) and two different five-ring phenylenevinylene oligomers (OPV5s), one without substituents (referred to as P5V4) and one that is 2-methoxy-5-(2⬘-ethylhexyloxy)-substituted (MEH-OPV5) (see insets, Figures 19 and 20, respectively) [63]. These oligomers can be used for optoelectronic applications and as model systems for the two most widely used corresponding conjugated polymers. The advantage of using these oligomers instead of the polymers is that their high chemical purity and relatively low mass enables us to prepare thin films in situ by deposition of the oligomer on a cleaned metal surface under UHV conditions. For each combination of metal and oligomer we have measured a misalignment between the vacuum levels of the two. The levels of the oligomers shift downward relative to the energy levels of the metal (a negative shift). The magnitude of this shift is between 0.4 and 1.2 eV, depending on the metal as well as on the oligomer. The rearrangement of energy levels

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Figure 20 UPS spectrum of MEH-OPV5 on Ag as a function of the deposition time (thickness). All spectra were recorded with a bias of ⫺4.0 V between sample and analyzer. The x axis represents the kinetic energy of the photoelectrons at the entrance slit of the analyzer. The dotted line is the spectrum from the polycrystalline Ag substrate. The plots in the right- and left-hand panels show enlargements of the central plot. After recording a spectrum, the thickness of the MEH-OPV5 layer was increased by approximately 0.5–1 nm, except for the last two spectra where the thickness was increased by several nanometers. The inset in the middle plot shows the structure of MEH-OPV5.

21.22 eV) and an effusion cell of our own design. A hemispherical analyzer in combination with a lens system was used as electron energy analyzer with an overall resolution of 150 meV. During measurement, a bias voltage of ⫺4.00 V was applied to the sample holder to clear the low kinetic energy cutoff of the spectrum. The metal substrates were cleaned by polishing and washing in an ultrasonic bath with toluene and acetone as solvents. Next, the substrates were placed in a sample holder on an x, y, z stage. Under UHV conditions, the substrates were Ar⫹ ion-sputtered to obtain UPS spectra without features from impurities or adsorbed overlayers. The oligomer–metal interfaces were formed by slow molecular beam deposition monitored by a calibrated thickness monitor, after degassing the effusion cell for several hours at 120⬚C. Typical sublimation temperatures range from 225⬚C to 275⬚C, and the pressure ranges from 0.5 ⫻ 10⫺7 to 2 ⫻ 10⫺7 mbar, giving a deposition rate of several monolayers per minute. The oligomer film thickness used in determining the energy level offset was typically 5–10 nm.

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145

Energy Levels at the Interface

Figure 19 illustrates how an interfacial energy diagram, including the presence of an electric field, is deduced from two UPS spectra. The figure shows the He I UPS spectra of Au (left) and P5V4 on Au (right). Both spectra were recorded while a bias voltage of ⫺4.00 V between sample and analyzer was applied. The energy diagram of the interface is shown in the middle of Figure 19. The y axis represents the kinetic energy of the photoelectrons at the entrance slit of the analyzer. The total spectrum consists of photoelectrons suffering no energy loss processes on their way out of the solid, providing information on the binding energies of the occupied states of the solid, and scattered electrons that have suffered energy loss, which contributes to a structured background. The high kinetic energy onset of the spectra corresponds to emission of electrons from the Fermi level of the metal [E max k (Au)] or from the highest occupied molecular orbital (HOMO) of the oligomer [E max The low kinetic energy cutoff [E min k (P5V4)]. k (Au), min E k (P5V4)] corresponds to electrons having suffered energy loss processes and ended up at the minimum kinetic energy in the solid required to escape into the vacuum. Thus, the position of the vacuum level of a material in the diagram is determined by adding 21.2 eV to the cutoff energies to give Evac(Au) and Evac(P5V4), respectively. Now one can determine the work function of Au using ⌽Au = h␯ ⫺ [E kmax(Au) ⫺ E kmin(Au)]

(4)

and the position of the HOMO level of P5V4 according to: min Is = h␯ ⫺ [E max k (P5V4) ⫺ E k (P5V4)]

(5)

By comparison of the two UPS spectra, we can construct an energy diagram as shown in the middle of the plot and determine the positions of the energy levels at both sides of the interface. We can determine the difference (⌬) between the vacuum levels of the metal and the oligomer layer using min ⌬ = E min k (P5V4) ⫺ E k (Au)

(6)

Since the vacuum level of the oligomer is lower than the vacuum level of Au, the electric field points from the oligomer (␦⫹) to the metal (␦⫺), making ⌬ < 0 (as defined in Ref. 62). In this case, the vacuum level shift leads to an increase of the barrier for hole injection (␧ Fv ) according to:

␧⬘vF = Is ⫺ ⌽Au ⫺ ⌬

(7)

This hinders hole injection from Au to P5V4 but facilitates electron injection from the Fermi level of Au to the LUMO.

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Dependence on Organic Layer Thickness

In Figure 20 we show the interface formation of MEH-OPV5 on a clean Ag substrate. In the center, the full spectra are given. The top line, with the largest offset, shows the Ag spectrum with the Fermi edge well resolved (at h␯ ⫺ ⌽Ag ⫹ Vbias = 20.6 eV). The bottom spectrum shows the UPS spectrum of the Ag substrate with a thin layer of MEH-OPV5 evaporated on top (in the order of 1 nm). The subsequent spectra show the photoemission of the surface with increasing MEH-OPV5 thickness (0.5–1 nm increase per deposition step). At the right-hand side, the gradual change in the HOMO region is depicted in more detail. As the film grows thicker, the features arising from the Fermi level of Ag give way to the pure MEH-OPV5 spectrum. The low kinetic energy cutoff region is plotted at the left-hand side of Figure 20. The shift in the onset (⌬) is complete after approximately 1 nm has been deposited on the substrate; it does not increase further with film thickness. This result indicates that one or very few molecular layers are involved in the process. The upper two spectra of MEH-OPV5 show a small shift of 0.1 eV toward lower kinetic energy. A slight charging of the sample may cause this shift, which is within our measurement accuracy.

E.

Consequences of a Dipole Field at the Interface for Charge Injection

The lowering of the vacuum level is caused by an electric field at the interface due to (partial) charge transfer in the interfacial region. The exact mechanism creating the electric field is unknown. Several processes can cause an electric field at an interface, e.g., electron transfer from a donor to an acceptor, image effects, tailing of the electron cloud of the metal towards the adsorbent causing dipoles or metal-induced gap states, or chemical interactions between metal and overlayer [65–67]. Charge transfer at the interface formed by similar PPV oligomers on a Ca substrate was observed by Park et al. [60,61]. The charge transfer caused an overall energy level bending of 0.5 eV in a space-charge region extending over 10 nm. The energy level bending in their experiment indicates electron transfer from the metal to the organic overlayer making ⌬ > 0, which is opposite to the values reported here. Image potential effects lower the ionization energy of insulators and semiconductors adsorbed on metal surfaces. This stabilizing effect is linearly proportional to 1/d, where d is the distance between the metal surface and the photoionized molecule [68,69]. In this study, such a distance dependence is not reflected in a dependence on the film thickness (see Figure 20), indicating that the image potential effect is either small for even the thinnest

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layer measured or compensated for by another electrostatic effect. Chemical interactions at the interface formed by evaporating Al on a sexithiophene layer have been reported [62] with a significant charge transfer from the metal to the oligomer, resulting in the formation of a bond between the oligomer and the metal. Such a charge transfer causes an electric field pointing from the metal toward the organic layer (⌬ > 0). For most metal–organic interfaces the dipole is negative (see, for example, figure 15 in Ref. 65), similar to our measurements. Table 1 summarizes the results of the UPS measurements. The shift between vacuum levels at the interface is negative in all cases, moving the energy levels of the oligomer downward relative to the levels of the metal. The magnitude of the shift is larger for Au than for Ag and also depends on the oligomer: the substituted oligomer causes a larger shift. The ionization energy seems to be independent of the metal substrate and compares well with previously reported values for similar OPV5s [70]. The hole injection barrier can be obtained in two ways: by measuring the energy difference between the high kinetic energy onset of the metal and that of the oligomer, as shown in Figure 19 (␧ Fv ), or by using Eq. (7) (␧ ⬘vF). The differences between the values for the same injection barrier (see fourth and fifth columns in Table 1) give an indication of the systematic error. The values clearly show that the misalignment of the vacuum levels strongly influences the hole injection barrier and acts as to keep the barrier almost independent of the work function of the metal substrate. Our UPS measurements of interfaces formed by evaporating PPV-type oligomers onto metals (Ag and Au) invariably reveal a misalignment between the vacuum levels of the metal and the organic overlayer. This shift of levels, presumably caused by an interfacial dipole layer, strongly influences the hole injection barrier in such a way as to keep this barrier nearly

Table 1 UPS measurements (eV) on P5V4 and MEH-OPV5 (⌽Au = 5.1 ⫾ 0.1 eV and ⌽Ag = 4.4 ⫾ 0.1 eV)a Interface P5V4/Au P5V4/Ag MEH-OPV5/Au MEH-OPV5/Ag

⌬

Is

␧ Fv

␧ ⬘v F

F ␧ v(⌬=0)

⫺1.0 ⫺0.4 ⫺1.2 ⫺0.5

5.6 5.6 5.2 5.3

1.4 1.7 1.2 1.3

1.5 1.6 1.3 1.4

0.4 1.3 0.0 0.8

⌬ represents the vacuum level shift (⫾0.1 eV); Is, ionization energy (⫾0.1 eV); ␧ Fv , measured hole injection barrier (⫾0.1 eV); ␧ v⬘F, calculated hole injection barrier using Eq. (7) (⫾0.3 eV); ␧ Fv(⌬=0), theoretical hole injection barrier for ⌬ = 0 (⫾0.1 eV).

a

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6 Studies of Organic Thin Films and Interfaces by Various Electron Spectroscopies Molecular Orientation, Electronic States, and Reactions Nobuo Ueno, Koji Kamiya Okudaira, and Satoshi Kera Chiba University, Chiba, Japan

Yasushi Azuma National Metrology Institute of Japan, Tsukuba, Japan

Shinji Hasegawa Institute for Molecular Science, Okazaki National Research Institute, Okazaki, Japan

I.

INTRODUCTION

There has been much interest in electronically functional organic materials with respect to various applications. The interest has been accelerated recently because of successful realization of organic electroluminescent (EL) devices that have given a good prospect for new functional device fabrication with organic molecules [1–4]. Organic devices involve various types of interfaces consisting of organic molecules, inorganic semiconductors, and metals, and their properties are largely influenced by electronic properties of interface region [4]. In Figure 1(a, b) basic principles of the organic EL and solar devices are illustrated for ease of understanding the importance of the interface. However, our knowledge on what is happening at the interface is still very limited even for simplified model interfaces. It is commonly recognized that electronic energy levels of two materials that form the interface should be known for the first step in order to understand the interface properties. Therefore, many studies have been done 153

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Figure 1 Schematic energy diagrams of organic electronic devices using interface electronic property (a, b) and schematic reaction between molecular orientation and overlap of wave functions at the interface (c1, c2). (a) EL device. (b) Organic solar cell. (c1) Wave function overlap for a flat-lying molecule. (c2) Wave function overlap for a largely tiled molecule.

concerning the determination of lowest ionization energies of organic thin films and the energy level alignment at interfaces [2,4–14]. However, it should be remembered that the charge transfer property through the interface is affected not only by the alignment of two energy levels concerned but also by the overlap of wave functions of corresponding energy levels. The spatial spread of wave functions is in general not isotropic, and therefore

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155

the overlap depends on the orientation of molecules in the interface region. In other words, if the wave function overlap is very small, effective charge transfer can not be expected between the two energy levels, even when the energy levels of two interface materials are located relevantly. Therefore information on the molecular orientation [2,15,16], the spatial distribution of the wavefunction and its character are of crucial importance when one wants to understand deeply the charge transfer properties through the interface. For the interface consisting of organic molecules, the important valence state, namely, the highest occupied molecular orbital (HOMO) state, in general consists of ␲ orbital derived from 2p atomic orbitals (AOs). For a planner molecule with ␲ electrons, the flat-lying orientation on a substrate will give more effective interaction with the substrate than upright orientation, since the ␲ orbital spreads normal to the molecular plane. These are schematically illustrated in Figure 1 (C1, C2). On the other hand, the location of the HOMO level depends on an electric dipole layer formed at the interface and degree of the charge separation near the interface. Furthermore, it is also affected by a change of the intermolecular interaction, namely, by the change of the molecular orientation. The change of the molecular orientation thus produces a band bending or energy level shift that plays an important role in the actual device properties. Although the shift of Fermi level (EF) from the center of band gap in an organic film due to introduction to low-density gap states can also produce the energy level shift as observed in long-chain alkane films [17], this type of energy level shift is still under investigation for organic solids. In this chapter we describe our recent results on three-dimensional molecular orientation, electronic states, and their wave function studied by various surface-sensitive electron spectroscopies. Furthermore, we will show an interesting thin-film growth observed for ClAl-phthalocyanine, and an example that clarifies the origin of a new electronic state produced by the chemical reaction by organic–metal contact. The chemical reaction at the interface would become more important in the near future, since new materials that are difficult to obtain by conventional synthetic methods can be obtained just by the contact of two materials under ultrahigh vacuum (UHV) conditions. Organic thin films discussed in this chapter involve a polymer thin film as well as epitaxial thin films of functional organic molecules. Our experimental methods described herein are angle-resolved ultraviolet photoelectron spectroscopy using synchrotron radiation (SR-ARUPS), metastable atom electron spectroscopy (MAES), low-energy electron transmission spectroscopy (LEET), high-resolution electron energy loss spectroscopy (HREELS), and low-energy electron diffraction (LEED) etc.

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EXPERIMENTAL METHODS AND THEIR CHARACTERISTICS

Here main experimental methods and their characteristics are briefly overviewed together with information obtained by them.

A.

Angle-Resolved UPS Using Synchrotron Radiation

ARUPS is probably the most powerful single experimental technique for studying the valence bands of solids [18,19]. By the use of linear polarization or high-energy tunability of synchrotron radiation (SR) in ARUPS measurements, full information on the valence band structure and geometrical structure near the surfaces can be obtained. However, until recently, the use of SR-ARUPS has been primarily limited to investigating the energy positions of valence bands along a few high-symmetry directions in the reciprocal space. This means that the intensity of SR-ARUPS features has in general not been discussed quantitatively. The main reason for this limited use of SR-ARUPS is the lack of any simple rules for explaining the intensity of such valence band spectra. Although Solterbeck et al. [20] recently performed the one-step calculation for the angular distribution of photoelectrons from GaAs (110) valence bands, such calculation is very complex and difficult to apply to various systems. Therefore, many SR-ARUPS results have been discussed with symmetry consideration [19,21–23] and the aid of calculations using a plane-wave final state because of its simplicity [24–27], although such simplified calculations are inevitably not very accurate [27]. We found that independent atomic center (IAC) approximation [27–31] and single scattering approximation (SS) [27,32–34] combined with molecular orbital (MO) calculation are very useful to explain the observed SR-ARUPS intensity for organic thin films, ranging from well-ordered thin films [27,30– 36] to disordered polymer systems [37–39]. The IAC/MO and SS/MO methods have been effectively used to analyze the photoelectron angular distribution from organic systems. As shown in Figure 2, three parameters to specify the molecular orientation are (a) the inclination angle ␤ at which the molecular plane is inclined to the substrate surface, (b) the molecular azimuthal angle ␾m , which specifies the rotational orientation of the molecule with respect to the direction normal to the molecular plane, and (c) the substrate azimuthal angle ␾s around the substrate normal. The substrate azimuthal angle ␾s corresponds to the experimental ␾. The theoretical formulae of the IAC/MO and SS/MO approximations and their plausibility were discussed in Refs. 33, 34, and 35. In the SS/MO

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Figure 2 Illustration of measurement parameters of ARUPS using synchrotron radiation (panel a) and parameters specifying the molecular orientation (panel b) adsorbed on a single crystal substrate. In panel a, the electric field vector E of h␯ and momentum vector of photoelectron Pe are on a plane perpendicular to the substrate surface. In panel b, the aximuthal angle ␾s around the surface normal (n) is measured from the plane of E and Pe . ␾m specifies the rotational orientation of a molecule around the molecular plane normal.

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approximation, the photoelectron intensity In(R) from the nth MO at the detector R is represented by In (R) ⬀ 兩A ntot (R)兩2,

冘冘 冘冘冘

n A tot (R) ⬀

a

⫹

Da C nXa e⫺ikn ⭈ Ra

Xa

冘 冘冘

ˆ LXa Y*(R)M L

L

n Db C Xa e⫺ikn ⭈ Rb ⫻

a

b≠a

Xa

L

ˆ l⬘ Y* L⬘ (R)t b (kn)

L⬘

GL⬘L(Rb ⫺ Ra)MLXa

(1)

where Da is the phenomenological damping factor for the photoelectron wave from atom a, C nXa is the nth MO coefficient of the Slater-type atomic orbital (AO) Xa, kn (= knR) is the wave vector of photoelectron, and Ra is the position of atom a. MLXa represents the matrix element including the phase shift and radial integral. t lb (kn) and GL⬘L (Rb ⫺ Ra) are the single-scattering vertex and the free propagator, respectively. The first term of Eq. (1) corresponds to the IAC/MO approximation that reflects multicentered photoemission waves scattered only by each atom from which the photoelectron is ejected, and the second term represents the waves scattered singly by all other atoms. For ease of understanding Eq. (1), the IAC waves (direct waves with phase shift due to potential scattering at each emitting cite) and selected two waves involved in the SS waves are schematically shown in Figure 3 for the photoemission from a model triatomic molecule. Here the wave emitted from each atomic site changes its phase by the attractive force field and these waves interfere with each other to give the final photoelectron wave. In the case of ‘‘large’’ organic molecules consisting of light elements, the direct IAC waves are stronger than the SS waves and the contribution of the SS waves usually modifies the IAC angular pattern slightly to give a rather futureless one with increase in the number of atoms in the molecule. Thus, the interference of the IAC waves mainly dominates the observed photoelectron angular distribution depending on relative positions of emitting sites. In any case, due to the interference, the angular distribution is very sensitive to the wavefunction of the initial state of photoemission, molecular orientation, and intermolecular geometry, and therefore we can obtain information on these by analyzing the observed angular distribution. B.

Metastable Atom Electron Spectroscopy

In MAES, the energy of electrons ejected from a target by the impact of metastable atoms is analyzed [40–42]. MAES selectively detects the outermost surface layer because metastable atoms do not penetrate the inner layers in contrast to the case of other surface-sensitive methods that use

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Figure 3 Schematic illustration of the IAC and SS waves in Eq. (1). The direct waves and scattered waves in the figure correspond to the IAC and SS waves, respectively. Only two scattered (SS) waves are shown.

photons or electrons as the probe. Therefore, MAES can be the most surfacesensitive electron spectroscopy, and electrons with wavefunctions extending outside the surface interact more effectively with the metastable atoms than those inside and give a stronger spectral band in MAES [41–43]. However, as shown in Figure 4, there are two main deexcitation mechanisms for the collison with solid surfaces: (a) Penning ionization (PI) or Auger deexcitation (AD), and (b) resonance tunneling followed by Auger neutralization (RI ⫹ AN) [40–42]. The former takes place for insulators, wide-gap semiconductors, and molecule-adsorbed surfaces, and provides a spectrum like UPS. The latter occurs for metals and narrow-gap semiconductors, and gives a featureless spectrum due to convolution of the electron energy distribution curves (EDCs) by many Auger neutralization paths [40–42]. When the PI process dominates the deexcitation of metastable atoms, MAES directly reflects the electronic structure of the outermost surface of

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Figure 4 Deexcitation mechanisms of metastable atoms at a solid surface. (a) Penning ionization (PI) or Auger deexcitation (AD) on an insulator surface, where the kinetic energy of the ejected electrons Ek is given as Ek ⬇ ⌬E(A*) ⫺ Eb. (b1, b2) Resonance ionization (RI) followed by Auger neutralization (AN) on a metal surface, where Ek ⬇ E⬘(A⫹) ⫺ 2(␰ ⫹ ␾).

the target solid and is often called Penning ionization electron spectroscopy (PIES). Using this characteristic, the orientation of the surface molecules was determined for organic thin films from the analysis of the relative band intensity of MAES as illustrated in Figure 5 [41–46]. An excellent advantage of this method is that qualitative molecular orientation can be obtained very easily just by comparing intensities of corresponding features in MAES and UPS. Furthermore, a monolayer formation of molecules on a substrate can be uniquely confirmed by measuring the deposition-amount dependence of MAES. That is, by a formation of a complete monolayer, the substrate surface is fully shielded for the metastable-substrate interaction to give little signal from the substrate in MAES [41–45]. In addition to the potential use of this method to investigate the molecular orientation and the ‘‘size’’ of wavefunction at the surface of an organic film as shown in Figure 5, it can also be useful for selective detection of surface electronic density of states (DOS) of an inorganic semiconductor, when the PI process dominates in the deexcitation of metastable atoms [47]. Another interesting and important characteristic of using metastable atom beam is that it can induce the outermost surface reaction selectively,

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Figure 5 Schematic illustration of Penning ionization electron spectra. The metastable atom of He (He*) ejects the electron at the outermost surface. When the orbital ␾1 is exposed outside, the intensity of electron from ␾1 state is observed more strongly than electron from orbital ␾2 (left). When the molecular orientation is reverted (right), electron from ␾2 is observed more strongly than electron from ␾1. Furthermore, the observed electron intensity reflects the degree of outspread of the wave function or the size of the wave function.

which cannot be realized by photon and electron [48]. Although we will not show the example of the surface selective reaction in this chapter, we point out that the outermost surface reaction induced by the metastable atom beam will be useful in preparing a molecular heterolayer using a single type of molecule by repeating the deposition/irradiation/deposition cycle. C.

High-Resolution Electron Energy Loss Spectroscopy

In HREELS [49,50], a monochromatic electron beam irradiates surfaces and scattered electrons are energy-analyzed to obtain high-resolution spectra of electron energy lost by excitations of surface phonons, vibrations, and electrons. Among these, measurements of vibrational excitations of organic surfaces offer information on chemical bonding and orientation of the bonds at surfaces. In general, the vibrational excitations mainly occur by two different scattering mechanisms in HREELS: dipole and impact scattering mecha-

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nisms [49–51]. These mechanisms can be distinguished by means of angular distribution of inelastically scattering electrons. The dipolar loss is strongly localized around the specular direction, whereas the impact scattering gives broad angular distribution [51]. In the dipolar excitation, the excitation probability (energy loss intensity) is enhanced for the dynamic dipole moment being perpendicular to the surface and is diminished when it is parallel to the surface due to the image dipole (Figure 6). Therefore, when the structure of target molecules is known, the molecular orientation can be estimated even for very thin films. Another useful characteristic of HREELS is that the scattering angle (␪s) dependence of the elastic electron reflects the roughness of the surface. For a very flat surface, the ␪s dependence of the elastic-electron intensity is very sharp peaked at the specular angle, whereas for an uneven surface a broad ␪s dependence is observed as illustrated in Figure 6. D.

Low-Energy Electron Transmission Spectroscopy

In LEET measurements, a monoenergetic electron beam impinges normally on a thin film deposited on an electrically conductive substrate, and the electron current transmitted through the film (It) is measured at the substrate as a function of the incident electron energy (Ei) (Figure 7). It uses an important characteristic of electron beam, namely, the high tunability of electron beam energy. From this characteristic, the electron beam is often called the ‘‘poor man’s synchrotron radiation,’’ and LEET is probably the simplest and cheapest form of electron spectroscopy. In LEET, the dIt (Ei)/dEi or ⫺d 2It (Ei)/dE2i curves are recorded as well as It (Ei), since features in It (Ei) curves are generally weak. LEET spectra provide unique information on the excess electron states (conduction bands) in dielectric thin films and the work function change of the surface [52,53]. Although complete assignments of spectral features appearing in LEET spectra have not yet been realized, the features mainly reflect the ‘‘conduction bands’’ above the vacuum level for thick films [52–59]. For example, the features correspond well with the conduction band DOS measured by secondary electron emission spectroscopy (SEES) [59]. Despite the extremely simple experimental setup, LEET spectra are very sensitive to the film thickness and the geometrical arrangement of molecules in the film [52– 55,59–62] because the excess electron state is basically determined by an interference of the electron which includes reflections at the film–vacuum and film–substrate interfaces, and periodic potential in the film as shown in Figure 7. Thus, LEET spectra involve information on the energy positions of critical points in the conduction bands [52,62,63], the effective mass of hot electron in the film, the inner potential, and the molecular orientation

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˚ ) and H2Pc (4 and 84 A ˚ ) films Figure 8 SR-ARUPS of CuPc films (3 and 50 A on MoS2 substrate at h␯ = 40 eV and ␣ = 0⬚. SR-ARUPS of CuPc films are observed at ␪ = 35⬚ and ␾ = ⫺38⬚. SR-ARUPS of H2Pc films are observed at ␪ = 33⬚ and ␾ = ⫺17⬚. The vertical bars shows calculated molecular orbital energies. The calculated density of states (DOS) by a Gaussian broadening (0.4 eV) of the molecular energy levels is included. Peak A represents the top band of Pc films. The shoulder indicated by arrows originates from MoS2 substrate. (From Ref. 35.)

calculation are shown. These calculated results were basically similar to those for H2Pc in the low binding energy region. In the low binding energy region, the SR-ARUPS of the 1-ML CuPc film is almost the same as that of the 1-ML H2Pc. The SR-ARUPS of the thick film is also similar to those of 1-ML film for both CuPc and H2Pc, except that the peak position is shifted by about 0.3 eV to the higher binding energy side. The binding energy shifts at the monolayer with respect to EF indicate that the band bending exist near the interface for both the films.

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Furthermore, it is interesting to see the full width at half maximum (FWHM) of band A in the 1-ML film and thick film. Band A is ascribed to the single HOMO (␲) band, which is distributed in a porphin-like ring in Pc molecule. We expect a much narrower bandwidth for the 1-ML film than the thick film, since the photoelectron bandwidth of a molecular solid has been considered to be due to the depth dependence of the relaxation energy [65,66]. However, the band A in the 1-ML H2Pc film is sharper only slightly than that in the thick film, and for CuPc film the sharpening of the band A is scarcely seen for the 1-ML film. The large bandwidth observed in the 1-ML film suggests that a strong intermolecular interaction may exist along the lateral direction, since it is difficult to consider that the bandwidth in the monolayer is due to a very short lifetime of the hole by the electron transfer from the substrate. ˚ -thick The ␪ and ␾ dependencies of SR-ARUPS spectra of the 3-A (about 1 ML) CuPc at the binding energy region of about 0–5 eV are shown in Figure 9. The intensity of the HOMO band A depends on ␪ and ␾. The ␪ and ␾ dependencies of the HOMO band intensities of the 1-ML CuPc films are plotted in Figure 10(a, b), where the ␾ dependence is shown after rotational shifts of 120⬚ and 240⬚ due to the six-hold symmetry of the MoS2

Figure 9 Examples of the take-off angle (␪) (a) and azimuthal angle (␾) (b) dependencies of the SR-ARUPS spectra measured at h␯ = 40 eV and ␣ = 0⬚ for CuPc ˚ ) on MoS2. The ␪ and ␾ dependencies were measured at ␾ = 8⬚ and ␪ = film (3 A 35⬚, respectively. Peak A represents the top band. Arrows show the shoulder originating from the top band of MoS2 substrate. The intensity is normalized to the incident photon flux. (From Ref. 35.)

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Figure 10 The take-off angle (␪) (at ␾ = 54⬚) and azimuthal angle (␾) (at ␪ = 36⬚) dependencies of photoelectron intensity of the highest occupied band of CuPc ˚ ) on MoS2 measured at h␯ = 40 eV and ␣ = 0⬚, and comparison monolayer films (3 A with calculated ones. The observed results are shown by 䡩. The calculated angular dependencies by SS/MO approximation at h␯ = 40 eV and ␣ = 0⬚ are shown for ␤ = 0⬚ ( ), ␤ = 6⬚ (– – –), and ␤ = 15⬚ (– ⭈ – ⭈ –). The calculated ␪ dependencies in panel a and ␾ dependencies in panels b and c are obtained by assuming that the three domain sizes are the same (see Figure 11). In panel d the ␾ dependence is calculated by use of the different domain sizes, S[A]/S[B]/S[C] = 1:1.5:1, where S[A] denotes the size of domain A, which is one of three equivalent domains (see Figure 11). (From Ref. 35.)

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surface. The ␪ dependence has a sharp maximum at ␪ = ␪max = 36⬚, and ␾ dependence shows a petal-like structure reflecting molecular structure and triple azimuthal orientation due to the rotational symmetry of the substrate. A typical LEED pattern and two-dimensional lattice structure of the CuPc on MoS2 are shown in Figure 11, indicating that there are triply oriented CuPc lattices, [A], [B], and [C]. In order to analyze the origin of the ␪ and ␾ dependencies, calculations of photoelectron angular distributions were fully performed for the experimental angular conditions used in the measurements. The theoretical ␪ dependencies were obtained by averaging those for the triply oriented lattices, [A], [B], and [C], shown in Figure 11, and including molecular rotational orientation at the lattice points, which were determined by the analyses of the observed ␾ pattern as described later. Since the scattered waves from neighboring molecules do not seriously affect final angular distribution [34], only single-scattering waves from all atoms constituting the molecule that emits the photoelectron were considered in the calculation. The calculated ␪ dependencies for the 1-ML CuPc film are compared with those observed in Figure 10a, where the calculated results using the SS/MO approximation are shown for ␤ = 0⬚, 6⬚, and 15⬚. It is seen that slight disagreement exists between the observed and calculated ␪ dependencies for ␤ = 0⬚, and the calculated ␪ dependence at ␤ = 6⬚ seems to be in better agreement with the observed one. Azuma et al. [67] observed that the ␪ dependence for ClAlPc on MoS2 is highly temperature dependent probably due to molecular vibrations. They showed that the ␪ dependence becomes sharper and the ␪max position shifts to the low-␪ side by decreasing

˚ ) on MoS2 at Ei = 30 eV. (b) Figure 11 (a) LEED pattern of CuPc film (6 A Theoretical LEED pattern calculated with the three symmetry-equivalent domain orientations of the proposed structure (c), where 䡩, 䉭, and ▫ correspond to diffraction spots from three domains [A], [B], and [C], shown in panel c. (From Ref. 35.)

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film temperature from room temperature (RT) to 120 K. By considering the temperature dependence of the ␪ dependence, the molecular inclination angle (␤) can be 0⬚ ⱕ ␤ < 6⬚ for the 1-ML CuPc film. The observed and calculated ␾ dependencies are compared in parts b and c of Figure 10. The observed ␾ pattern is rotated by 37⬚ for a best fit with the calculated one. It it noted that the number of peaks and their positions in the ␾ pattern calculated for ␤ = 0⬚ are in agreement with those of observed ones by the rotational fitting of ␾s = ␾ ⫹ 37⬚ (Figure 10b). The calculated ␾ dependencies for ␪ > 6⬚ are obviously different from those observed for any ␾m and any rotational fitting (Figure 10c). From these comparisons it is concluded that the Pc molecules in the monolayer film lie flat on the MoS2 surface. Similar results were obtained for the 1-ML H2Pc film on MoS2. Figure 10d shows the comparison between observed and calculated ␾ patterns, where the calculated results involve the size difference of the three domains [A], [B], and [C]. That is, the ␾ dependence is calculated for S[A]/S[B]/S[C] = 1:1.5:1, where S[A] denotes the size of domain [A]. The difference of the domain sizes was assumed from LEED measurements. The observed ␾ pattern shows excellent agreement with that calculated for ␤ = 0⬚ after rotational fitting of ␾s = ␾ ⫹ 37⬚ both in the relative intensities of peaks in the ␾ pattern and in the peak positions (␾max). From these comparisons and considering four-hold symmetry of the molecule and six-hold symmetry of the substrate, azimuthal orientation of CuPc and H2Pc molecules is determined to be ⫺7⬚ or ⫺37⬚ or ⫺67⬚ with respect to one of the surface crystal axes of MoS2. Here it is noted that the calculated ␾ patterns at azimuthal rotation of ⫹7⬚, ⫹37⬚, and ⫹67⬚ did not agree with the observed results. From the analyses of SR-ARUPS results, the inclination and azimuthal angles of CuPc and H2Pc molecules on MoS2 were determined as shown in Figure 12, where the inclination angle (␤) is 0⬚. Three azimuthal angles, ⫺7⬚ or ⫺37⬚ or ⫺67⬚, cannot be distinguished due to the four-hold symmetry of the molecules and the six-hold symmetry of the MoS2 surface. By joining the results obtained from the analyses of SR-ARUPS and LEED data, three possibilities for the full structure of the 1-ML film on MoS2 were determined experimentally [35]. Considering the smallest overlap of molecules, the structure shown in Figure 13 was finally obtained as the most preferable one. For thicker films, we found from SR-ARUPS studies that Pc molecules tilt about 10⬚ at the surface [35]. Thus, it is considered that for multilayer films molecular orientation changes depend on the distance from the substrate surface. Such a change in the molecular orientation produces a band bending in the films.

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Figure 12 The molecular rotational angles with respect to surface crystal axes of MoS2 determined by SR-ARUPS measurements. (From Ref. 35.)

Figure 13

Full molecular orientation of CuPc on MoS2. (From Ref. 35.)

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Bis(1,2,5-thiadiazolo)-p-quinobis (1,3-Dithiole)

As another example of complete determination of the molecular orientation using SR-ARUPS, results for bis(1,2,5-thiadiazolo)-p-quinobis (1,3-dithiole) (BTQBT) [68] on the MoS2 surface are described. The in situ determination of molecular orientation by SR-ARUPS enabled us to measure the energy band dispersion derived by intermolecular interaction. Figure 14 shows comparison between observed and calculated ␪ and ␾ dependencies of the photoelectron intensity of the HOMO band of BTQBT epitaxial thin films on MoS2, where experimental results are shown for 2-, ˚ -thick films and the calculated results by the IAC/MO and the 4-, and 8-A SS/MO are shown [34]. The observed ␪ dependencies agree better with the calculated one for ␤ = 0⬚ by the SS/MO than by the IAC/MO, indicating that molecules lie flat on the MoS2 surface and that the SS term in Eq. (1) cannot be neglected in this molecule, which contains S atoms. The experimental ␾ pattern agrees well with the calculated one for ␤ = 0⬚ in terms of the angular positions of the peaks as well as in terms of the number of peaks. In the case of ␾ dependence, both IAC/MO and SS/MO results show good agreement with the observed results because the difference of the ␾ dependencies in the two approximations is smeared out by the existence of six azimuthal orientations of the molecules due to the symmetry of the substrate surface and existence of two molecules with mirror symmetry with respect to each surface crystal axis of the surface. After similar procedure made for the CuPc/MoS2 system, a possible full structure of the ultrathin film of BTQBT on MoS2 can finally be obtained as shown in Figure 15, where the six domains of BTQBT coexist and the long molecular axes are rotated by 3 ⫾ 2⬚ with respect to the indicated line in the two-dimensional lattice [34]. The flat-lie orientation can also be obtained on an HOPG graphite [30,69], where the azimuthal orientation of the molecules cannot be observed because of the azimuthal disorder of the small graphite crystals in the sub˚ -thick films on HOPG substrates, the intermolecular band strate. For 30-A dispersion was obtained along surface normal from h␯ dependence of the normal-emission SR-ARUPS as shown in Figure 16, which corresponds to the band dispersion along the stacking axis of the molecules [69]. The reason we can determine the energy-band dispersion along the film normal using very thin films is summarized in Ref. 70. B.

Growth of Chloroaluminum Phthalocyanine Thin Films

As in the case of CuPc and H2Pc, the initial SR-ARUPS study indicated that chloroaluminum phthalocyanine (ClAlPc) molecules are oriented flat in the

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Figure 15 Two-dimensional lattice and molecular orientations at the lattice points determined for BTQBT on MoS2 from the analyses of the LEED and SR-ARUPS. The two domains of R = ⫾13.9⬚ are shown by solid and dotted lines, respectively. The molecules lie flat on the substrate surface and rotate about 3⬚ with respect to the indicated diagonal line. Although surface sulfur atoms of the substrate are shown by the filled circles, the exact location of the molecules with respect to the surface sulfur atoms is not determined from the experiments. (From Ref. 34.)

< Figure 14 Comparison between the calculated and observed angular dependencies for BTQBT on MoS2. The IAC/MO (– – –) and the intramolecular SS/MO ( ) results at h␯ = 40 eV and ␣ = 0⬚ are shown for ␤ = 0⬚. Experimental results are ˚ film). (a) Comparison between the ˚ film), ▫ (4-A ˚ film), and 䉭 (8-A shown by 䡩 (2-A observed and calculated ␪ dependencies at ␾ = ␾max, where the calculations were performed for the experimental conditions and for the molecular azimuthal orientations on MoS2 determined by the analyses of the LEED and the observed ␾ dependence. (b) Examples of calculated ␾ dependencies for a single azimuthal orientation of the molecule shown in the figure. (c) Comparison between the observed and calculated ␾ dependencies at ␪ = ␪max, where the calculations were performed for the experimental conditions and for the two molecular orientations (relative angle = 28⬚) with respect to each surface crystal axis of MoS2 determined by the analysis of the LEED pattern. The long molecular axes are shown by solid arrows, and the surface crystal axes of the substrate are shown by dotted arrows. The observed ␾ dependencies are shown after a rotation of ⌬␾ = 49⬚ for a best fit, where the sixhold rotation gives similar result. The angles (47⬚ and 19⬚) between one of the surface crystal axes and the long molecular axes corresponding to R = ⫾13.9⬚ (about ⫾14⬚) are shown in panel c. (From Ref. 34.)

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Figure 16 Photon energy dependence of SR-ARUPS at ␪ = 0⬚ and ␣ = 50⬚ for ˚ ) on HOPG (a), and intermolecular energy band dispersion (b) obBTQBT (30 A tained from the photon energy dependence of SR-ARUPS. The HOMO and next HOMO (NHOMO) bands are labeled by A and B. (From Ref. 69.)

1-ML film on MoS2, whereas they are oriented to tilt at about 10⬚ in thick films as shown in Figure 17 [71]. However, it was recently found that after the vacuum deposition of the monolayer equivalence amount on the MoS2 surface the film structure and molecular orientation change depending not only on temperature but on time even at RT, and showed an interesting thinfilm growth. This finding demonstrates that various experimental results on the ultrathin film change depending on the time after the vacuum deposition. Here we describe the thin-film growth of ClAlPc, which has an electric dipole, on MoS2. ˚ -thick (about Figure 18 shows the HREELS of a ClAlPc film of 5-A 1.5 ML) deposited on the MoS2 surface at the incidence beam energy of Ei = 3 eV with incidence angle ␪i = 53.6⬚ and scattering angle ␪s = 53.6⬚ relative to the surface normal (specular condition) [72,73]. In the figure, the spectrum measured after a heat treatment of the film at 100⬚C for 3 h is also shown for comparison. The spectrum of the as-grown film shows five vibrational loss peaks at 54, 92, 140, 171, and 384 meV (peaks A–E). The MoS2 substrate has an energy loss peak at 48 meV (peak F) corresponding to excitations of two surface phonons [74], which are not resolved here. The

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Figure 17 Take-off angle (␪) dependencies of photoelectron intensity of the HOMO band of 1-ML film (a) and 5-ML film (b) of ClAlPc on MoS2 measured at h␯ = 40 eV and ␣ = 0⬚. Calculated results for ␤ = 0⬚ and 10⬚ are compared in a and b, respectively. (From Ref. 71.)

˚ thick) Figure 18 HREELS of MoS2 substrate (a), as-grown film of ClAlPc (5 A (b), and heat-treated film at 100⬚C for 3 h (c). (From Ref. 72.)

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losses at 54, 92, 140, 171, and 384 meV for the film can be assigned to the Cl-Al stretch [␯ (Cl-Al)], the out-of plane C-H deformation [␥ (C-H)], the C— —N)], the in-plane C-H deformation [␦ (C-H)], and the —N stretch [␯ (C — C-H stretch [␯ (C-H)]. The ␯ (Cl-Al) peak was assigned by reference to the previous IR studies [75]. Furthermore, this assignment was confirmed by comparing the HREELS of ClAlPc with that of H2Pc, which has no ␯ (Cl-Al) peak [76]. The angular distributions of the elastic and inelastic electrons were found localized around the specular direction. Thus it is concluded that the vibrational excitations occur as a result of the dipole scattering for the ClAlPc/MoS2 system. Therefore, we apply ‘‘dipole selection rules,’’ which state that only the vibrational modes that have a dynamic dipole perpendicular to the surface are excited [49], to analysis of the spectra. When ClAlPc molecules lie flat on the MoS2 surface, we expect that only the peaks of ␯ (Cl-Al) and ␥ (C-H) can be detected in HREELS at specular condition because the dynamic dipoles of ␯ (Cl-Al) and ␥ (C-H) have large components perpendicular to the surface, whereas those of ␯ (C — —N), ␦ (C-H), and ␯ (C-H) have no perpendicular component. On the other hand, when the molecule lies tilted on the MoS2 surface, all of the peaks can be detected because all of the vibrations have the dipole components perpendicular to the surface. As seen in Figure 18, HREELS of the heat-treated film shows a remarkable difference when compared to the as-grown film in the peak intensities. After the heat treatment, the intensities due to the excitations of ␯ (C-H), ␦ (C-H), and ␯ (C — —N) become weak relative to those due to ␯ (Cl-Al) and ␦ (C-H). This finding indicates that the molecules are tilted to the substrate surface in the as-grown film and lie flat after the heat treatment. A similar change in the spectra was also observed by aging the as-grown film at RT. The scattering-angle dependencies of elastic peak intensity for the asgrown film and the heat-treated one are shown in Figure 19, where the angle is scaled using ⌬␪ = ␪s ⫺ ␪i . It is seen that the angular distribution for the as-grown film becomes sharper by the heat treatment. When the as-grown film is not well ordered, it also gives a broader angular dependence of the elastic electron. However, a clear LEED pattern was observed even for the as-grown films [71]. This indicates that the sharpening of the angular distribution due to the heat treatment does not originate from the ordering of molecules but to the decrease in the roughness of the film surface. It is thus expected that the as-grown film consists of ClAlPc islands and, upon heating, the molecules spread over the substrate surface to form a uniform film with flat surface.

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Figure 19 The scattering-angle dependence of elastic peak intensity for ClAlPc ˚ thick) on MoS2. ●, MoS2 substrate; ⽧, as-grown film; 䊱, heat-treated film film (5 A at 100⬚C for 3 h. (From Ref. 72.)

These results of HREELS measurements indicate that the as-grown film consists of islands in which the molecules are stacked with a tilt orientation to the surface. Upon heating, the film becomes a uniform monolayer in which the molecules are oriented flat. Figure 20 shows changes of the LEET spectrum of a 1-ML ClAlPc on the MoS2 substrate during aging at RT and by heating [72]. In these spectra, we note the spectral minima because they originate from the electron reflection at the surface [62]. The spectrum of the as-grown film strongly reflects the features characteristic of the substrate. It changes gradually wih the time to give new features with little contribution from the substrate. This indicates that the MoS2 surface eventually becomes fully covered by the molecules. The spectrum of the heat-treated film is similar to that of the well-ordered uniform monolayer with a flat-lie orientation, where Cl atoms protrude outside the film [77]. The spectrum recorded after the 89-h aging is slightly different from that of the heat-treated one. The difference is considered to be due to the small difference of molecular orientation; that is, the well-ordered uniform monolayer, in which Cl atoms protrude outside the film, was not fully obtained without heating. Figure 21a shows the change of the work function (⌬␾) of the 1-ML ClAlPc film upon aging and heating [72]. ⌬␾ is defined to be the difference of the work function of the film (␾film) and the substrate (␾substrate), ⌬␾ = ␾film ⫺ ␾substrate . The value of ⌬␾ was obtained from the change of the energy

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Figure 20 Change in the LEET spectrum of ClAlPc film (⬃1 ML) on MoS2 with the aging time after deposition of the film. The spectrum of the film heated at ⬃200⬚C for 2 h is also shown for comparison (uppermost curve). The origin of the electron energy scale corresponds to the vacuum level of the MoS2 substrate. (From Ref. 72.)

position of the leftmost peak in the dIt (Ei)/dEi curve, which corresponds to the vacuum level [52,60,77]. As seen in the figure, ⌬␾ shows a very gradual increase (region I) by aging at RT. Even after the 89-h aging, it shows a negative value. Upon heating at 200⬚C, it increases abruptly to give a positive value of about 0.1 eV (region II). The change of the work function can be correlated to that of the molecular orientation. On the other hand, we observed that the work function ␾ of wellordered ClAlPc films, which was obtained after a heat treatment, oscillates ˚ ), as shown when the film thickness increases from 0 to ⬃2 ML (8 ⬃ 10 A in Figure 21b [77]. The Cl atom in the molecule is negatively charged owing

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Figure 21 Change in the work function (⌬␾ = ␾film ⫺ ␾substrate) with the aging time (region I) and upon heating (region II) for the 1-ML film shown in Figure 20 (panel a) (from Ref. 72), and ⌬␾ for the heat-treated films as a function of the thickness (panel b) (from Ref. 77). In panel b, the oscillation of ⌬␾ in the film ˚ corresponds to the change of the molecular orientation. thickness region of 0–10 A In the monolayer, molecules are oriented with Cl atoms protruding outside. In the double layer, the second layer molecules are turned over to compensate the electric dipole in the first layer.

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to its electron-withdrawing nature, and therefore the molecule has an electric dipole perpendicular to the molecular plane, directed from the Cl to Al atom. By considering the electric dipole of ClAlPc, the oscillation of ␾ indicates that (a) in the well-ordered monolayer, the molecules lie flat with the Cl atoms protruding outside the film, and (b) in the double layer, the outer molecules are tuned over with the Cl atoms protruding into the film to compensate the electric dipole of the underlayer [77]. It is thus considered that the as-grown monolayer, which shows a negative value of ⌬␾ even after the 89-h aging (see Fig. 21a), does not consist of the monolayer with the Cl atoms protruding outside the film. In passing, it is of note that the molecules in the 89-h aged monolayer film completely shields the MoS2 surface. The abrupt increase of ⌬␾ upon heating indicates that all molecules become oriented with the Cl atoms protruding outside the film by the heating. The LEET can detect monolayer formation and molecular orientation very sensitively, and the results indicate that the heating is necessary to obtain well-oriented monolayer. Figure 22 shows the time and temperature dependence of the PIES (MAES) for the 1-ML ClAlPc deposited on the MoS2 surface [72]. The PIES of the film has six bands denoted by A–E and X. According to Pasinski et al. [44] and Aoki et al. [78], band A is related to the ␲-ring orbital, which is mainly distributed along the skeleton of the inner porphin-like ring. Band B is ascribed to MOs with large contribution from the outer four benzene (some ␲-ring orbitals are also responsible for band B). Bands C and D can be assigned to the nonbonding orbitals of the Cl atom, Cl(n㛳) and Cl(n⬜), respectively, which are distributed parallel and perpendicular to the molecular plane. B and E can be ascribed to the ␴ orbital. In the spectrum of the as-grown film, we reproducibly observed the substrate bands F and G, indicating that the substrate is not fully covered by the molecules. Furthermore, bands C and D are weak whereas band B is stronger in the as-grown film. From these results it is concluded for the as-grown film that the molecules form islands on the MoS2 surface in which the Cl atoms of the outermost molecules are mainly directed to the substrate. These results are fully consistent with those obtained by HREELS and LEET measurements. Even after aging for 45 h at RT, the spectrum shows little change except for the disappearance of the substrate peaks of F and G. On the other hand, upon heating, the intensities of bands C and D become strong, while band E disappears. The result clearly indicates that after the heat treatment the He* atoms interact effectively with the n㛳(Cl) and n⬜(Cl) orbitals, whereas they scarcely interact with the ␴ orbital, indicating that the molecules come to lie flat with Cl atoms protruding outside the film as a result of heating. From PIES studies, we can conclude that (a) the as-grown monolayer consists of islands and the outermost molecules orient with the Cl atoms protruding

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Figure 22 Time and temperature dependencies of He*(23S) PIES of a 1-ML ClAlPc deposited on MoS2 surface. (From Ref. 72.) The uppermost spectrum was obtained at 25⬚C after annealing the film at ⬃100⬚C for 23 h. The other spectra were measured during aging at 25⬚C and annealing at the indicated temperature.

inside the film; (b) the molecules move to cover the MoS2 surrface during aging at RT, where the molecules do not orient with the Cl atoms protruding outside the film; and (c) upon heating, the molecules form a uniform monolayer with a flat-lie orientation, where Cl atoms protrude outside the films. For thicker films, the PIES indicated that the molecules are tilted at the surface as observed by SR-ARUPS [71,78]. From the results of SR-ARUPS, HREELS, LEET, and PIES, the following conclusion can be consistently obtained for a structure of ClAlPc thin film of monolayer equivalence on the MoS2 surface: (a) ClAlPc molecules in as-grown film aggregate to form islands, where the outermost molecules orient with the Cl atom protruding inside the film; (b) in the film aged at room temperature, the molecules in the islands move to cover the substrate surface, where the molecules are oriented flat but are not fully oriented with the Cl atoms protruding outside the film; and (c) upon heating the film at about 100⬚C, the molecules spread over the substrate surface to

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form a well-ordered uniform monolayer with a flat-lie orientation, where all Cl atoms protrude outside the films. Using this controlled molecular orientation change, we can intentionally obtain molecular surfaces with different chemical properties. Figure 23 summarizes schematically the growth of ClAlPc thin-film vacuum deposited on MoS2. These results indicate that when one wants to prepare a well-oriented flat multilayer, careful control of the molecular orientation in the monolayer is required. C.

Molecular Orientation, Reaction, Band Gap State, and Wavefunction in Indium/PTCDA System

Hirose et al. [79] found that a new electronic state with very low ionization energy (band gap state) appears by a contact of perylene-3,4,9,10-tetracarboxylic dianhydride (PTCDA) and metals such as In, Al, Ti, and Sn in the band gap region of PTCDA. In an organic–metal interface such a state would play an important role, and therefore the study of the origin of the band gap state is of crucial importance. In this section, we show the origin of the band gap state appearing in In/PTCDA system. In Figure 24a, the ␪ dependence of the SR-ARUPS of PTCDA thin ˚ thick; about 1 ML) on the MoS2 surface is shown [80]. The ␪ film (3 A dependence indicates that PTCDA molecules lie flat on the MoS2 surface [80,81]. The assignment for each band in the spectra was made with MO calculation and the quantitative analysis of the photoelectron intensity [80,81]. Briefly, bands A (EB = 2.2 eV), B (4.0 eV), and D (5.7 eV), which show the intensity maximum at ␪ ⬵ 34⬚, 34⬚, and 46⬚, respectively, originate mainly from ␲ states. Among these, band A, which consists of single ␲ MO, and the feature C (EB = 4.6 eV), which appears at a larger value of ␪, originates from oxygen 2p nonbonding state n//(O) at C — —O parts with O 2px and 2py AOs. Figure 24b shows the dependence of the SR-ARUPS of an In–PTCDA ˚ -thick PTCDA interface that was prepared by evaporating In onto the 3-A ˚ . When we compare film [80]. The thickness of the In overlayer was 1 A parts a and b of Figure 24, some spectral changes can be seen. First, a new band X appears at EB ⬵ 0.8 eV in the PTCDA band gap after the In deposition. Second, at the region of EB = 1.5–5 eV, some new features arise in the spectra that are broader than those of the pristine PTCDA thin film in Figure 24a. Furthermore, the n//(O) band of PTCDA (band C in Figure 24a) was not resolved after the In deposition, while new band X⬘ (EB = 3.6 eV) is observed at ␪ = 62⬚. These spectral changes originate from a chemical reaction between the PTCDA and the In overlayer. The disappearance of the n//(O) band in Figure 24b suggests that the chemical reaction between the PTCDA and the In atoms takes place at the C — —O parts to change the

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Figure 23 Schematic illustration of thin-film growth of ClAlPc on MoS2. (1) and (2): After the deposition of 1 ML equivalence at RT, the molecules aggregate to form islands with multilayer. (3) By heating the islands, the aggregated molecules diffuse and spread over the substrate surface to form a uniform monolayer with the Cl atoms protruding outside. (4) By the second layer deposition, all molecules in the second layer turn over to compensate the electric dipole layer of the first layer. (5) In the multilayer formed by further deposition of molecules, paired molecules form layers, where the surface consists of slightly tilted molecules with the Cl atoms protruding inside. This means that there is no odd-numbered layer except for the monolayer. (From Ref. 77.)

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˚Figure 24 The take-off angle (␪) dependencies of SR-ARUPS spectra of a 3-A thick (⬃1 ML) PTCDA film on MoS2 before (a) and after (b) In overlayer deposition. ˚ . The intensity is normalized to the incidence The thickness of the In overlayer is 1 A photon flux. (From Ref. 80.)

n//(O) states. It is expected that band X⬘ originates from the AO which is distributed parallel to the surface, since band X⬘ was observed at a large value of ␪. On the other hand, HREELS for the In/PTCDA on MoS2 [82] showed that the intensities due to the excitation of C — —O stretching mode [␯ (C — — O)] become weak relative to those of C-C stretching mode [␯ (C-C)] and COC stretching mode [␯ (COC)], and only the peak position of ␯ (C — —O) shifts to the lower wavenumber side. These findings indicate that a strong interaction between the PTCDA molecules and the In atoms exists at the C— —O parts. Furthermore, Kera et al. recently performed MAES measurements of the In–PTCDA interface on the MoS2 surface and showed that four In atoms react with the one PTCDA molecule [83]. Figure 25 shows the ␪ dependence of the intensity of band X, where the angular distribution of the top ␲ band of the pristine PTCDA is also shown for comparison. The angular dependence of band X intensity shows a maximum at ␪max = 36–40⬚ with a broader angular distribution than the

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Figure 25 The take-off angle (␪) dependence of the band gap state of the In/ PTCDA interface (●) and the ␪ dependence of the top ␲ band of the PTCDA before In deposition (䡩) (panel a). The comparison between the observed (●) and calculated ␪ dependencies of the photoelectron intensities for the band gap state (panel b). The SS/MO results at h␯ = 40 eV and ␣ = 0⬚ are shown for ␤ = 0⬚ (– – –), 10⬚ ( ), and 20⬚ (– ⭈ –). The calculation was made by assuming azimuthal disorder of the reaction product, In4PTCDA. The tilt direction of the molecule is also shown. (From Ref. 80.)

top ␲ band of the pristine PTCDA. In general, when large ␲-conjugated planner molecules, such as PTCDA and phthalocyanine, lie flat on the substrate surface, the ␪ dependence of the photoelectron intensity from the single ␲ band, which spreads over the molecule, showed a value of ␪max = 32– 36⬚, with very sharp angular distribution due to the interface effects. The broader angular distribution for the X band suggests that the molecular plane becomes tilted by the reaction. In Figure 25b, the observed ␪ dependence of the band X intensity is compared with calculated one. The calculated results were obtained for the expected reaction product, In4PTCDA, by assuming random azimuthal orientation of the molecules for ␤ = 0⬚, 10⬚, and 20⬚, with the short molecular axis being parallel to the surface using the SS/MO method with ab initio (STO-6G) wavefunction for the initial state. The calculated ␪ pattern for ␤ = 10⬚ agrees better with the observed one in the value of ␪max and the width

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of the angular distribution. However, the observed angular distribution is broader than the calculated results for a single value of ␤, ␤ = 10⬚. Furthermore, the calculated result for ␤ = 20⬚ also gives a fair agreement when we look at the total width of the angular distribution. From this, we expect that the sample consists of some mixture of molecules with ␤ ⬵ 10–20⬚. This results indicates that the In4PTCDA molecules are tilted at ␤ ⬵ 10– 20⬚. In fact, SR-ARUPS experiments on the annealed In/PTCDA/MoS2 showed a broader ␪ dependence [82]. In Figure 26 the observed SR-ARUPS of the In/PTCDA is compared with the results of ab initio MO calculation (STO-6G) on In4PTCDA. The observed spectrum was shown after averaging the ␪ dependence in Figure 24b and the calculated DOS was obtained by 0.7-eV Gaussian broadening of the MO levels. From the calculation, band 1 consists of a single MO of ␲ character with large contribution from C 2pz , O 2pz and In 5pz AOs, and band 2 originates from an MO of ␲ character, which consist of C 2pz and O 2pz AOs. On the other hand, band 3 is related to four ␴ MOs with the contributions of O 2s, In 5s, 5px , and 5py AOs. Finally, band 4 involves four ␲ states distributed mainly at the perylene core. Considering the possibility that there exists some unreacted PTCDA molecules, the DOS curve agrees reasonably well with the measured angle-averaged UPS, indicating that X and X⬘ are assigned to the ␲ and ␴ orbitals of In4PTCDA, respectively. The agreement between the observed spectra and the DOS indicates that the electronic structure of the In/PTCDA is well simulated by that of In4PTCDA. The SR-ARUPS studies indicated that the band gap state involves contribution from large atomic wavefunction of In 5pz , and also the deeper lying band X⬘ has a large contribution from In 5s orbital. Therefore, it is interesting to see the relative intensity of MAES (PIES) features, since for nearly flat orientation of the In/PTCDA system on MoS2, these states should give stronger bands in the MAES due to larger spread of the wave functions. Such a comparison can provide experimental assignment of the observed spectral features and improve the assignment made using calculated MO levels, which depend on calculation methods. In Figure 27, the MAES and He I UPS of the In/PTCDA/MoS2 system are shown together with calculated results [83]. In this experiment, the sample was prepared at 100⬚C to obtain a more complete reaction. Here, we first note the main difference between the relative band intensities in the MAES and UPS. In the MAES, one can see that the gap state R1, which corresponds band X in Figure 24b, is stronger than the 2nd band R2. Moreover, the relative intensity of the third band R⬘3 , which may correspond to band X⬘ in Figure 24b, in the MAES is especially stronger than the corresponding band R3 in the UPS. The former result originates from the difference of MO distributions of R1 and R2 bands. The latter can be ascribed to

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Figure 26 Comparison between observed SR-ARUPS and the density of states (DOS) calculated for a model system of the reaction product, In4PTCDA. The vertical bars show the calculated MO energies obtained by ab initio method (STO-6G) for the In4PTCDA, where four In atoms are reacted with the four C — —O parts of PTCDA as shown by the inserted figure. The upper and lower bars indicate ␴ and ␲ states, respectively. The calculated DOS was obtained by a Gaussian broadening of the MO levels (see text). The observed SR-ARUPSs of the pristine PTCDA (– – –) and the In/PTCDA ( ) are shown for comparison, where both spectra were obtained by averaging of the ␪ dependencies. The MO coefficients of the band X in the reacted system are shown. (From Ref. 80.)

the difference of surface sensitivity between MAES and UPS; that is, an orbital extending outside the surface can be selectively detected by MAES and gives a stronger band in MAES than in UPS. Prior to discussion of the intensity, we first describe the assignment of the MAES and UPS features using density functional theory (DFT) calculations. The DOS by this method gave the best agreement with the observed spectra. The DFT result for the In4PTCDA indicates that the top band 1 in the DOS and selected DOS (SDOS) is due to the ␲ state consisting of C 2pz , O 2pz , and In 5pz AOs, and the second band 2 originates from a ␲ state

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Figure 27 Comparison between observed MAES (PIES) (He*, 23S), UPS (He I) ˚ )/PTCDA (1 MLE)/MoS2 system and calculated results (DFT spectra for In(1 A method) for the reaction product In4PTCDA. Here MLE indicates monolayer equivalence. Molecular energy levels are shown by bars for MO distributed in perpendicular (MO⬜) [␲ and perpendicularly distributed nonbonding state of oxygen n⬜ (O)] and parallel fashion (MO//) [␴ states, parallel distributed nonbonding state of oxygen n⬜ (O)] to the molecular plane. Selected DOS for MO⬜ states (SDOS) is shown by dashed curve. Molecular planes are nearly parallel to the substrate surface; therefore, the MAES reflects the SDOS rather than DOS. The negative of the second derivative of the MAES (⫺dI 2/dE 2) and MAES after the background subtraction are also shown, indicating that the HOMO band (R1) intensity is greater than that of the second band (R2) (next HOMO). (From Ref. 83.)

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derived from C 2pz and O 2pz AOs, distributed only in the perylene core (␲P). Here the SDOS corresponds to the DOS derived from MOs distributed perpendicularly to molecular plane (MO⬜) such as ␲ and n⬜(O). The third band 3 consists of five ␲ MOs, which involve a ␲ state with MO phase similar to the HOMO state of PTCDA and nearly degenerate four ␴ states with large In 5s AOs. The fourth band 4 is related to various MOs, such as ␲P, ␲CO, n⬜(O), n//(O), and ␴ states. From the fairly good agreements between calculated and observed results, the MAES and UPS bands can finally be assigned as follows: In the UPS, band R1 is ascribed to ␲P, involving In 5pz AOs, which is in good agreement with those obtained by the SR-ARUPS study [80]. Band R2 is related to ␲P, and band R3 to five ␲P MOs and four ␴ MOs involving In 5s AOs. Finally, band R4 is ascribed to various MOs, ␲, n⬜(O), n//(O), and ␴ states. As described before, MAES is sensitive to the molecular orientation and reflects MOs with large spatial distribution. It is noteworthy that expected characteristics of the MAES of the In/PTCDA system are (a) the MO⬜[␲ and n⬜(O)] were mainly observed because the molecules were oriented nearly flat to the substrate and (b) bands including large In AOs gave stronger intensities. From these, MAES band R1 is ascribed to ␲P, involving In 5pz AOs. Band R2 is related to ␲P MO, and band R⬘3 to ␲P MOs and ␴ MOs involving In 5s AOs. Finally, band R4 is ascribed to ␲P, ␲CO, and n⬜(O) MOs. Here, band R1 was observed strongly comparing with band R2 in the MAES (see the inset ⫺d 2I/dE 2 curve and background-removed spectrum in Figure 27). This is because the wave function for band R1 (HOMO) involves In 5pz AOs and spreads more outside than that for band R2 (␲P), which consists of only C and O 2pz AOs. The new electronic state, HOMO of In4PTCDA, has a large wave function that distributes near normally to the interface plane, and it may play an important role in the charge transfer properties through the interface by bridging the two interface materials. Furthermore, the especially strong intensity of the MAES band R⬘3 in comparison with corresponding band R3 in the UPS can be due to the electron emission from the In 5s–derived ␴ orbitals. These assignments considering MAES band intensities are consistent with those using the DFT calculation.

IV.

APPLICATION OF SR-ARUPS TO A DISORDERED POLYMER SURFACE: COMPARISON WITH NEXAFS

In a pendant group polymer, replacing a pendant chemical group largely controls chemical, physical, and/or electrical properties. In addition, due to molecular stability and facility of preparation, the pendant group polymers

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are widely used in practical applications. Furthermore, due to the flexibility of such polymers, they are promising materials that will be used in organic devices. Considering the high popularity of such polymers in industrial fields, there are relatively few investigations of molecular orientation at surfaces of their thin films because of lack of an appropriate tool to investigate such an unoriented molecular system. In general, the molecules in polymer solids are not oriented because of the large molecular weight distribution, mixed tacticity, and so forth; therefore, it is believed that the pendant groups are also not oriented at the surface. From an intuitive point of view, one supposes that such disordered system shows no properly dominated ‘‘orientation effects’’ of the pendant group in the disordered system. This may be one of the reasons that few basic scientists went into this area of ‘‘frontier surfaces.’’ However, we found from the SR-ARUPS study that three-dimensional random orientation of naphthalene pendant group at the surface of poly(2-vinylnaphthalene) thin film would be somewhat controlled by the naphthalene group tilted at large angles as described in this section. On the other hand, it is well known that polarization dependence of the near-edge X-ray absorption fine structure (NEXAFS), also called X-ray absorption near-edge structure (XANES), gives molecular orientation at the surfaces of various organic films [84–87]. The basis of NEXAFS is the excitation of electrons from core to unoccupied orbitals. The transition probabilities of these excitations are governed by dipole and symmetry selection rules; thus, molecular orientations can be determined by analyzing light polarization dependencies of spectral peak intensities. With linear-polarized synchrotron radiation, this is a powerful and convenient technique to examine molecular orientation. Therefore, it is interesting to compare the results obtained by NEXAFS and SR-ARUPS. Here we compare the pendant group orientations of the poly(2-vinylnaphthalene) thin-film surface determined by these two methods, and show an advantage of SR-ARUPS as well as a limitation of NEXAFS in studying molecular orientation. In Figure 29, NEXAFS spectra of the poly(2-vinylnaphthalene) (see Figure 28 for molecular structure) thin films at the CK-edge are shown for normal (␣ = 0⬚) and grazing (␣ = 72⬚) incidence conditions [39], where those of naphthalene thin films on Ag(111) with flat-lying orientation [88] are compared. In the naphthalene pendant polymer, as seen in Figure 29a, only weak polarization effect was observed. Judging from the polarization dependence of naphthalene [88] and polystyrene [86] NEXAFS spectra, the peak origins of the naphthalene polymer spectrum are assigned as follows. The first intense peak at 284.8 eV and the second one at 285.6 eV are attributed to transitions from the 1s orbital to the unoccupied ␲* orbital of the naphthalene pendant groups. Peak 5 at 293.5 eV and peak 6 at 300.0

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Figure 28 Molecular structure of poly(2-vinylnaphthalene) (a) and isopropylnaphthalene, which was used as a model unit in the SS/MO calculation (b).

eV correspond to transitions to unoccupied ␴* orbitals of the pendant groups. However, peaks 3 and 4 could not be assigned clearly. In NEXAFS, molecular orientation can be determined by analyzing the polarization dependence of the ␲* transition intensity. A dipole moment of the ␲* transition is oriented perpendicular to the molecule plane; hence, the transition intensity is largest when the light E vector is aligned parallel to the direction (main axis of atomic pz orbital) of the ␲* orbital. Quantitative description for the 1s → ␲* transition intensity f (␤, ␣) under a perfect polarized light condition has been given as [89,90] f(␤, ␣) ⬀ 2 cos2(␤)cos2(90 ⫺ ␣) ⫹ sin2(␤)sin2(90 ⫺ ␣)

(2)

where ␣ is the incidence angle of the linearly polarized light and ␤ is the tilt angle of the molecule with respect to the substrate surface. The calculated results of Eq. (2) for several tilt angles are presented as solid lines in Figure 30. This equation demonstrates that a large intensity change is expected when the molecule is oriented almost flat (␤ ⬵ 0⬚) or upright (␤ ⬵ 90⬚), whereas a relatively small intensity change is expected for intermediate angles. At ␤ = 54.7⬚, the magic angle, no intensity change is expected. The polarization dependence observed for the first ␲* transition intensity (peak 1 in Figure 29) of the naphthalene polymer is given as solid circles in Figure 30. The observed result indicates that the naphthalene pendant groups are tilted by about 57⬚ with respect to the polymer surface. In evaluating the tilt angle, it must be cautioned that the obtained angle of about 57⬚ could have resulted from a highly ordered orientation where all molecules are tilted at a single angle of 57⬚ or could be interpreted as an average angle of variously tilted naphthalene pendant groups. These possibilities cannot be distin-

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Figure 29 Carbon K-edge NEXAFS spectra of (a) poly(2-vinylnaphthalene) thin films at normal (solid line, ␣ = 0⬚) and grazing incidence (dashed line, ␣ = 72⬚) and (b) naphthalene thin films on Ag(111) surface at normal (solid line, ␣ = 0⬚) and grazing incidence (dashed line, ␣ = 70⬚) [88]. The insert shows the measurement geometry. (From Ref. 39.)

guished by NEXAFS because the absorption peaks for various orientations are located at a similar h␯ position in the spectra. SR-ARUPS of poly(2-vinylnaphthalene) are shown in Figure 31 [39]. The spectra were normalized to the incidence flux, and the binding energy is scaled to EF . The density of states (DOS) obtained by ab initio MO calculation (STO-6G) also is given in the same figure. The calculation was performed on the model molecule, isopropylnaphthalene (Figure 28), which can be regarded as a monomer of the naphthalene polymer terminated with methyl groups at the two ends. The calculated DOS peaks reveal excellent agreement with the measured ARUPS in spite of the fact that the MO calculation was made only on a single-model unit of the actual polymer. This agreement substantiates that the electronic structure of the naphthalene polymer is well represented by the model compound. From the MO calculation,

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Figure 30 Observed (●) and calculated (solid curves) incidence angle (␣) dependence of the 1s → ␲* transition intensity (peak 1 in Figure 29a) of poly(2-vinylnaphthalene) thin film. (From Ref. 39.)

the first two peaks (indicated by A in Figure 31) in the SR-ARUPS are assigned as originating from the ␲ states of the naphthalene pendant group. Peak B is also due to the ␲ orbital. Peaks C–F correspond mainly to the ␴ states from the main chain and the pendant group. In this binding energy region, there are two additional ␲ orbitals. One contributes to peak C and the other is located between peaks D and E. In the calculation of photoelectron intensity, the SS/MO method was applied to a single unit of the model polymer with three orientations (see below) of the tilt angle ␤ by assuming that the azimuthal orientation of the naphthalene pendant group is disordered. Photoelectron angular distributions for the three orientation models were thus calculated by summing over the two azimuthal orientations of the molecule, namely, azimuthal rotations with respect to normal to the molecular plane (␾m), and with respect to normal to the substrate plane (␾s) (see Figure 32 caption). 1.

2.

All naphthalene pendant groups at the polymer surface are assumed to be tilted at one unique angle (the single-␤ orientation model), ␤ = 57⬚, which was determined from the polarization dependence of NEXAFS measurements in the previous section. This corresponds to a highly oriented system. Naphthalene pendant groups are oriented randomly at the polymer surface. To clarify this model further, we define ‘‘random orientation’’ here. In Figure 32a, a schematic drawing to define orien-

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Figure 31 SR-ARUPS spectra of poly(2-vinylnaphthalene) thin film at h␯ = 40 eV, light incidence angle ␣ = 0⬚, and photoelectron take-off angle ␪ = 30⬚ ⬃ 62⬚. The vertical bars indicate MO energies obtained from an ab initio MO (STO-6G) calculation. The calculated binding energy scale was contracted by 1.35 and shifted to fit the observed spectrum. The dashed curve was obtained after DOS peaks were broadened by Gaussians (0.8 eV FWHM). The MO calculation was performed on the model unit shown in Figure 28. (From Ref. 39.)

tation for the naphthalene pendant group is presented. Assuming that P is an intersecting point between the molecular normal and the surface of the unit sphere, a molecular tilt angle ␤ can be redefined as a position of P. In the perfect random orientation for the tilt angles, positions of P must be located homogeneously on the sphere surface (the three-dimensional isotropic random orientation model). To achieve this condition, the number of pendant groups having tilt angle ␤ is proportional to sin ␤. This distribution

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Figure 32 Schematic drawing for the three-dimensional isotropic random orientation model (panel a) and tilt-angle distribution [N(␤)] of naphthalene pendant groups derived from the orientation models (panel b) (see text). In panel a, the molecule is located at the center of the unit sphere. P is an intersecting point between the molecular normal and the surface of the unit sphere. The molecular tilt angle ␤ can be defined as the location of P. The distribution curve (see solid line in panel b) provides a homogeneous distribution of P on the unit sphere surface. This condition is known as three-dimensional isotropic random orientation. ␾m and ␾s are rotational angles around these axes that are normal to the molecular plane and the substrate plane, respectively. In panel b, the dashed line is for model 3 and the solid line is for model 2 (three-dimensional isotropic random orientation) (see text). The mean tilt angles derived from the distributions are 57.3⬚ and 45⬚ for models 2 and 3, respectively. (From Ref. 39.)

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function N(␤) is given by the solid line in Figure 32b, revealing that the number of pendant groups increases sinusoidally with increasing tilt angle ␤. The mean value of ␤ for this distribution is calculated to be 57.3⬚, which is very close to the mean tilt angle value of about 57⬚ obtained from the NEXAFS measurement. In the third model, the distribution function N(␤) is constant as a function of the tilt angle, as shown by the dashed line in Figure 32b. Tilt angle ␤ varies between 0⬚ and 90⬚ with the flat distribution; therefore, the mean value of ␤ becomes 45⬚. This model corresponds to a random orientation only in ␤ and is considered to be an intermediate orientation between model 1 (highly ordered orientation) and model 2 (random orientation).

Measured ␪ dependence of the peak A intensity in Figure 31 is presented in Figure 33 as the solid circles, where the intensities were obtained by summing the two ␲ states contributing to the peak A. The observed angular pattern is quite broad as compared to those in which all molecules are lying flat on a surface. The calculated ␪ dependencies of the abovedescribed orientation models are also given in Figure 33 as the dotted (model 1), solid (model 2), and dashed (model 3) lines. The most striking feature

Figure 33 The observed and calculated take-off angle (␪) dependencies of the photoemission band A of poly(2-vinylnaphthalene) thin film. The experimental results are plotted as solid circles. The results of the SS/MO calculations for the three orientation models are given as a dotted line for the single ␤ orientation (model 1), a solid line for the three-dimensional isotropic random orientation (model 2), and a dashed line for model 3 (see text). (From Ref. 39.)

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of the calculated results is that the different molecular orientations are clearly distinguishable in the photoelectron intensity angular pattern. This was not the case in the NEXAFS polarization measurements. In NEXAFS, the single-␤ orientation (model 1) and the three-dimensional isotropic random orientation (model 2) are indistinguishable because the mean tilt angle of both orientations is the same (about 57⬚). As a result, the same polarization dependence would be observed in either case. By comparing the calculated results and the observed angular dependence, the three-dimensional isotropic random orientation (model 2) represented by the solid line in Figure 33 gives the best agreement with the observed angular pattern. Therefore, it is concluded that the naphthalene pendant groups are oriented randomly according to the distribution given by Figure 32b (solid line) at the polymer surface. Since the distribution has larger populations in the larger tilt-angle region, properties of the polymer surface are considered to be dominated by the ␴ states arising from C-H bond (naphthalene ring edge). As seen in the calculated angular patterns for models 2 and 3, where molecules are oriented randomly but differently, each angular pattern is clearly different, suggesting a unique capability of the ARUPS-SS/MO technique to investigate molecular orientation in disordered systems.

V.

CONCLUDING REMARKS

As understood from many recent works on enegy level alignment at organic–inorganic and organic–organic interfaces, the investigation of electronic states has gained a central position in developmental works of actual organic devices. This situation could not be expected when we started electron spectroscopic studies on the electronic structure of organic systems in the early 1970s because such studies on electronic states were considered to be not useful for developing actual organic devices in industrial field. Steady progress with various electron spectroscopic techniques and theoretical analysis methods was undoubtedly the basis for the successful use of these techniques in the analyses of organic materials. On the other hand, one should remember that major aspects of these investigations, namely, development of organic devices, have been performed by chemists, with only a limited number of physicists taking part in these. Generally, organic molecules, which have already been used in various fields, have very complicated molecular structures, thus posing a serious barrier for physicisits wishing to do research on organic systems. There are frontier fields of science and technology in organic systems that are still sleeping, and we believe that they are waiting for ‘‘morning call’’ by chal-

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lenging people in various fields with new ideas and new experimental methods. We hope that such new-frontier fields will appear in the near future through challenging research on organic interfaces.

ACKNOWLEDGMENTS The authors are grateful to Prof. Y. Harada of Seitoku University and to Prof. K. Seki and Dr. H. Ishi of Nagoya University for useful discussions about the electronic structure of organic films and their supports in various ways. The authors also thank Prof. E. Morikawa of Louisiana State University, Dr. T. Miyamae of the National Institute for Materials and Chemical Research, Dr. M. Aoki of the University of Tokyo, and Dr. D. Yoshimura of the Institute for Molecular Science for their help in SR-ARUPS experiments, and to the staff of UVSOR for their support during SR-ARUPS measurements.

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7 Organic–Metal Interfaces From Physisorption to Covalent Bonding Norbert Koch and Gu¨nther Leising Institute of Nanostructured Materials and Photonics, Joanneum Research Forschungs GmbH, Weiz, Austria

Aparna Rajagopal Rutgers University, Piscataway, New Jersey

Jean-Jacques Pireaux Faculte´s Universitaires Notre-Dame de la Paix, Namur, Belgium

I.

INTRODUCTION

The possible interactions between conjugated organic materials (COM) and metals cover the whole range from physisorption to covalent bonding and charge transfer in terms of ionic bonding. This makes it virtually impossible to predict a certain energy level alignment at a given interface, based on individually determined materials properties such as metal work function, ionization potential, and electron affinity. Without doubt, investigations on the electronic properties at metal–organic interfaces with photoemission spectroscopy (PES) techniques can provide the information required for a better understanding of the nature of the interaction and the energy level alignment, which determine to a large extent the stability and performance of organic light-emitting devices, photovoltaic cells, or field effect transistors. As will be outlined below, great care must be taken when employing PES on COMs. Often, concepts valid for inorganic semiconductors are not appropriate for these organic materials. Many of the electroactive COMs used for organic light-emitting devices (OLEDs) have an energetic spacing between the highest occupied mo205

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lecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) larger than 2 eV, defining the family of wide band gap semiconductors. In such materials it is possible to change the intrinsic work function (i.e., the energetic distance of the Fermi level from the vacuum level) through a very wide range by even small amounts of dopants, or charged defects in general. One of the advantages of PES experiments is the need for ultrahigh vacuum conditions. It is possible to grow thin films of COMs (oligomers or relatively large molecules, or polymers by vapor deposition polymerization [1]) in situ by sublimation techniques, which enables the reduction of all kinds of impurities and contaminations to a minimum. However, sometimes it can be difficult to determine intrinsic work functions with the help of PES. Usually, some metal or semiconductor serves as substrate for the organic film. The first monolayer of the COM will have a certain interaction with the substrate material, defining the alignment between the observable substrate Fermi level (metal) or valence band (semiconductor) and the molecular orbitals of the adsorbed organic material. Certainly, this is true only if the interaction between substrate and COM is relatively weak and does not lead to a chemical decomposition of the organic molecule. For the following discussions we exclude any diffusion effects. Molecules deposited subsequent to the first monolayer will ‘‘feel’’ the surface potential of the first monolayer. Naturally, the molecular orbitals of additional layers will line up with the ones of the underlying COM. As the intermolecular interaction within one COM is mainly of the van der Waals type, a so-called vacuum level alignment will occur within the organic layer, and flat band conditions will prevail for thick films. If the level of contamination in the COM is low, there will be no (free) charge carriers, which would allow the development of band bending, as observed for many inorganic semiconductor–metal interfaces [2]. As a consequence, the work function of a thick conjugated organic layer on a conductive substrate measured in PES will be determined by the work function of the substrate corrected for an optional interface dipole, located mainly at the first organic monolayer. The contribution of the interface dipole depends on the individual chemical interaction at the considered interface and can cover a range from 0 eV to over 1 eV for some common pairs of COMs [3–7] and metals. Often the reverse case, i.e., the energy level alignment occurring when a metal is deposited on top of a COM layer, is of more practical interest for OLEDs. A different energy level alignment might be established compared to the organic-on-metal scenario because single metal atoms or small clusters impinging on the organic surface can possess chemical properties different from those of a solid metal surface. Again, if diffusion is neglected, the energy level alignment will be determined by the nature of the chemical interaction of the metal atoms/clusters with the organic material, including

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interface dipoles. The final alignment will only be established when the metal layer has reached sufficient thickness to develop true metallic properties. The work function of the pristine organic layer, determined by the substrate material, has no influence on this energy level alignment. The work function measured in PES will gradually attain the value expected for the metal as the overlayer grows thicker and becomes metallic. As the interfacial region is confined to a few molecular/atomic layers, no band bending can occur in a pure wide-band gap semiconductor. The situation becomes more complex as metal atoms diffuse into the organic film. If the interaction of the metal with the COM is of ionic nature, the metal ions indeed can form a gradient within the organic layer and an effect similar to band bending may be observed [8,9]. The information obtained in PES studies of organic–metal interfaces often becomes obsolete for ‘‘real’’ devices, as high-vacuum conditions are more common in the device production process than the ultrahigh vacuum (UHV) sustained during PES experiments. In high vacuum the rapid formation of a contamination layer of unwanted adsorbates (O2, H2O, hydrocarbons, etc.) may occur. These change the chemistry at the organic–metal interface; for example, there might be oxidation of the first metal layer, and therefore the energy level alignment may differ substantially from the one derived from photoemission data obtained under UHV conditions. However, it ought to be mentioned here that there are examples where the introduction of an oxide layer (intentionally or unintentionally) increases the performance of OLEDs [10,11]. This is by no means the only method to increase charge injection properties of metal–organic interfaces, and there is an ongoing effort to explore methods to accomplish this task [12–16]. The following PES studies of organic–metal interfaces should demonstrate the wide variety of different interactions between metals and conjugated organic materials. We have chosen one oligomer (para-sexiphenyl) and one polymer {poly(2,5-diheptyl-1,4-phenylene-alt-2,5-thienylene) [17], PDHPT} that exhibit blue electroluminescence and are therefore promising candidates for application in OLEDs. The chemical structure of para-sexiphenyl (6P) and PDHPT are shown in Figure 1a and b, respectively. The interface of these two COMs has been studied with a variety of metals [7,18–22]. In this chapter we discuss the interaction of 6P with Al, Sm, and Ca, and that of PDHPT with Ca, as they are representative of the spectrum of possible interactions at such interfaces. For aluminum, a widely used cathode material in OLEDs, on para-sexiphenyl we obtain evidence of physisorption. Another common cathode material is calcium, due to its low work function. For the Ca–6P interface an ionic interaction is proposed. As an alternative low work function metal in OLEDs we suggest samarium, where we find indications for a physisorption-like interaction with 6P and a very

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Figure 1 Chemical structures of (a) para-sexiphenyl and (b) poly(2,5-diheptyl1,4-phenylene-alt-2,5-thienylene).

favorable energy level alignment for electron injection. Onto the polymer PDHPT we deposited Ca, for the same reason as for 6P. In this case, where sulfur is a constituent in the polymer’s chemical structure (in contrast to 6P), we find indications for the formation of new covalent bonds. In order to investigate the above-mentioned interfaces, all experiments were conducted in UHV. Typically, in most cases, a thick film of a few hundred angstroms of 6P was deposited on a polycrystalline gold substrate; this was then characterized in situ using ultraviolet photoelectron spectroscopy (UPS) and X-ray photoelectron spectroscopy (XPS). Subsequently, the metal was deposited incrementally in small steps from a Knudsen-type cell in UHV (base pressure 1 ⫻ 10⫺10 Torr). Thickness of the 6P and metal overlayers was monitored with a quartz microbalance. However, no correction for sticking coefficient differences was made. UPS measurements and 6P and metal film depositions were conducted in a multichamber UHV system, either at the laboratory or at the FLIPPER II beamline at HASYLAB in Hamburg; details of the experimental setup have been described elsewhere [23]. Sample transfer from preparation to the analysis chamber was carried out without breaking UHV conditions. The pressure in the analysis chamber was maintained below 1 ⫻ 10⫺10 Torr. Changes in the electronic structure at the interface of the organic film and metal were monitored with spectroscopy after each incremental step of in situ deposition of 6P/metal. In some cases, the curves obtained by photoemission were analyzed using a curve fitting procedure. The curve fitting was done using a mixture of Lorentzian and Gaussian contributions to account for the intrinsic linewidths and instrumental broadening.

Organic–Metal Interfaces

II.

INTERFACES OF PARA-SEXIPHENYL WITH METALS

A.

Aluminum on p-Sexiphenyl

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Aluminum is widely used as a cathode in OLEDs, and it has proved to be a satisfactory choice, particularly in OLEDs of p-sexiphenyl (6P). Here we discuss a case where aluminum was evaporated onto a pristine 6P film in a stepwise manner, and interface change with increasing aluminum coverage was monitored using XPS. Figure 2 shows the evolutions of the Al2p (Figure 2a) and C1s (Figure 2b) spectra for increasing coverage of the 6P film. The thickness of the aluminum deposit was estimated from a quantitative analysis of the C1s/Al2p signal ratio, taking into account individual cross sections, and assuming a uniform layer. There is an apparent shift of the Al2p peak ˚ coverage to 72.75 eV for 15 A ˚ of alumaximum from 73.2 eV for 0.5 A minum on 6P. The value for the peak position after the final deposition step already corresponds to metallic aluminum and does not change for thicker metal films. In Figure 2b, where the C1s peak for each evaporation step is presented, no shift in the position is visible, and the curve fitting procedure yields the same value for all peaks to within 0.1 eV. Remarkable is the fact that even the lineshape of the C1s spectrum remains unchanged as we go from pristine 6P up to the maximum aluminum coverage. Using the curve fitting procedure outlined in the introduction and keeping all the fitting parameters constant, except for intensity, results in a fit that is in perfect agreement with the experimental data. Since this C1s peak is a convolution of the contributions from the different carbon atoms within the 6P molecule, this indicates that—at the actual instrumental resolution limit—no change in charge distribution at a preferred atomic site takes place upon aluminum deposition. Because there is no variation in the distribution of electron density for the carbon atoms, it can be concluded that no chemical reaction (i.e., formation of new bonds with charge transfer) occurs between 6P and aluminum, and that their interaction is of the pure van der Waals type (physisorption). The 0.5-eV shift of the Al2p peak toward lower binding energy for higher coverage can be explained by the fact that for very low coverage only few metal atoms are in intimate contact; thus, the electronic relaxation of the photoemission-induced core hole is not as efficient as for the infinite metal surface, and the measured binding energies for small clusters are higher. It is only after the last evaporation step, corresponding to a thickness ˚ , that truly metallic aluminum is formed to give the observed binding of 15 A energy of 72.75 eV. This behavior for thin metal films has been reported previously [24,25]. To gain more insight into the energy level alignment at the 6P–aluminum interface and the change in the electronic structure of the oligomer

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Figure 3 UPS spectra of a 6P film for increasing aluminum coverage: The topmost curves exhibit the Fermi edge of Al; the photon energy is 80 eV.

due to aluminum deposition, a study that is analogous to that presented ˚ ) film of 6P was evaporated above was performed with UPS. A thick (⬇200 A ⫹ onto an Ar -sputtered stainless steel plate, and aluminum was deposited stepwise onto it. A photon energy of 80 eV was used for the excitation of photoelectrons (for spectra obtained at 32 eV photon energy, see Ref. 7). The growth study was carried out until the Fermi edge of metallic aluminum was clearly visible. Figure 3 shows the low binding energy spectra obtained in this experiment. The binding energies are presented relative to the Fermi level. The lowest curve in Figure 3 shows the UPS spectrum of pure 6P [26]. No shift for the whole 6P–derived spectral features is visible for increasing aluminum coverage. Also, the spectral shape does not change upon the deposition of Al; this is consistent with what is observed in XPS. Again, this indicates a very weak interaction of aluminum with 6P. Furthermore, from the absence of a rigid shift of the C1s and the valence band spectra we can conclude that no changes in the sample work function occur. In addition, after each evaporation step the sample was biased negatively with respect to the electron spectrometer for the measurements of the

< Figure 2 (a) XPS Al2p core level spectra and (b) XPS C1s core level spectra for increasing coverage (␪) of 6P with aluminum.

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facial chemical reactions makes it difficult to define the electronic structure at the interface; only a qualitative description can be given [31]. Because no such interaction takes place at the Al–6P interface, a quantitative energy level diagram of the aluminum–6P contact can be proposed (Figure 4). Since the electronic structure of 6P is practically not influenced by the metal, we assume that (a) the bulk work function for aluminum (literature value of about 4.2 eV [32]); (b) the first ionization potential of 6P constant at 6.1 eV for any aluminum coverage, as determined with UPS by the low kinetic energy cutoff, and (c) the optical energy gap of 6P unchanged at 3.1 eV [33]. The energy level alignment at the interface yields an almost common vacuum level for aluminum and 6P, the estimated difference being smaller than 0.1 eV (within the experimental accuracy for this value of 0.2 eV). It can be seen from this estimation that there remains a barrier ⌬ = 1.3 ⫾ 0.2 eV for electrons to be injected at the Al–6P interface.

B.

Interfaces of Samarium and para-Sexiphenyl

1.

para-Sexiphenyl on Samarium

With a work function of about 2.7 eV [34] samarium would be expected to be a promising metal for use as a cathode material in OLEDs. To test this hypothesis we investigated the energy level alignment at the interface between 6P and Sm. First, 6P was evaporated stepwise onto a metallic polycrystalline Sm film (measured work function: 2.8 eV), and the respective UPS spectra were recorded. The spectra obtained in this way can be observed in Figure 5, with the Fermi level as binding energy reference. For a ˚ we observe a strong attenuation of the emission 6P coverage of less than 8 A ˚ 6P we can distinguish the photoemission charfrom Sm. Starting from 8 A acteristics of the valence levels of the organic material on this energy scale, and even HOMO and HOMO-1 (at 3.5 and 4.0 eV binding energy, respectively) can be discerned. Subsequent deposition of 6P results in an increase of its photoemission intensity and the signal in the 6P energy gap from the substrate vanishes completely. Notably, there is absolutely no binding energy ˚ ) up to shift of the 6P levels from the beginning of their appearance (8 A ˚ ). On an extended the maximum film thickness in this experiment (130 A binding energy scale, the emission from the localized 6P orbitals (at a binding energy of 5 eV and more) can be seen from the very beginning of the ˚ ). Not shown is the secondary electron cutoff, deposition sequence (1 A which was also recorded for each step. Its energy value does not change within the experimental error for this position (less than 0.2 eV). The ionization potential of 6P determined under these experimental conditions was 5.8 eV. These results are interpreted as follows:

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Figure 5 UPS spectra of a metallic samarium film for increasing coverage (␪) with 6P; the photon energy is 32 eV.

Given that the sticking coefficient of 6P on Sm is similar to that on 6P (the quartz microbalance was covered with 6P in previous experiments), we can conclude that the first monolayer of 6P is lying flat on the samarium film, as, for instance, on gold [35]. This is indicated by the very strong ˚ 6P (one monolayer of standing attenuation of the Sm photoemission at 8 A ˚ ) and by a complete 6P molecules would correspond to approximately 30 A ˚ screening of the Sm signal for more than 32 A of the overlayer. The exact nature of the interaction between Sm and the first 6P monolayer cannot be derived from UPS alone, but the appearance of clear 6P features in the ˚ indicates a rather weak interaction. This is supcoverage range below 8 A ported by the fact that the energy position of the secondary electron cutoff does not change notably, i.e. the work function of the sample remains constant. This also rules out the formation of a significant interface dipole, again pointing to a weak interaction (physisorption). Furthermore, increasing the 6P thickness does not change the energy level alignment. These experimental findings are a clear demonstration that the substrate work function and an optional interface dipole alone determine the energy level alignment at a COM–metal interface. In the present case, due to a very weak interaction and the lack of an interface dipole, we have a case of practically ‘‘vacuum level alignment’’ between the two materials. This results in the rather un-

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expected small work function for the 6P film (2.8 eV). Consequently, the HOMO onset of 6P is found far from the Fermi level of the metal substrate: 3.1 eV. Interestingly, this value coincides with the energy spacing between the electronic ground state and first optically excited state of 6P (in an approximation corresponding to the HOMO-LUMO separation). This could indicate that the lowest unoccupied molecular orbital of 6P interacts with the delocalized Sm electrons close to EF. More important for the application in devices is the fact that the LUMO of 6P is found at (or very close to) the Fermi level of samarium. This leaves a negligible injection barrier for electrons into the COM. However, what is true for the case of 6P on Sm could differ for the reverse case. In the next section, we show that the weak interaction is also sustained when Sm is deposited onto 6P film. 2.

Samarium on para-Sexiphenyl

˚ -thick 6P film on Sm from Figure 5, increasing Starting from the 130-A amounts of Sm were deposited stepwise and UPS spectra recorded. These are shown in Figure 6, with EF as binding energy reference. In principle, we simply observe the reverse sequence as depicted in Figure 5. The photoemission from 6P does not exhibit any shift or lineshape change, and is attenuated as the Sm deposit grows thicker. In the empty energy gap of 6P

˚ thick 6P film on Sm for increasing coverage Figure 6 UPS spectra of a 130 A (␪) with samarium; the photon energy is 32 eV.

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˚ Sm coverage), which gains we see an emission growing (even at 0.2 A ˚ intensity with the amount of Sm present on the sample. Eventually, at 23 A ˚ Sm on 6P we observe an emission at EF, and for 60 A Sm and more, a clear ˚ coverage onward no photoemisFermi edge from Sm is visible. From 23 A sion signal from 6P is detected anymore. Once more, the secondary electron cutoff shows no shift within experimental error, indicating no change of the sample work function. Because the valence emission from 6P undergoes no changes due to the Sm deposition (except attenuation), we conclude that the interaction is very weak (i.e., Sm physisorbs on 6P). The broad emission in the binding energy region below 3 eV is believed to stem from Sm, as it eventually develops into the metallic signature of samarium. We can rule out any ionic (doping-like) interaction between 6P and Sm because in such a case we would expect a clear attenuation of the HOMO relative to the HOMO-1, plus a measurable change in the secondary electron cutoff position (cf. Ca on 6P). Presumably, Sm does not grow in a layer-by-layer mode ˚ coverage only a faint emission at EF is detected. This because at even 23 A would point to the formation of small Sm clusters at lower coverages, which have not yet developed real metallic character. However, as in the case of 6P on Sm, the HOMO onset of 6P is found 3.1 eV below the Sm Fermi level. The work function of the thick samarium film was determined to be 2.7 eV. The difference of 0.1 eV with respect to the Sm film from the previous section can easily be explained by a slight change in the film morphology. (The work function of a metal surface is extremely sensitive to even small changes in the surface structure [36,37,38].) We would like to mention that in another Sm deposition experiment we started with a 6P film on native silicon oxide as substrate. There we arrived at essentially the same energy level alignment as described here. The energy level diagram of the Sm–6P interface is shown schematically in Figure 7.

Figure 7 Schematic energy level diagram of the Sm–6P interface; vacuum level (Evac) alignment is achieved within the experimental error (⫾0.1 eV); the injection barrier for electrons from Sm into the 6P LUMO is estimated to be close to 0 eV.

Organic–Metal Interfaces

C.

217

Calcium on para-Sexiphenyl

For systems with a nondegenerate ground state the formation of singly charged (polaron) and doubly charged (bipolaron) states has been investigated extensively in many conjugated polymers and oligomers, and the transition from polarons to bipolarons has been reported [39,40]. This last issue is important because one would expect to have polarons in low doping regimes, whereas bipolarons often seem to be the most stable configuration for higher doping levels [41]. There is also evidence for the formation of charged states higher than the bipolaron, i.e., more than two electrons on one sexiphenyl molecule [42]. By definition, a negative polaron results from filling the former lowest unoccupied molecular orbital (LUMO) of a conjugated chain with one electron, leading to a half filled orbital. The consequence is a destabilization of the former highest occupied molecular orbital (HOMO) and a stabilization of the former LUMO, yielding new electronic states in the normally forbidden energy gap. A finite density of valence states (DOVS) is expected to be found at the Fermi energy EF [43,44]. In the case of a bipolaron, the former LUMO is completely filled with two electrons, and EF is expected to be located well above the occupied electronic levels. Quantum chemical calculations show that the two new intragap states of the bipolaron are located deeper in the gap than the polaronic states [45]. A schematic energy level diagram for pristine, singly and doubly charged nondegenerate ground state

Figure 8 Schematic energy level diagram for a nondegenerate ground-state conjugated organic molecule in the neutral state (N), with one excess electron (negative polaron, nP) and with two excess electrons (negative bipolaron, nBP). The binding energy scale only reflects the general trend in such systems. The dashed line indicates the expected position of the Fermi level for a homogeneous ensemble of the individual species.

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conjugated molecules is shown in Figure 8. A finite density of states at EF in photoemission experiments performed with doped samples is generally attributed to the presence of polaron states. Ca is one of the low work function metals that is commonly used as a cathode material in OLEDs. It has been demonstrated that if Ca is used as the electron-injecting contact in OLEDs with poly(para-phenylenevinylene) (PPV) as the active material, the performance of the device is very sensitive to the presence of oxygen at the metal–polymer interface [46]. Without the formation of a calcium oxide interface layer the performance is rather poor. The reason seems to be that in the case of pure Ca the metal diffuses significantly into the active material and dopes the polymer in the interface region. In an ultraviolet photoemission spectroscopy (UPS) study of Ca on PPV the newly emerging intragap states were attributed to the formation of negative bipolarons [29,47], even at very low concentrations of Ca on the polymer surface. On the other hand, in another experiment rubidium was deposited onto PPV and a model oligomer thereof [40], and evidence was found that for the low doping regime polarons are formed, whereas for higher doping levels the intragap emissions are attributed to bipolarons. This assignment was based on the observation that the Fermi level was located at the position of the observed new DOVS for the low Ca concentration and that EF was found below this DOVS position at higher doping levels. This leads to an interpretation in terms of a polaron–bipolaron transition. A similar experimental result was obtained by Steinmueller et al. [39]. When Cs was evaporated onto a bithiophene thin film, the evolution of two energetically well-separated intragap emissions on the low binding energy side was clearly observed. Their relative intensities were strongly dependent on the amount of Cs present on the sample, clearly indicating a transition from one doping regime to another, presumably polarons and bipolarons. One drawback of these experimental observations is that, even in the polaron domain, EF is located below any DOVS, which is attributed to an incomplete screening of the hole left behind by the emitted photoelectron. The valence electronic region of a pristine 6P film on Au is seen in the lowest curve in Figure 9. The binding energies in this figure are given with respect to the vacuum level, set at zero. As before, we find this value to be 6.1 eV for pristine 6P films. Ca was deposited stepwise onto this clean 6P/Au film, and the photoelectron spectra recorded after each subsequent step are presented in Figure 9. A close-up of the frontier orbitals is plotted in Figure 10. The first deposition step of Ca on 6P was chosen to be ex˚ . This corresponds roughly to one Ca atom per tremely small, i.e., 0.05 A 100 6P molecules, assuming standing molecules and no Ca diffusion. Subsequent Ca depositions were done with an increasingly larger step size. In the survey plot (Figure 9) one can observe that after having deposited 1.5

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Figure 9 UPS spectra of a 6P thin film on Au with increasing amounts of Ca (␪) deposited on it. The binding energy is given with respect to the vacuum level, set at zero; the photon energy is 32 eV.

˚ Ca on the 6P surface new occupied electronic states appear with signifiA cant intensity in the formerly empty gap. The new low binding energy emission is centered at 3.75 eV; the binding energy for the other new state cannot be estimated accurately because it is located only slightly below the HOMO peak of the pristine 6P. These emissions continue to become more intense ˚ of Ca coverage, the overall for higher coverage. In addition, after 10 A spectral features of 6P become broader and the whole spectrum shifts to lower binding energy by about 0.2 eV (with respect to the pristine organic ˚ Ca and more, the emission from Ca begins to dominate film). With 20 A the spectrum. However, it is not yet in a metallic state because no clear ˚ is the coverage Fermi edge can be observed. Here it must be noted that 20 A indicated by the microbalance. This does not necessarily mean that a uniform film of this thickness is formed. A close-up of the gap region of 6P (Figure 10) provides further insight into the evolution of both the intensity and energy position of the newly occupied electronic states due to the deposition of Ca on the organic surface. A bar indicates the position of the Fermi energy ˚ Ca coverage the upper new for each spectrum. Now, even at only 0.05 A electronic state centered at roughly 3.9 eV becomes visible (albeit with a poor signal/noise ratio). With increasing coverage this emission grows in

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Figure 10 Close-up of the low binding energy region of Figure 9. The bar indicates the position of the Fermi level as determined in the UPS experiment.

intensity, gradually shifting to reach its final binding energy of 3.6 eV at a ˚ . Throughout the experiment, up to 10 A ˚ Ca on 6P, the coverage of 3 A ˚ Ca and more, there is lineshape of this emission does not change. For 10 A an observable change in the lineshape, and this change can be attributed to a superposition of the photoemission from the Ca-6P complexes and small Ca clusters, which will finally form the metallic Ca film. The increasing amount of Ca, leading to the new intragap states, has a significant effect on the position of the Fermi energy (EF) of the 6P/Au sample. For the pristine 6P/Au film EF is found 4.2 eV from the vacuum level (i.e., the work function is 4.2 eV), or 1.9 eV below the HOMO onset. After the first deposition of ˚ ), the work function of the sample decreases by a mere 0.05 eV. Ca (0.05 A Any further deposition of the metal causes a continuous shift of EF to lower ˚ covenergies, reaching a final value of 2.6 eV, which is achieved at 10 A erage. Since at all stages the new emission in the former energy gap of pristine 6P is present and it only shifts by roughly 0.3 eV (on the binding energy scale with respect to vacuum), EF is initially found within this peak, and then gradually moves out of it, until it is found well below any density ˚ . The metallic Fermi edge of occupied states for coverages higher than 5 A ˚ , and we of Ca becomes visible only for Ca coverage higher than 50 A measure a work function value of 3.2 eV for this metallic film.

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In another set of experiments, a thin film of 6P was evaporated onto metallic Ca and subsequently covered with increasing amounts of Ca. Figure 11a shows the photoemission spectra of this series. The Fermi level is chosen as energy reference, set at zero. For comparison, spectra of the series with Au as substrate are also shown in Figure 11b, on the same energy scale and for similar Ca coverage. We find that when 6P is evaporated onto Ca, EF is located 2.7 eV below the HOMO onset (i.e., a work function of 3.4 eV). This must be attributed to a different interaction of the first 6P monolayer with metallic Ca as compared to that with metallic Au. In Ref. 35 it is shown in detail that 6P physisorbs on Au. We observed that photoemission spectra ˚ and less) on Ca (spectra not shown) exhibit marked of thin 6P films (20 A differences compared to thicker ones or those formed on Au. This clearly indicates a strong chemical interaction between 6P and a metallic Ca surface. The 6P molecules in intimate contact with the Ca surface determine the position of EF, and due to the low intrinsic defect concentration in 6P this position is pinned for the thicker film also, as is to be expected for large band gap materials [48], and as explained in the introduction. The fact that for 6P/Ca the work function of the pristine 6P is already small, the shift of

Figure 11 Close-up of the low binding energy region of Ca deposited onto (a) 6P/Ca and (b) 6P/Au. The binding energy is given with respect to the Fermi level, set at zero.

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the Fermi level for increasing Ca deposition shows a noticeably different ˚ Ca coverage EF behavior than when Au serves as substrate. Even at 0.5 A is found well below the new doping-induced DOVS; and for increasing Ca deposition the whole spectrum shifts slightly to higher binding energies (cf. Figure 11a). As one can see from Figure 11b, when the work function of the pristine 6P film is higher (Au as substrate), the new DOVS are found at ˚ . We propose the following interpreEF for even higher coverage than 3 A tation of the experimental observations for the two different substrates: When the former LUMO of 6P is filled with one electron a negative polaron (nP) is formed. This leads to stabilization (increased binding energy) of this newly occupied state and destabilization (lower binding energy) of the former HOMO. Due to the fact that this molecule now has a half-filled orbital, DOVS at the Fermi level could be expected [44,49,50]. If a second electron is added to the molecule, the half-filled orbital will become a completely filled one and EF will be expected to be found between this orbital and the new LUMO (LUMO⫹1 in the neutral molecule), well below any DOVS. For this negative bipolaron (nBP) the new completely filled orbital is stabilized even more than in the case of the nP; the opposite holds for the former HOMO, i.e., it would be more destabilized. First we consider the experiment with Au as substrate. From the perspective of the movement of the Fermi level alone, one could readily conclude that we observe negative polarons during the initial stages of Ca deposition, since EF is found well within occupied electronic states (Figure 10). For a higher amount of Ca on the 6P surface EF moves toward lower binding energy, and is found well below the new intragap emission, indicating that the observed species might be negative bipolarons. What we observe then could be interpreted as a transition from nP’s in our sample with low Ca concentration to a situation with only nBP’s being present at high coverage with Ca, passing through intermediate stages where nP’s and nBP’s are present at the same time. However, considering the same experiment with Ca as substrate makes the validity of an nP-nBP transition questionable in the case of doping 6P with Ca, and offers another probable explanation. For high Ca coverage on 6P, we assign the intragap emission to a bipolaron for the following reasons: Since for high Ca coverage the Fermi level is found below any occupied electronic states, it is reasonable to assume that this emission is coming from the HOMO of negative bipolarons. The experimentally determined shift of EF and the observation of DOVS at EF for low coverages for 6P/Au could have an origin different from the formation of nPs. This becomes evident when investigating the spectral shifts observed for 6P/Au and 6P/Ca. By comparing the position of EF relative to the new doping-induced emissions for the two different experiments shown in Figure 11a and b, one must conclude that the presence of DOVS

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at EF in doped, wide-band-gap–conjugated materials does not automatically ˚ Ca is deposited onto 6P/ imply the presence of polarons. When up to 3 A Au a finite DOVS is observed at EF, and the presence of nP’s could be concluded. However, for similar doping concentrations no DOVS are found at the Fermi energy when Ca is used as substrate. It appears that for 6P/Ca EF of the pristine film is already very close to the position obtained for high ˚ Ca, and more for 6P/Au). Even at very low Ca coverage doping levels (5 A EF attains a position corresponding to that expected for nBP’s. The observation of EF within the new DOVS when Au is used as substrate must be attributed to the way the Fermi level is determined in photoemission spectroscopy for an inhomogeneous sample. The conclusion from our experimental results is that from the very beginning in both doping series negative bipolarons are formed. Only the shift of EF is different, due to the variation of initial work functions of 6P/Au and 6P/Ca. This is supported by the absence of a change in the lineshape of the new intragap emission throughout all doping levels. In addition, the formation of nBP’s is expected to be energetically favorable compared to the formation of nP’s. Here we also note that in a solid-state cyclovoltammetry experiment performed with 6P the first observed reduction wave was due to a two-electron transfer, indicating the formation of dianions (bipolarons) [51]. Furthermore, we can rule ˚ out a layer-by-layer growth of Ca on a 6P film. At a Ca coverage of 20 A we still can see a contribution to the photoemission spectra from the Ca-6P complexes. The elastic mean free path of the photoelectrons is very low in ˚ should considerably attenuate UPS, and hence a uniform Ca film of 20 A the Ca-6P emission. As this is not observed in the experiment, we can suppose that Ca either diffuses into the 6P film or grows in islands. Combining all the measured energy values for the interface of 6P with Ca from above allows us to draw a quantitative energy level diagram (shown in Figure 12), which enables the prediction of some charge injection properties of this interface in OLEDs. Due to the doping-like interaction of Ca with 6P, this interface consists of three different regions that must be considered: (a) the ‘‘pristine’’ 6P, (b) a region of Ca-6P complexes (doped region), and (c) metallic Ca. For the doped region we see that the peak maximum of the new intragap peak (the HOMO of the 6P bipolaron, BP-HOMO) is about 0.8 eV below EF. We can assume that the energetic separation of the BP-HOMO to the BP-LUMO (the LUMO of the 6P bipolaron) is of the order of a few tenths of an electron volt. Thus, the BP-LUMO would be very close to the Fermi level of Ca. Therefore, the injection of electrons from Ca into the doped region is expected to be rather efficient. From Figure 12 we also see that the BP-LUMO (not drawn, but most likely just above EF) and the LUMO of the pristine 6P molecule are at approximately the same energy. Consequently, the transport of electrons from the doped region

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Figure 12 Schematic energy level diagram of the Ca–6P interface. There are three regions to be considered: pristine 6P, Ca-doped 6P, and metallic Ca. For explanation, see text.

to the pristine 6P region is expected to be highly efficient as well. Since the overall injection barrier for electrons from Ca into 6P appears to be a few tenths of an electron volt, the use of Ca instead of Al (cf. Figure 4: 1.3 eV) as cathode material in OLEDs seems to be favorable. However, the possible diffusion of Ca into the bulk 6P film could result in a very extended interface region, where the electroluminescence could be significantly quenched. This would require some additional engineering of this interface (e.g., insertion of a diffusion barrier).

III.

INTERFACE OF Ca AND POLY (2,5-DIHEPTYL-1,4PHENYLENE-ALT-2,5-THIENYLENE)

Another completely different interaction between a metal and a COM is found for calcium on poly(2,5-diheptyl-1,4-phenylene-alt-2,5-thienylene) (PDHPT). PDHPT, with a chemical structure as shown in Figure 1b, is a blue light–emitting polymer with a band gap of approximately 3 eV and an ionization potential of about 6.2 eV. A thin film of this polymer can be obtained by spin-coating a solution of PDHPT in CHCl3. As mentioned earlier, due to its low work function, Ca is expected to be a good choice for a cathode as electron injection can be more efficient in low work function metals. However, there have been instances in which Ca has not proved to be a good cathode despite its low work function, e.g., an OLED consisting of ITO/PDHPT/Ca was found to exhibit poor performance [52]. This is one motivation to investigate the chemistry at such interfaces. An earlier investigation of Ca/polydiheptylphenylenevinylene (PDHPV) by Lazzaroni et al. [29,53], showed that the interface with Ca was ionic, and charge transfer from metal to the polymer was similar to ‘‘classical’’ n-doping of conjugated systems. Contrary to this result, we demonstrate here that the Ca/PDHPT

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interface shows no evidence of doping; instead, there is evidence of Ca-S bond formation and a disruption of the ␲-conjugated structure. This stark difference in the character of the interface of Ca with these two polymers could be attributed to the presence of sulfur in the backbone of the otherwise similar conjugated structures. A.

Pristine PDHPT Characterization

Figure 13 shows S 2p and C 1s spectra for a clean PDHPT film. A leastsquares fitting procedure has been used to find the energy positions of the component peaks in each of the spectra. To obtain an accurate picture of shift in the peak position of spectra taken subsequently during the experiment, we do not calibrate the binding energy positions with respect to alkyl carbon at 285 eV. It can be seen that the C 1s spectrum has three components: phenyl carbon at 285.66 eV, alkyl carbon at 285.95 eV, and carbon atoms bound to sulfur at 286.46 eV. Spin splitting of the S 2p state can be clearly seen with 2p 1/2 peak at 166.1 eV and 2p 3/2 peak at 165 eV. The relative energy positions of all peaks are in excellent agreement with published data [52]. A UPS spectrum of the pristine PDHPT film can be seen in Figure 14 (UPS data have been plotted on a kinetic energy scale). The

Figure 13 XPS spectra of S and C core levels of a pristine PDHPT film showing the individual components obtained by a least-squares fit.

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Figure 14 UPS spectrum of a pristine PDHPT film. The left side is the secondary electron cutoff, and the right side shows the valence level features; photon energy is 82 eV, and the sample is biased by ⫺17.6 V.

spectrum exhibits a sharp ␲-band edge and well-defined valence features, which are fingerprints of chemically pure ␲-conjugated polymers [31]. The lineshape of the spectrum is in very good agreement with both experimentally obtained and DOVS-derived energy bands (VEH quantum chemical calculations) [52]. However, the HOMO and the HOMO-1 are not as clearly resolved, which could be attributed to differences in order and thickness of the deposited film. The emission at lowest binding energy is attributed to states delocalized along the polymer backbone, giving rise to a dispersed ␲ band. A more detailed explanation of the valence level features can be found elsewhere [52]. The ionization potential (IP) obtained from our experiment at incident photon energy of 82 eV varies between 6.29 and 6.4 eV, which is larger than the calculated value of 5.29 eV [52]. We rule out charging in our experiments on the basis of our well-resolved XPS and UPS spectra of clean PDHPT films. In addition, in the XP spectrum there was no trace of oxygen or any other impurity, indicating a very clean pristine PDHPT film.

Figure 15

Continued

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of Ca. Based on this, along with a plot of the intensity of Ca and C core levels with respect to increasing Ca thickness (Figure 17), we speculate that Ca grows in the form of clusters. Upon heating the sample of PDHPT covered with a thick Ca layer, it was seen that there was a decrease in asymmetry and widths of the core levels and the ratio of Ca to C increased with respect to that before heating (Figure 18). Taking all these in conjunction we conclude that heating causes ‘‘wetting’’ of the PDHPT by Ca, i.e., we have a more uniform surface Ca layer [19]. These results are very different from those obtained with Ca on PDHPV, where it was observed that deposition of Ca caused ‘‘doping’’ of the polymer. This difference in interaction of Ca with these similar polymers could be attributed to the fact that PDHPT has a thiophene unit that is far more reactive than the vinyl unit of PDHPV. This change in ␲-conjugation along with other chemical changes is induced by interaction of Ca with S atoms. Diffusion of Ca into the polymer would render Ca an unsuitable candidate as the cathode material for a PDHPT-based OLED. This explains the observed short lifetime of such a device investigated in Ref. 52.

IV.

INTERFACES IN OLEDs

As we have shown in the preceding sections, the electronic structure of interfaces between metals and COMs can be well characterized by PES studies, which are carried out in UHV. However, almost exclusively, the fabrication process of OLEDs involves the evaporation of the top cathode (low work function metal) under less defined conditions, i.e., high vacuum (HV, typically 10⫺6 mbar). Naturally, these conditions can result in formation of an interface composed not only of the COM and the metal; water, oxygen, or any other contaminant from the residual gas in the recipient can also be incorporated at this interface. This will lead to a change in the electronic structure, and therefore charge injection properties, at this interface compared to the one composed of only two constituents. In order to investigate if incorporation of residual gas components into the interface takes place we prepared simple OLEDs with 6P as active material and Al as cathode (anode: ITO/glass). The first hint that oxygen actually is found at the Al/6P interface came from Rutherford backscattering (RBS) measurements (Figure 19) on such OLEDs [54]. Figure 19b shows the as-recorded RBS spectrum taken on a spot of an OLED prepared in HV in the following sandwich structure: glass/ITO/6P/Al. For interpretation of RBS data it is necessary to compare the experimental results with numerically simulated spectra. With certain assumptions on the geometrical and compositional properties of the

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Figure 17 Change in the intensity of C1s and Ca2p3/2 core levels as a function of Ca thickness as measured on the microbalance.

device, in our case a stacking as shown at the bottom of Figure 19, it is possible to obtain simulated spectra fitting the experimental data, shown in Figure 19a. Within this fitting procedure, peak F corresponds to ␣ particles backscattered by the In and Sn atoms in the ITO electrode. The 12-keV resolution of the detector is not able to resolve the two structures. The square-like structure E is assigned to the particles backscattered by the Al film on top of the 6P. The sharp step labeled D is the fingerprint of the particles backscattered by Si atoms in the glass substrate, while the extended edge around 360 keV comes from the oxygen atoms in both ITO and glass substrate. Finally, peak A originates from the backscattered particles by the carbon in the 6P region. Feature C corresponds to aluminum oxide on top of the Al contact. Within this fitting procedure, the feature appearing around 400 keV (labeled B) corresponds to oxygen localized at the interface between aluminum and the 6P layer. Apparently, due to HV conditions during the evaporation of Al onto the 6P film, oxygen is incorporated at this interface, most probably changing the electronic structure, and consequently electron injection properties, compared to the UHV study presented above. As the low signal/noise ratio in Figure 19 does not allow for a quantitative analysis of the oxygen content, a depth-profiling Auger spectroscopic investigation was performed on similar devices, including a variation in the Al evaporation conditions.

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Figure 18 XPS spectra of C1s (top), Ca2p (middle), and S2p (bottom). The spec˚ Ca coverage as read on the microtra labeled (a) are for a PDHPT film with 15 A balance, (b) the same film left in UHV (1 ⫻ 10⫺10 Torr) for 12 h, and (c) after this film has been annealed at 100⬚C for 30 min in UHV (1 ⫻ 10⫺10 Torr).

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Figure 19 Top: Simulated (a) and as-measured (b) RBS spectra of a conventional ITO/6P/Al LED; the peak assignment is explained in the text. Bottom: Side view of the LED model used to calculate the simulated RBS spectrum above.

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Single-layer OLEDs were produced by evaporating thin films of 6P (200 nm) onto ITO-covered glass substrates. The aluminum cathode was evaporated by resistive heating from a filament at a pressure in the vacuum system below 5 ⫻ 10⫺6 mbar. Keeping all parameters constant, except the evaporation rate of Al, we produced two different types of devices. For the ˚ /s while first type, the evaporation rate was chosen to be greater than 10 A ˚ for the second it was 1 A/s. Right after production, the OLEDs were characterized for their performance, i.e., current–voltage (I-V) and electroluminescence–voltage (E-V) curves recorded under ambient conditions. The curves for the two different device types are presented in Figure 20. Figure 20a represents the curves for a device with the fast evaporated Al cathode, and Figure 20b for a device with the slowly evaporated metal. The most striking difference is the enormous reduction of the onset voltage for electroluminescence from roughly 15 V down to 3.5 V for the device with the slowly evaporated cathode (type 2). At this voltage, the blue 6P electroluminescence emission could be seen by the naked eye in a dark room. In addition, the behavior of the current vs. driving voltage changes significantly. The onset of the current is around a bias of 10 V for the device of type 1, far away from the onset voltage of the electroluminescence. This finding supports the common consideration that one species of charge, i.e., holes, are injected (from the ITO) and transported more efficiently in conjugated materials of this type. At higher voltages electrons are injected as well and can travel to the zone of radiative recombination in the active layer. But for the device where the Al cathode was evaporated slowly, the current and electroluminescence start to rise practically simultaneously. Therefore, one can assume that in this device carrier injection and transport properties of electrons and holes do not differ significantly. As a consequence, the power efficiency of the device of type 2 is increased. For instance, at a power consumption of 50 mW the luminance of devices with the slowly evaporated cathode is seven times higher than for the devices with Al deposited at a faster rate. The maximum external quantum efficiency of these single-layer devices is increased by the careful deposition of aluminum, namely, from 1% to 2%. Therefore, the only difference in these two OLED types must be located at the Al–6P interface. To investigate this interface, elemental depth profiling was carried out on devices with a scanning Auger microprobe sputtering with 3-kV Ar ions, while recording Auger electron spectra (C KLL, O KLL, Al LMM). In Figure 21 elemental depth profiles of OLEDs of type 1 and 2 are presented. Figure 21a shows the depth profile of a device with a fast deposited cathode and reveals that there is a rather broad region (of about 57 nm within a rough estimation) at the organic–metal interface where aluminum, carbon from the 6P, and oxygen are present simultaneously (marked as IF in Figure

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Figure 20 Current–voltage (open squares) and electroluminescence–voltage (solid circles) curves for (a) normal LED and (b) LED with the slowly evaporated Al cathode. Insets show the EL onset on a logarithmic scale.

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Figure 21 Elemental depth profiles (from Auger electron spectra) of the AL–6P interface (marked IF) in the (a) LED with Al deposited at a high rate and (b) LED with the slowly evaporated Al cathode: thin dashed line, indium; thick solid line, oxygen; dotted curve, aluminum; dash-dotted curve, carbon.

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21). Such an intermixed region can also be found in a device where the top electrode was evaporated slowly. As can be seen in Figure 21b, the thickness of the interface region (corresponding to 5–10 min sputter time) is smaller by far (approximately 30 nm), resulting in a higher effective thickness of the pure active material, and in a higher oxygen concentration. The presence of the interfacial oxygen cannot be attributed to oxygen adsorbed on the 6P surface during transfer in air before metallization. It is known that when a 6P film is exposed to air and introduced into vacuum subsequently, only small amounts of oxygen remains adsorbed on the surface of 6P [54]. The diffusion of oxygen through the aluminum layer from atmosphere or through the 6P from ITO can also be ruled out as the cause; in a series of controlled experiments the interfacial oxygen content was found to be independent of time or atmospheric exposure of the devices over a time scale (10⫹ days) much greater than that of the experiments reported here. Another difference between the two device types is the oxygen content within the aluminum layer: while no oxygen is found in the Al electrode for device type 1, there exists a residual O signal throughout the cathode for the slowly evaporated metal (type 2). In addition, the atomic ratios (i.e., Al:O:C) within the interface region differ significantly. In the device where the cathode was deposited with a high rate, the maximum oxygen content at the interface was found to be approximately 5%. However, in the device with the slowly evaporated cathode the ratio of oxygen at the maximum exceeds 20%. It thus appears that the reason for the O signal within the Al layer, and the higher interfacial oxygen concentration, is the slow evaporation rate in a moderate vacuum (5 ⫻ 10⫺6 mbar), giving enough time for oxygen to be incorporated during Al deposition, which is not the case for the fast evaporation. Although oxygen concentrations as high as 20% were found at the interfacial region, the Al LMM Auger spectra clearly indicate the almost total absence of oxidized Al. This unexpected finding can be rationalized as it is known that coverages greater than 30% are required before true aluminum oxide is formed when oxygen is adsorbed on Al(111) [55], while at lower coverages oxygen is present in a chemisorbed state. Therefore, it is not unreasonable that oxygen/aluminum ratios of the magnitude of those found in the devices do not give rise to oxide formation. However, the presence of chemisorbed oxygen can have a significant effect on the electronic structure of Al–organic interfaces [56]. It can be concluded that a major part of the voltage drops over region IF (Figure 21). It has been shown that the deliberate introduction of an insulating layer at the organic– cathode interface improves the overall electron injection efficiency [10,11]. However, in our case the thickness of the layer found to be present at the interface is very large, in comparison with the optimum values given in Refs. 10 and 11, and the Al at the interface is still metallic.

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OUTLOOK

The examples of possible interactions between conjugated organic materials and metals presented here demonstrate that by choosing appropriate material combinations one can control the chemistry and tune the energy level alignment at interfaces over an extremely wide range. Undoubtedly, organic molecules will be employed as the functional units in future electronic and optoelectronic devices. Recent reports of an electrically pumped organic laser [57], highly efficient organic solar cells achieving 4.7% [58] up to 7.3% [59] conversion efficiency, and the observation of fractional quantum Hall states in an organic molecular crystal [60], as well as superconductivity [61], have catalyzed research on similar systems. At the same time, the size of devices will shrink to the subnanometer level and thus ‘‘interface-only’’ structures will result. Therefore, a complete understanding of processes and properties of the interfaces involved is essential for this newly evolving field. Photoemission spectroscopy will be one of the important tools to elucidate interfacial properties. ACKNOWLEDGMENTS The authors thank Prof. K. Kaeriyama for providing the polymer and A. Pogantsch, Dr. J. Ghijsen, Dr. E. Zojer, Dr. R. I. R. Blyth, Prof. M. G. Ramsey, Prof. F. P. Netzer, and Prof. R. L. Johnson for their support. A. R. acknowledges the financial support of the Interuniversity Research Projects (PAI/IUA) on ‘‘Sciences of Interfacial and Mesoscopic Structures’’ sponsored by the Belgian Prime Minister’s Office (Federal Services for Scientific, Technical and Cultural Affairs). REFERENCES 1. 2. 3. 4. 5. 6. 7. 8.

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8 Electronic Structure of Interfaces Between Organic Molecules and van der Waals Surfaces Toshihiro Shimada and Atsushi Koma University of Tokyo, Tokyo, Japan

I.

INTRODUCTION

Recent technological advances in organic electroluminescence devices have triggered intensive research on the electronic structure at the interfaces between organic molecules and inorganic materials. Since the charge injection characteristics in the devices are governed by the electronic level alignment at interfaces, much effort in the surface scientific studies has been focused on the determination of the electronic structures. It has been done by various electron spectroscopic methods such as ultraviolet photoelectron spectroscopy, X-ray photoelectron spectroscopy, inverse photoemission, and scanning tunneling spectroscopy. Through experimental studies by several groups it has been found that dipole layers are formed at various organic–inorganic interfaces [1]. However, the precise mechanism has not been elucidated so far, and it is still impossible to predict the strength of dipole layers of a given interface. Suggested mechanisms include charge transfer, molecular orbital hybridization, polarization, and influence of metal-induced gap states, which are interesting by themselves as subjects in quantum chemistry. To study the mechanism of the dipole formation, it is essential to prepare atomically well-defined ultrathin films, hopefully on a series of structurally similar materials with different electronic properties. One of the most effective guidelines to obtain such interfaces is the epitaxial growth of molecules on surfaces without dangling bonds, which will be referred to as 241

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‘‘van der Waals surfaces’’ in the following discussion. Since there are no strong site-selective atomic interactions on van der Waals surfaces, various molecules can be grown with their own lattice constants, with the lattice orientation aligned with that of substrate surfaces [2]. In this chapter, we describe the preparation and electronic structures of organic molecular films on van der Waals substrates. Three materials systems are described to show typical examples of analysis scheme: C60 on MoS2 as HREELS application to highly symmetrical molecules, copper phthalocyanine (CuPc) on layered materials as a systematic study using substrates with different work functions, perylene tetracarboxylic dianhydride (PTCDA) on MoTe2 and TaSe2 as an application of ultraviolet photoemission spectroscopy (UPS), and quantum chemical calculation to clarify the precise mechanism of dipole layer formation.

II.

EPITAXIAL GROWTH ON VAN DER WAALS SURFACES

Most organic molecular crystals have lattice constants much larger than inorganic substrate materials. Thus, it is very difficult to find lattice-matched substrates, which is the usual requirement for good heteroepitaxial growth. This has been one of the major obstacles in epitaxial growth of organic thin films. The lattice matching condition is especially severe in heteroepitaxial growth between tetrahedrally bonded covalent semiconductors such as GaAs and Si. This comes from the fact that dangling bonds on the surface of a substrate material can be connected only to the atoms of well lattice-matched materials because the length and the angle of the covalent bonds cannot be changed easily. The lattice matching condition, however, can be relaxed dramatically when heteroepitaxial growth is mediated by van der Waals interaction. A typical example is the growth of a layered material onto a cleaved face of another layered material having no dangling bonds. It has been shown that an ultrathin film with its own lattice constant can be grown by van der Waals epitaxy even under the influence of lattice mismatch as large as 300% [3]. It has also been shown that a very abrupt interface with small numbers of defects can be fabricated by van der Waals epitaxy because of the absence of the dangling bonds [4]. Since the molecules in the organic crystals are held by weak van der Waals forces, the idea of van der Waals epitaxy can be naturally extended to the epitaxial growth of organic thin films, with which the problem of the lattice matching can be resolved. We have attempted molecular epitaxial growth of organic molecules on various van der Waals surfaces with different lattice constants and made a systematic study of the epitaxial growth features [2]. In this chapter we focus on the growth of molecular films on the layered materials with dis-

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tinctive electronic structures to study the interaction at interfaces by electron spectroscopies. Readers are referred to the literature in the ‘‘Tips for Beginners’’ section about the organic epitaxial growth on van der Waals surfaces.

III.

EXPERIMENTAL AND COMPUTATIONAL TECHNIQUES

A.

Preparation of Substrates

Van der Waals surfaces, namely, single crystalline surfaces on which no dangling bonds exist, are classified into two categories. The first one is the cleaved faces of the layered materials, such as graphite, mica, and metal chalcogenides. Most of them have sixfold symmetry in the atomic arrangement on topmost surfaces. In particular, layered metal chalcogenides are a group of materials possessing various electronic properties ranging from insulator to superconductor [5]. They have various hexagonal lattice constants from 0.3155 nm of WS2 [6] to 0.405 nm of InSe [7] and they show variety of physical properties and work functions [8,9]. A single crystal of MoS2 can be obtained as a natural mineral molybdenite, while those of other layered metal chalcogenides can be prepared by halogen vapor transport in a sealed quartz tube [10]. They can be cleaved easily to bare its (0001) surfaces. Surfaces can be cleaned by cleavage in ultrahigh vacuum (UHV) or, depending on the materials, cleavage in air followed by heating in UHV. The second category is surfaces of three-dimensional materials with the dangling bonds terminated. The termination is sometimes accomplished in solution simultaneously with etching (e.g., hydrogen-terminated Si(111)), and sometimes done by dosing terminating atoms onto the surfaces already cleaned in UHV. The details of termination technique will also be given in ‘‘Tips for Beginners’’ section. In the experiments described in the present chapter, layered materials [highly oriented pyrolytic graphite (HOPG), MoS2, MoTe2, and TaSe2] are used as the substrates. The lattice constants, work functions measured by the present authors, and electronic properties are summarized in Table 1.

Table 1

Parameters of substrates used in the present experiments

Parameter

(␣-)MoTe2

(2H-)TaSe2

Graphite (HOPG)

Lattice const. (in-plane) a Work function

0.352 nm 4.1 eV

0.344 nm 5.5 eV

0.245 nm 4.5 eV

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Film Growth and Structure Characterization

The molecules used in the present experiments are shown in Figure 1. Those molecules are evaporated from ‘‘Knudsen cells,’’ or heated crucibles with diameter smaller than length. The Knudsen cells used in the present experiments are equipped with a load-lock mechanism, which enables one to retract the cell from the main growth chamber during the bake-out of the vacuum system. The growth rate and thickness of the film were measured by quartz crystal microbalance located near the substrate. The substrate was introduced via an entry chamber separately pumped by a turbomolecular pump. The experimental setup is shown schematically in Figure 2. The atomic and molecular arrangement on the substrate surface and films are characterized by reflection high-energy electron diffraction (RHEED). Basically, streak patterns are obtained if the topmost surface is flat and crystalline. Due to the different symmetry of the unit cells of the molecular film and of the substrates, equivalent rotated domains can exist in many of the organic–inorganic heteroepitaxial systems. Thus, precise examination of in-plane ordering of domain structure is necessary, which can be done by changing the azimuth of incident electrons. In extreme cases, only one axis perpendicular to the surface is aligned, which gives ‘‘columnar’’ RHEED patterns unchanged by the sample rotation around the surface

Figure 1 Molecules and substrates presented in this chapter. (a) C60, (b) CuPc, (c) PTCDA, (d) TaSe2, (e) MoTE2, (f) graphite.

van der Waals Interfaces

Experimental setup. (a) HREELS system; (b) UPS system.

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Figure 2

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normal. Aligned sot patterns are obtained when the films have ordered crystalline island structures, whereas ring patterns like Debye-Scherrer ring in X-ray diffraction is observed when the film consists of totally disoriented polycrystals. A halo pattern is observed when the surface is amorphous. We have found that the use of microchannel plate (MCP) to reduce the incident current less than 1 nA is particularly useful for the RHEED observation of organic films because it drastically reduces the electron bombardment effects [11]. C.

Electron Spectroscopy

Several electron spectroscopic techniques were used in the present experiments. High-resolution electron energy loss spectroscopy (HREELS) was used to measure the vibrational excitation of the surfaces. Ordinary commercial HREEL spectrometer (Eiko Engineering EES-10) is equipped with two sets of double monochrometers by which incident electrons are monochromertized and the electrons reflected from surfaces are detected with energy resolution of ⬃10 meV. The incident and exit angle of electrons can be changed by rotating the monochrometers and the samples. The term ‘‘specular’’ is used when the incident and exit angles are the same, while ‘‘off-specular’’ is used when they are not the same. In the specular arrangement, only dipole-allowed vibrational modes can be observed. In the offspecular arrangement, the dipole selection rule is broken and all the vibrational modes can be observed in principle. Thus it is possible to identify the vibrational peaks by comparing the spectra taken with specular and offspecular arrangements. UPS is used to measure the occupied states of the substrate and organic molecules. The equipment used in the present study was Vacuum Generators ADES 500 spectrometer with He discharge lamp. He II line (h␯ = 40.8 eV) was used in the present UPS measurement. X-ray photoemission spectroscopy (XPS) was measured by a double-pass CMA (PHI 15-255G) with Mg K␣ X-ray source (h␯ = 1253.6 eV). D.

Work Function Measurement

Work functions of the samples including that of substrates were determined by the low kinetic energy cutoff in the UPS spectra using He I photons (h␯ = 21.2 eV). In order to avoid the spectrometer threshold effect, the sample was biased electronically from the ground level. Values for work functions measured from identical samples were consistent within ⫾0.05 eV when the bias voltage was changed from 1.5 to 7.5 V. Photovoltaic effects can be ignored in our UPS measurements because no binding energy shift occurred

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when the photon intensity of the discharge lamp was changed by a factor of 3. E.

Observation of Surface Morphology by Atomic Force Microscopy (AFM)

Morphology of the films is important to analyze the results of electron spectroscopy. An atomic force microscope (Seiko SPI-3800) was used in contact mode in ambient atmosphere to observe the surface morphology of the grown films. F.

Quantum Chemical Calculation

Recently the computer power available for the quantum chemical calculation has been drastically increased. It is now possible to calculate the electronic structures of a moderate-sized molecule such as PTCDA in laboratories with tolerable accuracy which can be compared with the experiments. We have used Gaussian 98W [12] for the computation.

IV.

ELECTRONIC STRUCTURES AT VAN DER WAALS INTERFACES

A.

Characteristic Examples of Electronic Interaction

We will describe three typical experimental results on electronic interaction between molecules and substrates in the following. First, we examine C60 grown on MoS2(0001). Owing to its highly symmetrical shape, intramolecular vibrational modes of C60 have characteristic selection rules which can be determined by the measurement of HREELS. The validity of the selection rule can be a good measure of the molecular environment at the interface. The amount of charge transfer will be derived by comparing the vibrational energy in the films with infrared spectra of alkali-doped C60. Second, we show a combined analysis of XPS, UPS, and work function measurement using van der Waals substrates with different electronic properties. CuPc was chosen as the molecule because of the well-known electronic structure and the epitaxial growth features on layered materials. The third example is PTCDA grown on two layered materials (MoTe2 and TaSe2) with similar lattice constants (0.352 nm and 0.344 nm, respectively) and different electronic properties. Since the size of the molecule allows us to perform various quantum chemical calculations, detailed comparison with angle-resolved UPS has become possible.

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HREELS Study of C60 Ultrathin Films Grown on MoS2: Symmetry Lowering Due to the Molecule–Substrate Interaction [13–15]

Fullerenes have a symmetrical structure and characteristic vibrational modes. From the selection rules in specular scattering it will be possible to detect the distortion of the molecules which is difficult in other systems and techniques. It is also known from infrared (IR) spectroscopy of alkali-metaldoped fullerenes that the vibrational energy is shifted when charge transfer occurs in fullerenes. It is therefore expected that systematic study of HREELS of fullerenes will reveal the interaction between molecules and surfaces. C60 grows epitaxially on MoS2(0001) almost in layer-by-layer growth mode [16,17]. Since thick films are expected to have no interaction with the substrate, HREELS spectra of 20-nm-thick films were taken as reference (Figure 3). The spectrum taken with specular configuration (a) represents the IR-allowed modes [18], while that with off-specular configuration (b) agrees well with neutron scattering data which represents total density of states of phonons [19]. In order to characterize the interaction between C60 and MoS2, HREELS of ultrathin C60 films were measured. Figures 4 and 5 show onemonolayer-equivalent (MLE)-thick and 2-MLE-thick samples, respectively. As for the specular configuration, strongly observed peaks at 66, 147, and 177 meV also exist in the spectrum from 20-nm-thick C60 (Figure 4a) and the intensity ratio are nearly the same. It indicates that IR-allowed modes of C60 do not change their energies and dipole strength at the interface. Spectra taken with off-specular configuration also show similar shapes between 20 nm and 1 ML, which suggests no strong chemical bonds due to double bond breaking or large charge transfer. However, the spectrum of the 1-ML-thick sample taken with the specular configuration shows a large peak at 158 meV. It should be noted that this peak is not observed in that of the 2-MLE- and 20-nm-thick samples taken with the same configuration. It means that the peak at 158 meV is dipole forbidden in the symmetrical molecule and appeared in the 1-MLE-thick film because the symmetry of the molecule is lowered at the interface. The vibration frequency of C60 – MoS2 interaction is too low to couple with the intramolecular modes of C60 because the former is of weak van der Waals nature while the latter are due to covalent bonding. Therefore, the mechanism of symmetry lowering is not the vibrational coupling with the substrate but must be due to static distortion of the C60 molecule caused by the electronic interaction between the molecule and the MoS2 surface. Since no strong charge transfer is suggested because no shift in vibration energy is observed in the off-specular spectra

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Figure 3 specular.

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HREELS of 20 nm-thick C60 grown on MoS2. (a) Specular; (b) off-

of 20-nm-thick and 1-MLE-thick films, it does not come from charge transfer. Polarization of C60 molecules on surfaces has been pointed out in the case of another van der Waals system of C60 /GeS(001) [20,21] from partial shift of XPS C 1s spectra. It can contribute to the symmetry lowering of C60. Although it has been attributed to the external electric field on the substrate surface [21], we consider the symmetry lowering as a result of hybridization of wavefunction at the interface in the case of C60 /MoS2(0001) for the following reasons. The strength of the van der Waals interaction can be estimated using the ‘‘wetting’’ behavior of C60 grown on surfaces. C60 / MoS2 shows good wetting (layer-by-layer growth) while C60 /alkali halides and C60 /CaF2(111) show poor wetting property (island growth). It is there-

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Figure 4 specular.

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HREELS of 1-ML-thick C60 grown on MoS2. (a) Specular; (b) off-

fore considered that the external electric field, which is stronger on alkali halides and CaF2(111) than on MoS2, is not related to the interface energy, which determines the wetting properties. Here it should be noted that the wetting properties involve the balance between interface energy and bulk and surface energies of constituent materials. Since the bulk and surface energies of C60 and the substrates are not expected to be much changed by the formation of the interfaces, the wetting property is a good measure to estimate the interface energies or the strength of van der Waals interactions. Thus, it is concluded that the molecule–surface interaction is not caused by external static electric field but by some hybridization of wavefunctions. This mechanism suggests that symmetry lowering of C60 is a result of van der Waals interaction between C60 and MoS2.

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Figure 5 specular.

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HREELS of 2-ML-thick C60 grown on MoS2. (a) Specular; (b) off-

In summary, static distortion of the molecular shape at the interface was detected by HREELS in monolayer C60 /MoS2 system owing to the highly symmetrical shape of C60. Hybridization of wavefunctions directly related to van der Waals attraction is suggested as the mechanism. Precise measurement by other spectroscopic technique using various substrates is under way. C.

XPS and UPS Study of CuPc/Layered Materials: Hybridization at a van der Waals Interface [22]

We can expect that molecule–surface interaction will be strongly observed when the molecule has a planar shape and lies flat on a substrate surface.

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Figure 6 XPS and UPS of CuPc films grown on HOPG. (a) HOPG substrate. (b) to (h) are spectra of CuPc films on HOPG. CuPc thickness is (b) 0.15 nm, (c) 2 nm, (d) 5 nm, (e) 15 nm, (f) 0.3 nm, (g) 0.6 nm, (h) 15 nm.

Metal phthalocyanine is a group of such molecules with electronic structures modifiable by changing central metal atoms. We have chosen copper phthalocyanine (CuPc) as the first example because its electrochemical characters and epitaxial growth conditions are well understood [23,24]. We have used UPS and XPS to look at the electronic interaction in detail. For substrate materials, we selected three kinds of materials which have distinctly different electronic properties with various values of work function (␾); MoTe2 is semiconducting with ␾ = 4.0 eV, highly oriented pyrolytic graphite (HOPG) is semimetallic with ␾ = 4.5 eV, and TaSe2 is metallic with ␾ = 5.5 eV. The growth of CuPc films was carried out in a growth chamber with a base pressure of 3 ⫻ 10⫺7 and equipped with RHEED apparatus. The single crystalline substrates were grown using halogen vapor transport method [25] and cleaved under UHV conditions. CuPc was commercially obtained (Tokyo Kasei Co.) and purified twice by vacuum sublimation. It was evaporated from a Knudsen cell kept at the temperature between 380⬚C and 420⬚C, which resulted in a growth rate of 0.1–0.5 nm/min. The film thickness was monitored by a quartz crystal microbalance. Immediately after the epitaxial growth, the samples were transported to an analysis chamber without breaking UHV.

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All samples of thickness less than 15 nm showed streaky RHEED patterns, indicating epitaxial relation of the films with the substrates. The separation between streaks that had strongest intensities corresponded to the distance of about 2 nm. This was consistent with a distance of the second nearest neighbors in the two-dimensional quasi-square lattice of CuPc observed by STM [25]. It indicates that CuPc molecules are lying with molecular plane nearly parallel to (0001) surfaces of the substrates. The sharpness of the RHEED streaks also revealed that the crystallinity of CuPc films was excellent for all cases. When the thickness was more than 15 nm, the RHEED pattern became spotty but the crystallographic orientation was unchanged. It was found from AFM observation that the growth proceeds nearly in Stranski-Krastanov mode in which the first molecular layer covers the substrate almost completely and islands are formed on it. Figures 6, 7, and 8 show the thickness-dependent UPS and XPS spectra obtained from the films grown on HOPG, TaSe2, and MoTe2, respectively. The core levels of the substrates measured by XPS (Ta 4f, Se 3d, Mo 3d, and Te 3d) did not show any dependence on the CuPc film thickness. When HOPG or MoTe2 was used as the substrate, all peaks in UPS and XPS did not change their binding energies regardless of the thickness of overlayers. Note that the C 1s of CuPc could not be resolved from that of HOPG

Figure 7 XPS and UPS of CuPc films grown on TaSe2. (a) TaSe2 substrate. (b) to (h) are spectra of CuPc films on TaSe2. CuPc thickness is (b) 0.15 nm, (c) 3 nm, (d) 5 nm, (e) 40 nm, (f) 1.2 nm, (g) 1.8 nm, (h) 20 nm.

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Figure 8 XPS and UPS of CuPc films grown on MoTe2. (a) MoTe2 substrate. (b) to (h) are spectra of CuPc films on MoTe2. CuPc thickness is (b) 0.15 nm, (c) 2 nm, (d) 3 nm, (e) 40 nm, (f) 4 nm, (g) 8 nm, (h) 45 nm.

substrate. For MoTe2, all peaks were shifted upward by 0.4–0.5 eV as compared to the case of HOPG substrates. For TaSe2, the Cu 2p, N 1s, and C 1s XPS peaks were shifted by the same amount toward the higher binding energy side as the film thickness was increased. Total shift was about 1.0 eV when the thickness was changed from 1.2 nm to 20 nm. Similar trend in the valence levels measured by UPS was observed but their behavior was not simple; two peaks at 0.8 eV and 2.3 eV appeared uniquely independent of the thickness of CuPc thin film up to a nominal thickness of 3 nm (Figure 7b, c). Beyond this thickness, those peak heights were diminished and taken over by bulk CuPc peaks (e). Work function was also measured for each case (Figure 9). Work function for CuPc films grown on either HOPG or MoTe2 did not change their magnitude from the initial value of the substrate (4.5 eV for HOPG and 4.0 eV for MoTe2). When TaSe2 was used as the substrate, work function of the samples decreased from the substrate value (5.5 eV) steeply at first, and gradually then, as the CuPc thickness was increased. Work function approaches a value almost identical to that of CuPc grown on HOPG beyond 5 nm. Band bending of the semiconducting substrate, MoTe2, was negligible in the present case, since the core levels of the substrate elements did not

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Figure 9 Work functions of CuPc grown on (a) HOPG; (b) MoTe2; (c) TaSe2 at various thicknesses.

show any shift in XPS. From the results of photoelectron spectroscopy (PES) and work function measurements, we propose models for three respective substrates (Figure 10). From our measurements, a main feature can be accounted for by the difference of work function in the substrates. Since the binding energy positions for PES and work function values were nearly

Figure 10 Band alignment models for (a) HOPG, (b) TaSe2, (c) MoTe2 substrates. HOMO and LUMO represent the positions of maximum density of states. HOMO is obtained from UPS and LUMO is estimated from HOMO energy plus the optical gap in Ref. 28.

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reports that CuPc is a p-type semiconductor [24]. It is probably because our films are made in UHV with the different crystal structure. In summary, electronic structures of the interfaces between epitaxially grown CuPc films and layered materials were studied using UPS, XPS, and vacuum level measurements. Three types of behavior were observed depending on work function of the substrates. A flat-band condition was achieved on HOPG, whereas a physical adsorption model seems to be valid for CuPc grown on MoTe2. Interface dipole layer formation was observed at the epitaxial interface between CuPc and TaSe2, and the molecular orbital energies involved were identified. D.

PTCDA/TaSe2; Analysis by UPS and Quantum Chemical Calculation

In the previous section, the hybridization between the orbital wavefunctions of CuPc and TaSe2 is suggested from UPS results, but the precise mechanism was not determined because of the Stranski-Krastanov growth mode of the system and the difficulty of quantum chemical analysis due to the heavy atom in the molecule (Cu). In this section, we report on a similar study of PTCDA ultrathin films grown on cleaved surfaces of two kinds of layered materials: n-type semiconducting MoTe2 (lattice constant = 0.352 nm, ␾ = 4.1 eV) and metallic TaSe2 (lattice constant = 0.344 nm, ␾ = 5.5 eV). PTCDA is one of the best studied molecules as the model system of organic–inorganic interfaces, judging from the publication numbers (more than 120 papers) to date. In spite of the extensive studies on the molecular arrangement, electronic structure, and reaction at the interfaces, it seems that the molecular mechanism of the dipole layer formation is not understood even on this well-studied molecule. We here present the results on the interface between PTCDA and layered materials, in which no complication is expected due to the strong chemical bonding between molecules and dangling bonds on the substrate surfaces. The effect of the tunneling of free electron from metal surfaces and that of the chemical potential difference will be clarified by the combination of electron spectroscopy and quantum chemical calculation. RHEED was used to characterize the molecular arrangement in the films. Angle-resolved UPS was applied to measure the band alignment and molecular orbital levels at interfaces. Ab initio quantum chemical calculation was performed for the determination of molecular levels and to estimate the extent of charge transfer. PTCDA was purified by sublimation in a Pyrex tube pumped by a turbomolecular pump to 5 ⫻ 10⫺6 Pa. It was found that the vacuum range higher than ordinary purification for the organic sublimation was essential to obtain reproducible results. The green-brown shiny crystalline powder was

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characterized by mass spectrometry, differential thermal analysis, and elemental analysis to be confirmed as PTCDA without detectable impurities. The crystalline powder was charged in a Knudsen cell and thoroughly degassed by heating at 250⬚C in a UHV growth chamber. The substrate single crystals were mounted on a sample holder and cleaved in a vacuum better than 1 ⫻ 10⫺6 Pa, and quickly transferred to the growth chamber with a base pressure of 6 ⫻ 10⫺8 Pa. The amount and molecular arrangements on the substrate surface and films are characterized by RHEED during the growth. The substrates were kept at certain growth temperatures during the irradiation of PTCDA molecular beam. The thickness of the films and the growth rate were estimated by quartz crystal microbalance. Knudsen cell of PTCDA was kept at 290⬚C during the growth, at which time the growth rate was about 0.1 MLE/min. Streak patterns were observed after about 1 MLE of PTCDA was irradiated from both of the MoTe2 and TaSe2 substrates. The patterns observed from 1-MLE-thick PTCDA/MoTe2 indicated two types of nonequivalent lattices existing on the surface when the substrate was kept at room temperature during growth. However, when higher substrate temperature was used (100⬚C), one of the lattices was not formed, which means one was less stable than the other. The detail of the structural analysis of the films will be described elsewhere [31]. On the other hand, the streak patterns from PTCDA/TaSe2 did not change when the incident azimuth of electron changed. This means that only one axis of the molecular lattice perpendicular to the surface is aligned. This result was not changed by elevating the substrate temperature during the growth. Considering that the lattice constant of MoTe2 and TaSe2 differs by only 2.3%, this result suggests that the surface potential, which determines the lattice orientation of the films, is more uniform on metallic TaSe2 than semiconducting MoTe2. When the thickness of the films was increased to 8 MLE, spots appeared on the streaks in the RHEED pattern for both substrates without changing the lateral distance of the streaks. The spots are aligned on parallel lines. The lines are inclined and the angles from the surface normal changed when the incident azimuth of the electron beam was varied. The distance between spots along the streaks corresponded to 0.36 nm for both substrates. This indicates that the distance between the molecules along surface normal is 0.36 nm, which is close to the distance between the molecular plane (0.3231 nm) in the bulk PTCDA [32]. It suggests that the PTCDA is stacked nearly parallel to the substrate surfaces. Figure 11 shows work functions of the films as functions of the thickness. While the films on MoTe2 do not show any difference after the growth of 1 MLE, those on TaSe2 gradually changed the work function. The converged value was 4.8 eV for both substrates.

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Figure 11 Work functions of PTCDA grown on (a) TaSe2, (b) MoTe2 at various thicknesses.

Figure 12 UPS of PTCDA/MoTe2 at normal emission. The film thickness is (a) 0 ML (=MoTe2), (b) 1 ML, (c) 2 ML, (d) 4 ML, (e) 8 ML, (f ) 100 ML.

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Figure 13 UPS of PTCDA/TaSe2 at emission angles 35⬚ from surface normal. The film thickness is (a) 0 ML (=TaSe2), (b) 0.5 ML, (c) 0.5 ML (substrate subtracted), (d) 1 ML, (e) 2 ML, (f) 3 ML. The photoelectrons from ␲ orbitals are pronounced if the molecules lie flat on the surface.

energy (h␯ = 40.8 eV). The subtracted spectra are shown in Figures 13c and 14c. We note that the overall feature of the spectra is similar to that of the thicker film, but the continuous shift is observed by increasing the thickness for both of the polar angles of electron emission. Another thing to notice is that the spectra around EB ⬃ 3–5 eV show different shapes when the thickness is changed. The intensity of the peak next to the highest occupied molecular orbital (HOMO) is decreased for both polar angles when the thickness is reduced. Three of the possible models proposed by Ishii et al. [1] seem applicable to D2h symmetrical PTCDA molecule lying flat on the metal substrate:

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Figure 14 UPS of PTCDA/TaSe2 at emission angles 70⬚ from surface normal. The film thickness is (a) 0 ML (= TaSe2), (b) 0.5 ML, (c) 0.5 ML (substrate subtracted), (d) 1 ML, (e) 2 ML, (f) 3 ML. The photoelectrons from nonbonding orbitals are pronounced if the molecules lie flat on the surface.

(1) change in the electron distribution on the metal surface due to the orbital hybridization between the molecule and the metal substrate (without net charge transfer); (2) polarization of the molecular ␲ orbitals; and (3) charge transfer between the molecule and the substrates. In order to analyze the result of the electron spectroscopy, we have calculated the electronic structure of neutral, anionic, and cationic PTCDA and a neutral molecule polarized perpendicular to the molecular plane. The electric field for the polarization was 5 ⫻ 108 Vm⫺1 (0.5 V/nm), which gives the potential drop at 3 MLE in the present experiment. First we optimized the structure of each chemical species using Hartree-Fock energies with 3-21G basis functions. The optimized structures were used to calculate the electronic structures

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using density functional method B3LYP with 6-31G(d) basis functions. The results of density-of-states (DOS) curve are shown in Figure 15 by summing up the Gaussian with FWHM 0.7 eV located at each molecular orbital level. The UPS spectrum from a thick PTCDA/MoTe2 is shown in Figure 15e for comparison. Calculated DOS curves are shifted as shown in the caption of Figure 15 to adjust the HOMO position with the experiment. It should be noted that neutral and cationic PTCDAs show DOS similar to the experimental curve (Figure 15e) whereas anionic PTCDAs show a different DOS shape. The difference is mainly due to the change of molecular orbitals involved with carbonyl group, as typically seen at the large peak at EB ⬃ 4–5 eV. We now examine the origin of the interface dipole formed at PTCDA grown on metallic TaSe2. Four features should be considered in the present results: (1) The shift in the valence levels in UPS is observed between 1 MLE, 2 MLE, and ⱖ3 MLE. Almost no shift is observed between 0.5 MLE and 1 MLE. (2) UPS at 0.5 MLE coverage after subtraction of the substrate signal shows the peak shape around HOMO different from that of the thicker films. It indicates that the spectrum cannot be reproduced by the simple summation of normalized substrate spectra and that of PTCDA. (3) The second peak from the low binding energy side in the UPS spectra becomes less pronounced near the interface (0.5–1 ML). This holds true independent of the polar angles of electron emission. If this is due to the change in the orientation of the molecular plane accompanied by the increase of the thickness, the intensity of the peak will behave differently depending on the polar angle of electron emission. It therefore comes from the change in the electronic structure of PTCDA. (4) In quantum chemical calculation, the molecular orbital level corresponding to the diminished peak in feature 3 is not observed when the molecule has a negative charge. The shift of the calculated HOMO levels between PTCDA and PTCDA⫺ (3.7 eV) corresponds to the vacuum level shift between a molecule and an anion floating in vacuum. Although the value of the shift is far larger than the observation, the direction of the shift is consistent with the observed shift in UPS. The difference in the value is accounted for by the existence of dielectricity around the molecules and small percentage of the PTCDA⫺. Thus, the calculation seems to suggest that observations 1 and 3 can be accounted for by the existence of PTCDA⫺ mixed in neutral PTCDA at the interface. We now attempt to construct a model consistent with findings 1–4. First, the existence of orbital hybridization between HOMO of PTCDA and electrons of TaSe2 is strongly suggested by observation 2. This conclusion has the drawback that 1-MLE-thick film does not show the change in the HOMO peak shape. We consider this is caused by the fact that the estimate of the thickness is slightly erroneous because of the different sticking co-

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Figure 15 Quantum chemical calculation of PTCDA at various chemical states under (a) electric field of 109 V/m perpendicular to the molecular plane, (b) PTCDA⫹, (c) neutral PTCDA, (d) PTCDA⫺, and (e) UPS of 20 nm PTCDA/MoTe2. Molecular structures in (b)–(d) are optimized in the respective chemical states, whereas that in (a) is optimized as a neutral molecule without the electric field [same as (c)].

efficient between quartz crystal microbalance and TaSe2. The shift in the valence spectra (1) cannot be explained by the orbital hybridization alone above 1 MLE because hybridization between PTCDA layers does not modify the electronic structure strongly. Charge transfer or polarization of molecular layers must be considered. A simple rigid band picture considering PTCDA as a doped semiconductor suggests the electron donation from PTCDA to TaSe2 substrate or the existence of cationic PTCDA (PTCDA⫹) because of the difference in the work functions between thick PTCDA and TaSe2. However, the UPS (feature 3) and calculation 4 suggests the existence of anionic PTCDA (PTCDA⫺). We consider this controversy can be explained as follows. Electrons spill out from the surface of metallic TaSe2 by tunneling for a distance of about 0.5–1 nm [35]. Those electrons must be in PTCDA

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when the film is formed on the surface, which create chemical species resembling anionic PTCDA observed in the UPS. At the same time, this is not energetically favored because the electron motion is contrary to the direction expected from work function difference. The electron spill-out from the metal surface is therefore, suppressed as proposed recently [36]. It leads to the dipole layer formation consistent with the results observed in our work function measurement. In summary, we propose a model of the dipole formation at the PTCDA/TaSe2 interface from angle-resolved UPS measurements and quantum chemical calculations. The suggested model involves the hybridization of HOMO level of PTCDA with TaSe2 at the first layer and charge transfer from TaSe2 to PTCDA which continues up to 2 ML. We expect that the present approach using UPS measurement of molecular film grown on layered metals will lead to the full understanding of the microscopic nature of interface dipoles.

V.

CONCLUDING REMARKS

We have described the results of our recent study on electronic structure of the van der Waals interfaces between molecular film and layered materials. HREELS is particularly useful to detect the symmetry lowering of highly symmetrical molecules such as C60 and can be a probe to characterize the molecular–surface interaction as model cases. Planar molecules such as CuPc and PTCDA are sensitive to the orbital hybridization when they lie flat on the substrate surface. Systematic study by combination of well-defined interfaces, various electron spectroscopies, and quantum mechanical calculation will elucidate the nature of van der Waals interaction and the mechanism of interface dipole formation.

ACKNOWLEDGMENT The authors are grateful to the graduate students in our laboratory, Dr. Y. Fujikawa, K. Hamaguchi, and I. Sugiura, who conducted the experiments of the present research. They also wish to thank Prof. F. S. Ohuchi, K. Saiki, Dr. K. Ueno, and K. Iizumi for helpful discussions. The financial support by Grants-in-Aid for scientific research from Ministry of Education, Culture and Sports is acknowledged. A part of the present work was supported by NEDO for International Joint Research Grants.

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13. 14. 15. 16. 17. 18. 19. 20.

21. 22. 23. 24.

H. Ishii, K. Sugiyama, E. Ito, K. Seki, Adv. Mater. 11:605, 1995. A. Koma, Prog. Cryst. Growth Charact. 30:129, 1995. A. Koma, Thin Solid Films 216:72, 1992. A. Koma, K. Yoshimura, Surf. Sci. 174:556, 1986. J.A. Wilson, A.D. Yoffe, Adv. Phys. 18, 193 (1969). J.C. Wildervanck, F. Jellinek, Z. Anorg. Allg. Chem. 328:309, 1964. E. Bjerkelund, A. Kjekshus, Acta. Chem. Scand. 21:513, 1967. T. Shimada, F.S. Ohuchi, B.A. Parkinson, Jpn. J. Appl. Phys. 33:2696, 1994. A. Aruchamy (ed.), Photoelectrochemistry and Photovoltaics of Layered Semiconductors, Kluwer Academic, Dordrecht, 1992. R.M.A. Lieth (Ed.), Preparation and Crystal Growth of Materials with Layered Structures, D. Reidel, Dordrecht, 1976. T. Shimada, M. Nagahori, A. Koma, Surf. Sci. Lett. 423:285, 1999. Gaussian 98 (Revision A.7), M.J. Frisch, G.W. Trucks, H.B. Schlegel, G.E. Scuseria, M.A. Robb, J.R. Cheeseman, V.G. Zakrzewski, J.A. Montgomery, R.E. Stratmann, J.C. Burant, S. Dapprich, J.M. Millam, A.D. Daniels, K.N. Kudin, M.C. Strain, O. Farkas, J. Tomasi, V. Barone, M. Cossi, R. Cammi, B. Mennucci, C. Pomelli, C. Adamo, S. Clifford, J. Ochterski, G.A. Petersson, P.Y. Ayala, Q. Cui, K. Morokuma, D.K. Malick, A.D. Rabuck, K. Raghavachari, J.B. Foresman, J. Cioslowski, J.V. Ortiz, B.B. Stefanov, G. Liu, A. Liashenko, P. Piskorz, I. Komaromi, R. Gomperts, R.L. Martin, D.J. Fox, T. Keith, M.A. Al-Laham, C.Y. Peng, A. Nanayakkara, C. Gonzalez, M. Challacombe, P.M.W. Gill, B.G. Johnson, W. Chen, M.W. Wong, J.L. Andres, M. HeadGordon, E.S. Replogle and J.A. Pople, Gaussian, Inc., Pittsburgh PA, 1998. Y. Fujikawa, K. Saiki, A. Koma, Surf. Sci. 357/358:176, 1996. Y. Fujikawa, M. Sakurai, A. Koma, Jpn. J. Appl. Phys. 34:L188, 1995. Y. Fujikawa, Doctoral thesis, University of Tokyo, 1995. M. Sakurai, H. Tada, K. Saiki, A. Koma, Jpn. J. Appl. Phys. 30:L1892, 1991. M. Sakurai, H. Tada, K. Saiki, A. Koma, Chem. Phys. Lett. 208:425, 1993. R. Meilunas, R.P.H. Chang, S. Liu, M. Jensen, M.M. Kappes, J. Appl. Phys. 70:5128, 1991. R.L. Cappelletti, J.R.D. Copley, W.A. Kamitakahara, F. Li, J.S. Lannin, D. Ramage, Phys. Rev. Lett. 66:3261, 1991. G. Gensterblum, K. Hevesi, B.-Y. Han, L.-M. Yu, J.-J. Pireaux, P.A. Thiry, R. Caudano, A.-A. Lucas, D. Bernaerts, S. Amelinckx, G. Van Tendeloo, G. Bendele, T. Buslaps, R.L. Johnson, M. Foss, R. Feidenhans’l, G. Le Lay, Phys., Rev. B 50:11981, 1994. P.A. Gravil, Ph. Lambin, G. Gensterblum, L. Henrard, P. Senet, A.A. Lucas. Surf. Sci. 329:199, 1995. T. Shimada, K. Hamaguchi, A. Koma, F.S. Ohuchi, Appl. Phys. Lett. 72, 1869 (1998). H. Saijyo, T. Kobayashi, N. Ueda, J. Cryst. Growth 40:118, 1977. J.D. Weight, Prog. Sur. Sci. 31:1, 1989.

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9 Indium-Tin-Oxide Thin Films for Characterization of Electrochemical Processes in Molecular Assemblies Surface Characterization and Surface Modification Carrie L. Donley, Darren Dunphy,* Rebecca A. P. Zangmeister, Kenneth W. Nebesny, and Neal R. Armstrong University of Arizona, Tucson, Arizona

I.

INTRODUCTION

Transparent metal oxide thin films (indium-tin-oxide, ITO; and F-doped or Sb-doped SnO2) are of interest as the anode in several active display applications, e.g., organic light-emitting diodes (OLEDs) and light-emitting polymer thin films [1–4]. They have also found widespread use as at least one of the electrodes in electrochromic displays [5], and as a convenient indicator electrode in the spectroelectrochemical characterization of various solution electrochemical reactions and the redox processes of conductive polymer films [6,7]. We have also recently shown that 25 to 50 nm ITO films can be used as a conductive overlayer on a step-index integrated optic waveguide (EA-IOW), providing the sensitivity to follow solution redox processes of about 1% of a monolayer of a surface-confined redox active molecule or molecular assembly [8,9]. The initial composition of the metal oxide surface is critical to the performance of these metal oxides, whether as anodes in a display device or as indicator electrodes in a spectroelectrochemical sensor. There is a rich *Current affiliation: Sandia National Laboratories, Albuquerque, New Mexico.

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tradition of modification of the surfaces of these oxide electrodes with various molecular monolayers, including silanes, cyanuric chloride, and various functionalized phosphonates, carboxylates, etc., to tailor their electrochemical or photoelectrochemical properties [10–13]. The goal in many of those studies has been to create a non-cross-linked modification layer, which would improve the chemical compatibility of the surface with molecular redox active species and accelerate charge transfer. There are several recipes in the literature that apparently lead to successful silane attachment to ITO surfaces, although it is not always clear whether these are extensively crosslinked layers, attached at relatively few points to the ITO surface, or true ‘‘brush-phase’’ silane layers [13,14]. Marks and coworkers recently showed how silane chemistries can be readily extrapolated to ITO surfaces, where the goal is to achieve a fully cross-linked layer, to regulate charge injection in an OLED [15,16]. There are substantial variations in microstructure of ITO films [17,18], which undoubtedly lead to variations in surface chemical composition. These variations depend on the commercial source of the ITO and the deposition procedure, solution and plasma cleaning processes, etc., which likely lead to differences in chemical reactivity and modifier coverage. The presence of variations in apparent work function of these ITO surfaces attests to the significant effects such variability may produce on the electronic properties of the ITO surface [2,19–22]. In addition to the OLED applications of ITO electrodes, we have become interested in the use of this material as a transparent substrate for the characterization of charge transfer events of low coverages of molecular assemblies based on columnar phthalocyanine aggregates or polymers [23,24]. As an example, extremely coherent columnar assemblies can be obtained from phthalocyanines (Pcs) based on 2,3,9,10,16,17,23,24-octakis(2-benzyloxyethoxy)phthalocyaninato copper and its metal-free analogue (CuPc(OC2OBz)8 and H2Pc(OC2OBz)8) (Figure 1) [25]. These Pcs possess both ethylene oxide side chains, which enhances their solubility in nonaqueous solvents, and benzyl termination of these side chains, which provides an unusual stiffness to bilayer thickness films (5.6 nm) prepared by compression on a Langmuir Blodgett trough. These Pc films consist of coherent columns of cofacially aggregated molecules tilted at angles with respect to the molecular plane (Figure 1) reminiscent of the ␤ phase of the unsubstituted crystalline Pc. Once compressed to bilayer thickness, Pc films consisting of closely packed, parallel columnar aggregates can be ‘‘stamped’’ off the subphase by a horizontal transfer approach to appropriately modified surfaces (see below). If the Pc is diluted into a monolayer of an electroinactive amphiphile, isolated Pc aggregate columns can be formed on an ITO

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surface, over the integrated optic waveguide, and the electron transfer processes of individual ‘‘molecular wires’’ explored [23,24]. The efficacy of transfer of these two types of Pc film can be critically dependent on the surface chemical functionality of the substrate materials [26]. We have seen that modification of gold and oxide covered silicon wafers, with phenyl-terminated thiols or silanes, respectively, can increase the amount of Pc transferred in the first bilayer by up to 100%, relative to the bare gold or silicon oxide surface. It is of interest to extrapolate these modifications to ITO surfaces and ultimately to determine their effect on electron transfer rates at such interfaces. However, the composition of the ITO surface strongly affects the efficiency of silane attachment. In this chapter we outline some of our recent studies of ITO surface composition, the changes which these surfaces undergo after a piranha etch, and, finally, the changes indicated by the successful attachment of a phenylterminated trialkoxysilane. In aqueous media the apparent electron transfer rate for reduction–oxidation of a charged molecule such as Fe(CN)⫺4/⫺3 is 6 significantly lowered by the presence of the phenylsilane surface modifier, and the surface becomes electrochemically heterogeneous. In nonaqueous media the electrochemical response of probe molecules like ferrocene and 4,4⬘-bis(m-tolylphenylamino)biphenyl (TPD) is not appreciably affected. At low supporting electrolyte concentrations, however, the rate for electron transfer at a phenylsilane-modified ITO surface is enhanced. The presence of the phenyl-terminated silane on the ITO surface enhances the transfer efficiency of a closely packed Pc bilayer, as evidenced by the increased total charge attributed to the electrochemical oxidation of these assemblies.

II.

EXPERIMENTAL STUDIES

The ITO used in these studies was obtained from Donnelly Corporation with a surface resistivity of 13 ⍀/▫. Preliminary cleaning of the ITO consisted of ultrasonication in the following solutions for about 10 min each: a Triton X-100 detergent solution, Millipore water (18 M⍀), and ethanol. ITO was then stored in pure ethanol until it was ready to be processed further. Several different established recipes were evaluated that had been developed for attachment of silanes to ITO or SnO2 surfaces [14–16,27]. The method used for most of these studies involves a procedure developed by Wilson and Schiffrin [28]. First, the ITO was heated to approximately 80⬚C in a solution of 10 mM NaOH for 4 h. Then the surface of the ITO was etched with a piranha solution (4:1 sulfuric acid/hydrogen peroxide) for 1 min. Longer exposure times etched significant amounts of the ITO as was evident by the partial or complete loss of conductivity. Finally, the ITO was

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rectangular bilayer thickness regions whose stability and integrity allow for rapid horizontal film transfer by the Schaefer method (‘‘stamping’’) without disturbing the rest of the film. All electrochemistry was carried out with an Cypress Systems CySy 2Ra potentiostat, a platinum counter electrode, and a Ag/AgCl reference electrode. XPS/ultraviolet photoemission spectroscopy (UPS) was conducted using a Kratos AXIS-ULTRA spectrometer, with a base pressure of about 5 ⫻ 10⫺9 torr. The X-ray source was monochromatic Al(K␣) and the approximate size of the analysis area on each sample was 300 ⫻ 700 ␮ms. The photoemission spectra were recorded in a constant-analyzer pass energy mode (pass energy = 20 eV). Samples were generally floating rather than electrically grounded to the spectrometer, with charge compensation applied to optimize photoelectron count rate and minimize peak width. These conditions were directly compared with a situation in which the samples were grounded to the spectrometer to ensure that no artifact in peak shape was introduced by this charge compensation procedure. Charge compensation did introduce a shift in the peak positions to lower binding energies, but charge shifting of the photoemission peaks was further corrected by including a fully reduced indium species on the sample holder, which in the fully charge compensated spectra showed an In(3d5/2) binding energy of 443.2 eV, and subsequently adjusting the binding energies of the rest of the photoemission peaks. This method of binding energy determination was checked against procedures involving the use of adventitious carbon on the sample surface (observed in our studies to be at 284.9 eV) and procedures computing the binding energy differences [O(1s) ⫺ In(3d5/2)] [32] to ensure that the binding energies reported here are consistent with previous studies [32,33]. Following linear background subtraction the photoemission peaks were curve-fitted using summations of 70% Gaussian and 30% Lorentzian peaks. The indium metal standard for XPS was obtained commercially (AlfaAesar) and was mechanically scraped to expose bare metal before analysis. The In2O3 and the In(OH)3 powders were obtained from Fluka Chemika and Aldrich, respectively, and used as received. As a way of producing a mixture of both hydroxide and oxide states, the In(OH)3 powder was heated to 200⬚C for about 72 h, producing a partially dehydrated oxide state at about 50 mol %. The partially hydrolyzed indium oxide, ‘‘InOOH,’’ was electrochemically grown on indium metal in a 0.1 M borate buffer as recently described by Metikos-Hukovic and Omanovic [34]. This electrochemically grown oxide is an intermediate on the way to electrochemical production of the stoichiometric oxide and is presumed to consist of a disordered film of polynuclear indium oxide/hydroxide species. The indium metal surface was first

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polarized to ⫺1.8 V (vs. Ag/AgCl) for 5 min, to remove surface oxide. The indium foil was then scanned from ⫺1.5 to ⫹0.85 volts creating an ‘‘InOOH’’ film that was about 3.3 nm thick, and removed from the cell under potential control, rinsed, and inserted into the vacuum environment for analysis. After initial XPS analysis of this surface the sample was sputtered to expose the underlying zero-valent indium metal and this peak was used as a reference to determine the position of the In(3d5/2) and O(1s) peaks in the InOOH state.

III.

RESULTS AND DISCUSSION

A.

Oxide Surface Chemistries

Figure 2 shows a tapping-mode atomic force microscopy (AFM) image of a typical cleaned ITO sample. Typical root mean square roughness for these samples and for silane-modified ITO was about 0.8 nm and grain sizes were about 30 nm. These small grains of ITO are similar to ‘‘subgrain’’ features observed within larger grain structures of some previously reported ITO films [17,18]. However, our AFM images of the Donnelly ITO on glass do not show this larger (about 250–750 nm in diameter) grain structure. Sputtered ITO films are primarily composed of the bixbyite In2O3 lattice structure with predominantly the 具100典 and 具111典 faces lying parallel to the substrate in the microcrystalline grains [17,35]. X-ray diffraction studies of the films used in these studies showed the (400) and (222) reflections in the approximate ratio of 1:3, which has been attributed to the presence of random distribution of 具100典 to 具111典 crystal faces [17,35]. A schematic view of the types of oxides to be found in this thin film material, including the accepted nomenclature to describe this oxide (In2-y SnyO3ey) (36), and another schematic view of the types of chemistries likely to occur during exposure to water vapor, solutions, and oxygen-rich and oxygen-poor environments are also included in Figure 2. A minimum of two forms of oxygen are expected at an oxide surface, including those bridgebonded into the lattice, and those arising from disruption of those bonds due to hydroxylation ( — OH) and/or ionization ( — O⫺) or protonation ( — OH⫹ 2) of the hydroxide. ITO thin films also typically contain a substantial fraction of oxygen defect sites to optimize both the electrical and optical properties of this material, and it is presumed that the oxygen ions surrounding these defect sites will experience a sufficiently altered environment as to be spectroscopically distinct [1,17,18,37]. Hydrolysis of indium and tin oxides can occur as a result of contact with aqueous environments or even atmospheric water:

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Figure 2 (a) AFM image of one of the ITO films which are the subject of this chapter, showing the grain microstructure of the ITO. (b) Schematic of the ITO surface at the approximately 10 nm scale, showing the columnar grain structure, the bulk of which is In2O3, with a heterogeneous capping layer of ‘‘InOOH’’ and In(OH)3. (c) Schematic of the nanometer scale structure of the ITO surface, showing bridging and nonbridging oxygen sites, hydroxide sites, and certain acid–base chemical interactions.

In2O3 ⫹ 3H2O ⇒ 2In(OH)3

(1)

Of course, if those environments contain chloride ion, the hydrolysis process is facile enough to produce soluble indium chlorides, etching the ITO surface, which is the basis of the lithographic processing steps commonly used to produce patterned ITO films [38,39]. As revealed by stability diagrams of these materials, calculated from potentiometric electrochemical studies [39], hydrolysis of indium oxides is generally more favored than for tin oxides, at any pH. The solubilities of In(OH)3 and polynuclear indium hydroxides Inx(OH)x⫹3 are quite low, so it is anticipated that the products of such hydrolysis reactions will remain physisorbed to the ITO surface until

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processing steps provide a pathway for their removal. It is possible that certain silane modification steps, which are not preceded by removal of the polynuclear metal hydroxides, simply attach to those species, not to the electrochemically or electronically active ITO surface, which then impacts the efficiency of the surface modification process. These surface chemistries are under further exploration. Lastly, there is the possibility of incomplete hydrolysis of an indium surface oxide, creating an intermediate hydroxylated species: In2O3 ⫹ H2O ⇒ 2InOOH

(2)

This species may arise as part of the incomplete hydrolysis of the stoichiometric oxide and by the nucleophilic attack of water on the exposed oxide defects, adding — OH at metal sites adjacent to the defect site [40–43]. Recent studies of the electrochemical formation of indium oxides on indium surfaces, which proceeds through a stable oxyhydroxide intermediate (‘‘InOOH’’) [34], suggest a means by which such a species could be produced for XPS characterization, and such a surface is included in our standard materials discussed below. B.

Indium Oxide Standards

Figure 3 shows the In(3d5/2) and O(1s) spectra for the ultrasonically cleaned ITO surface and for a series of standard materials that are representations of the possible products of ITO surface reactions. Table 1 summarizes the peak positions for the most important forms of indium, oxygen, and tin on the ITO and standard oxide surfaces explored to date. O(1s)-In(3d5/2) binding energy (BE) differences, also included in Table 1, are useful in that they are insensitive to charge shift corrections or charge compensation procedures [32,44]. These BE difference values correlate well with those reported previously for other standard indium oxides. The In(3d5/2) peak on the mechanically cleaned indium foil, and the corresponding higher binding energy In(3d5/2) peak due to the residual surface oxide, were at positions consistent with several previous studies [32,37]. The In(3d5/2) peak in the In(OH)3 sample was shifted about 1.6 eV to higher BE versus the pure metal, and 1 eV higher BE than for indium in In2O3, as had been previously reported [32,37]. The annealed In(OH)3 sample showed a broadening of the In(3d5/2) peak to lower BE, indicative of the formation of some In2O3 in the near surface region, and the O(1s) peak showed the clear formation of a new low-BE (In2O3) form of oxygen, also consistent with this surface dehydration process. The surface of the ‘‘InOOH’’ sample showed an In(3d5/2) peak position whose BE was intermediate between that seen for In2O3 and that seen for

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Figure 3 In(3d5/2) and O(1s) XPS data for ITO mechanically scraped indium foil, with residual oxide layer, In2O3 powder, an ‘‘InOOH’’ film, electrochemically grown on indium foil (before and after argon ion sputtering), In(OH)3 powder (before and after annealing to convert a portion of the near-surface region back to the In2O3 form).

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Table 1 BE values for photoemission peaks from ITO and related indium oxide/ hydroxide standards In 3d5/2 (ev)

O 1s (ev)

Sn 3d5/2 (ev)

O (1s)–In (3d5/2) (ev)

␣⬘ (ev)b

485.8a 486.8 (d) 487.7 (e) —

851.6

—

854.3

In2O3

443.8

—

85.7

850.8

‘‘InOOH’’

444.3

—

86.9

—

In(OH)3 Heated In(OH)3

444.8 444.9 444.3

529.5a 530.6 (b) 531.8 (c) 531.5c 529.4c 529.5a 531.4 531.2a 529.1 531.3 531.6 529.8

85.6

In0

443.9a 444.9 (a) 445.9 443.2

— —

86.5 86.7 85.5

850.3 851.1

Species ITO

a

Denotes the main peak, when only one value for O (1s)–In (3d5/2) is shown, this value was used. b ␣⬘ = KE (Auger) ⫹ BE (photoelectron) (a) Due to In(OH)3 (see text). (b) Oxygen adjacent to oxygen deficiency sites in the ITO lattice. (c) Oxygen due to In(OH)3. (d), (e) Other forms of tin as seen for indium. c These forms of oxygen were due to oxide contamination in the In0 surface.

In(OH)3, with an In(3d5/2) BE = 444.3 eV. The O(1s) spectra for this material showed broad peaks consistent with the presence of both oxide-like species and hydroxide-like species. Argon ion sputtering of this sample surface produced an O(1s) signal more consistent with In2O3, suggesting that the electrochemical growth of this ‘‘InOOH’’ film produced a layered product with the more stoichiometric oxide closer to the indium metal interface. C.

Forms of Indium and Oxygen on the ITO Surface

The 0.35 eV bandwidth of the monochromatic Al(K␣) source in these XPS experiments allows a clear distinction between In2O3 on the ITO surface and other indium oxide forms which contribute to the asymmetry in the In(3d5/2) peak. We cannot completely discount final state effects as contributing to this asymmetry; however, the high binding energy shoulder in the In(3d5/2) peak for ITO coincides closely with the In(3d5/2) signal for the In(OH)3 standard. Reasonable fitting of the In(3d5/2) peaks with or without contri-

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butions from the ‘‘InOOH’’ standard was possible, but for the reasons stated above, we include its contribution to these spectra [34,39]. The O(1s) spectrum for the ITO surface (Figures 3 and 4a) is composed of a minimum of three peaks (A, B, and C). Several explanations have been given for the chemical origins of these three O(1s) peaks. Ishida and coworkers attributed peak A in Figure 4a to a crystalline form of ITO; peak B was assigned to an amorphous phase of ITO; and peak C was assigned to hydroxide oxygens [45]. X-ray diffraction data on most commercial ITO films does not indicate the presence of an amorphous phase, making this assignment unlikely. Nelson and Aharoni reported these two forms of oxygen as those bound to indium (A) and oxygen bound to tin (B) [46]. These O(1s) peaks would therefore need to correlate in relative peak area to the Sn/In

Figure 4 O(1s) XPS data for as-received, strong base/piranha-etched, and phenylsilane-modified ITO at both a 0⬚ take-off angle (top) and a 75⬚ take-off angle (bottom). All peaks show the fitted data, and the percentages in the corners represent the relative oxygen percentages of that species. Species A: In2O3-like oxygen; species B: oxygen adjacent to oxygen defect sites; species C: In(OH)3-like oxygen; species D: In-O-Si-, Si-OH-, or Si-OCH3-like oxygens.

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ratio of these thin films, which they do not; therefore, we reject that assignment as well. The more plausible assignments are that peak A in Figure 4 is due to In2O3-like oxygen, peak B is due to oxygen near defect sites in the near surface region, and peak C is due to In(OH)3-like oxygen from the hydrolyzed oxide and from adventitious surface water [1]. Oxygen deficiencies are intentionally introduced during processing of ITO thin films to enhance electrical conductivity and optical properties [36], and their concentrations can be as high as 30% in the near surface region. The presence of these oxide defect sites in ITO does not generally appear as an additional In(3d5/2) peak at lower BE than the dominant peak due to the In2O3 form. However, we recently observed that ITO samples extensively plasma etched during cleaning or argon ion sputtered to provide a clear Fermi edge photoemission in UPS studies do show traces of a partially reduced form of indium, on the low-BE side of the In(3d5/2) peak. These results are reported elsewhere [43]. The Sn(3d5/2) photoemission peaks also indicate that there are at least two forms of tin present in the surface region of these ITO samples (Table 1), which are likely due to SnOx and Sn(OH)x. Some recent studies have suggested that the tin sites are the point of attachment for certain metal complexes (zirconium-based and organotin species) on the ITO surface [47,48]. In the surface modification studies we have conducted to date, we do not see significant changes in lineshape for either the Sn(3d) or In(3d) spectra as the surface pretreatments described above are conducted. D.

Angle-Resolved XPS

Figure 5a and b shows the O(1s) and In(3d5/2) peaks for ITO at both the O⬚ and 75⬚ sample/analyzer take-off angles, the latter angle being used to accentuate photoemission from the uppermost surface region. It is clear that there is an increase in the higher BE components relative to In2O3, indicating that both the oxide defect sites and the hydroxylated oxide components are closer to the surface than the In2O3 phase. We also show in this figure the ratio of photoemission peak intensities at 0⬚ relative to 75⬚ (R0/75 = N(E)0-degrees/N(E)75-degrees), vs. BE, where the peak intensities have been normalized to 1.0 at the In(3d5/2) peak maxima. The ratio R0/75 accentuates the influence of photoemission from species near the surface. In computing this ratio we treat the contribution to the photoemission signal from the In2O3 species as though it is the same throughout the XPS sampling region at both take-off angles, accentuating the contributions from other species at the surface [43]. At higher BE values above those for the main peak R0/75 diminishes for both the O(1s) and In(3d5/2) peaks, suggesting that both the In(OH)3 and ‘‘InOOH’’ species dominate the composition of ITO in the near-surface re-

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gion, as shown schematically in Figure 2. Consistent with this finding, the O(1s)/In(3d) ⫹ Sn(3d) peak area ratios increase significantly at the 75⬚ takeoff angle, as expected if the near surface region is rich in hydroxylated indium species [1].

E.

ITO Surface Modification

Several studies have been conducted on ITO surface modification with both chloro- and alkoxysilanes, which show apparently good coverage of the ITO surface with the silane modifier [14]. Silane chemistries that we initially used to modify SnO2 electrodes [27,49,50] have not been as successful in recent modification studies with ITO. When significantly cross-linked silanes are desired, it is clear that good coverages can be obtained but that this may significantly hinder the electrochemical response of the ITO substrate [15,16]. In addition, the use of chlorosilanes appears likely to liberate HCl as the product of silane attachment, which will accelerate the decomposition of the ITO surface [38,39], so that addition of small amounts of base during those reactions is advisable [15,16]. Recently reported pretreatment of the ITO surface by soaking in strong base followed by a brief exposure to a piranha solution [28], annealing, and reaction with alkoxysilanes does appear to provide adequate coverage of the phenylsilane, without extensive cross-linking. This treatment appears to be most effective at removing adventitious carbon on these ITO surfaces, incorporated into the ITO surface during fabrication or storage, which we have found difficult to remove by other chemical/mechanical cleaning steps. The C(1s) signal due to this adventitious carbon is decreased by about 50% following this piranha treatment protocol. At the same time, the static contact angle for water was 30–35⬚ for the piranha-treated ITO and increased to an average of about 65⬚ for the phenylsilane-treated ITO surfaces. Similar silane modifications of oxide layers on Si(100) wafers, and on silica surfaces, generally yield contact angles of about 85⬚ [26], so we must conclude that the phenylsilane coverage is not as high on these ITO samples and that hydrophilic regions are still present in trace quantities. Figure 4 (c and d, center column) and Figure 5 (c and d) summarize the changes induced by the strong base/piranha treatment to the oxygen and indium near-surface species. The O(1s) and In(3d5/2) spectra clearly indicate the loss of the oxygen near defect sites (see above) and the increase of — OH-like oxygen and indium species as a result of this pretreatment process. The RB/A values (RB/A = N(E)before etch/N(E)after etch) for both the O(1s) and In(3d5/2) spectral regions, at a 75⬚ take-off angle, are especially sensitive to these chemical changes (Figure 5, c and d).

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Figure 5 (a) O(1s) and (b) In(3d5/2) spectra for ITO at both 0⬚ and 75⬚ take-off angles. The difference between the two peaks in the high binding energy region is clearly noticeable. The ratio R0/75 is also shown, computed from the ratios of the peaks at the two different take-off angle (see text). R0/75 shows a strong decrease at the binding energies corresponding to ‘‘InOOH’’ and In(OH)3. (c) O(1s) and (d) In(3d5/2) spectra for as received ITO and strong base/piranha etched ITO at 75⬚ takeoff. The ratio of RB/A is also shown computed from the ratios of the peaks before and after the etch process. RB/A shows a slight increase in the O(1s) region assigned to photoemission from oxygen defects, and a decrease in the O(1s) region assigned to hydroxide forms. Little significant change is seen in the In(3d5/2) signal as a result of this etch process.

After silane modification of the ITO the In2O3-like oxygen and defectrelated oxygen forms (peaks A and B) appear to remain unchanged, although their intensities are reduced because they now lie under the silane layer. A broad peak D was assigned to all higher BE forms of oxygen on these surfaces, which may supplant most of the — OH-like oxygen from before silanization, and may include oxygen from — Si-O-In- species, and -Si-OHlike surface groups from hydrolyzed Si-OCH3 that did not react with the ITO surface. In addition, Si(2p) and Si(2s) peaks were observable at inten-

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sities consistent with a monolayer of attached material [27]. No discernible change was observed in the In(3d5/2) signal as a result of these pretreatment and silanization steps. The contact angles obtained from this treatment process are likely to be increased further by using monofunctional silanes, possessing only one attachment site to the ITO surface, and terminated with — CH3 groups or additional phenyl groups. Additional studies are under way with other silanes to probe the details of these ITO modification chemistries. F.

Electrochemistry on Modified ITO

The voltammetric reduction–oxidation of millimolar solutions of ferricyanide [Fe(CN)⫺4 6 ] (Figure 6, top) with both plain ITO and phenylsilane-modified ITO was carried out in aqueous media with 0.1 M potassium hydrogen phthalate (KHP) as the supporting electrolyte. The total current flowing on both the reduction and oxidation cycles is lower by about 40–50% on the phenylsilane-modified ITO surface, suggesting blockage of the ITO surface by phenylsilane, with regard to Fe(CN)⫺4 6 , with remaining electrochemical activity likely proceeding through small patches of unmodified ITO [51– 54]. In addition, the apparent rate of electron transfer, ks,app, as revealed by the separation of anodic and cathodic peaks is significantly lower. For plain ITO ⌬Epeak = 140 mV vs. ⌬Epeak = 291 mV for phenylsilane-modified ITO at a voltammetric sweep rate of 100 mV/s. The apparent electron transfer rates, ks,app, not corrected for the inherent resistance within the ITO film, decreases from about 9.7 ⫻ 10⫺5 cm/s to less than 10⫺6 cm/s, which is our lowest rate detectable under these conditions [55]. Actual electron transfer rates are expected to be about one order of magnitude higher, but still affected by the presence of the silane on the ITO surface. The broadening of the reduction peak for these ferricyanide solutions on phenylsilane-modified ITO also suggests that this surface contains multiple sites for redox activity, each with its own unique charge transfer rate. This is an expected result from a charged molecule such as ferricyanide if its redox chemistry occurs through strong interactions with the ITO surface and the hydrophobic surface modifier is heterogeneously distributed. The electrochemical oxidation of ferrocene and TPD were also probed in 50:50 acetonitrile/toluene/0.1 M tetrabutylammonium hexafluorophosphate (TBAHP) solutions. There was no significant increase or decrease in the electron transfer rates of these molecules at the phenylsilane-modified ITO (Figure 6, lower panel). The supporting electrolyte concentration was next lowered to 1 mM for the TPD solutions and the voltammetric oxidation process reexamined on both the clean and silane-modified ITO. Lowering the supporting electrolyte concentration increases the solution resistance and decreases the electric field gradient at the electrode surface, making electron

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Figure 6 Cyclic voltammograms for reduction of about 10⫺3 M aqueous potassium ferricyanide solutions (top) and about 10⫺3 M toluene/acetonitrile solutions of TPD (bottom). Both graphs show the electrochemistry on both plain ITO (solid line) and phenyl-modified ITO (dotted line). The electrochemistry of ferricyanide on phenylsilane-modified ITO shows less total current passed. In addition, the peaks have broadened, indicating multiple types of sites for redox activity.

transfer more difficult at either surface. The voltammetric responses on both clean and modified ITO surfaces are quite distorted due to the higher solution resistance; however, significantly more charge is passed on each portion of the voltammogram on the silane-modified ITO surface and the peak potential separation is slightly less for each redox process, suggesting higher rates of electron transfer for this molecule on the phenyl-modified ITO sur-

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face, in a solution of significantly lowered polarity. The impact of such modifications on charge injection rates in condensed-phase thin-film devices is under further exploration. Electrochemical oxidation–reduction of single-bilayer Pc films, as shown in Figure 1, was also explored as a means to quantify differences in uptake of these films on the ITO surface (Figure 7). Deposition of more than three to four or more bilayers of the Pc film on most surfaces generally results in an equivalent coverage of the dye in these later transfers, suggesting that the first deposited layers help to facilitate the transfer of the next hydrophobic layers. We are, however, interested in the spectroelectrochemical study of electron transfer processes involved in single bilayers and in isolated Pc aggregate columns, in association with electroactive waveguide and attenuated total reflectance (ATR) devices, so that 100% transfer efficiency of these layers to the ITO surface is critical. Previous studies have shown that gold and silica surfaces modified with phenyl-terminated alkanethiols or silanes, respectively, provide for a higher transfer efficiency of these bilayer Pc films than do modifiers terminated with alkyl groups, or no modification at all [26]. This same philosophy has driven our ITO surface modification procedures. On both the silane-modified and unmodified ITO surfaces the electrochemistry of the Pc films is qualitatively the same, with oxidation of about 20–25% of the Pc rings occurring at between ⫹0.4 and ⫹0.6 V vs. Ag/

Figure 7 Electrochemistry of one bilayer of CuPc(OC2OBz)8 transferred from a LB trough onto both plain (solid line) and phenyl-modified (dotted line) ITO showing the increased current passed for the Pc film on phenylsilane-modified ITO.

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AgCl. A higher fraction of Pc rings are oxidized at potentials more positive than ⫹0.7 V, and the potential at which these processes occur, and the peak shapes, are strongly dependent on the nature of the solution counter ion, which must be incorporated into the Pc film to maintain electroneutrality during this electrochemical doping process [6,23,31]. Interestingly, these voltammetric processes are better defined into two distinct potential regions on bilayer thickness films vs. 3–15 bilayer thickness films. Counter ion incorporation is more difficult in the thicker films, due to diffusive transport limitations, which impacts the shape of these voltammetric peaks [10,23]. The total charge transferred on the oxidative sweep in these voltammograms is correlated to the total surface coverage of Pc molecules in these thin films [31]. We estimate that on phenylsilane-modified ITO the surface coverage of Pc increased by about 10% compared to clean ITO as calculated from the area under the voltammetric oxidation peak, from about 0.1 to 1.0 V, which is consistent with spectroscopic characterization of Pc transfer to both gold and silica surfaces, modified with phenyl-terminated alkyl or alkoxy chains [26]. Using this coulometric data we estimate the surface coverage of Pc on the unmodified ITO to be about 9.0 ⫻ 10⫺11 mol/cm2 vs. 1.0 ⫻ 10⫺10 mol/cm2 on phenylsilane-modified ITO. Based on a projected area per ˚ 2 per molecule (accounting for the packing and tilt molecule of about 120 A of the Pcs within each columnar aggregate), we estimate that the electrochemically active surface coverage of Pc should be about 3.0 ⫻ 10⫺10 mol/ cm2 if every Pc is electrochemically accessible. It is not surprising that our estimated coverage is lower than this number since we do not expect perfect transfer of the Pc film, even onto the modified surface, and we anticipate domains within these columnar aggregate assemblies which are not electrically connected to the ITO surface and/or not in electrochemical equilibrium with the remaining portion of the Pc film [6,23,31]. The improvement in the electrochemical response of these Pc films on the phenylsilane-modified ITO suggests that these impediments to efficient charge injection from such a molecular assembly can be partially ameliorated by the modification process.

IV.

CONCLUSION

Despite their widespread use in a variety of emerging and established technologies, ITO films remain poorly understood, especially in the attempt to correlate the surface chemical composition with the microstructure of the ITO film and its device properties. Our characterization of just this one commercial source of ITO, and the variability in chemistries that is apparent from the standard oxides we have studied, suggest that substantially more study is needed before predictable relationships can be established. From

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our own experience it is also clear that modification chemistries developed in one laboratory, e.g., for the monolayer attachment of a silane to the ITO surface, may not translate easily to another laboratory because of subtle differences in ITO composition and microstructure. Because of the reactivity of defects in the oxide structure, the quality of the vacuum system used to produce the ITO film and the type of laboratory ambient the oxide surface sees when first removed from the ITO fabrication chamber may also be critical in determining surface composition and reactivity. It has been our general experience that more aggressive treatments are needed to prepare the ITO surface for covalent bonding to a surface modifier than was necessary for similar studies on SnO2 electrode materials [10,27,29,49]. The origins of these differences are an area of active exploration for us, and elaboration on the modification chemistry shown here is under way. Clearly, enhancing the wettability of ITO substrates toward hydrophobic molecular assemblies, providing for electrocatalytic sites that enhance rates of charge injection, and creating molecular layers that ‘‘template’’ the self-assembly of new mesogenic materials are areas needing further exploration in terms of these and related transparent conducting oxides.

ACKNOWLEDGMENTS This research was supported in part by grants from ONR (Center for Advanced Multifunctional Polymers and Molecular Assemblies), the National Science Foundation (Chemistry), and the Materials Characterization Program–State of Arizona.

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10 Energy Level Alignment at Organic–Metal Interfaces Hisao Ishii and Kazuhiko Seki Nagoya University, Nagoya, Japan

I.

INTRODUCTION

Until around 1950, organic materials have been recognized as insulators. The first observation of weak electric conductance of an organic material was reported in 1948 [1]. From systematic studies, Akamatu and Inokuchi found that a series of aromatic hydrocarbons with condensed aromatic rings have weak conductivity and named such a class of organic materials ‘‘organic semiconductors’’ [2]. They and Matsunaga also found that doping halogens into aromatic hydrocarbons gives highly conducting organic materials [3]. After their pioneering work, various semiconducting and conducting organic materials have been investigated [4] as well as conducting polymers [5] and superconductors [6]. In analogy to inorganic semiconductors, applications of organic semiconductors to electronic devices, such as light-emitting diodes and photovoltaic cells, have been extensively studied. In electronic devices, electronic signal or electric power should be exchanged between a device and an external circuit; electric carriers must move across the interface between the device and the electrode attached to it. Thus, the understanding of the interface between an organic semiconductor and the electrodes is a key issue for organic devices. For example, in organic electroluminescent (EL) devices, electrons and holes are injected from the cathode and the anode, respectively. The related processes have been discussed using a conventional energy diagram like that in Figure 1a. 293

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Figure 1 Conventionally used energy diagrams of organic devices. (a) Organic electroluminescent (EL) devices. (b) Organic solar cells.

The lowering of the barrier height for carrier injection at organic–metal interfaces is crucial to the performance of the device. The influence of interfaces is not only limited at the interface, but the properties of bulk semiconductor are affected in the region near the interface. For example, in photovoltaic cells, the space charge layer with band bending has been assumed in discussing the performance of the device as shown in Figure 1b. The gradient of the potential curve in the space charge layer is important for the efficiency of the separation of electron-hole pair generated by photon absorption. Thus the question ‘‘How do the energy levels of an organic solid and a metal align at their interface?’’ is a key in the interface business. Until recently, a naive model assuming a common vacuum level at the organic– metal interface had often been applied to estimate the interfacial electronic structures of organic devices. This traditional model is in analogy to the Mott-Schottky model, which is a widely known textbook model for inorganic–semiconductor/metal interfaces. The application of the model to organic systems was without serious discussion. The validity was not examined experimentally or theoretically. In 1995 we have published the first paper on the energy level alignment problem of organic–semiconductor/metal interface [7], where the invalidity of the assumption of a common vacuum level was clearly demonstrated by the results of ultraviolet photoemission spectroscopy (UPS). Our systematic studies for various organic solids revealed that the model is not valid for

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most organic–metal interfaces [8–24]. The breakdown of the assumption of a common vacuum level means that the vacuum level does not align but shifts at organic–metal interfaces. After our first work, other research groups all over the world, e.g., at Princeton [25–30], Rochester [31–36], Arizona [37–41], Hong-Kong [42–44], and Linko¨ping [45–48], also investigated the energy level alignment problem. These studies covered not only the energy level alignment problem but also other important factors such as the structures and the reactivity of organic–metal interfaces. Large amount of data on organic–metal interfaces are now accumulating. In 1999, we published a review paper on interfacial electronic structures, with stress on the energy level alignment [22]. Now further progress in the basic understanding of the above-mentioned factors is highly desired for the development of organic devices. In this chapter, we focus our attention mainly on the energy level alignment at organic–metal interfaces. The traditional model including misunderstanding about the vacuum level has been widely used for a long time, partly because various concepts for inorganic semiconductors are believed to be valid even for organic semiconductor without serious discussion of the difference between the inorganic and organic semiconductors. Thus, in Section II.A, first we reexamine the basic aspects for the electronic structures of organic semiconductors. The field of organic–metal interfaces is really interdisciplinary, and locates at the interface region between surface science and materials science of organics. Thus the basic understanding for the electronic structures of metal surfaces are also described in Section II.B. Then the energy level alignment model for semiconductor–metal interface is described in Section II.C. In Section III, our UPS studies on energy level alignment at various organic–metal interfaces including model interfaces of organic electroluminescent devices are described. In Section IV, we discuss the trends in vacuum level shift and the origin of interfacial dipole layer. Such discussion is important for estimating the energy level alignment at organic–metal interfaces of real organic devices. In real devices, ambient atmosphere can affect the electronic structure at organic–metal interfaces. Section V.A discusses such effects. Most UPS data in this chapter were measured for the organic film deposited on metal substrate. On the other hand, the reversely formed interface, i.e., metal-onorganic system, is also important in real organic devices. Section V.B briefly discusses the problem on this deposition sequence. Finally, Section V.C briefly discusses the band bending behavior at organic–metal interfaces. The chemical structures of many organic compounds described in this chapter are summarized in Figure 2.

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Figure 2 Chemical structures of the compounds discussed in this chapter. TPD, N,N⬘-diphenyl-N,N⬘-(3-methylphenyl)-1,1⬘-biphenyl-4,4⬘-diamine; Alq3, tris(8-hydroxyquinolino)aluminum; DP-NTCI, N,N⬘-diphenyl-1,4,5,8-naphthyltetracarboxylimide; BCP, bathocuproine; CBP, 4,4⬘-N,N⬘-dicarbazolylbiphenyl; ␣-NPD, N,N⬘diphenyl - N , N⬘ - bis(1-naphthyl) - 1,1⬘-biphenyl - 4,4⬘ - diamine; Gaq3, tris(8-hydroxyquinolino)gallium; TTN, tetrathianaphthacene; TCNQ, tetracyanoquinodimethane; TTC, tetratetracontane; CuPc, copper-phthalocyanine, ZnTPP, 5,10,15,20-zinc-tetraphenylporphyrin; H2TPP, 5,10,15,20-tetraphenylporphyrin; H2T(4-Py)P, 5,10,15,20tetra(4-pyridyl)porphyrin.

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Figure 2

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II.

CONCEPTUAL ASPECTS OF THE ELECTRONIC STRUCTURE OF METAL AND ORGANIC SEMICONDUCTOR

A.

Electronic Structures of Organic Semiconductors

Before discussing organic–metal interfaces, we will reexamine the basic features of the electronic structures of inorganic and organic semiconductors. Figure 3 illustrates the evolution of the electronic structure from an isolated atom to a semiconducting solid. The typical electronic structure of an isolated atom is shown in Figure 3a. Within the potential well, various atomic orbitals (AOs) are formed. The orbitals with small binding energies are va-

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lence AOs, whereas deep AOs localized around the core nucleus with large binding energies are core AOs. The flat part of the potential curve is the vacuum level, which is the energy of a rest electron outside of the atom. In the case of an inorganic crystalline semiconductor, where a large number of atoms (typically ⬃1021/cm3) come together to form a crystal, the individual atomic potential wells merge with each other, forming a periodic potential as shown in Figure 3b. The core levels of the AOs are still localized around the nucleus. On the other hand, the valence AOs form delocalized wavefunctions distributed over the atoms. They are characterized by wave vector (k) or linear momentum due to the translational symmetry of the system. Because the number of the energy levels is of the order of 1021/cm3, these levels can be regarded as quasi-continuous in energy, and the set of such wavefunctions is called an energy band. The energy band with occupied valence electrons is called a valence band, and the upper unoccupied band is called a conduction band. There is an energy gap between the two bands in semiconductors. The flat part of the potential is again the vacuum level. The details of the definition of the vacuum level will be discussed later. The energy diagram is often simplified as Figure 3c, where the potential curves at the internuclei region are omitted and the set of the energy levels is denoted by a box. These energy bands are filled by electrons subject to Fermi-Dirac statistics. The Fermi level, which is determined from the density of states and the number of electrons in the solid, corresponds to the chemical potential of the system. The energy separation between the Fermi level and the vacuum level is called work function. The energy band structure can be characterized with so-called E ⫺ k relation or energy band dispersion. Because of the translational symmetry of the system, the wave vector k is a good quantum number of the system. Each electronic state is characterized by its energy (E) and wave vector (k), as shown in Figure 3d. In the case of organic semiconductors, a set of atoms, possibly of several elements such as H, C, N, and O, forms a molecule. The potential well of the molecule is the superposition of atomic potential wells as shown in Figure 3e. The valence AOs interact with each other and mix to form

< Figure 3 Evolution of the electronic structures of inorganic and organic semiconductors from isolated atoms. (a) An isolated atom. (b) An inorganic crystalline semiconductor. (c) Simplification of (b). (d) An energy band dispersion relation of (c). (e) A polyatomic molecule. (f ) Molecular solids. (g) Simplification of (f ). (h) Typical energy band dispersion relation of organic solids. EF, Fermi level; ⌽, work function; ␹g and ␹, electron affinities in gas phase and solid phase, respectively; Ig and I, ionization potentials in gas phase and solid phase, respectively.

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valence molecular orbitals (MOs), while the core AOs remain localized around the nuclei. The top of the occupied orbital is the highest occupied molecular orbital (HOMO), and the bottom of unoccupied orbital is the lowest unoccupied molecular orbital (LUMO). Energy separation between the LUMO and the vacuum level (VL) corresponds to electron affinity in gas phase (␹g). Ionization potential in gas phase (Ig) is equivalent to the energy difference between the HOMO and the VL. When molecules are condensed to form a solid, the molecular energy diagrams are combined as shown in Figure 3f. In most organic solids, molecules are combined with van der Waals interaction. The outer part of the potential well of an isolated molecule does not merge at the condensation. Thus, valence electrons are often well delocalized over a molecule, but they are still localized within each molecule. The top of occupied state is the highest occupied valence band and the bottom of the unoccupied one is the lowest unoccupied conduction band in the terminology of band theory. However, they are usually called HOMO and LUMO using a molecular orbital model. Electrons occupy these energy levels subject to Fermi-Dirac statistics again. Thus, the concept of the Fermi level can be operative for organic solids, if thermal equilibrium is established in a system with particle exchange among the molecules. This type of diagram is also often simplified as in Figure 3g. The energy separation between VL and LUMO is electron affinity in solid phase (␹), while that between VL and HOMO is ionization potential in solid phase (I). Work function of the solid (⌽) corresponds to the energy difference between the Fermi level and the VL. In a molecular solid, the interaction among molecules is weak; hence the energy band dispersion is usually flat, as shown in Figure 3h. Since this type of diagram is very similar to that for inorganic semiconductor, we often tend to discuss organic semiconductors on the basis of similar model for inorganic semiconductors. However, as mentioned above, some aspects of the electronic structures is quite different between them. Therefore, we have often had to remind ourselves of the basic differences in the electronic structure between them in discussing the semiconducting nature of organics. Various concepts that have been developed for inorganic semiconductors should be examined in terms of their validity for organic materials both experimentally and theoretically. B.

Surface Electronic Structure, Vacuum Level, and Work Function of Metal

Next let us see the basic features of electronic structure of metals. Figure 4 illustrates the energy diagram of a metal. The energy bands are filled up to the Fermi level EF according to Fermi-Dirac statistics. Since the conduction

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Figure 7 Schematic view of an electric double layer. (a) Uniformly charged plates face to each other with the spacing d. The electric field (b) and the potential (c) for electrons are also described along the z axis perpendicular to the plates.

formed at the surface of jellium metal induces potential shift, which forces electrons to remain in the crystal. The shift that corresponds to the difference of the potential outside [␾(outside)] and inside [␾(inside)] the solid gives the surface term ⌽s, as shown in Figure 8. It should be noted here that this dipole layer model is valid only if the electrodes can be regarded as infinitely spread plates. For an electron far on the scale of the dimensions of the macroscopic crystal faces, such a simple potential breaks down, leading to the conclusion that a VL infinitely far from the sample is not suited to define work function as mentioned above. As shown in Figure 8, the bulk term ⌽B corresponds to the difference between ␯xc and បk 2F/2m. The former is the correlation and exchange potential of electron. In the jellium model, only this part remains as potential energy in the Hamiltonian due to the cancellation of core-electron and electronelectron potential. The latter is the kinetic energy of the electron in the Fermi level. The quantities ប, kF , and m are Plank’s constant divided by 2␲, Fermi wavevector, and mass of an electron, respectively.

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Figure 8 The potentials at the surface region of a metal. ␾(inside) and ␾(outside) are the electrostatic potential inside and outside the metal, respectively. ␯xc denotes the total effective one-electron potential. ⌽m can be factored to the bulk term, ⌽B, and the surface term, ⌽S.

In Figure 9, the calculated values of ⌽B, ⌽s, and ⌽ are plotted as functions of rs [49]. For metals with larger rs value, such as an alkaline metal, the contribution of ⌽B is dominant. With decreasing rs, ⌽B decreases and changes from positive to negative, while ⌽s becomes dominant. Work function as the sum of the two terms increases with decreasing rs. This result

Figure 9 The relation between ⌽ and rs derived from the jellium model, including the observed values of the work functions of typical metals [49]. With decreasing rs, the contribution of ⌽S becomes dominant.

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Table 1 Work functions (in eV) of metals related to electrodes of organic devices [53,54] Element Cs K Li Ca Mg Al Ag Cu Pd Au Pt

Polycrystal

(110)

(100)

(111)

2.14 2.30 2.9 2.87 3.66 4.28 4.26 4.65 5.12 5.1 5.65

1.7 1.65

1.7 1.78

1.6 1.85

4.06 4.52 4.48

4.41 4.64 4.59

4.24 4.74 4.98

5.37

5.47 5.84

5.31 5.93

indicates that the work function is substantially affected by the degree of the penetration of electron into vacuum in the case of metals with fairly large work function. On the other hand, the work function is determined only by bulk properties in the case of metals with lower work function. Work function of a metal surface depends not only on rs but also on the surface structure. Surface orientation dependencies of work function for several metals are listed in Table 1. For each metal element, the work function increases with increasing surface density of the metal atoms. For example, (111) surface of an fcc crystal, which has the most dense surface, has the largest values. This effect can be explained as follows. At surfaces, the electron density not only spills beyond the surface along the surface normal direction, creating the dipole layer, but also redistributes parallel to the surface. If a surface has corrugation as shown in Figure 10, the energy can be lowered when the electron distribution moves from the ‘‘hills’’ to the ‘‘valleys’’ resulting in a smoothed charge density surface as shown by dotted curve [50]. This mechanism leads to the electron

Figure 10 Surface roughness induces decrease of work function (Smoluchowski electron smoothing effect). Cases with large corrugation (left) and step (right) are shown.

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deficiency on vacuum side. Thus, the dipole layer tends to decrease the work function of the surface. This effect is called the Smoluchowski electron smoothing effect. According to this effect, a rough surface, such as that of evaporated metal film, has a lower work function than a well-ordered surface in Table 1. As described above, work function is defined as the minimum energy required to remove an electron from a solid. In the case of metal, it corresponds to the binding energy of the electron on a Fermi surface. In the case of semiconductors and insulators, this type of definition is not applicable because they have no Fermi surface. Thus, the definition must be extended. The energy required to remove an electron far from a solid with N electron is

冉 冊

⫺(EN ⫺ EN⫺1) = ⫺

⭸EN ⭸N

= ⫺␮

(1)

T

and equals the negative value of the chemical potential ␮. EN is free energy of N electron system. By using this value, work function of a solid with band gap is defined as

⌽ = ⫺␮

(2)

In other words, the work function is defined as the energy separation between the VL and the Fermi level of the semiconductor. C.

Energy Level Alignment at Semiconductor/Metal Interfaces

In the preceding subsections, each electronic structure of organic semiconductors and metals was discussed. Now we are ready to consider the energy level alignment problem: the question is how energy levels align at organic– metal interfaces. For this purpose, the following simple model has so far been applied to organic interfaces until recently without serious discussion. (As for the basics of the energy level alignment at inorganic semiconductor interfaces, textbooks such as those given in Refs. 51 and 52 are available.) The Mott-Schottky model is a well-known textbook model for inorganic-semiconductor/metal contact. Figure 11 illustrates this model for an n-type semiconductor. When a semiconducting solid and a metallic solid are isolated, the VL at infinite distance from the surface is a common energy level for both solids. When the two solids become in contact, the model requires two assumptions. The first assumption is that the VLs of the semiconductor and the metal coincide (vacuum level alignment). According to this assumption, the barrier height for hole (or electron) injection, ⌽ pB(⌽ Bn ) is determined as follows.

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Figure 11 Energy level alignment models at semiconductor–metal interfaces. Top: Mott-Schottky model. Bottom: the case with vacuum level shift. The energy levels are bent by the charge redistribution in the semiconductor to achieve the electrical equilibrium with the alignment of the Fermi levels of the two solids. This leads to the buildup of built-in potential, Vbi, within a space charge layer.

⌽ pB = I ⫺ ⌽m

(3)

⌽ nB = ⌽m ⫺ ␹ = ⌽m ⫺ (I ⫺ Eg) = Eg ⫺ ⌽ Bp

(4)

where I, ␹, and Eg are ionization potential, electron affinity, and energy gap of a semiconductor, respectively. The assumption of VL alignment is often invalid for inorganic–metal interfaces, where some interface states exist. On the other hand, this assumption has been widely applied for estimation of injection barrier height of organic devices. This is probably due to two reasons. One is that weak interaction at organic–metal interfaces can be ex-

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pected from the van der Waals nature of organic solids. The other is that the VLs of a metal and an organic solid are believed as a good common energy reference between the two solids. However, as pointed out already, the VL at an infinite distance from the surface is not adequate for the definition of work function. The second assumption is the Fermi level alignment of the two solids with band bending. In most cases, the work function is different between a semiconductor and a metal, and there is offset of the Fermi level, which corresponds to chemical potential for this situation. Thus, if the electron (and also hole) can be exchanged across the interface, the total system should go to the thermoequilibrium state with the alignment of Fermi levels between the two solids. In the present example, the semiconductor is charged positively and its energy level is shifted downward since negatively charged electron is stabilized, while the metal is charged negatively and its levels go upward. The amount of positive charge is equivalent to that of the negative charge for the neutrality condition. For the metal, the transferred charges are located just at the surface region, and they are rapidly screened by abundant free electrons in the metal (e.g., 8.5 ⫻ 1022/cm3 for Cu) [51]. The screening distance is typically 0.055 nm (the Thoms-Fermi screening length). On the other hand, the amount of the mobile carrier in a semiconductor is much lower than that of metal. Thus, the charges necessary to the Fermi level alignment distribute near the interface region, forming the so-called space charge layer. Since the electrostatic potential is varying in this space charge layer according to Poisson’s equation, the energy of the electronic levels of the semiconductor is also changing along the potential. This is called ‘‘band bending.’’ In the Mott-Schottky model, the built-in potential Vbi, in the space charge layer corresponds to the difference of the work functions between a metal and a semiconductor, i.e., Vbi = ⌽m ⫺ ⌽s

(5)

The above model is a very naive one, although it has long been applied to discussions of the interfaces of organic devices. However, as will be shown in many examples in this chapter, a common VL is not achieved at most organic–metal interfaces. The observed results are deviated from that expected from the model. Here the deviation is called vacuum level shift, ⌬. This ‘‘shift’’ is actually hypothetical: the concept of vacuum level does not make sense at the interface because there is no vacuum space where a VL is defined. In this sense, an interfacial VL is a virtual concept for conveniently discussing the energetics of the interfaces. When we take account of ⌬, the barrier height can be expressed as

⌽ pB = I ⫺ (⌽m ⫹ ⌬) ⌽ = ⌽m ⫹ ⌬ ⫺ ␹ = ⌽m ⫹ ⌬ ⫺ (I ⫺ Eg) = Eg ⫺ ⌽ n B

(6) p B

(7)

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The existence of ⌬ also affects the behavior of band bending, where Vbi can be expressed as ⌽m ⫹ ⌬ ⫺ ⌽s. Thus, the understanding of energy level alignment, especially the VL shift, is indispensable for estimating and controlling the interfaces of organic devices.

III.

UPS STUDIES ON ENERGY LEVEL ALIGNMENT AT VARIOUS ORGANIC ON METAL INTERFACES

A.

Principle of UPS

When a solid surface is irradiated by light, an electron is emitted by absorbing the photon. The emitted electron is called photoelectron. Photoemission spectroscopy is a method for analyzing the kinetic energy distribution of photoelectrons emitted by monochromatic light. When vacuum ultraviolet light is used, it is suited for investigation of valence electronic structures and called ultraviolet photoemission spectroscopy (UPS). According to the energy conservation rule, the binding energy of the electron to be excited, EB, can be derived from EB = h␯ ⫺ Ek

(8)

Here, h␯ and Ek are the photon energy and the kinetic energy of the photoelectron, respectively. Thus, the electronic structure of occupied states can be experimentally determined. It should be noted that the observed binding energy EB includes many-body effect. When a system with n electrons is considered, energy conservation gives h␯ ⫹ E(n) = E(n ⫺ 1) ⫹ Ek

(9)

where E(n) is total energy of the system with n electrons and E(n ⫺ 1) is that of the ionized system with n ⫺ 1 electrons. Then, EB can be expressed as EB = E(n ⫺ 1) ⫺ E(n)

(10)

This equation demonstrates that photoemission is a method for probing the energy of ionized states. Thus, when the ionized final state is energetically relaxed by screening of the surrounding medium, the term of E(n ⫺ 1) decreases, leading to the decrease of EB. As is described later, this effect becomes often important at the interfacial region. For other details, good references are available regarding this method [54,55]. The electronic structure of organic–metal interfaces can be experimentally studied by forming an interface either by depositing an organic material onto a metal substrate or by depositing a metal onto an organic

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solid. Here the principle of application of UPS to an interface study is explained for the former case. First we measure the UPS spectrum of a clean metal surface as shown in Figure 12a. The valence electrons of the metal are excited by incident UV light and emitted out into vacuum. The kinetic energy (Ek) distribution of the emitted electrons (photoelectrons) is called the UPS spectrum and roughly reflects the density of the occupied states of the solid. The schematic spectrum of UPS is illustrated at the upper part of the energy diagram, where the kinetic energy axis is set vertically. The photoelectron with the maximum kinetic energy [E max k (metal)] corresponds to the excitation from the Fermi level of the metal, while the low-energy cutoff (the position at Ek = 0) corresponds to the VL. By Eq. (8), the work function of the metal, ⌽m, can be determined as

⌽m = h␯ ⫺ E max k (metal)

(11)

As an example, the UPS spectrum of Au substrate is shown in the lower part of Figure 12c for the low kinetic energy part and the lower part of Figure 12d for the high kinetic energy part. The step at the maximum kinetic energy in Figure 12d is due to Fermi edge labeled by E mF , while the onset in Figure 12c corresponds to the vacuum level of the Au substrate. Next we deposit organic layer on the metal substrate in a stepwise manner and then measure UPS at each deposition step. For a sufficiently thick organic layer, the spectrum becomes that of the organic layer, since the escape depth of electrons is a few nanometers at most. This situation is shown in Figure 12b. In this case, the onset at the maximum kinetic energy [E max k (org)] corresponds to the emission from the highest occupied molecular orbital (HOMO), and the low-energy cutoff corresponds to the VL at the surface of the organic layer. Here threshold ionization of organic material (I) can be determined as I = h␯ ⫺ E max k (org)

(12)

By comparison with the UPS spectra, we can determine the relative positions of the energy levels at the interface. As an example, the spectrum of 20-nm-thick Alq3 film on Au substrate is shown in the upper part of Figure 12c and d. From the energy difference of the high kinetic energy onset between the two spectra, we can determine the energy of the top of the valence states ␧ Fv relative to the Fermi level of the metal as shown in Figure 12d. Also the energy of the VL of the organic layer relative to the Fermi level of the metal ␧ Fvac is obtained as

␧ Fvac = I ⫺ ␧ vF

(13)

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Figure 12 Principle of the UPS study of an organic–metal interface. (a) Photoemission from the metal. (b) Photoemission from the organic layer deposited on the metal substrate. (c) High binding (low kinetic) energy region of UPS of Au and Alq3 film on Au. The change of the cutoff corresponds to the vacuum level shift ⌬. (d) Low binding (high kinetic) energy region of UPS of Au and Alq3 film on Au. The change of the threshold gives to the HOMO energy, ␧ Fv .

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It should be noted here that ␧ Fvac is defined relative to the Fermi level of the ‘‘metal’’ substrate but not that of the ‘‘organic’’ layer. Basically, energies of various levels are determined by UPS relative to the Fermi level of an electron energy analyzer, which is equal to the Fermi level of the electrically conducting substrate or the sample holder. Thus, the primary energy origin becomes the Fermi level of the sample holder or substrate. Only if Fermi level alignment is achieved between the organic layer and the metal, ␧ Fvac corresponds to the work function of the organic layer. However, this condition is often not met at least for the experiments in ultrahigh vacuum condition. The observed quantities of ␧ Fv and ␧ Fvac often depend on the film thickness ␪. The VL shift can be considered as the difference of the VLs between clean metal surface and organic monolayer on it. The VL shift ⌬ can be obtained as

⌬=

lim

␪ → monolayer

(␧ Fvac ⫺ ⌽m)

(14)

⌬ thus defined corresponds to the work function change by monolayer adsorption in surface science. If the film does not grow in layer-by-layer mode, we cannot get the data for a monolayer. In such a case, some ambiguity remains in determining the value of ⌬.

B.

Model Interfaces of Organic–Metal Interfaces in Organic Electroluminescence Devices

In the remainder of Section II, we describe the general aspects of the UPS studies for organic interfaces, taking examples mostly from systems in relation to organic EL devices depicted in Figure 1a. Also other related systems were studied for elucidating the details of interfacial electronic structures. If the vacuum level shift at an organic–metal interface can be predicted, it will be useful to estimate and control the barrier height to optimize the performance of organic devices. Unfortunately, the mechanism of the VL shift has not yet been so well understood as to predict the shift at the interface of real devices under low vacuum or ambient conditions. However, as for ‘‘ideal interfaces,’’ prepared by depositing organic material on clean metal substrate and also measured in ultrahigh vacuum, accumulation of data is now in progress for qualitative understanding of the VL shift. These results not only give the general outlook about the interfacial electronic structure but also hint at the origin and mechanism of factors dominating the interfacial electronic structures, which will be discussed in the next section.

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HTL–Metal Interfaces

At first we will examine the hole-injecting interfaces, where a hole transport layer (HTL) contacts a metal as the anode. As the first example for holeinjecting interfaces, UPS spectra of TPD film on Au substrate are shown as a function of film thickness from 0.05 to 10 nm in Figure 13 [14,56]. The bottom spectrum is for the Au substrate evaporated in the pressure range of ⬃10⫺7 –10⫺8 torr, which is a typical pressure for fabrication of real devices. The abscissa is the electron binding energy relative to the Fermi level of the Au substrate. Because of strong emission from the Au substrate, the spectral features can be regarded as the overlap of those of TPD and Au in the low-thickness region. With increasing thickness, the contribution from TPD becomes dominant. The top UPS spectrum reflects the bulk electronic structure of TPD. The cutoff at the lower binding energy side (left) corresponds to the vacuum level of the system and its position clearly depends on the film thickness as indicated by the dotted lines. Figure 14 shows the thickness dependence of the energy of the VL relative to the Fermi level of the metal substrate (␧ Fvac) of the TPD film derived from the cutoff, together with the onset energy of TPD (␧ Fv ) determined from the high-energy onset (right side) of the spectra. Initial deposition of TPD induced abrupt decrease of ␧ Fvac in the thickness range up to around 0.5 nm. The vacuum level was shifted slightly downward by the further deposition up to 2 nm thickness and then became almost constant in the larger thickness region. This variation of the vacuum level energy clearly indicates that the assumption of a common vacuum level is not valid for the interface. The initial steep decrease of ␧ Fvac suggests that the potential drop across the interface is confined to the interfacial region. When we use more clean Au substrate, the VL shift becomes larger. The observed VL shift increased to 1.25 eV with increasing the work function of Au substrate from 4.7 to 5.1 eV. This result suggests that the VL shift is very sensitive to surface cleanliness. The HOMO energy, ␧ Fv , also exhibited film thickness dependence as shown at the bottom of Figure 14. The observed values of ␧ Fv shows downward shift up to the thickness of about 1.5 nm, and then the slope becomes smaller. One of the possible reasons for the small variation up to 1.5 nm thickness is the change of the relaxation energy of the final state of the photoemission event: In photoemission spectroscopy of molecular systems, the final state is a photoformed molecular positive ion. The ionic final state is relaxed by the electronic polarization of the surrounding molecules and metal surface by the positive charge of the ion. As the thickness of the organic layer increases, the metal substrate, which is more polarizable than the organic layer, becomes far from the outermost region of the film, which

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Figure 13 UPS spectra of TPD film on a Au substrate as a function of the TPD thickness. The shift of the left-side cutoff corresponds to the vacuum level shift ⌬ in Figure 12.

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Figure 14 Plot of the energies of the vacuum level and the HOMO level of TPD film on Au determined from the spectra of Figure 13 as a function of TPD thickness.

mainly contributes to the UPS spectrum. This induces the decrease of the relaxation energy for the hole left by photoemission. Since the Au substrate is polycrystalline film, the inhomogeneous film growth on it also affects the observed thickness dependence. The final state of the hole injection process contains a hole in a molecule as in the case of photoemission. Thus, the barrier height for hole injection can be estimated to be about 0.7 eV, which is the value of the observed ␧ Fv around monolayer thickness, when we neglect the additional relaxation energy due to organic layers over the first layer. Similar shift of vacuum level was also observed at TPD/indium-tinoxide (ITO) interface, which is a typical anode interface of organic EL devices. Figure 15 shows the observed energies of the vacuum level (␧ Fvac) and HOMO (␧ Fv ) of TPD film on ITO as a function of the film thickness. The ITO electrode was cleaned by UV ozone treatment. The vacuum level was gradually shifted downward by about 0.3 eV at the deposition of TPD up to 1 nm and showed no significant change by further deposition, indi-

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Figure 15 Plot of the energies of the vacuum level and the HOMO level of TPD film on an indium-tin-oxide (ITO) substrate as a function of TPD thickness.

cating the lack of significant change in final-state screening. This is probably due to the lower dielectric constant of ITO relative to metal because it is not an intrinsic metal but a degenerate semiconductor. The direction of the VL shift at the ITO interface is similar to that for the Au interface; i.e., lowering of the VL was observed for both cases. On the other hand, the value of 0.3 eV was smaller than that at the Au interface. Such a small value of ⌬ was also observed at the interfaces formed between ITO and starburst molecules [16], which form a newly developed class of compounds for HTL. These small values may be due to the surface contamination or due to the low concentration of mobile carrier as discussed later. Because of such small ⌬, the HOMO level comes closer to the substrate Fermi level of ITO than that at the Au interface, and the barrier height for hole injection into TPD is decreased to 0.5 eV for the ITO interface. This difference suggests that the ITO electrode is better than Au from the

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viewpoint of the barrier height for hole injection, although they have similar work functions. 2.

ETL–Metal Interface

Next we examine the interfaces between electron transport layer (ETL) and cathode metals. In this type of interface, an organic material with large electron affinity and a metal with lower work function are often used. Here we show the results for Alq3 /Al interfaces. Figure 16a shows the energy diagram at the interface constructed from UPS measurements of Alq3 film deposited on clean Al substrate. The energy of LUMO was derived assuming

Figure 16 (a) Energy diagram of Alq3 /Al constructed from UPS results. (b) Energy diagram of Alq3 /Al estimated with an assumption of vacuum level alignment. (c) and (d): Diagram of Alq3 /LiF/Al interfaces. The thickness of LiF layer is 0.5 nm and 2 nm for (c) and (d), respectively.

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that HOMO-LUMO gap equals to the energy of optical absorption edge (⬃2.7 eV). This assumption includes two possible errors; the first is due to an exciton binding energy. In the case of organic materials, the absorption edge of HOMO-LUMO transition is often dominated by the exciton absorption to form Frenckel-type exciton, where HOMO → LUMO transition occurs accompanied by the Coulomb interaction between the electron in the LUMO and the hole in the HOMO. Thus, the actual HOMO-LUMO gap must be larger by the excitonic binding energy than that of the absorption edge energy. The determination of the actual band gap from optical spectra is difficult because the intensity of the transition into unbound electron-hole pair is generally weak. The second possible error is due to the screening of injected electron in real device interface. While the final state of the optical absorption is still neutral in charge, the final state for the electron injection is a negatively charged ion molecule. At that case, the injected electron is screened by the polarization of the surrounding Alq3 molecules and the surface region of the cathode metal. Hence, the actual HOMO-LUMO gap must be smaller by such relaxation energy. From these two factors it is noted that the estimation of HOMO-LUMO gap from absorption edge should not be so accurate. Very recently, the energy location of the LUMO of Alq3 was directly measured with inverse photoemission by Kahn et al. [57]. This method is free from the above-discussed errors. From the spectra, the gap can be determined to be larger than 4 eV if the energy separation of the thresholds of UPS and IPES is assumed to correspond to the gap. The LUMO position estimated from this value is shown in Figure 16 as LUMO*. Both LUMO and LUMO* are denoted by boxes, emphasizing the possibility of the error. As shown in Figure 16a, the VL shift of ⫺1.4 eV was observed and the LUMO is very close to the Fermi level of Al. This is consistent with the favorable electron injecting nature of this interface. It should be emphasized here that the traditional assumption of a common VL leads to the wrong picture; Figure 16b is the energy diagram derived from the traditional assumption of a common VL, which neglects ⌬. Here the Fermi level is assumed to locate around the midpoint of the HOMO-LUMO gap. In this case, the barrier height for electron injection becomes larger and we cannot expect effective electron injection at this interface in contrast to the observation for the real interface. The observed large ⌬ of ⫺1.4 eV is probably related to the strong interaction between Alq3 and Al as follows. Figure 17 shows the valence top region of UPS spectra of Alq3 /Al interfaces [18]. The bottom spectrum is for clean Al surface, and others are for 0.5-, 1-, 2-, and 5-nm-thick Alq3 film on the Al substrate, respectively. The lineshape of the spectrum of 5nm Alq3 film is similar to that of 20-nm-thick Alq3 film, where the onset peak labeled A is due to the HOMO level. At the spectrum of 0.5 nm Alq3,

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Figure 17 HeI UPS spectra of Alq3 on clean Al surface. Peak A is due to HOMO of the original Alq3 film. For the low-thickness region, a gap state labeled X was observed.

on the other hand, an extra structure labeled X was observed at lower binding energy side than the HOMO. A similar gap state has been reported for an Al-on-Alq3 system [36]. The formation of such a gap state indicates that Alq3 is chemisorbed on the Al surface. The observed large ⌬ is probably due to the chemical interaction at the interface. A similar gap state has been reported for Mg/Alq3 and Ca/Alq3, although the behavior of the peak is different from that for Al [33,36]. Recently, it was found that the insertion of an insulating layer such as poly(methyl methacrylate) (PMMA) [58], LiF [59], and aluminum oxide [60] improves the performance of organic EL devices. Figure 16c is the energy diagram of Alq3 /LiF (0.5 nm thick)/Al interface, which was reported to show the optimum EL performance. VL shifts of ⫺1.0 eV and ⫺0.8 eV were observed at Al/LiF and Alq3 /LiF interfaces, respectively. The overall shift is ⫺1.8 eV, whereas it is ⫺1.4 eV for the direct contact of Alq3 on Al. This result indicates that the insertion of LiF decreases the barrier height for electron injection. Metastable atom electron spectroscopy (MAES) revealed that the surface of Al substrate was partially covered by LiF due to an island

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growth of LiF on this substrate [61]. At 20 nm thickness, the surface was completely covered by LiF, and the energy diagram of the interface at this stage is shown in Figure 15d. The value of ⌬ at the LiF/Al interface was ⫺1.7 eV, whereas it is ⫺0.1 eV at the LiF/Alq3 interface. The potential drop across the interface occurred only at the LiF/Al interface, and the magnitude of the overall ⌬ is still larger than that for the Alq3 /Al interface. Similar results were reported for Alq3 /LiF/Al [62] and Alq3 /Al2O3 /Al [44] interfaces, and the decrease of the barrier height was proposed as a reason for the efficiency improvement. Besides the change in the barrier height, the chemical interaction between Alq3 and Al is suppressed by the insertion layer. Figure 18 shows the spectra of Alq3 (0.5 nm)/LiF/Al system with increasing thickness of LiF (d). With 0.5-nm-thick LiF layer, the extra structure X almost disappears. The interface discussed here was fabricated on Al substrate, which is in reverse to the real device structure. Similar suppression of the interaction was re-

Figure 18 HeI UPS spectra of Alq3 (0.5 nm) of LiF-deposited Al surfaces. The gap state X disappears with increasing thickness of LiF.

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ported for the interface that was formed by depositing Al metal onto the organic layer, as in the real device structure [36]. As another example of ETL interface, the energy diagrams of DPNTCI/(Al or Au) interfaces obtained by UPS are shown in Figure 19. The film thickness of the organic layer is 2.5 nm. At the DP-NTCI/Au interface, the Fermi level of gold is much closer to the LUMO than to the HOMO, corresponding to the electron injecting character of the interface [63]. It is interesting to examine the DP-NTCI/Al interface because electron injection may be made even easier by the use of Al with a lower work function than Au. Actually, the results in Figure 19 shows that the VL is shifted upward at DP-NTCI/Al interface (⌬ = ⫹0.4 eV). This is in contrast to the case of other interfaces, where downward shifts were observed. Thus, the barrier height for electron injection increased despite the fact that the work function is decreased from Au to Al. The positive sign of ⌬ suggests the formation of interfacial dipole layer with DP-NTCI side negatively charged. This can be explained by the occurrence of electron transfer from Al to DP-NTCI, which is a medium-strong electron acceptor. This transfer does not occur for Au with larger work function. This interpretation is confirmed by our result of TCNQ/Au interface in Figure 20, where upward VL shift of ⫹0.2 eV was also observed. TCNQ is a stronger electron acceptor than DP-NTCI and can extract electrons even from Au. Thus, the popular assumption that cathode material with lower work function is preferable for decreasing the barrier height for electron injection is not the case for DPNTCI/metal interfaces, indicating that caution must be paid in the selection of organic–metal combination to achieve good carrier injection.

Figure 19 The interfacial energy diagrams of DP-NTCI/Al (left) and DP-NTCI/ Au (right) interfaces obtained from UPS results.

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Figure 20 Plot of the energies of the vacuum level and the HOMO of TCNQ film on a Au substrate as a function of TCNQ thickness.

3.

Organic–Organic Interfaces

To construct the whole-energy diagram of organic EL devices, the energy level alignment at organic–organic interfaces is necessary for multilayer devices. Although this is not the main subject of this chapter we will discuss it briefly. For several organic–organic interfaces, such as Alq3 /TPD [14,56], Alq3 /BCP, ␣-NPD/BCP, CBP/BCP, CuPc/BCP [29], Alq3 /CBP [25], and CuPc/␣-NPD [30], no significant shift of VL was observed, indicating that the traditional assumption of a common VL is feasible at organic–organic interfaces. This is probably because van der Waals interaction is expected to be dominant at organic–organic interfaces. However, finite VL shift can still be induced for some interfaces. Figure 21 is the energy diagram of TTN/TCNQ interfaces measured by UPS [14,56]. TTN is a strong electron donor, and TCNQ is a typical acceptor. The deposition of TCNQ on TTN film exhibited an upward shift of vacuum

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Figure 21 Plot of ␧ Fvac of TCNQ layer deposited on TTN film against the film thickness of TCNQ.

level of ⫹0.2 eV. This behavior can be explained by the formation of an interfacial dipole layer with the TCNQ side negatively charged. This polarity is consistent with the charge transfer nature of TTN-TCNQ combination. Fairly large VL shift was also reported for PTCDA/Alq3 and ␣-NPD/Alq3 interfaces (⌬ ⬃ 0.5 and 0.25 eV, respectively) where the origin of the interfacial dipole formation is not clearly understood [26]. From the results described above, we can construct the whole-energy diagram of a typical organic EL device, i.e., Al/Alq3 /TPD/ITO system as shown in Figure 22. We note that the energy level alignment at Al/Alq3 interface was estimated using the value of ⌬ for Alq3 on Al. Of course, the data for ⌬ of Al on Alq3 film is more desirable. However, it is not easy to precisely determine the energy level alignment for such metal-on-organic system as described later. This energy diagram clearly demonstrates that the barrier height for carrier injection is significantly modified by the VL shift at the cathode and anode interfaces.

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Figure 23 Plot of the energy of the vacuum level of TTN film on a Au substrate as a function of TTN thickness. The interfacial energy diagram is shown in the inset. The position of the HOMO is estimated using the ionization potential of TTN for the thick film.

across the organic–metal interface seems to induce the interface dipole with proper polarity as expected from the charge transfer nature of the combination of the organic material and an electrode. D.

Organic Wide-Gap Insulator–Metal Interface

In the examples described above, we see that strong chemical interaction such as charge transfer can be a possible factor for the interface dipole formation. In the case of organic insulator with a wide HOMO-LUMO gap, the frontier orbitals of organic insulator are far from that of metal electrode (i.e., Fermi level). Thus, we can expect a very small chemical interaction or charge transfer at the interface. It is interesting to see whether the VL shift can ever occur even for such an inert interface. A long-chain alkane has a wide gap of about 9 eV [64] and large I of 8.5 eV [65], with the LUMO even above the VL. From angle-resolved UPS [21] and infrared reflection absorption spectroscopy of n-C44H90(TTC)/ Au(111) [66], the physisorptive nature of this system was concluded. Figure 24 shows the energy diagrams of the TTC/Au interface derived from conventional estimation (a) and UPS measurements (b) [15]. UPS results clearly demonstrate the occurrence of VL shift of ⫺0.7 eV. Similar results were

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Figure 24 Energy diagrams of TTC/Au interfaces. (a) Energy diagram estimated assuming a common vacuum level. (b) Diagram constructed from UPS results.

obtained at TTC/Ag (⌬ = ⫺0.5 eV), TTC/Pb (⌬ = ⫺0.3 eV) [15], TTC/ Cu(100) (⌬ = ⫺0.5 eV) [21], and TTC/Au(111) (⌬ = ⫺0.7 eV) [20] interfaces. These findings clearly indicate that the traditional model does not work even at insulator–metal interface. The occurrence of such VL shift at a weakly interacting system suggests the contribution of physical effect to ⌬, as described in the following section.

IV.

TRENDS IN VACUUM LEVEL SHIFT AND THE ORIGIN OF DIPOLE LAYER

In the last section we described that the vacuum level shift at organic–metal interfaces is an important factor governing the interfacial electronic structures. In this section, the origins of the formation of interfacial dipoles and the factors affecting the value of ⌬ are discussed in detail. At first we describe the trend of the value of ⌬ and discuss the possible origins of the dipole formation. A.

Trend of the Vacuum Level Shifts

Systematic studies of ⌬ for various combinations of organics and metals were carried out in the groups at Nagoya and Princeton. Based on these data, the values of ⌬ for various organic materials are plotted as functions of ⌽m in Figure 25. At all the interfaces, nonzero values of ⌬ are observed. For some interfaces, the magnitude of ⌬ reaches 1 eV, which is comparable to the HOMO-LUMO gap (typically ⬃2–3 eV), suggesting that the effect

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Figure 25 The relation between the observed ⌬ and ⌽m from our UPS measurements of various org–metal interfaces. The substrate metals are denoted beside the data points, and the symbols for organic materials are shown in the right panel. Data with * and ** are from Ref. 27 and 39, respectively. For the structures and the full names of the compounds, see Figure 2 and caption.

of the VL shift can be significant in discussing the interface of organic devices. As seen in the plot, there is no analytical relation between ⌬ and ⌽m, which can express the total behavior of all data points. However, phenomenologically we can find two trends in the magnitude and polarity of ⌬, which corresponds to the trends of the error of the traditional estimation method. First, the polarity of ⌬ is negative except for strong acceptor–metal interfaces such as DP-NTCI/Al, TCNQ/Au, and PTCDA/Mg,In,Sn interfaces. This means that the vacuum level of the metal substrate tends to be

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lowered by the deposition of organic material, indating the formation of dipole layer with the vacuum side positively polarized. In other words, the traditional method of estimation tends to underestimate (or overestimate) the barrier height for hole (electron) injection at organic–metal interfaces. The positive values of ⌬ for acceptor/metal interfaces can be explained that electron transfer from metal to acceptor molecule forms interfacial dipole layer with the vacuum side negatively polarized, leading to the increase of the work function. Second, the magnitude of ⌬ seems to be monotonic function of ⌽m, i.e., the value of ⌬ tends to decrease or remain constant with increasing the work function of the metal substrate. For some interfaces, good linear relationship was observed. This trend suggests that (1) ⌽m of metal substrate is one of the physical parameters to determine the VL shift, and (2) the error induced by the traditional estimation often becomes larger for the anode interface, where ⌽m is large, than the cathode interface. Such linearity against the metal work function has been also observed for inorganic semiconductor–metal systems. The slope parameter defined as S=

d⌽ nB d⌽m

(15)

is used for characterizing such relation. It is known that S depends on the type of the semiconductor [52]. Since ⌽ nB can be expressed as ⌽ nB = ⌽m ⫹ ⌬ ⫺ (I ⫺ Eg), from Eq. (7), we obtain S=1⫹

d⌬ d⌽m

(16)

indicating that the slope of the VL shift against ⌽m corresponds to S ⫺ 1. In the field of inorganic semiconductors, the following are known. When Mott-Schottky rule holds, i.e., there is no interface state at the interface (Schottky limit), ⌬ = 0 and the value of S should be unity. On the other hand, when there is interface state with sufficient density to pin the Fermi level of the semiconductor, S should be zero (Bardeen limit). The reported values of S in most inorganic-semiconductor/metal junctions are between zero and unity. For example, S is about 0.1 for Si, GaAs, and Ge, and about 0.5 for compound semiconductors such as ZnSe and ZnS, suggesting the effect of the interface states [67]. The values of S obtained from Figure 25 are listed in Table 2. S in the table ranges from 0 to 1 as in the case of inorganic semiconductor–metal interfaces. The value of S strongly depends on material. The values of S for acceptor-type molecules, such as DP-NTCI, C60, and PTCDA, are close to zero, suggesting that the barrier height at the interface is almost independent

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Table 2

Interface parameter S reported for several organic–metal interfaces

Organic materials ZnTPP H2TPP H2T(4-Py)P TTC(C44H90) DP-NTCI C60 PTCDA Alq3

␣-NPD CBP

Substrate metal

Interface parameter, S

Ref.

Mg, Al, Ag, Au Mg, Al, Au Mg, Al, Au Pb, Ag, Au Cu(100), Au(111) Al, Au Mg, Bi, Cr, Ag, Au Mg, In, Sn, Au Mg, Ag, Au Mg, Al, Au Mg, Au Mg, Ag, Au

1.0 0.5 0.4 0.6 0.5 ⬃0 ⬃0 0 0.9 ⬃1 0.49 0.6

7, 9–11 9–11 9–11 15 20, 21 12 69 27 27 70 27 27

of the work function of metal as in the case of the Bardeen limit. On the other hand, S is close to unity for ZnTPP interface, indicating that the vacuum level shift is almost independent of ⌽m. This case is similar to that of the Schottky limit, although the ⌬ is not zero. Recently, Ramsey et al. reported on the energy level alignment at the interfaces between bithiophene and well-defined surfaces such as Ni(110) and Al(111) [68]. They observed no linear relationship between ⌬ and ⌽m. As described in the next section, there are several origins of the formation of interfacial dipole, including chemical and physical ones. Depending on these origins, the actual relationship between ⌬ and ⌽m can be linear or can be deviated from monotonic function. To clarify this point, additional data on ⌬ and ⌽m are needed. B.

Possible Origins of Interfacial Dipole Layer

The clarification of the origin of the formation of dipole layer at organic– metal interfaces is indispensable for estimating and controlling the barrier height. Although the origin is not yet fully clear, we have proposed that the VL shift is caused by the possibly combined contributions from the factors described below. 1.

Charge Transfer and Bond Formation by Chemisorption

When a molecule is adsorbed on a metal surface with strong chemical interaction between the adsorbate and the surface, it is called chemisorption.

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The first layer of organic solid at the organic–metal interface is critical for understanding the interfacial dipole layer, and the layer can be regarded as an adsorbed system. The interaction between the first layer and other organic layers is usually weak van der Waals interaction. Thus, the perturbation from other layers can be approximately neglected. Work function change by molecular adsorption has been widely reported for various chemisorption system. The work function change has usually been ascribed to the redistribution of the electrons by (1) charge transfer and (2) hybridization of wavefunctions of adsorbate and surface. Charge transfer between an adsorbate and the surface can be expected for the interface between a strong acceptor (or donor) molecule and a low (high) work function electrode, inducing anion (cation) formation, as shown in Figure 26a1 and a2. Anion formation causes the dipole layer with the vacuum side to be negatively polarized, which raises the VL of organic side relative to metal side (positive ⌬). Such positive values of ⌬ have been reported for TCNQ/Au, DP-NTCI/Al [14], and PTCDA/Mg,In,Sn [27] systems. C60 /metal is also the case. Ohno et al. [69] reported an almost constant value of ␧ Fv for various metals, which corresponds to an interface parameter S of almost zero. They ascribed this to the charge transfer from metal into LUMO through weak mixing of the metal wavefunction and the LUMO wavefunction, even claiming that this leads to Fermi level alignment. A similar trend was reported for PTCDA–metal systems by Hill et al. [27]. The value of ␧ Fv is also almost constant for various metals. These phenomena like a pinning may be related to the acceptor-like nature of this type of materials. In the case of cation formation, the dipole layer is formed with the vacuum side polarized positively, making the vacuum level at the organic side lower than that at the metal side. This mechanism should give a negative ⌬. The observed negative ⌬ for TTN/Au [56], TPD/Au [14], and ␣-NPD/ Au [27] can be ascribed partly to this effect. As for negative ⌬, other factors described below can also induce the VL shift with the same polarity as the cation formation. Thus, it is not easy to distinguish the contribution from charge transfer effect among them. Examples for hybridization of wavefunction at adsorption can be widely seen in the case of adsorption of small molecules, such as CO, ethylene, and benzene, to metal surfaces. Parts of their molecular orbitals mix with the wavefunction of metal surface, leading to chemical bond formation at the interface [70]. Such interaction accompanied with charge redistribution at the interface can contribute to the work function change of the surface. At the interface between metal and organic semiconductor with ␲ system, such an interaction can be expected. The behavior of Alq3 and ZnTPP interface may be due to this kind of chemical effect. Actually, a mid-gap state

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Figure 26 Possible origins of the formation of electric double layer at organic– metal interfaces. Chemical effects: (a1 and a2) Charge transfer across the interface. (b) Charge redistribution by the adsorption bonding. Physical effects: (c) Shift of electron cloud of the molecule toward the metal surface by image force. (d) Rearrangement of electron cloud of the metal surface with the reduction of tailing into vacuum. (e) Existence of interface state serving as a charge reservoir. (f ) Orientation of polar molecules or functional groups.

was reported for Alq3 /Al systems [18], indicating strong chemical interaction at the interface. In general, a chemical bond has ionic bond character due to both charge transfer and covalent bond character. The relative ratio of the two characters depends on the situation. In the case of strong donor–acceptor systems, CT character may be dominant. However, both factors will be operative for most cases. There seems to be a trend in the literature to explain all the observed VL shifts by charge transfer. There are two points to be claimed about such a trend.

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First, chemisorption formation can contribute to the VL shift not only through charge transfer but also by charge redistribution, as shown in Figure 26b. Thus, the dipole layer formation does not directly mean the occurrence of charge transfer across the interface. We have to pay attention to the character of the chemisorption bonding before using the term of ‘‘charge transfer.’’ Second, VL shift can occur even in the case without strong chemical interaction. As described above, significant shift of VL was observed at the TTC/Au interface, where strong chemical interaction cannot be expected. This indicates that a dipole layer must be formed by a physical mechanism, as described below. 2.

Image Effect and Surface Rearrangement by Physisorption

When a molecule is adsorbed on metal surface without strong chemical adsorbate–surface interaction, it is called physisorption. In this scheme, van der Waals interaction is dominant in the adsorbate–surface interaction. Since molecular semiconductors form molecular solids with van der Waals interaction, it is suggestive to refer to the rare gas–metal system when discussing the origin of the dipole layer in such cases. At the first glance, van der Waals interaction seems to be unable to induce significant charge redistribution because the molecular electronic structure is basically preserved in physisorption. However, it is well known in the field of surface science that physisorption of Xe atoms onto clean metal surface induces a fairly large change in the work function, although these systems are typical physisorption systems, reflecting the nature of rare gas Xe atom. In Figure 27, the observed values of work function change at Xe adsorption on clean metal surfaces, which corresponds to ⌬ at Xe–metal interfaces, are shown [71]. The plot is superimposed on the plot of Figure 25. We can see that the polarity of ⌬ is negative for all cases. Although the data points are somewhat scattered, ⌬ seems to increase with increasing ⌽m. Significant shift was observed for metals with higher work function, while ⌬ is almost zero for lower work function metals such as alkaline and alkaline-earth metals. These tendencies are very similar to those in organic–metal interfaces described above. One explanation for the work function change for physisorption system is adsorption-induced dipole in the adsorbate [72]. Consider a Xe atom on metal surface. Because of the motion of the electrons in Xe atom, the atom has instantaneous multipole moments that generate dynamic long-range fields with which the metallic electrons interact. If the metal is assumed to be a perfect conductor, its electrons respond instantaneously to the fields

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Figure 27 The relation between ⌽m of metals and the work function change ⌬ at Xe adsorption on the metal surfaces after Ref. 71. Metal substrate is denoted beside each data point. The plot was superimposed on the gray plot of Figure 25.

generated by the atomic charges and completely screen them. This leads to an attractive interaction between the adsorbate and the metal surface, i.e., van der Waals interaction. The interaction of the fluctuating dipole moment with the screening charge can derive to a net displacement of the electron distribution in the adsorbate. Such displacement is predicted from the result of local density approximation (LDA) of a Xe/jellium system [73,74], as shown in Figure 28. The left part shows the contours of constant charge density, and the right denotes the difference in charge density between Xe/ jellium and a superposition of the isolated jellium metal and Xe atom charge densities. The charge density at the region between the jellium and the Xe atom is increased by adsorption. Although it seems to indicate the bond formation in the usual chemical sense, bonding and polarization occur actually because the energy of the valence electrons becomes minimal when surrounded by their exchange-correlation hole as much as possible. In a simple physical picture, this effect can be described as follows: consider the

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Figure 28 Contours of constant charge density of Xe on jellium (rs = 2); total density (left) [74] and difference density (right) [73]. In the right, solid lines indicate that electronic density increased by adsorption, and dotted lines denote the opposite.

electron cloud of the valence electron of adsorbed an Xe atom. An electron closer to the metal substrate is attracted more strongly by its positive image charge than an electron far from the substrate. Thus, the distribution of the electron cloud is shifted toward the metal surface [75]. In this case, the adsorption-induced polarization always leads to the formation of dipole layer with its vacuum side positively polarized, leading to the lowering of the vacuum level. Note that in this case the rearrangement of the electron cloud inside the adsorbate is important. A second explanation is the rearrangement of the electron cloud at the metal surface. As described in the introduction, the electron spills out into the vacuum region, forming a surface dipole layer. When Xe atom is adsorbed on metal surface, the spilled electron feels repulsive interaction from the closed shell of the Xe atom. Thus, the tailing part of the electronic cloud is pushed back, causing a decrease in the potential drop across the dipole layer. The net change by adsorption corresponds to the formation of an additional dipole layer with the vacuum side positive. This model also leads to lowering of the vacuum level by adsorption. Note that the main change in this case is not in the adsorbate but at the metal surface. As shown in Figure 9, the contribution of surface dipole to the work function of metal is increased with decreasing rs. that is, ⌽m is determined mainly by the bulk term for low ⌽m metal such as alkaline metals, while the surface term is dominant for other metals with higher ⌽m. Thus, we can expect that the decrease of surface dipole by push-back effect becomes more important for metal surfaces with higher ⌽m. This picture is consistent with the observed trend of monotonic relation between ⌬ for Xe systems. Now we will go back to the TTC/metal system. The observed finite shift of vacuum level can be explained by these two mechanisms described

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above for Xe. As for the image-effect model, the static polarizability of typical organic molecules is comparable to that of Xe. For example, the polarizabilities (in units of 10⫺24 cm3) of methane, ethane, n-hexane, and benzene are 2.593, 4.47, 11.9, and 10.0, respectively, while that for Xe is 4.044. This suggests that electron cloud of an organic molecule can be polarized as well as in Xe. The push-back effect will be also operative, and will explain the increase of 兩⌬兩 with ⌽m. Since these mechanisms are operative for all adsorption systems, we can expect that they can be applied for all organic–metal interfaces. They explain the general trend of VL lowering and the increase of 兩⌬兩 with ⌽m in some case. We can also think about the combination with chemical interaction, e.g., charge transfer. For example, in the case of cation formation, VL lowering is caused by the concerted combination of image and pushback effects. On the other hand, for anion formation for acceptor adsorbates such as TCNQ, PTCDA and DP-NTCI, upward shift of VL by anion formation is partly compensated by these effects. 3.

Other Origins

There are also possible mechanisms contributing to the dipole layer formation. One is the possible existence of interface states. For some organic– metal interfaces, we observed almost linear relation between the ⌬ (also ␧ Fv ) and ⌽m. Phenomenologically, this trend is similar to that observed for inorganic semiconductor–metal interfaces. In the latter case, it has been explained by the existence of interface states. If these states have sufficient densities of state, the Schottky barrier height becomes independent of ⌽m (Bardeen limit, S = 0), while S becomes unity if there is no interface state (Schottky limit). In most cases, S is between zero and unity, which has been ascribed to mediumly dense interface states. The population of such midgap states changes when the relative position of the Fermi level of metal electrode to the gap is changed by the variation of ⌽m, and their change can induce the change of the interfacial dipole. As the origin of the interface states, many models have been proposed, e.g., metal-induced gap state (MIGS) [52] and defect-induced gap state (DIGS) [76]. Among them, MIGS, which is formed by the penetration of metal wavefunction into semiconductor layer [52], may be also applied to organic system. For polar organic molecules, a net dipole layer can be also formed if the molecules form an oriented ordered layer as shown in Figure 26f. Possible contribution of this factor was pointed out for merocyanine dye molecules on silver halide surfaces [77]. Campbell et al. reported the control of the work function of metal electrode by depositing polar alkane thiol derivatives by self-assembled monolayer technique [78,79]. By using three com-

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Table 3 Summary of the possible origins of the formation of the interfacial dipole layer at organic–metal interfaces Type

Origins

Chemical Charge transfer Chemical bond formation Physical Image effect Push-back effect Dipole orientation

Nonpolar Polar mol. mol. 䡩 䡩 䡩 䡩 ⫻

䡩 䡩 䡩 䡩 䡩

Sign of ⌬ ⫺(donor), ⫹(acceptor) ⫾ ⫺ ⫺ (兩⌬兩 increases with ⌽m) ⫾

pounds of various magnitude and direction of dipole moment, they could change the work function by more than 1 eV, and also the injection current into a polymer layer. Table 3 summarizes the possible origins of the interfacial dipole. Although they seem to explain many aspects of ⌬ for organic layers, further systematic studies will be necessary for a deeper understanding of such phenomena.

V.

FACTORS TO BE CONSIDERED FOR INTERFACES OF REAL DEVICES

In the preceding sections, energy level alignment of organic–metal interfaces has been discussed on the basis of the experimental results for organic film on metal substrate formed under ultrahigh vacuum (UHV) condition. These systems are simple and convenient for discussing the basics of energy level alignment. However, the environment of real interfaces used in organic devices is quite different. First, the existence of ambient atmosphere must modify, more or less, the surface of a metal electrode. Since vacuum vapor deposition is usually performed in high (not ultrahigh) vacuum region (typically ⬃10⫺6 Torr) to make organic devices, the surface of electrode can be contaminated by oxygen, water, or hydrocarbons. Second, the order of the sequence of evaporation of an organic material and a metal is also an important factor. A.

Effect of Atmosphere

During fabrication of organic devices, the surface of a metal electrode tends to be oxidized. When the device is in ambient atmosphere, oxygen in air

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can penetrate to the interface and can oxidize the metal surface. Such oxidation also can be expected in the preparation process of evaporation of metal under a high vacuum condition of typically 10⫺6 Torr. In such vacuum, typically the main component of the residual gas in the vacuum chamber is water, and the number of the water molecules that impinge on the electrode surfaces during 1 s is close to the number of molecules in the monolayer of water on it. For cathode materials such as Al and Mg, adsorbed water molecules also can induce surface oxidation. Figure 29 shows the change of UPS spectra of ZnTPP film on Mg, Al, Ag, and Au by oxygen exposure [7]. Solid circles denote data for samples prepared and measured under UHV conditions. After the measurement in UHV, the sample films was exposed to oxygen (4 Torr, 5 min), reevacuated, and the UPS spectrum was measured again. The latter is shown by open circles. Peaks A and B show energy shift by the exposure, and the direction and the magnitude are dependent on the metal. This result demonstrates the necessity of taking the effect of atmosphere into account. Vacuum level shift is also affected by atmosphere. Figure 30 shows the work function of Ag, Cu, and Au and the observed VL shifts of TPD

Figure 29 UPS spectra in the low binding energy region for ZnTPP films on various metals (Au, Ag, Al, and Mg) in UHV (●) and after exposure to O2 of 4 Torr for 5 min (䡩). Vertical lines indicate the onsets of the HOMO band (peak A).

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Figure 30 Effect of exposure to air to the work function of metal substrate and the vacuum level shift. Top: Work functions of Ag, Cu, and Au substrates before and after exposure to air. Bottom: The vacuum level shift at TPD–metal interfaces using clean and air-exposed metal substrates.

on these metals [80]. The hatched bar denotes the data for clean substrate, and the gray one indicates the data for air-exposed substrate. By exposing to air, the work function of the metal electrode itself changed due to oxidation or contamination, and the magnitude of the VL shift tends to decrease. This decrease is probably because of the reduction of the interaction between organic layer and metal surface by the surface oxidation or contamination. It should be noted that the magnitude of the shift is significant even after such a change. This clearly demonstrates that we have to pay attention to VL shift even for considering interfaces of real organic devices. A similar effect of surface cleanliness can be seen at the interface between HTL and ITO interfaces. For relatively clean ITO surface prepared by UV-ozone treatment, the work function as well as the VL shift were found to be increased from those of uncleaned ITO [16]: In the case of TPD/ ITO interfaces, the work function is typically 4.4 eV and ⌬ is almost zero for dirty ITO surface cleaned only by ultrasonic bath treatment, while the work function is changed to around 4.7 eV and ⌬ reaches about ⫺0.3 eV for UV-ozone–cleaned ITO. To predict the energy level alignment of real interfaces, such effects of atmosphere and contamination should be further investigated.

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Metal on Organic System

All of the data shown in this chapter so far were taken for organic-on-metal interfaces where the interface is formed by depositing organic molecules on metal substrate. However, in real devices, the reverse system, metal-on-organic also is important. In such systems, deposited metal atoms often penetrate and diffuse in organic film, forming diffuse interface. Such diffusion also often enhances the interface reaction between organic and metal. An example of the difference of the structure between metal-on-organic and organic-on-metal interfaces is shown in Figure 31. This figure shows the MAES spectra of (1) 6P film deposited on Au film and (2) Au film deposited on 6P film. In MAES, only the outermost layer of the sample film is probed because of the extreme surface sensitivity of this method resulting from the fact that the excitation source (excited He* atom) cannot penetrate the sample [81]. In the case of 6P on Au, the spectrum of Au substrate is drastically changed by 0.3 nm deposition of 6P. The spectra show no more change by further deposition, and the final spectrum resem-

Figure 31 The metastable atom electron spectroscopy (MAES) spectra of (a) 6P film deposited on Au film and (b) Au film deposited on 6P film.

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bles that of thick 6P film. This indicates that 6P film grows on the Au substrate in nearly layer-by-layer manner. On the other hand, in the case of Au on 6P, the deposition of Au up to 110 nm does not induce spectral change. (The spectra for 0.3 and 1 nm seem to be broadened and quite different. This is probably due to a special charging effect of ‘‘buried Au cluster’’ [82].) The thickness of the initial 6P film was 20 nm. This means that the outmost layer is still 6P, although the amount of deposited Au is 5 times larger than that of the 6P film, suggesting that it is hard to form a top electrode of Au on 6P film. This phenomenon may be due to (1) the strong diffusive nature of Au on 6P film or (2) quick diffusion of 6P molecules on the Au surface. In the case of Mg on 6P, large amount of Mg was necessary to cover 6P film due to diffusion of Mg atoms into 6P film, but finally a top electrode can be formed in contrast to Au on 6P [82]. Similar diffusive nature (often with chemical reaction) was reported for metal-on-polymers [45,46] and metal-on-molecular materials [28,83], but the present case of Au diffusion is extraordinary large. When such complex interface structure occurs, studies of interfacial electronic structure also give complex results. Combined studies by various techniques is necessary to get a reasonable unified picture of the physical structure and electronic structure. C.

Band Bending

In this chapter, we have focused our attention on the examination of VL alignment of MS model. The examination of the second assumption of MS model, i.e., band bending and Fermi level alignment, is also critical for understanding organic devices such as photovoltaic cells. Although the situation is not yet so clear as in the case of VL shift, we will briefly describe our recent efforts to examine the validity of the band bending picture below. In Figure 32, the VL energy of TPD film on various metal substrates (Au, Cu, Ag, Mg, Ca) is plotted as a function of ⌽m of the substrates and the thickness of the TPD film. The measurements were performed with Kelvin probe method (KPM) instead of UPS. In the case of UPS, it is often difficult to distinguish the spectral shift by band bending from artificial shift due to sample charging and/or photovoltaic effect during the measurement. KPM is more suitable for the measurement for thicker film because it is free from charge-up and photovoltaic effect. At all the interfaces in Figure 32, an abrupt shift of the vacuum level was observed within 1 nm thickness, while further deposition of TPD up to 100 nm led to no change of the location of the VL. According to the MS model, the energy levels of semiconductor should show gradual shift with the increase of the thickness leading to the Fermi level alignment between the semiconductor and the metal, where the energy position of the levels of the semiconductor relative to the

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Figure 32 Plot of the vacuum level energy (␧ Fvac) of TPD film on various metal substrates (Au, Cu, Ag, Mg, Ca) as a function of ⌽m of the substrates and thickness of the TPD film.

Fermi level of the metal becomes independent of the work function of the substrate metal. The observed flat behavior of the UV indicates that band bending does not occur or is very small within 100 nm thickness under UHV conditions. These results are probably due to the low concentration of the impurity doped in TPD film formed in UHV. For the four metal substrates, the energy position of the VL at the thicker region is dependent on the substrate metal, indicating that the Fermi level alignment is not achieved. The difference of the VL energy between Au and Ca is about 2 eV. In spite of such a large difference in VL position, no band bending feature was observed. This clearly implies that the picture of MS model is not adequate for TPD–metal interfaces formed in UHV. In contrast to the results in UHV, some band bending feature and Fermi level alignment have been reported on the measurements under atmospheric condition. For example, Iwamoto et al. reported Fermi level alignment by measuring surface potential of polyimide Lagmuir-Blodgett films as a function of number of layers [84]. Their experiments were performed in ambient atmosphere and the sample was heated before the measurements. A doping

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effect, probably due to oxygen in air and activation by heating, may assist the alignment of Fermi level. In order to understand the band bending behavior, further study with controlled doping is necessary. Band bending has been reported also by UPS [37–41], although it is often difficult to distinguish the superficial shift due to band bending from that due to sample charging. For example, the HOMO energy of TPD in Figure 13 showed a gradual shift like band bending. As already mentioned, this shift can be interpreted as final state screening effect of UPS, and not directly related to band bending in the MS model. If we extend the definition of the term band bending, however, this effect may be accepted as effective band bending because the hole injected from anode to organic layer can be affected by this screening effect. Thus it is critical to distinguish between intrinsic band bending leading to Fermi level alignment and extrinsic band bending in the case of photoemission spectroscopy.

VI.

CONCLUSION

In this chapter, our recent studies on the energy level alignment at organic– metal interfaces were discussed. Experimental data clearly demonstrated that the traditional model, which assumes a common vacuum level at organic– metal interfaces, is not valid. The breakdown of the traditional model is not surprising if we consider the basic nature of organic solids and the properties of surface electronic structures of metals. One reason for the long use of the traditional model is the interdisciplinary nature of the field of organic– metal interfaces. By exchanging the knowledge between surface science and materials science of organics, the basic concepts related to organic–metal interfaces could be reexamined to lead to a unified view of such interfaces. Since the vacuum level shift across the interface due to interfacial dipole directly modifies the barrier height for carrier injection, neglect of this shift may lead to an incorrect prediction of the carrier injecting nature of the interfaces. When the traditional model is adopted as a practical way to estimate the interfacial electronic structure, the possible error should be taken into account in discussions of the interfacial phenomena. The trend regarding the vacuum level shift we pointed out can be a guideline for estimation of such errors. In order to precisely predict the energy level alignment at the interfaces in real organic devices, metal-on-organic interfaces should be investigated further. The effect of ambient atmosphere should be also examined. The number of the studies for such subjects is still limited, and the accumulation of data is highly desired. For the former subject, a combination of several

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surface science techniques are necessary for clarifying the complex nature of the interface, such as diffusivity of metal atoms and chemical reaction. For the latter subject, the application of a technique to detect a small amount of the impurity in organic devices is necessary. We hope that serious and continuous efforts in this direction will soon reveal the physics and chemistry at organic–metal interfaces to enable us to predict and tailor the energy level alignment at organic–metal interfaces.

ACKNOWLEDGMENTS This work was performed in collaboration with Profs. K. Yamashita and Y. Harima, Drs. S. Narioka, K. Sugiyama, D. Yoshimura, E. Ito, H. Oji, S. Hasegawa, T. Miyazaki, and T. Miyamae, and Messrs. T. Imai, N. Hayashi, and T. Yokoyama. We gratefully acknowledge these individuals. This work was supported in part by Grant-in-Aids for Scientific Research from the Ministry of Education, Science, Sports and Culture of Japan (nos. CE072004, 10440205, 10146102, and 12304041). This work was performed as a Joint Studies Program of the UVSOR facility of Institute for Molecular Science (No. U-606).

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11 Organic Molecular Interfaces Investigations of Electronic Structure, Chemistry, and Carrier Injection Properties Chongfei Shen and Antoine Kahn Princeton University, Princeton, New Jersey

Ian Hill Sarnoff Corporation, Princeton, New Jersey

I.

INTRODUCTION

The contributions included in this book underscore the importance of the structural, chemical, and electronic properties of interfaces in organic devices. The progress made in the past decade in developing new multicolor organic light-emitting diodes (OLEDs) [1–7], optically and electrically pumped lasers [8,9], thin-film transistors [10–13], and solar cells [14,15] is advancing the field of organic polymers and small molecule materials at an astounding pace. The vast majority of these devices have a thin-film architecture, which comprises multiple organic–inorganic and organic–organic interfaces [1–7]. Because of the very nature of thin-film devices, these interfaces are generally located within a few molecular planes from active regions. Their electronic and chemical properties determine the characteristics of charge carrier injection into and transport through the device. OLEDs require that injection of electrons and holes be balanced in order to maximize device quantum efficiency (Figure 1). The injection of these charge carriers must be efficient and stable under operation. Metal–organic contacts undergo complex and spatially extended chemical interactions, which can dominate the electrical properties of interfaces. Organic–organic heterojunctions in OLEDs control the transport of carriers between layers 351

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Figure 1 Schematic energy diagram of a heterojunction OLED under bias. Electrons (holes) are injected from a low (high) work function cathode (anode) into the electron (hole) transport layer. The heterojunction is designed here to favor exciton formation and radiative recombination in the electron transport layer.

and the region of the device where recombination takes place. Understanding the mechanisms that control the interface electronic structure, including the impact of interface chemistry on carrier injection, is therefore of paramount importance for a more directed engineering of organic devices. Considerable progress is being made in the area of device efficiency and reliability, although much of it has so far been achieved by empirical means. The drive to develop technologically viable and marketable devices has often left behind the science required to fully understand the mechanisms at play. Over the past decade, however, surface science has begun to play a key role in developing a deeper fundamental understanding of organic interfaces [16], particularly with respect to the way electronic structure and chemistry relate to charge carrier injection. As a consequence, the level of theoretical and experimental understanding of organic interfaces has improved to the point where some predictability is possible. However, this

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understanding is still limited in comparison to what has been achieved for interfaces of inorganic semiconductors. The molecular nature of organic films makes the investigation of their interface properties both experimentally more flexible and theoretically more challenging than for inorganic semiconductors. The flexibility derives from the fact that organic molecular films with good electronic and optical properties can be formed on a variety of substrates under a broad range of deposition conditions. Inorganic semiconductors must be grown epitaxially or must be passivated to reduce the density of electronically active defects associated with unsaturated covalent bonds (e.g., hydrogenated amorphous silicon). Lattice matching and temperature compatibility put strict constraints on the choice of substrate and growth conditions for epitaxial layers. Organic molecular films, on the other hand, are not bound by the same restrictions. Molecular solids are characterized by strong covalent intramolecular bonds but weak van der Waals (vdW) intermolecular bonds. The overlap between wavefunctions centered on neighboring molecules is small, leading to weak intermolecular electronic coupling. Therefore, the electronic and optical properties of organic molecular films derive in first approximation from those of the closed-shell molecular moieties. They are much less affected by structural defects or lack of long-range molecular order than in inorganic solids, adding considerable flexibility in device design and processing. A large variety of molecules can be used to realize specific optical or transport applications. Electronic and optical grade films can be deposited from liquid phase (polymers) or by vacuum evaporation (small molecules) on a variety of crystalline or amorphous, organic or inorganic, rigid or flexible substrates without regard for structural match. Multiple layers can be superposed and different molecular materials can be mixed to engineer complex structures and introduce new functionalities while keeping the properties of the constituents essentially unchanged. This must be contrasted to the crystalline perfection required for most optical and optoelectronic devices based on inorganic semiconductors. Of course, the flexibility afforded by molecular films applies to surface and interface studies of these materials as well. For example, the sequence of deposition of various materials can be reversed, e.g., organic-on-metal vs. metal-on-organic, without necessarily deteriorating the intrinsic properties of the materials, thereby increasing the range of meaningful interfaces and phenomena that can be investigated. Yet the flexibility enjoyed at the experimental and device processing level comes at a price of considerable complexity at the theoretical level [17,18]. This complexity is inherent to molecular materials and is due to (1) the nature of the constituents of these molecular films, i.e., the large neutral molecular moieties; (2) the relaxation and polarization phenomena that take

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EXPERIMENTAL APPROACH

The experiments described in this review, including the deposition of all organic and metallic thin films, were performed in ultrahigh vacuum (UHV) at pressures of 10⫺9 –10⫺10 Torr or below. These conditions are sometimes regarded as remote from the conditions under which organic thin films are being and will be processed for commercial applications. However, UHV processing permits investigations of fundamental interface properties designed to extract information related to the intrinsic nature of the materials under investigation rather than to extrinsic factors such as contamination, known to affect organics and metals used in devices. A specific example of the importance of looking at ‘‘clean’’ interfaces in order to understand the physics of organic–metal interfaces is discussed at the end of Section III.C in the context of the Mg–Alq3 interface [19]. This chapter also focuses exclusively on vacuum-evaporated thin films of ␲-conjugated molecules that can be easily sublimated and condensed in UHV. These organic materials are purified via multiple cycles of sublimation prior to being loaded in the UHV deposition chamber. The chemical structures of the molecules discussed in this review are given in Figure 2. The principle group of investigation techniques used in this work is photoemission spectroscopy. Ultraviolet photoemission spectroscopy (UPS) is used to investigate the valence states of the organic and inorganic films, in particular the evolution of the electronic structure around the highest occupied molecular orbital (HOMO) upon formation of interfaces. Photons are produced by synchrotron radiation or by a helium discharge lamp that provides light at 21.22 eV (HeI) and 40.8 eV (HeII). A typical HeI valence spectrum of an organic thin film, measured here from PTCBI (3,4,9,10perylenetetracarboxylic bisimidazole) [20] is shown in Figure 3. The spectrum consists of a broad, low-energy, secondary electron peak, which results from inelastically scattered electrons, and several well-defined features related to the density of states of the ␲ system and ␴ bonds of the molecular solid. The highest kinetic energy (lowest binding energy) feature corresponds to the HOMO. The energy of the HOMO is generally referred to the Fermi level, measured separately on a metallic electrode. The low-energy edge of the secondary electron peak, i.e., the photoemission cutoff, corresponds to the vacuum level (Evac) of the molecular film. The ionization energy (IE) is defined as the energy difference between Evac and the leading edge of the HOMO peak. It is given by IE = h␯ ⫺ W where h␯ is the photon energy and W is the total width of the valence spectrum measured from the leading edge of the HOMO peak to the pho-

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toemission cutoff (Figure 3). The shift of the photoemission cutoff upon formation of an interface, corrected for eventual ‘‘band bending’’–induced shift of the molecular levels, leads to the determination of the dipole barriers (see next section). The typical energy resolution achieved in our UPS system is 150 meV. The other photoemission technique extensively used in our work on organic interfaces is X-ray photoemission spectroscopy (XPS). XPS probes the core level binding energy of atoms at and near the interface and provides information on chemistry. We use the K␣ line from an aluminum anode (1486 eV) and the M␨ line from a zirconium anode (151 eV). The low kinetic energy of electrons photoexcited by the latter provides for a very surface-sensitive chemical probe. Both photon lines produce spectra with energy resolution of the order of 0.5 eV. An important aspect of our work concerns the study of empty states in the energy region of the lowest unoccupied molecular orbital (LUMO) in organic films and at interfaces via inverse photoemission spectroscopy (IPES). IPES is performed in the isochromat mode, by varying the electron energy while keeping the energy of the detected photons fixed. The system, which has been described in detail in previous publications [21], has a resolution of 0.5 eV. One of the main concerns with IPES is whether electron beam–induced deterioration and/or charging occurs in the organic film. Typical experiments reported in this chapter were performed using small incident electron current densities in the approximate range of 10⫺7 –10⫺4 A/ cm2. Decreasing the current at the expense of the count rate is generally found to result in superior data. The analysis of our IPES data involves standard nonlinear least-squares fitting techniques and background subtraction. The interested reader is referred to previous publications for more details [21–23]. Finally, the relationship between electronic structure, chemistry, and charge carrier injection characteristics at organic interfaces under investigation is established via current–voltage (I-V) measurements on small devices grown and measured in UHV. The complete in situ process eliminates

< Figure 2 Chemical structures of molecules investigated: Alq3, tris(8-hydroxyquinoline)aluminum; ␣-NPD, N,N⬘-diphenyl-N,N⬘-bis(1-naphthyl)-1,1⬘biphenyl-4,4⬙ diamine; CuPc, copper phthalocyanine; F16CuPc, hexadecafluoro copper phthalocyanine; PTCDA, 3,4,9,10-perylenetetracarboxylic dianhydride; PTCBI, 3,4,9,10perylenetetracarboxylic bisbenzimidazole; CBP, 4,4⬘-N,N⬘-dicarbazolylbiphenyl; BCP, bathocuproine; ␣-6T, ␣-sexithiophene.

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Figure 3 He I spectrum of pristine PTCBI showing the valence states with the HOMO and the photoemission cutoff. The Fermi level EF is measured separately on a Au electrode. (From Ref. 20, with permission.) Inset: Chemical structure of the PTCBI molecule.

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eventual effects of cathode oxidation and deterioration of the organic layer in air on the transport characteristics.

III.

ELECTRON AND HOLE BARRIERS AT ORGANIC INTERFACES

A.

Organic-on-Metal Interfaces: Molecular Level Alignment and Interface Dipoles

Understanding the mechanisms that control the energetics of interfaces, in particular the magnitude of energy barriers between the Fermi level (EF) and the HOMO (LUMO) levels across interfaces, has been one of the main challenges to surface and interface studies of organic thin films over the past few years. These energy barriers determine the injection of carriers across interfaces. In that respect, it is useful to briefly reconsider the situation encountered with inorganic semiconductors. Inorganic semiconductor interfaces are controlled by intrinsic and extrinsic electronic gap states. Intrinsic states, like metal-induced gap states at metal–semiconductor interfaces [24], are the natural consequence of the loss of three-dimensional periodicity on the band structure of the two materials and of the boundary conditions between them, such as interfacial bonds and atomic structure. On the other hand, extrinsic states result from structural defects or chemistry-induced defects that are formed at interfaces [25,26]. Both types of states affect the electronic structure of the interface, as seen, for example, in the complex dependence of the Schottky barrier height on the metal work function at inorganic semiconductor interfaces [24]. The relative positions of energy levels across inorganic interfaces are controlled by the occupation of these interface states, which gives rise to charge exchange and dipole barriers, as depicted in Figure 4b for a metal–semiconductor interface. Because of the closed-shell nature of molecules and their weak vdW intermolecular bonding, organic molecular semiconductor interfaces had been expected to form weakly interacting boundaries with weak interface bonds and no electronically active defects, leading to interfaces without gap states. Under such conditions, the relative position of molecular levels across the interface would be obtained by simple alignment of the vacuum levels of the two solids, as depicted in Figure 4a. This situation, which is the equivalent of the Schottky-Mott limit [27] in metal–semiconductor interfaces and of the Anderson rule [28] for semiconductor heterojunctions, was widely assumed to be correct in the literature on organic devices. Recent investigations have now firmly demonstrated that this is rarely the case at metal–organic interfaces [29–38]. Interface dipole barriers with magnitudes greater than 1.0 eV have been seen to introduce discontinuities in the vac-

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Figure 4 Schematic energy diagram of a metal–organic semiconductor interface (a) without and (b) with a dipole barrier ⌬. ␾Be and ␾Bh are the electron and hole barriers, and Evac(O) and Evac(M) are the organic and metal vacuum levels, respectively.

uum level across interfaces and cause substantial departure from the simple Schottky-Mott limit (Figure 4b). Figure 5 shows UPS spectra taken upon formation of the interface between ␣-NPD [N,N⬘-diphenyl-N,N⬘-bis(1naphthyl)-1,1⬘biphenyl-4,4⬙ diamine] and Au, which clearly demonstrates this point. The schematic energy diagram of the interface (bottom panel) shows (1) the hole injection barrier measured from the energy difference between the HOMO of the organic film at the interface and the EF of the substrate in the UPS spectrum (top right panel); (2) the 1.2-eV interface dipole barrier deduced from the abrupt shift of the onset of photoemission (top left panel) upon deposition of the first layer of organic molecules. Neglecting the dipole in this case would introduce a rather large error in the evaluation of the electron or hole injection barriers.

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Figure 5 Soft X-ray photoemission spectra recorded as a function of deposition of ␣-NPD on Au. (a) Valence states, showing the position of the leading edge of the HOMO with respect to the Au Fermi level. (b) Abrupt shift of the photoemission onset indicating the presence of an interface dipole. (c) Schematic interface energy diagram deduced from the photoemission measurements.

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Figure 6 summarizes (1) the position of EF and (2) the dipole barriers vs. metal work function measured in our laboratory via photoemission spectroscopy for a series of organic-on-metal interfaces. A positive dipole in Figure 6b indicates a vacuum level drop from the metal into the organic film (identical to the situation depicted in Figure 4), suggestive of an electron transfer from the organic film to the metal. Alternatively, a negative dipole corresponds to a rise in the vacuum level, indicative of an electron transfer

Figure 6 (a) Position of EF vs. HOMO (and LUMO) as a function of metal work function for six molecular materials. The data points are obtained via PES for organic-on-metal interfaces. The dashed lines correspond to the Schottky-Mott limit and the vertical lines give the magnitude of the measured interface dipole barriers. (b) Sign and dependence of dipoles vs. metal work function.

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from the metal to the organic. The dipole barriers also appear in Figure 6a as the vertical segments around each measured point. The dashed lines in panel a represent the position that EF would occupy if these interfaces did follow the Schottky-Mott limit. Also indicated on each plot is the S parameter, which gives a measure of how much the energy barrier varies as a function of metal work function (S = 1 for the Schottky-Mott limit). Two remarks should be made here. First, the interfaces represented in Figure 6 were formed by deposition of the organic film on the metal surface, i.e., organic-on-metal (O/M). This deposition sequence generally leads to a more abrupt interface than the reverse sequence, i.e., metal-on-organic (M/ O). The O/M interface allows a more straightforward investigation of the fundamental mechanisms that operate at these interfaces, in particular with respect to molecular level alignment. The reverse deposition sequence is exceedingly important as well, especially in the context of devices, but often leads to diffused interfaces that are more difficult to model. More will be said on this subject in the next section. The second remark concerns the

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energy gap of each organic film. The energy difference between the two horizontal lines marking the HOMO and LUMO levels was arbitrarily set equal to the optical absorption gap of the material and not to the more appropriate ‘‘transport’’ gap (discussed in detail in Section IV), which is to be used when discussing energy barriers for charge carrier injection. However, this approximation does not affect the following discussion. The obvious conclusion from these data is that the rule of vacuum level alignment breaks down for most of these interfaces. This conclusion was reached early on by Seki and coworkers [29–31] and was subsequently confirmed in a number of experiments by our group [19,20,33–36] and others [37,38]. It raises two key questions: (1) Why is the Schottky-Mott rule not respected, in particular in seemingly simple cases where the work function of the metal falls deep between the electron affinity and ionization energy of the semiconductor? (2) What is(are) the mechanism(s) that control molecular level alignment and formation of interface dipole barriers? In cases where the work function of the metal is smaller than the electron affinity (EA) of the organic solid, like PTCDA on Mg, In, or Sn, PTCBI on Mg and Ag, and F16CuPc on Mg and Al, a simple vacuum level alignment would lead to an interface where EF is very close to or above the LUMO of the organic, a clearly untenable situation. Electron transfer from the metal to the organic solid occurs in order to raise the potential of the organic molecular solid and eventually stop the net flow of electrons. In equilibrium, EF is stabilized below the LUMO at the interface. The electron transfer from the metal to the interface molecules gives rise to a negative dipole (Figure 6b) and to occupied states in the gap of the organic semiconductor. All but one of the seven interfaces mentioned above are reactive, and the gap states originate from chemistry as well as charge transfer, as will be discussed below in the case of Alq3 /Mg. However, PTCBI/Ag displays no evidence of chemical reaction. Bauschlicher et al. [39] have studied the interaction of transition metal ions with benzene and found that the Ag⫹ ion forms a weak ␲ complex. The optimized structure consists of a virtually undisturbed benzene ring with the Ag⫹ ion relatively far from the C atoms. The aromaticity of the benzene ring is not disturbed, consistent with our observations of bulk-like PTCBI molecular features at the lowest coverages and with the fact that molecular level separations remain constant ⫾0.1 eV as a function of film thickness (Figure 7). In this case, the gap states cor-

< Figure 7 (a) He I spectra of interfaces formed by incremental deposition of PTCBI on Ag. (b) Top of the valence density of states at the PTCBI–Ag interface. Vertical lines are intended as guides to the eye. (From Ref. 20, with permission.)

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respond to an electron transfer from the metal to the LUMO level of the molecule, which induces a reconformation of the molecule. The result is referred to as polaron (singly occupied) or bipolaron (doubly occupied). At the PTCBI–Ag interface, the two levels that peak above the HOMO at about 0.6 eV and 1.6 eV below EF are tentatively attributed to bipolaron-like and polaron-like states, respectively [20]. The second type of interface, for which the metal work function ranges between EA and IE of the organic solid, is more complex and diverse. At these interfaces, the molecular level alignment depends on various bulk and interface characteristics, such as the possibility of moving EF across the gap of the organic material, the appearance of interface electronic states, and the formation of chemical bonds across the interface. A relatively simple situation emerges when gap states or other mechanisms inherent to the organic material restrict the position of EF in the gap in the bulk and at the surface of the organic film. In order to reach thermodynamic equilibrium, the interface develops a dipole barrier equal to the difference between the metal work function and the organic work function. This is precisely the situation encountered with PTCDA and PTCBI on Au. If the Schottky-Mott rule was valid, the 5.2-eV Au work function should lead to an EF position deeper in the gap at the interface than experimentally observed. However, UPS measurements on PTCDA [40–43] and PTCBI [20] films of various thicknesses deposited on a variety of metallic and semiconducting substrates consistently show that the Fermi level is ‘‘pinned’’ at a fixed position, roughly 1.5 eV and 1.8 eV above the HOMO for PTCBI and PTCDA films, respectively. These organic molecular materials are nominally undoped and EF should occupy a much deeper position in the gap. The reason for this ‘‘pinning’’ is unknown, although a small density of impurities or molecular defects could play a role. The consequence is that EF is consistently pinned high in the gap of these two materials, leading to a positive dipole at interfaces with Au. Besides PTCDA and PTCBI on Au, the majority of interfaces presented in Figure 6 exhibit positive dipoles that are perhaps more difficult to justify. In particular, interfaces like CBP and Alq3 on Ag and Au exhibit dipole barriers that are nearly independent of the metal work function, as shown in Figure 6b. Image-force effects, whereby the charge density of the interface molecule is polarized by the presence of the metal, and redistribution of the metal charge density by the adsorbed molecules, leading to a decrease in the metal work function and an effective drop of the vacuum level upon adsorption of the molecules, have been invoked to explain the sign and fixed magnitude of these barriers. These mechanisms have been recently reviewed by Ishii et al. [32]. Another very important mechanism, which is clearly orthogonal to the notion of weakly interacting interfaces and, as such, is reminiscent of mech-

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anisms active at inorganic semiconductor interfaces [24–26], must be brought up to explain the behavior of interfaces like Alq3-on-Mg or Alq3Al. It is based on chemistry-induced formation of interface defects, which pin the Fermi level and results in the formation of a dipole. Our group reported 3 years ago the formation of gap states upon deposition of Mg on Alq3 [44]. We recently reported the formation of similar gap states for the reverse deposition sequence, i.e., Alq3 on Mg and Al [45] (Figure 8a). These states result from the exothermic formation of an organometallic complex in which the metal atom attaches to the pyridil side of the quniolate ligand of the molecule [46]. The reaction is favorable due to the tendency of Mg and Al to form covalent bonds with carbon. More will be said on this reaction in the next section dedicated to metal-on-organic interfaces. In the organic-on-metal case, metal atoms do not diffuse into the organic film and the reaction is limited to the first molecular layer. Yet the reaction product and resulting electronic states are similar to those obtained for Mg or Al on Alq3. These chemistry-induced states are responsible for establishing the interface energetics, i.e., the position of the HOMO relative to EF, early in the deposition sequence. They are filled and force EF in the upper part of the gap. The positive dipole barrier, i.e., 0.8 eV and 0.9 eV for Alq3 on Mg and Al, respectively, results from an electron transfer from the interface organometallic complex (i.e., the gap states) to the metal (Figure 8b). A ˚ metal–molecule distance and a dielectric simple calculation based on a 10-A constant of about 3 shows that about 10% of the gap states need to be ionized to sustain the observed dipole. All other states remain occupied and thus detectable via UPS (Figure 8a). It is likely that such defect- or chemistry-induced states play a predominant role in determining the molecular level alignment at reactive interfaces. The next section will emphasize the role of these states in inducing band bending and establishing barriers at metal-on-organic interfaces. B.

Metal-on-Organic (M/O) Interfaces: Chemistry, Interdiffusion, and Barriers

Thin-film organic devices like OLEDs comprise several organic films sandwiched between an anode and a cathode. Future devices will likely include stacked structures in which multiple O/M and M/O interfaces coexist. M/O interfaces are typically formed by vacuum evaporation or sputtering of a metal on the organic film. In general, M/O interfaces are structurally and chemically more complex than their O/M counterparts. Whereas O/M interfaces are abrupt, with the metal–molecule interaction limited to the first molecular layer, isolated metal atoms evaporated on an organic surface are chemically more active and capable of diffusing more deeply into the mo-

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Figure 9 Soft X-ray photoemission spectra as a function of In deposition on PTCDA: (a) In(4d) core level showing the oxidized state on In. (b) C(1s) core level showing the loss of the carbonyl C peak (B). (c) HOMO and gap region, showing the In-induced gap states. (From Ref. 48, with permission.) The pristine PTCDA molecule and the molecule with In attached to the end groups are represented.

portant M/O interfaces, i.e., Mg and Al on Alq3. Alq3 is one of the most extensively used electron transport and emissive layers for OLEDs. Mg and, to a lesser degree, Al are low work function metals used for electron-injecting cathodes on Alq3 films. Several groups have investigated these interfaces [51–55] and have reported various interactions ranging from simple metal-to-molecule charge exchange to metal–O coordination at the center of the molecule. Our work, summarized below, shows how complex the chemical interaction resulting in the formation of an organometallic complex can be. Also discussed are the results from carrier injection experiments

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designed to investigate the respective behaviors of Mg–Alq3 and Alq3 –Mg interfaces. The tendency of Mg and Al to form covalent metal–C bonds [56] and the affinity of these metal atoms for O lead to considerable interface chemical interaction upon deposition on Alq3. The reaction results in a complex set of chemical shifts in the XPS data measured for Mg deposited on a 100˚ Alq3 film (Figure 11). Mg is ‘‘oxidized’’ during the initial stages of depA osition, resulting in a dominant high binding energy (BE) component of Mg(2p). The metallic Mg component signaling accumulation of the metal on the organic film becomes dominant only following the deposition of 32 ˚ and 64 A ˚ . The C(1s) undergoes a 0.6-eV shift to higher BE with the first A Mg deposition, and its position and shape remain basically unchanged thereafter. The same shift is observed on the N(1s) and O(1s) core levels. It represents a rigid shift analogous to band bending induced by the deposition of a metal on an inorganic semiconductor surface. However, unlike the C(1s)

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Figure 10 Combined UPS/IPES spectra of PTCDA as a function of In deposition. Particularly notable are the position of EF near the LUMO on pristine PTCDA; the appearance of gap states upon In deposition; and the concomitant decrease in LUMO signal.

peak, the N(1s) and O(1s) peaks broaden considerably. The N(1s) core level develops a low BE component at 399.1 eV (⬃1.8 eV chemical shift), which corresponds to an excess electronic charge on and around N. On the other hand, the O(1s) core level develops a high BE component. The 1.4-eV chemical shift corresponds to a net withdrawal of negative charge from oxygen. Finally, the Al(2p) undergoes a 1-eV shift to high BE, which corresponds

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Figure 11 Mg(2p), C(1s), N(1s), O(1s), and Al(2p) core levels measured as a function of Mg deposition on Alq3. (From Ref. 45, with permission.)

to a net chemical shift of 0.4 eV to higher BE over and beyond the ubiquitous 0.6-eV band-bending shift. The peak areas of the O(1s), C(1s), N(1s), ˚, and Al(2p) remain basically unchanged up to a nominal coverage of 32 A confirming that Mg does not initially accumulate on the surface of the organic film. The increase in the total area of the Mg(2p) peak demonstrates that the overall metal atom concentration does increase, and thus indicates that the metal atoms diffuse into the film. Eventually, reaction sites are depleted, Mg begins to accumulate at the surface, and the intensity of the various Alq3 signals is attenuated. The key result in terms of interface chemical structure is the opposite sign of the O(1s) and Al(2p) vs. N(1s) chemical shifts, which is incompatible with a simple charge donation from the metal atom to the molecule. It signals the formation of an organometallic complex. In that regard, the results are quite different than for Ca, K, or Li on Alq3. These latter elements form considerably weaker bonds with C [57] and undergo only a charge exchange with the molecule, leading to a metal cation–molecular anion complex, consistent with XPS chemical shifts observed experimentally for the Ca-Alq3 system [51]. In the Mg/Alq3 case, Mg forms a covalent bond

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Figure 11

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with C at the pyridyl ring of the quinolate ligand. Mg donates a charge to the Alq3 LUMO, which is maximized on the pyridyl ring and at N [22]. The result is a high-BE component of Mg(2p) (oxidized species) and a low-BE component of N(1s). The Mg atom further coordinates as an electrophile to a quinolate O of the same or a neighboring molecule, and in the process pulls charge away from O. This results in high-BE components on O(1s) and on the 2p level of the central Al atom of the molecule to which O is directly bound. The model of the Mg-Alq3 complex is presented in Figure 12. The deposition of Al on Alq3 leads to a very similar set of chemical shifts, consistent with the similar tendency of Al to form covalent Al-C bonds. The chemical model developed above therefore applies to the latter interface as well [45]. The UPS data provide information on the electronic structure of the Mg–Alq3 interface. The initial deposition of Mg attenuates the Alq3 features, consistent with the chemistry-induced ‘‘smearing’’ of the top molecular layers of the organic film (Figure 13). All the molecular levels shift by 0.6 eV toward higher BE and remain at the same energy thereafter, in agreement with the 0.6-eV band bending shift deduced from the XPS data. The origin of this shift becomes clear when considering the gap states induced by the Mg deposition at about 1.9 eV above the original HOMO. These filled states, which correspond to the organometallic complex formed in the top layers of the film, push EF upward in the gap by 0.6 eV with respect to its original position at 2.1 eV above the leading edge of the HOMO. The shift corresponds to the band bending depicted in Figure 13. The spatial extend of the band bending is not known in the present case, although the width of the reacted interface region is believed to be limited to 5–10 molecular planes. In relation to the discussion given in the previous section, it is interesting to note that the final positions of the Alq3 HOMO with respect to EF

Figure 12 Structural representation of the organometallic complex resulting from Mg-Alq3 reaction. (From Ref. 46, with permission.)

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are nearly identical at the O/M and M/O interfaces, suggesting that the electron and hole injection barriers are identical ⫾0.1 eV as well. Given this identity, it is important to understand whether the presence of a wider reacted interface in the M/O case plays a role in the injection process. It is generally believed that a depth distribution of the type of gap states observed at these interfaces could facilitate current injection via a ‘‘stepping-stone’’ effect [58], leading to a higher injection current from the metal-on-top interface. Such an asymmetry was believed to occur in nominally symmetrical Mg:Ag/Alq3/ Mg:Ag structures (Figure 14a), based on I-V measurements showing that the current injected from the metal-on-top electrode was two to three orders of magnitude larger at identical bias than the current injected from the bottom electrode [59] (Figure 14b). However, this asymmetry was later shown to be due to contamination effects. Mg/Alq3/Mg and Al/Alq3/Al sandwich structures fabricated and tested in UHV (10⫺9 –10⫺10 Torr) showed nearly perfect symmetry of the currents injected from the top or bottom electrodes [19] (Figure 14c, d), consistent with the identical barriers measured via photoemission and reported above. The previously reported asymmetry was shown via XPS to be due to a blocking oxide layer accumulated on the reactive Mg electrode prior to deposition of Alq3 in a 10⫺6 Torr vacuum. The symmetry exhibited by these interfaces does not invalidate the notion that stepping-stones, provided by chemical reaction or other means, can help the injection process. The lack of effect in the two examples discussed above is merely due to the fact that the states that could play a role in the extended M/O reacted region are filled states and thus do not participate in the electron injection process. The situation is clearly different if the interface states are relevant to conduction. For example, we have demonstrated a remarkable increase in hole injection from indium-tin-oxide ˚ ) layer of a compound (i.e., CuPc) (ITO) into ␣-NPD using an ultrathin (30 A exhibiting a hole conduction level at intermediate energy between the EF of ITO and the HOMO of ␣-NPD [60] (Figure 15). The stepping-stone effect is clearly active in such a case. As a final remark, Figure 14e shows that electron injection into Alq3 obtained from both top and bottom Mg/Ag electrodes is superior to that obtained from pure Mg electrodes. Although recognized early on by OLED practitioners [1] and believed to be due to improvement in morphology and stability of Mg by Ag, this behavior has not yet been fully explained.

< Figure 13 UPS valence spectra measured as a function of increasing Mg coverage on Alq3 and illustration of corresponding Mg-induced shift of molecular levels, or ‘‘band bending,’’ at the surface of Alq3.

Figure 14 (a) Mg/Alq3/Mg structure. Forward (electrons injected from top electrode) and reverse (electrons injected from bottom electrode) current–voltage characteristics of a (b) Mg:Ag/Alq3/Mg:Ag device grown at 10⫺5 –10⫺6 Torr. (From Ref. 59, with permission.) (c) Mg/Alq3/Mg device; (d) Al/Alq3/Al device; (e) Mg:Ag/ Alq3/Mg:Ag device. The latter three were grown at 10⫺9 Torr. (From Ref. 45, with permission.)

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Figure 15 (a) Comparison between current–voltage characteristics of ITO/␣-NPD and ITO/CuPc/␣-NPD devices showing the interlayer-induced increase in hole injection. (b) Molecular level alignment diagram constructed using HOMO offsets measured via UPS. (From Ref. 60, with permission.)

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Organic-Organic Heterojunctions: Molecular Level Alignment

The electronic structure of organic–organic (OO) heterojunctions is exceedingly important for multilayer organic devices. High-efficiency OLEDs made of vacuum-deposited molecular films often comprise an electron transport layer (ETL), a hole transport layer (HTL), and an emissive layer. Energetically efficient transport of electrons and holes across boundaries between these organic layers presupposes a favorable interface arrangement of molecular levels. Hole- or exciton-blocking layers are often inserted near one of the contacts of the device to restrict exciton formation and recombination to one layer of the device or to prevent nonradiative exciton recombination at the metal interfaces. The proper function of these various blocking layers in the device depends on the energy barriers they introduce in the path of the carriers. Understanding molecular level alignment at OO interfaces is therefore highly relevant to modeling device behavior. Figure 16 shows the evolution of the UPS spectra of the HOMO energy region during the formation of the ␣-NPD-on-CuPc heterojunction [60,61].

Figure 16 He I (21.22 eV) UPS spectra showing the evolution of the top of the valence spectra for the interface formed by depositing ␣-NPD on CuPc. (From Ref. 61, with permission.)

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The two HOMOs coexist at intermediate coverage, with the intensity of the CuPc HOMO diminishing as the ␣-NPD HOMO intensity increases at higher coverage. The CuPc base layer is not completely suppressed even at a nom˚ , indicating some degree of three-dimensional inal ␣-NPD thickness of 48 A island growth. This allows a direct and accurate measurement of the HOMOHOMO offset up to an ␣-NPD coverage of several molecular layers, sufficient to ensure that the electronic structure of the overlayer film is fully developed. A summary of the alignment of organic levels at 12 organic–organic heterointerfaces is presented in Figure 17. Within the experimental error of 0.1 eV, the molecular level alignment follows the vacuum alignment rule in a majority of cases [61–63]. This is not surprising, considering the closedshell nature of the molecules that comprise both constituents of the heterointerface. Charge exchange, chemical bonds, and chemistry-induced defects of the type encountered at metal–organic interfaces are not expected to occur at OO heterointerfaces. However, some of these OO interfaces exhibit larger dipoles: CuPc/PTCDA (0.4 eV) [36], PTCDA/Alq3 (⫺0.5 eV) [42–43], ␣NPD/Alq3 (⫺0.25 eV) [43], and BCP/PTCBI (0.4 eV) [61] where the dipole is defined as positive when the vacuum level of the second material is at a higher energy than the first. Ishii et al. [64] recognized that dipoles occur at heterointerfaces between materials of greatly differing ionization energies and electron affinities, which are referred to as donor–acceptor interfaces. In such cases, a partial charge is expected to be transferred from the low ionization energy constituent to the high electron affinity molecule. This argument, which is also at the basis of the work on molecular doping [65,66], may be used to explain the interface dipoles measured between the high electron affinity molecule, PTCDA, and the lower ionization energy molecules, CuPc and Alq3. However, the same argument would predict that the PTCDA/␣-NPD interface should also contain a strong dipole, given that ␣-NPD has a lower ionization energy and electron affinity as compared to Alq3. This prediction contradicts the observation of a negligible dipole at this interface. Three of the materials leading to interfaces with significant dipoles also exhibit limited ranges of EF movement at heterointerfaces (PTCDA, PTCBI, BCP). As noted in the previous sections, EF may be ‘‘pinned’’ due to defects or impurities, and the dipoles are the result of (de)populating a high density of impurity states near the interfaces. Yet the 0.5-eV dipole observed at the PTCDA–Alq3 interface cannot be explained

> Figure 17 Alignment of molecular levels at organic–organic heterointerfaces. (From Ref. 61, with permission.)

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on this basis, given the above observation that the Fermi level is relatively free to move within the Alq3 gap (Figure 6). Considering the mechanisms suggested above and the counterexamples, one must conclude at this point that neither of the suggested mechanisms for dipole formation at organic– organic heterointerfaces is sufficient to explain all of the observed behaviors. Other OO interfaces, such as PTCDA/ZnPc (zinc phthalocyanine) [67], PTCDA/ClInPc (chloroindium phthalocyanine) [67], CuPc/NPB (naphthylsubstituted benzidine derivative) [68], Alq3/NPB [69], and Alq3/TPD [70] have been investigated and found to exhibit dipoles ranging between 0 and 0.25 eV. The overwhelming conclusion is therefore that interface dipole barriers at OO heterojunctions are, if significant at all, smaller than at OM interfaces. The assumption of vacuum level alignment, which is widely used at present in the literature for device modeling, is not expected to cause the same level of error than at OM interfaces. The situation is somewhat different at organic–inorganic semiconductor interfaces. Figure 18 shows energy level alignment measured for ␣-NPD and CuPc on InP(110) and for PTCDA on GaAs(100). The dipole barrier is substantial (magnitude ⱖ 0.4 eV) at each interface [71]. Several mechanisms contribute to this apparent difference with OO interfaces. First, the InP(110)

Figure 18 Energy level diagram of the ␣-NPD/InP(110), CuPc/InP(110), and PTCDA/GaAs(100) interfaces. VBM and CBM denote the valence band maximum and conduction band minimum, respectively, in the inorganic semiconductor.

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and GaAs(100) surfaces are unknown to exhibit rather strong near-mid-gap EF pinning following deposition of almost any ad-atom or ad-molecule. The alignment of EF across these interfaces is therefore more constrained than at OO interfaces. Second, the GaAs and InP band gaps are smaller than those of all three organics. The ionization energy of GaAs (5.56 eV) is significantly smaller than that of PTCDA, placing the inorganic valence band top relatively close to the PTCDA LUMO, and the electron affinity of InP (4.45 eV) is smaller but close to the ionization energy of CuPc, leading to possible mixing of states and charge transfer. Indeed, the dipole barriers measured at all three interfaces tend to realign the energy levels to decrease mixing and charge transfer. Third, dangling bonds at the inorganic semiconductor surfaces lead to stronger chemical bonding with the molecules than exist across OO interfaces, and thus allow more readily charge transfer and formation of dipoles. In that sense, organic–inorganic semiconductor interfaces are closer to MO interfaces than to OO interfaces.

IV.

TRANSPORT GAP AND EXCITON BINDING ENERGY IN MOLECULAR FILMS

This final section focuses on the difficult task of reconciling energy levels and gaps measured via PES and IPES, which we will argue to be relevant to carrier injection and transport in molecular solids, with energy levels and gaps measured via optical absorption and other means. This task, which is still largely ignored in today’s literature on devices, is key to the development of realistic interface transport models. The complexity of the relaxation and polarization effects, and electronic and vibrational excitations in the basic molecular unit of the organic films, which we alluded to in the introduction, brings an important question with regard to the assignment of energy levels of charge carriers. Energy levels and charge carrier transport in standard inorganic semiconductors are well described in terms of delocalized states and the single-band approximation. Because of delocalization and efficient screening, exciton binding energies are small. The transport gap Et, also known as the single-particle gap, defined as the minimum energy of formation of a separated free electron and hole in the solid, is within a few meV of the optical gap Eopt. Eopt corresponds to the onset of optical absorption. The situation is fundamentally different in molecular solids, in which localization and polarization dominate the physics of charged excitation and transport [17,18,72]. These solids exhibit small intermolecular overlaps and large polarizabilities involving charges and induced dipoles [18,73]. The optical gap, Eopt, corresponds to the formation of a Frenkel or charge transfer

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HOMO level of the isolated molecule. This polarization includes an electronic component, which consists of the response of the electronic charge density of the central and surrounding molecules to the formation of the hole. This response is fast (⬃10⫺16 s) and accounts for most of the polarization energy [18]. The polarization also includes a vibronic component, which corresponds to the new equilibrium geometrical configuration of the molecular cation. This component switches on a time scale set by the ‘‘breathing’’ (ag) molecular modes (⬃10⫺15 –10⫺14 s). The UPS spectrum of the HOMO includes both polarization components and therefore provides a measure of the nearly fully relaxed positive polaron (the missing piece is the lattice relaxation component, which is too slow to be included in the PES process, but is also very small, i.e., a few tens of meV), which is the relevant hole transport level in the molecular solid. The measured energy level of the nearly fully relaxed positive polaron is labeled E⫹ P in Figure 19. Similarly, the IPES spectrum shows the LUMO shifted by the polarization P⫺ induced by the presence of the molecular anion. P⫺ includes components that are conceptually identical, although not necessarily equal in magnitude, to those of P⫹. The energy level of the nearly fully relaxed negative polaron is labeled E⫺ P in Figure 19. A comparison between occupied and unoccupied electronic levels measured via PES/IPES and calculated with a semiempirical Hartree-Fock intermediate neglect of differential overlap method is given in Figure 20 for ␣-NPD [22]. Similar agreement was obtained for PTCDA and Alq3, demonstrating in the process that IPES can be used to investigate empty states of a number of organic films without significant electron beam–induced deterioration of the film. Figure 21 gives the composite PES/IPES spectra for four organic films. From these spectra, the determination of the transport gap Et requires the consideration of several points. The first addresses the choice of the peak-to-peak HOMO-LUMO gap, rather than the edge-to-edge gap, to evaluate Et. The rationale for this choice is that the HOMO and LUMO peaks measured on the molecular films derive from narrow HOMO and LUMO peak of the gas phase molecule shifted by polarization and broadened by solid-state effects and random disorder. Thus, the relevant HOMO-LUMO gap is between the peak centroids. Vibrational levels, too closely spaced to be resolved, and perhaps some delocalization of molecular states due to intermolecular wavefunction overlap, contribute to the widths, but to a lesser extent [83]. Vibrational excitation in PES/IPES tends to shift both the measured HOMO and LUMO peaks away from the Fermi level. The measured peak-to-peak separation increases by the Franck-Condon maxima. To correct for this, a 100-meV shift is estimated for each peak and 0.2 eV is subtracted from the PES/IPES gap. Second, the difference between surface and bulk polarizations must be understood and corrected for in the

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Figure 20 (a) Comparison between UPS and IPES spectra measured from ␣-NPD thin film (top) and INDO-simulated UPS and IPES spectra (bottom). The energy scale is referenced to the Fermi level. A compression factor of 1.3 is used to simulate the UPS spectrum and the FWHM is set to 0.5 eV prior to compression. The vertical bars at the bottom of the graph refer to the calculated energies of the molecular orbitals. (b) LCAO patterns of the HOMO and LUMO levels, calculated at the INDO level ␣-NPD. (From Ref. 21, with permission.)

PES/IPES gap. PES and IPES are surface-sensitive techniques and, as such, measure hole and electron levels located predominantly in surface molecules. Polarization induced by a surface molecular ion is reduced due to the reduced number of neighboring molecules, leading to a ‘‘surface transport gap’’ that is in fact larger than the bulk transport gap. It is interesting to

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Figure 20

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Continued

note that the evaluation of polarization energies using the available gas phase PES data on PTCDA [84], Alq3 [85], and CuPc [86] and our condensed phase PES data lead to identical P⫹’s equal to 1.15 ⫾ 0.3 eV for all three materials. The correction to obtain the bulk from the surface polarization was experimentally estimated by Salaneck to be around 0.3 eV for a molecule like anthracene [87]. This difference may in fact change in layers of molecules like PTCDA whose in-plane polarization is significantly larger than out-of-plane polarization, thereby increasing the weight of in-plane neighbors relative to the polarization from the lane below the molecule. Understanding the film structure and surface morphology is of course necessary to quantify these differences more precisely. Short of additional information, it is therefore assumed that the surface to bulk difference in polarization is 0.3 eV for the HOMO and 0.3 eV for the LUMO. The total correction from measured HOMO-LUMO gap to the bulk Et, including the

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Table 1 Bulk transport gap, Et, obtained by correction of the UPS-IPES gap, as explained in Ref. 23a Organic material CuPc PTCDA ␣-6T ␣-NPD Alq3

Et ⫾ 0.4 (eV)

Eopt (eV)

Et ⫺ Eopt ⫾ 0.4 (eV)

2.3 3.2 3.4 4.5 4.6

1.7 2.6 3.0 3.5 3.2

0.6 0.6 0.4 1.0 1.4

Eopt is the absorption peak energy; the charge separation, or exciton binding, energy is Et ⫺ Eopt .

a

0.2 eV for vibrational excitation, is therefore 0.8 eV, leading to the numbers given in the second column of Table 1. Note that the uncertainty on Et is large (⫾0.4 eV), given the approximations stated above. The third column gives Eopt, defined as the energy corresponding to the maximum of the absorption peak. The last column gives Et ⫺ Eopt, the charge separation energy or exciton binding energy in the molecular film. The more delocalized ␲ systems like PTCDA exhibit bulk electronhole binding energies around 0.6 eV, which is in relatively good agreement with the 0.80-eV result reported for pentacene [88]. The 0.4-eV result for ␣-6T is also in line with the 0.2 eV reported for ␣-6T [75]. The latter ˚ vs. 23 A ˚ ) and sustains a molecule is significantly longer than ␣-6T (32 A ˚ ˚ more extended exciton (28 A vs. 14 A) with lower binding energy. The binding energy in ␣-NPD and Alq3 is significantly larger and equal to 1.0 and 1.4 eV, respectively. Note that the Alq3 result is incompatible with the 0.3-eV value based on internal photoemission [88]. The reason for this discrepancy is unclear at this point. It is somewhat surprising, however, given the good agreement between the resulting electron barriers at this interface (see below). These results for Et are for conjugated molecules with transport or device applications as thin films. As expected, the transport gap is more accurately known in organic crystals than in amorphous films. Crystalline polymers based on substituted polydiacetylenes (PDAs) have Et = 2.5 eV and O-O peak of Eopt = 2.0 eV, with small variations with substituents and temperature [89]. These binding energies are consistent with the above estimates for PTCDA, ␣-6T, and CuPc; their O-O peak just above or below 2 eV suggests comparable delocalization and polarizabilities. Anthracene crystals have Et = 4.1 eV and Eopt = 3.1 eV [17], in accord with a smaller ␲-system and cation–anion separation. Comparison with ␣-NPD is rough,

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however, because that molecule is not planar. The binding of naphthalene is large [18] and comparable to that of Alq3, whose quinolate ligands also contain two rings. The anion and cation charges in adjacent Alq3 are largely localized on individual ligands. We note that these encouraging comparisons with crystal data do not extend to amorphous films of conjugated polymers, where Eopt is known but different techniques for extracting Et yield binding energies ranging from a few kT to more than 1 eV [74]. Beyond the evaluation of the exciton binding energy, the impact of this work on interfaces of interest here should be to provide more accurate barriers for electron and hole injection. Indeed, the position of the hole state relative to the Fermi level at an MO interface or to another hole state at an OO interface, i.e., the hole barrier, is known from PES. As argued above, the determination of the electron barrier cannot be inferred from the hole barrier and the optical gap, and requires IPES or internal photoemission measurements or knowledge of the transport gap Et. The sum of the electron and hole barriers is the single-particle or transport gap of the material. Figure 22 gives a tentative example of gap correction for the Mg–Alq3 interface. EF is measured at about 2.7–2.8 eV above the leading edge of the HOMO peak (Figure 6a). Using the optical gap and neglecting the exciton binding energy (Figure 22, left) makes the electron injection barrier (leading edge of the LUMO peak ⫺EF) extremely small (ⱕ0.1 eV). Et, defined above, is the energy gap between the centroids of the HOMO and LUMO peaks. The gap between the edges of the HOMO and LUMO peaks is smaller than Et by roughly twice the base width of the peaks (Figure 22, right). The composite Alq3 PES/IPES data [23] give an edge gap of about 3.7–3.8 eV. Using the more appropriate transport gap increases the energy difference between the EF and the edge of the LUMO peak to about 0.9 eV, more in line with the 0.6-eV value based on internal photoemission [82] and with current thinking that transport at this interface is injection limited.

V.

CONCLUSIONS

Surface and interface techniques have been shown time and again to be extremely powerful for unraveling fundamental problems related to semiconductor heterointerfaces. The work done over the past decade on interfaces of molecular films certainly confirms the importance and validity of a surface–interface approach. The overwhelming conclusions of our work and that of others are that organic interfaces are chemically and electronically more complex and varied than initially perceived. In the context of an emerging technology based on thin-film devices, interfaces continue to play a disproportionately important role, as they often constitute the bottleneck

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Figure 22 Energy diagram and electron injection barrier at the Mg:Ag/Alq3 interface using (a) the optical gap Eopt and (b) the transport gap Et (gap between the HOMO and LUMO centroids, as defined in the text).

to the injection and transport of charge carriers. In the same time, the flexibility afforded by organic molecular films in terms of modes and conditions of deposition, chemical modification of materials, molecular mixing, and doping is opening a number of avenues for modifying and improving interfaces in ways that cannot be applied to inorganic semiconductors. The role of surface–interface science in taking advantage of this flexibility will be extremely important.

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ACKNOWLEDGMENTS Several people participated in the experimental work and in the theoretical analysis of the data reviewed in this chapter. Many thanks are due to Drs. Yutaka Hirose, Aparna Rajagopal, and Chih-I Wu and Ms. Weiying Gao for their PES and IPES work on organic interfaces, to Prof. Jean-Luc Bre´das and Drs. Jeroˆme Cornil and Denis dos Santos for their collaboration on filled- and empty-state calculations, to Prof. Zoltan Soos for his collaboration on the problem of electron-hole separation energy, and to Prof. Steve Forrest and his group for many helpful discussions and generous help in providing organic materials. Support of this work by the MRSEC program of the National Science Foundation (DMR-9809483) and by the New Jersey Center for Optoelectronics is gratefully acknowledged.

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12 Metal–Polyfluorene Interface and Surface Structures and Stability Liang-Sheng Liao, C. S. Lee, and Shuit Tong Lee City University of Hong Kong, Hong Kong, China

M. Inbasekaran and W. W. Wu The Dow Chemical Company, Midland, Michigan

I.

INTRODUCTION

Since the first report of polymer light-emitting diodes (PLEDs) by Burroughes et al. in 1990 [1], rapid progress has been made in improving device brightness, efficiency, and operational lifetime [2–7]. These developments have brought PLEDs to the stage where they are ready for commercialization [4]. To date, poly(p-phenylenevinylene) (PPV) and its derivatives are by far the most popular light-emitting layers used in PLEDs and perform very well [5]. Recently, the highest luminescence efficiency (22 lm/W) in PLEDs has been achieved in devices using yet another kind of polymer—fluorene-based polyer—as the active light-emitting layer [5,6]. Fluorene-based polymers, such as poly(9,9-dioctylfluorene) (PFO), are now attracting much interest because the PFO-based PLEDs have low turn-on voltage, high brightness, and high efficiency [6–12]. In PLEDs, the interface between the metal electrode and the active light-emitting polymer is an important factor in determining device perfor401

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mance [13]. Calcium is usually used as the cathode in PLEDs because of its low work function, which leads to a low barrier for electron injection [14]. Several groups [15–18] have studied extensively the interface formation between Ca and PPV (or substituted PPVs). Several useful conclusions have been drawn, such as the following: (1) no ideal metal-on-polymer contact can be formed and there is always some chemistry at the interface [15]; (2) for a very clean polymer surface, Ca diffuses into the near-surface region ˚ [16], and to form an interface with a thickness in the range of 20–40 A bipolaron states within the band gap can be generated [15,16]; (3) for an oxygen-containing polymer surface, an interfacial layer of Ca oxide is formed first, which is then followed by a layer of Ca metal [17]; and (4) the performance of PPV-based PLEDs can be improved, if Ca is deposited in a background O2 partial pressure of ⬃10⫺6 mbar [18]. These results provide guidance in optimizing the performance of PPV-based PLEDs and are useful references for studies on other kinds of light-emitting polymers. In contrast to PPV, there is no double bond in between the phenylene groups along the backbone of polyfluorene. As a result, it is expected that the characteristic of polyfluorene will be different from that of PPV. Therefore, characterization of the electronic structure of the polyfluorene film, understanding of the interface formation between metal and polyfluorene, investigation of the oxygen effect on the metal/polyfluorene film, investigation of surface modification of polyfluorene film, and stability study of the film are important for the fabrication of polyfluorene-based PLEDs. In this chapter, we take PFO as an example and report our systematic investigations on the metal–PFO interface and surface mainly by using ultraviolet photoemission spectroscopy (UPS) and X-ray photoemission spectroscopy (XPS).

II.

EXPERIMENTAL

A.

Instrumentation

The UPS and XPS measurements were performed in a VG ESCALAB 220iXL photoemission spectroscopy (PES) system, consisting of an analysis chamber, a multiport carousel chamber, and a deposition chamber. The base pressures in the three chambers were lower than 8.0 ⫻ 10⫺11, 3.0 ⫻ 10⫺10, and 3.0 ⫻ 10⫺10 mbar, respectively. The analysis chamber is equipped with XPS with fast imaging capabilities, UPS, a field emission gun, Auger scanning microscopy, scanning electron microscopy, and optical imaging. We used the He I excitation line (21.2 eV) or the He II excitation line (40.8 eV) from a He discharge lamp for the UPS measurements, and the monochromated Al K␣ excitation line (1486.6 eV) for the XPS measurements. The

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absolute energy resolution of the analyzer was set at 0.018 eV for the UPS measurement and 0.36 eV for the XPS measurement. The UPS spectra were recorded with a sample bias of ⫺4.00 V to allow the observation of the inelastic electron cutoff. A clean Au film on Si substrate was used for the Fermi level (EF) and binding energy calibration. Photoluminescence (PL) and photoluminescence excitation (PLE) spectra were obtained using a Perkin Elmer LS 50B luminescence spectrometer. Surface morphology and surface roughness were observed using Scientific Autoprobe CP atomic force microscopy under ambient conditions. B.

Material and Sample Preparation

The chemical structure of PFO is shown in the inset of Figure 1. The polymer was synthesized via a Suzuki coupling reaction [19]. Possible impurities introduced during synthesis were carefully removed. The PFO polymer used in this experiment has an Mw of 46,000 and an Mn of 33,000, respectively, as determined by multiangle laser light scattering measurements. The PFO was spin-coated from a toluene solution (0.15 wt%) to form ˚ films on indium-tin-oxide (ITO) glass substrates for approximately 1000-A

˚ on ITO Figure 1 XPS survey scan of PFO film with a thickness about 1000 A substrate. The inset is the chemical structure of PFO.

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˚ films for metal deposition and surface cleanness evaluation, and about 70-A interface study either on clean ITO or clean Si substrates (or Au-coated Si substrates). The ITO substrates were further treated by ultraviolet (UV) ozone before spin-coating the PFO. The work function of the ozone-treated ITO substrate was ⌽ = 4.7 eV, as determined by UPS. The root-mean-square˚ and average roughness of the ITO surface and the Si surface were 9.25 A ˚ 1.9 A, respectively, as determined with atomic force microscopy. These substrates were flat enough for the experiments. C.

Sample Treatments

The freshly spin-coated PFO film sample was immediately loaded in a fastentry air lock, which was connected to the multiport carousel chamber in the PES system. For metal deposition, Ca and Ag were evaporated by resistively heated tungsten baskets in the deposition chamber. The Ca layer, with nominal thickness monitored by a quartz crystal microbalance, was ˚ /min and under a pressure of ⬃3 ⫻ deposited at a constant rate of ⬃0.25 A ˚ /min 10⫺9 mbar. The Ag layer was deposited at a constant rate of ⬃0.6 A and under a pressure of ⬃2 ⫻ 10⫺7 mbar. For oxygen exposure experiments, the exposure was performed in the fast entry air lock (with a base pressure of 2.0 ⫻ 10⫺7 mbar) by filling it with pure (99.99%) oxygen gas. For ion irradiation experiments, 1.0 keV Ar ion irradiation with a dose of ⬃6.5 ⫻ 1014 ions/cm2 was performed in the analysis chamber, with half of the sample surface covered by a movable mask for later comparison. Before the above experiments, the possible spectral changes due to irradiation of x-ray and UV sources on PFO film were monitored, but no discernible spectral change was detected.

III.

RESULTS AND DISCUSSIONS

A.

Characterization of PFO Films

Shown in Figure 1 is the XPS survey scan of the PFO film with a thickness ˚ on ITO. Only one intensive C1s core level is seen in the of ⬃1000 A spectrum and no oxygen contamination is detected. This indicates that our samples for surface analysis are pure in material and clean on their surface. It also implies that, unlike some other kinds of light-emitting materials, very clean PFO film can be fabricated using the solution processing method under ordinary ambient conditions. The XPS C1s core level with its shake-up feature is shown in Figure 2. There are three peaks in the spectrum. The main peak (named as peak ␲) has a full width at half maximum (FWHM) of 0.91 eV and is positioned at

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Figure 2

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XPS C1s core level spectrum of a PFO film.

the binding energy (BE) of 285.0 eV. The two shake-up peaks [␲*(1) and ␲*(2), respectively] are easily discernible after 20 times enlargement. These peaks correspond to a shake-up process involving core-electron excitation and a ␲ – ␲* transition. The energy differences of ␲ – ␲*(1) and ␲ – ␲*(2) are estimated as 3.55 ⫾ 0.10 eV and 5.75 ⫾ 0.10 eV, respectively. Shown in Figure 3 is the UPS (He II) spectrum of the PFO film. Several peaks are well resolved in the spectrum. Although the valence band structure is a combined contribution from both the phenylene backbone and the alkyl side chains, the first two peaks (called A and B, respectively) are attributed to the delocalized ␲ states of the backbone, since they are similar to that of poly(p-phenylene) [20]. We notice that the energy difference between peak A and peak B is about 1.65 eV, and the interval between the maximum of peak A and its onset is about 0.45 eV. Figure 4 shows the He I UPS spectrum of the PFO film. The inelastic electron cutoff and the highest occupied state (HOS) of the highest occupied molecular orbital (HOMO) are enlarged as well. According to the values of the cutoff and the HOS, the ionization potential (Ip) can be obtained to be Ip = 21.22 eV ⫺ (⫺1.40 eV ⫹ 17.00 eV) = 5.62 eV. In the UPS measure-

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Figure 3

UPS spectrum of a PFO film obtained with He II excitation.

ments, we set the energy resolution of the system at 0.036 eV and used an energy scanning step of 0.02 eV. From more than five measurements for each of seven PFO samples prepared, we obtained the Ip with the value of 5.60 ⫾ 0.05 eV. Figure 5 shows the PL (curve a) and PLE (curve b) spectra of the PFO film. The PL spectrum has a maximum at 2.83 eV, a subpeak at 2.66 eV, and a shoulder at 2.91 eV. The PLE spectrum, similar to the absorption spectrum, has a maximum at 3.13 eV (peak 1), a small peak at 4.24 eV (peak 2), and a shoulder at about 5.36 eV (peak 3). The onset of peak 1 is at 2.85 eV. The present spectra are similar to the corresponding spectra in Ref. 7, except that the maximum absorption peak (peak 1) is 0.1 eV less in the present case. These slight differences may be due to the use of different solvents in the preparation of the PFO solutions [21]. (We used toluene instead of xylene). B.

Energy Band Structure of PFO Film

Usually, UPS is the single most applicable method that gives the amount of information needed to reveal the electronic structures of materials [15]. The

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Figure 4

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UPS spectrum of a PFO film obtained with He I excitation.

ionization potential (Ip) and the energy level of the highest occupied molecular orbital (HOMO) can be measured with UPS, as described above. However, information on the lowest unoccupied molecular orbital (LUMO) is usually not readily available from the PES techniques. Fortunately, since conjugated polymers have their delocalized electronic ␲ and ␲* states along their backbones, their ␲ – ␲* transition may be detected from the shake-up peaks on the high binding energy (BE) side of the main core level peaks in the spectra obtained by XPS [16,22]. Specifically, for PFO film, since its XPS C1s core level has a FWHM of less than 1.0 eV, the shake-up peak close to the C1s main peak is well resolved [peak ␲*(1) in Figure 2] without severe overlapping. Thus, both the Ip and the electron energy band gap (Eg) could be measured by PES in an identical experimental setting.

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Figure 5 (a) The photoluminescence (PL) spectrum and (b) photoluminescence excitation (PLE) spectrum of a PFO film on a quartz substrate.

Here we describe a method of determining the Eg of PFO film. Although it may not be an exact way, it provides a practical way to obtain Eg. The electron energy level alignment of Figure 6a is obtained from the UPS measurement (from Figures 3 and 4). We did not consider the possible band bending at the interface of PFO/Au in this comparison. Figure 6b is derived from Figure 2. The ␲-state level (the baseline) was assigned to align with A(HOMO). After settlement of this alignment, ␲*(1) is obviously the LUMO of the PFO film. Combining Figure 6a and b, we can then obtain the energy band structure of PFO film, as shown in Figure 6c. According to this energy level alignment, we tentatively define Eg as the energy difference between LUMO [␲*(1)] and the HOS, which is 3.10 ⫾ 0.10 eV. This value is similar to the single-particle energy gap (Eg,s = 3.20 ⫾ 0.10 eV) of the PFO film obtained by Alvarado et al. using a scanning tunneling microscope [23]. As a comparison, the optical energy band gap Eg,opt (Eg,opt = 2.85 eV), i.e., the onset of peak 1 in the PLE spectrum in Figure 5, is also shown in Figure 6c. As is known, using Eg,opt as Eg is a common practice in determining Eaffi. But Eg,opt is usually ⬃0.3–0.4 eV less than Eg [23]. Therefore, the Eg we measured is a better approximation of real Eg than Eg,opt.

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Figure 6 Diagrams for energy level alignment. (a) As derived from the UPS spectra in Figures 3 and 4, (b) as derived from the XPS spectrum in Figure 2, (c) a summary of (a) and (b) of this figure. Evac is the vacuum energy level and ⌽ is a work function.

C.

Interface Formation in Metal/PFO Film

Although the interface formation between metal and PPV (or substituted PPVs) has been well studied [15–18], few studies have reported on metal– PFO film [24]. Since the chemistry that occurs at the metal-on-polymer interface varies with the nature of polymers as well as metals [5], understanding of the interface formation between metal and polyfluorene is important for the fabrication of polyfluorene-based PLEDs. In this section, we will take Ca/PFO and Ag/PFO as examples to discuss the interaction at the interface of metal–PFO film. Figure 7 shows the UPS spectra with increasing Ca coverage. The pristine sample (PFO film on ITO) shows two peaks at the high kinetic energy side near the EF. Peak A is the HOMO of PFO. The energy difference between the HOMO edge or HOS and the EF is about 1.1 eV in this exper-

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Figure 7 Evolution of UPS spectrum with increasing Ca coverage. The ionization potential of the pristine sample is Ip = 5.60 eV.

iment. As the Ca coverage increases (this is only a nominal Ca coverage since some Ca atoms diffuse into the PFO film), two obvious spectral changes can be observed. First, the inelastic electron cutoff moves to the ˚ . The difference lower kinetic energy side until the Ca coverage reaches 16 A ˚ Ca in the cutoff energy between the pristine PFO sample and the 16-A covered sample is 1.95 eV. This means that the vacuum energy level (Evac) shows a 1.95-eV decrease after the Ca coverage. Since further Ca deposition does not change the cutoff energy, this suggests that the Evac has changed ˚ Ca deposition. Second, from that of PFO to that of Ca metal after about 16-A peaks A and B also move away from the EF, with a decrease both in intensity and in peak sharpness. At the same time, new gap states emerge in the former energy gap. To compare the evolution of the new gap states more conveniently, we aligned the UPS spectra according to peaks A and B. Since valley C (shown

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in Figure 7) had the same shift as peaks A and B, we chose this as the reference point for alignment. Shown in Figure 8a are the enlarged UPS spectra in the region near EF after the peak alignment. Figure 8b is the difference spectra taken from Figure 8a. The evolution of the gap states with increasing Ca coverage is clearly illustrated in Figure 8. This is especially evident in the difference spectra in Figure 8b, which shows that the gap ˚ states were first formed at the HOMO edge when the Ca coverage was 1 A

Figure 8 Enlarged UPS spectra in the region near the EF after aligning peak B in each spectrum of Figure 7. (a) Evolution of the UPS spectrum and (b) the evolution of the UPS difference spectrum with increasing Ca coverage. The energy difference between the peaks of the bipolaron states is about 1.8 eV.

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˚ , two-peak gap states or below. When the Ca coverage was higher than 1 A emerged in the former gap. Metal valence band structure would be formed by filling the valley between the two peaks if Ca coverage higher than 16 ˚ was deposited. The energy difference between the two peaks was about A ˚. 1.8 eV when the coverage was about 16 A The two-peak gap states shown in Figure 8 are the bipolaron states similar to those reported by Salaneck et al. in the interface analysis of Ca/ poly(2,5-diheptyl-p-phenylenevinylene), and the energy difference between the two peaks is the same in the two cases [16]. Moreover, in our case, we have shown a transition from one-peak gap states to two-peak gap states. This might be an indication of a transition from polaron states in a lightly doping stage to bipolaron states in a heavily doping stage, since a singly charged polaron should have gap states that are closer to band edges than is the case for a bipolaron [25]. Figure 9 shows a set of C1s XPS spectra, as well as their enlarged shake-up portion. The pristine PFO sample has the C1s core level with a FWHM of 0.91 eV. The energy difference between ␲ and ␲*(1) is 3.55 eV. The energy band gap (Eg) is estimated as ⬃3.1–3.2 eV, as discussed in Section B. Moreover, in the set of enlarged shake-up peaks, we can see that the two shake-up peaks gradually broaden, diminish, and eventually vanish at the later deposition stage. Since the HOMO peaks (peaks A and B in Figure 7) of the PFO were affected and broadened with increasing Ca coverage, the evolution of the shake-up peaks suggests that the LUMOs of the PFO were also affected and broadened with increasing Ca coverage. Based on the data obtained from the figures above, a schematic energy diagram of Ca/PFO/ITO structure can be drawn as in Figure 10. This diagram shows details in the interface region of Ca/PFO. This interface region is featured with bipolaron states, broadened HOMO, and broadened LUMO. The thickness of the interface region (⌬) was estimated to be more than 40 ˚ by angle-dependent XPS measurement (not shown here), which was A thicker than that in the case of Ca/PPV [16]. This is reasonable because there is no oxygen element in PFO, and Ca cannot form any compound as a diffusion barrier. Figure 10 is also an energy structure of a single layered PLED. By considering the operational situation, this energy diagram suggests that (1) efficient electron injection from Ca electrode can be achieved due to low or no injection barrier at the interface of Ca/PFO; (2) a thicker interface region with bipolaron states is detrimental to band–band radiative recombination because electrons from cathode and holes from anode are subject to recombination in this interface region, especially when the PFO film is thinner and when the hole mobility of the PFO is much higher than electron mobility in this film. Therefore, it is useful for the improvement of device perfor-

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Figure 9 Evolution of XPS C1s core level of the PFO film with increasing Ca coverage. The inset is the enlarged shake-up portion. The FWHM of the C1s peak of the pristine sample is 0.91 eV and the energy difference between peaks ␲ and ␲*(1) is about 3.55 eV.

mance by modifying the interface and stopping the formation of interface region with bipolaron states; and (3) a hole transport layer is of course necessary for the reduction of the hole injection barrier and for improvement of the interface at the anode side. Shown in Figure 11 is the evolution of UPS spectra with increasing Ag coverage (as indicated by nominal thickness). The spectral shape is identical but with little shift to higher energy. The spectral shift is about 0.3 eV ˚ . Figure 12 is the enlarged portion of when the Ag coverage is about 5 A

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Figure 10 Energy diagram of Ca/PFO/ITO based on the results of Figures 7, 8, and 9. The energy unit is eV.

HOMO (peak A) in Figure 11. By comparing it with Figure 8, we noticed that no significant change in HOMO and no gap states were induced upon Ag deposition. The evolution of C1s core level with increasing Ag coverage, shown in Figure 13, indicates that there is no obvious spectral change including the shake-up peaks upon Ag deposition. Therefore, unlike Ca on PFO surface, Figures 11, 12, and 13 show that there is no obvious interaction between Ag and PFO at the interface. Upon Ag deposition, both HOMO and LUMO do not change but show a slight energy shift. Accordingly, the energy ˚ -thick Ag deposition can be drawn in Figure band structure of PFO upon 5-A 14, which is quite different from that in Figure 10. Moreover, we found that the Ag diffusion region, or the interfacial region (⌬), is much thicker than that of Ca because no Fermi edge can be observed in the UPS spectra (shown ˚ -thick Ag deposition. This means metal contact has in Figure 12) upon 5-A not yet formed and the Ag atoms have diffused into the PFO layer. In this section, we have discussed the interface formation between Ca and PFO and between Ag and PFO. It is obvious that the interface formation between metal and PFO is dependent on the nature of the metal.

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Enlarged UPS spectra in the region near the EF in Figure 11.

strates, taking Ca/PFO as a study case. We use UPS to measure occupied molecular orbitals and gap states, and XPS to measure the shake-up peaks of the C1s core level of the PFO films, and then extract information about the LUMOs. We present a strong correlation between oxygen exposure and the carrier injection barrier at the interface of Ca/PFO. Shown in Figure 15a is the UPS (He I) spectral evolution of the PFO film with increasing Ca coverage and with increasing O2 exposure. Enlarged views of the cutoff at the high binding energy side and HOMO edge at the low binding energy side are shown in Figure 15b and c, respectively. In these figures, curve a is the spectrum of the as-prepared PFO film, and curves b and c are the spectra after Ca deposition with nominal coverage of 0.5 ˚ , respectively. Curves d, e, and f are from the 6-A ˚ Ca/PFO sample and 6 A upon exposure to O2 for 103, 105, and 107 LM (1 LM = 1 ⫻ 10⫺6 Torr⭈s), respectively. Four peaks in curve a, called A, B, C, and D, are discernible. ˚ (the Ca/C atom ratio is about Upon Ca deposition with a coverage of 0.5 A 1 at. %; i.e., about three repeat units of PFO have only one Ca atom), the spectral shape changed obviously. Except for peak B, other peaks were no

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Figure 14 Energy diagram of Ag/PFO/ITO based on the results of Figures 11, 12, and 13. The energy unit is eV.

between the two peaks. After exposing the Ca-deposited PFO to O2 (curves d, e, and f), the spectral shape was recovered partially, with peaks C and D becoming discernible again. The cutoff began to shift back to the lower binding energy side. Most importantly, the gap states were removed from the forbidden energy gap (peak 1 totally disappeared and peak 2 was reduced gradually). Figure 16 shows the evolution of the XPS C1s core level of the PFO film, with increasing Ca coverage and with increasing O2 exposure. The inset is the enlarged portion of shake-up peaks with the same binding energy scale. The curves, from a to f, have the same denotation as in Figure 15. After Ca deposition or O2 exposure, the C1s core level of the PFO film had not only a peak shift but also a change in its FWHM. The as-prepared PFO film had an FWHM of 0.91 eV in its C1s peak. Upon the Ca deposition of ˚ and 6 A ˚ , the FWHM was changed to 1.49 eV and 1.65 eV, respec0.5 A tively. This indicated that a charge transfer from Ca atoms to C atoms and chemical interaction between Ca and C had occurred. After the Ca-deposited PFO film was exposed to O2 of 103 –107 LM, the FWHM was reduced from 1.40 eV to 1.23 eV. This suggested that the reaction between Ca and C be reduced to some extent with increasing O2 exposure. However, the present results suggest that the interaction between Ca and C upon Ca deposition will not be fully eliminated by the oxygen exposure.

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Figure 15 Evolution of the UPS spectrum with increasing Ca and O2 exposure. (a) The full-range spectrum. (b) Enlarged view of the high binding energy portion. (c) Enlarged view of the low binding energy portion. Curve a: As-prepared PFO ˚ . Curve c: Ca sample. Curve b: Ca-deposited with a nominal thickness of 0.5 A ˚ . Curve d: With 103 LM O2 exposure. deposited with a nominal thickness of 6 A Curve e: With 105 LM O2 exposure. Curve f: With 107 LM O2 exposure.

We now examined the evolution of the shake-up peaks ␲*(1) and ␲*(2). These peaks are the ␲ – ␲* (HOMO-LUMO) transitions of the delocalized electronic ␲ and ␲* states along the backbones of the conjugated PFO [12,16,22]. Peak ␲*(2) is the common shake-up peak for any molecules with aromatic group, but peak ␲*(1) is specifically for the phenylene chains [28]. With this information, the evolution of LUMO can be revealed. Upon Ca deposition, peak ␲*(1), as well as peak ␲*(2), of the PFO film was affected by the broadened C1s peak and the Ca-induced bipolaron states. (If ˚ , the two shakethe nominal Ca coverage increases gradually to about 24 A

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Figure 16 Evolution of the XPS C1s core level of the PFO film with increasing Ca coverage and O2 exposure. The inset is the enlarged shake-up portion. The denotation of each curve is the same as in Figure 15.

up peaks will be diminished and eventually disappear, as shown in Figure 9.) Upon 103-LM-O2 exposure, as shown in curve d of the inset, ␲*(1) emerged again, due to the increase of FWHM of the C1s main peak and removal of the bipolaron states in the former forbidden energy gap. However, further O2 exposure resulted in the removal of ␲*(1), as shown in curves e and f of the inset. We have repeated this experiment several times, with the same result, which may suggest that the phylene chains in PFO were broken upon extra O2 exposure and some phenoxide groups might be formed. The ␲ – ␲*(2) transition began to dominate the band gap in the Ca-

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doped region. When the sample was further exposed to ambient for 60 s, the intensity of ␲*(2) increased further (with a slight shift in peak position) and ␲*(1) was no longer visible (not shown here). Figures 15 and 16 provided us with enough data to draw the energy band structures. The extracted data from Figures 15 and 16 are listed in Table 1. Figure 17 is a schematic energy band diagram drawn according to the data in Table 1. Figure 17a is the as-prepared PFO film on n-type silicon. Here, we estimated the energy band gap Eg = HOS ⫺ LUMO = 3.1 ⫾ 0.1 eV as discussed in Section B. Upon Ca deposition, Ca atoms diffused into the near surface to form an interfacial region (or Ca-doped region). The ˚ , according to our depth thickness (⌬) of the region was more than 40 A profile measurements (not shown here). The diagram in Figure 17b was drawn based on the assumption of Fermi level (EF) alignment, taking into account that bipolaron states were formed in the forbidden energy gap and the vacuum energy level (Evac) was moved downward by 1.6 eV. This structure (Figure 17b is obviously not good for radiative recombination of the

Table 1

Energy band parameters extracted from Figures 15 and 16 (eV)

Parameter Ipa ⌽b LUMO-2 peak [␲*(2)]c LUMO-1 peak [␲*(1)] EF Gap states (peak 1) Gap states (peak 2) HOMO edge HOMO peak (␲)d Cutoff a

As-prepared (curve a)

Ca-deposited (curve b)

103 LM O2 exposure (curve d)

5.6 4.3 3.9 1.7 0 None None ⫺1.4 ⫺1.9 ⫺17.0

5.5 2.7 22.5 Overlapped 0 ⫺1.0 ⫺2.8 (⫺2.9) ⫺3.4 ⫺18.6

5.5 3.2 3.0 0.5 0 None Little ⫺2.3 (⫺2.8) ⫺18.1

107 LM O2 exposure (curve f) 5.6 3.7 3.1 None 0 None Little/none ⫺1.9 (⫺2.4) ⫺17.6

Ip = photon energy (21.22 eV)–(HOMO edge-cutoff). ⌽ (work function) = Ip ⫺ (EF ⫺ HOMO edge). We use cutoff position (edge) at the high binding energy side of UPS spectrum to calculate ⌽. If the midpoint value on the slope of the high binding energy side would be used for the calculation, the ⌽ then would be increased by ⬃0.2–0.4 eV. c From LUMO-2 peak down to cutoff are readings either converted from XPS data or obtained from UPS data. d According to our UPS (He II) spectra or curve fitting, the energy difference between HOMO peak and HOMO edge is usually about 0.45 eV. The values in the brackets are the estimated values by considering the difference of 0.45 eV. b

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Figure 17 Schematic energy diagram of Ca/PFO/Si based on the data of Table 1. ˚ in nominal thickness). (c) With 103(a) As-prepared. (b) Ca-deposited (with 0.5 A 7 LM-O2 exposure. (d) With 10 -LM-O2 exposure.

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carriers because a portion of holes and electrons will recombine nonradioactively in the gap states area (the shadowed area). Gap states can be removed by exposing the sample to O2. Figure 17c shows the energy band diagram upon 103-LM-O2 exposure. In this diagram, gap states were mostly removed and Evac was increased by 0.5 eV. Although there was a 0.5-eV barrier for electron injection, the nonradiative recombination path was broken, due to the removal of gap states. This is a good structure for radiative recombination. Conceptually, if the O2 were just enough to convert all doped Ca atoms to CaO, all the gap states would be removed. However, if Ca(OH)2, CaCO3, or other kinds of compounds were formed at the same time, much more O2 would be needed. In our experiment, 103-LM-O2 exposure means that the sample is exposed to 10⫺6 Torr of oxygen partial pressure for 1000 s, or to ambient for about 10⫺5 s. However, if the O2 exposure is more than 105 LM in the experiments, the barrier height will be too high for electron injection. As shown in Figure 17d, the original LUMO-1 peak was broken due to extra oxygen exposure and the electron injection barrier height was increased to about 3.1 eV, which made it very difficult to inject electrons into the thick interfacial region. Therefore, according to the experiments, too ‘‘clean’’ Ca will induce gap states and too ‘‘dirty’’ Ca will produce a very high injection barrier in the interfacial region. The optimal amount of oxygen exposure lies between the two extreme situations. This very high injection barrier formed upon 105-LM-O2 exposure was not due to the formation of a full CaO layer at the interface because the Ca ˚ had already atoms deposited on the surface with a nominal thickness of 6 A ˚. A diffused into the film and formed an interfacial region thicker than 40 A full layer of Ca had not yet formed before oxygen exposure in this experiment. Therefore, it appears that this barrier would be formed much earlier than the oxidization of Ca electrode itself, if there were indeed a Ca layer already on the PFO surface. The breakdown of the LUMO-1 peak in the interfacial region in Figure 17d, or the disappearance of peak ␲*(1) in the shake-up portion of C1s core level in Figure 16, upon more than 105-LM-O2 exposure may indicate that the structures of some PFOs in the Ca-doped region or interfacial region were changed due to extra oxygen exposure. Possibly the C — C bond connecting the two benzene rings in one repeat unit, or between two repeat units of PFO, was attached by extra oxygen, and phenoxide groups might be formed. However, without Ca deposition, the structure of PFO is very stable. For example, the two shake-up peaks did not show any significant change, even when the film was exposed to ambient conditions for 4 weeks, as will be shown in a later section. In this section, we have investigated the interface formation in the Cadoped interfacial region in PFO film upon oxygen exposure using UPS and

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XPS. The UPS results indicated that the deformed UPS spectrum after Ca deposition can be partially recovered and the Ca doping–induced bipolaron states in the former forbidden energy gap can be removed upon O2 exposure. The XPS results indicated that the interaction between Ca and C can be reduced and the LUMO can be recovered after O2 exposure. However, too much O2 exposure will break the original LUMO and result in a very wide band gap, which might be due to the change in the chemical structures of PFO in the Ca-doped region after extra O2 exposure. By combining both XPS and UPS data, we can explain the evolution of oxygen effect on the carrier injection and recombination in the Ca-doped region.

E.

Improvement of Interface Formation Between Ca/PFO Film by Surface Modification Using Mild Ion Irradiation

As mentioned above, Gao et al. reported that nonradiative decay channels were created by deposited metal atoms on PPVs. Even a submonolayer of Al, Ag, and Ca will severely quench the luminescence from the polymers [26,29]. We can also deduce from their results that these metal-induced gap states are also detrimental to the performance of polymer transistors and photodiodes. Therefore, interfaces with less indiffusion and sharper contacts between metal electrodes and polymers are needed for all high-quality polymer-based electronic devices. In PLEDs, interface formation can be improved by inserting a very thin insulator layer, or by forming a metal oxide thin layer in between metal electrode and polymer [26]. This layer is effective in stopping the indiffusion of metal atoms and in reducing the gap states in the polymer layers, resulting in an enhancement of light emission efficiency. However, the quality and the thickness of the insulation layer are not easily controlled. Too thick a layer would handicap the carrier injection, and too thin a layer would not stop the indiffusion of metal atoms efficiently. Moreover, from a circuit speed point of view, especially for polymer transistors, the inserted thin insulator introduces additional capacitance, which may reduce the device speed. Surface modification techniques are commonly used to optimize the thick polymer surface for some applications, such as to enhance adhesion between polymer and metal coating, and to improve the electrical conductivity of some conducting polymers [30–32]. However, these techniques, such as ion beam modification, are not suitable for use in electronic devices with very thin organic films as functional layers because the thin films are subject to severe surface damage [33]. Up to now, few studies have been

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reported on polymer surface modification in PLEDs and in other polymerbased electronic devices. However, in our experiments we found that thin polymer films are much more stable than some small molecular organic films when exposed to energetic ions and their surfaces can be modified with mild ion irradiation. Therefore, as an alternative way to improve the interface between metal electrode and polymer, we used 1.0-keV Ar ions to irradiate the surface of PFO film before Ca electrode deposition. In this section, we report our recent results obtained from XPS and UPS. Shown in Figure 18 are sectioned UPS spectra of the PFO film with or without Ar⫹ irradiation, which focus mainly in the secondary electron cutoff and the HOMOs near the EF. Surface modification was achieved by damaging the molecular structures at the surface after irradiation. The sec-

Figure 18 Sectioned UPS spectra of PFO film in nonirradiated area (dotted line) and in 1000-eV Ar⫹-irradiated area (solid line).

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ondary electron cutoff was moved about 0.2 eV toward the lower kinetic energy (or the work function of the PFO film on the irradiated surface was decreased by about 0.2 eV). Thus, the surface was positively charged after the irradiation. It may indicate that some C — H and C — C bonds were broken and the interface structures became more disordered. In this experiment, the HOMO also shifted by about 0.2 eV and was broadened after irradiation. However, this surface damage was very mild compared to our previous observation of the surface of small molecular organic film irradiated by 100 eV Ar⫹ [33]. Since the surface modification can be controlled by varying the irradiation energy and the dosage, higher irradiation energy or a higher irradiation dosage can be used if more severe surface damage is needed. Figure 19 shows the evolution of XPS spectra with increasing Ca coverage (nominal thickness values obtained from the crystal microbalance are shown in the graph). The XPS core levels from the substrate (Si2p, Si2s), the PFO film (C1s), and deposited Ca metals (Ca2p, Ca2s) can be seen in the figure. Both spectra from the Ar⫹-irradiated and the nonirradiated areas are grouped according to Ca coverage. Before Ca deposition, both areas reveal almost identical spectra, as shown in curves a and f, respectively (but the FWHM of the C1s peak is slightly increased from 0.90 eV to 0.97 eV upon the 1-keV Ar⫹ irradiation, as will be shown in Figure 20b). As the nominal Ca coverage increases, we find that the Ca2p peaks from the irradiated area are higher than those from the nonirradiated area. Since the spectroscopy can ‘‘see’’ more Ca in the irradiated area than in the nonirradiated area, more Ca atoms deposited on the irradiated area stay on the uppermost surface, while more Ca atoms deposited on the nonirradiated area have diffused into the film. This is evidence that the irradiated surface can form a sharper contact upon Ca deposition. Shown in Figures 20a–c are the evolutions of the Ca/C atomic ratio, the FWHM of C1s peak, and the work function of the PFO film with increasing Ca coverage, respectively. Similar to Figure 19, the Ca/C ratio in the irradiated area increases much faster than that in the nonirradiated area, as is further shown in Figure 20a. For example, at the nominal coverage of ˚ , the Ca/C ratio in the irradiated area is almost twice that in the non24 A irradiated area. Since C atomic concentration is basically the same in both areas, the detected surface Ca atomic concentration [Ca(irrad.)] in the irradiated area is then considered twice that of Ca(nonirrad.) in the nonirradiated ˚ , the FWHM of C1s peak area. In Figure 20b, at a nominal coverage of 16 A stabilized in the irradiated area. This indicates that further increase of Ca coverage will not lead to the additional interaction between Ca and PFO, which may suggest that a full Ca layer begins to build up on the PFO surface. In contrast, FWHM increases monotonically in the nonirradiated

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Figure 19 XPS spectral evolution of PFO film with increasing Ca coverage. Curves a–e, taken from Ar⫹-irradiated area. Curves f–j are taken from nonirradiated area. The curves are grouped with nominal Ca coverage.

area. This indicates that the interaction between Ca and PFO in the nonirradiated area continues, which may suggest that Ca continues to diffuse into PFO film with increasing Ca coverage. In Figure 20c, we noticed that the work function decreases rapidly with increasing Ca coverage at low cover˚ in the irradiated area. This is conage and becomes constant beyond 16 A sistent with the change of FWHM in the irradiated area (shown in Figure 20b.

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Figure 20 Ca coverage dependence of (a) Ca/C ratio, (b) FWHM of C1s core level, and (c) work function, of PFO film in Ar⫹-irradiated area (—●—) and in nonirradiated area (— 䊱 —).

Figure 21a and b are portions of the spectral evolutions near the Fermi edge with increasing Ca coverage in the irradiated area and in the nonirra˚ Ca deposition, the diated area, respectively. In the irradiated area, upon 4-A cutoff edge was discernible at the EF, and no doping-induced bipolaron states were formed in the former forbidden energy gap, which indicates a metal contact on PFO surface. Further deposition of Ca atoms continued to build ˚ and 8up the metal contact. However, in the nonirradiated area, upon 4-A

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Figure 21 Evolution of UPS spectra as increasing Ca coverage. (a) Ar⫹-irradiated area and (b) nonirradiated area.

˚ Ca deposition, only bipolaron states were found near the EF, and metal A ˚ deposition. Therefore, Ar⫹ irradiation contact could not be formed until 16-A can form a sharper contact between Ca electrode and PFO film. As mentioned in Figure 20a, Ca(irrad.)/Ca(nonirrad.) ⬇ 2. Based on the data, we now estimated the thickness ratio of Ca diffusion region, or interfacial region, in both areas. For the convenience of simple calculation, we assumed that the Ca atomic concentration in the irradiated area, Ci(x), is rectangularly distributed: Ci(x) =

再

C0 0

0 ⱕ x ⱕ ti x > ti

(1)

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and assume that in the nonirradiated area, Cn(x), is triangularly distributed:

再

冉 冊

␤C0 1 ⫺

Cn(x) =

x t2

0

0 ⱕ x ⱕ tn

0<␤<1 (2)

x > tn

and

冕

ti

冕

tn

Ci(x)dx =

0

Cn(x)dx

(3)

0

where ti and tn are the thickness of the Ca diffusion region in the irradiated area and the nonirradiated area, respectively. By using the following equation,

冕

t

It(A) =

B ⭈ C(x) ⭈ e⫺x/␭A cos ␪ dx

(4)

0

where It(A) is the photoelectron intensity of a core level of an element (A) in a layer with thickness of t; B is a constant including X-ray photon intensity, geometrical factor for the electrons, atomic photoemission cross section, etc; ␭A is the inelastic mean free path of the photoelectrons from a specific core level of element A; ␪ is the polar angle between the normal of the sample and the detector direction; and x is the depth perpendicular to the sample surface, we can roughly estimate that tn/ti > 4, when Ca(irrad.)/ ˚ (deduced from Ref. 34). If Ci(x) is also Ca(nonirrad.) ⬇ 2 and ␭A = 20 A triangularly distributed, tn/ti will increase further. In our experiment, the ˚ (measured by ellipsometer); thus, the thickness of PFO film was about 80 A thickness of the interface region in the Ar⫹-irradiated area was possibly 20 ˚ at most. A Upon Ca deposition on the Ar⫹-irradiated area, the broken C — C or C — H bonds may form C — Ca and other bonds with high density in the top surface. This disordered top layer with high density of C — Ca and other bonds may efficiently stop the diffusion of later deposited Ca atoms. Therefore, further Ca deposition will easily form a Ca layer on the top. It is also likely that this metal layer has better adhesion to PFO film due to surface modification. In the fabrication of polyfluorene-based PLEDs, the emission layer is usually 200 nm in thickness [7], which is twice or three times thicker than that in typical small molecular OLEDs [35]. Moreover, PFO is more stable than small molecular organic films during irradiation [36]. Therefore, mild ion irradiation on PFO surface may not have a negative effect on the performance of the devices. However, further investigation is in progress.

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In this section PFO surface was modified by 1-keV Ar⫹ ions with a dose of about 6.5 ⫻ 1014 ions/cm2 before Ca deposition. This modified surface can effectively stop Ca diffusion into PFO film and form a sharper metal contact on the surface of the PFO film upon Ca deposition by comparing with the surface without Ar⫹ irradiation. Therefore, suitable surface modification of polymer surface may be useful to improvement of the devices. F.

Energy Alignment of PFO-Based PLEDs

Recently, increasing numbers of papers related to PFO-based LEDs have been published. However, to the best of our knowledge, no energy level alignment for the devices has been reported, apart from schematic energy level structures [7]. In fact, information on the energy level alignment is crucial to improve the performance of PFO-based devices. Usually, in PFObased devices, a polymer named as poly{9,9-dioctylfluorene-co-[bis-N,N⬘(3-carboxyphenyl)-bis-N,N⬘-phenylbenzidine]} (BFA) is used as a hole transport layer (HTL) (its chemical structure is shown in the inset of Figure 22) [37]. Thus the device configuration usually consists of Ca/PFO/BFA/ ITO. Therefore, in this section, we use UPS and PLE to measure the energy level structure of the device with the above-mentioned configuration. In this experiment, clean ITO was treated with ozone and BFA solution (⬃0.2 wt %), spin-coated on the ITO surface, and heated overnight at 70⬚C to form an approximately 6-nm-thick layer. After the UPS measurement on the BFA in the UHV, an approximately 6-nm-thick PFO film was formed on the BFA layer by spin-coating, as mentioned in the experimental section of this chapter. After characterizing the PFO layer with UPS in the UHV, ˚ -thick Ca was deposited on the PFO layer followed again by UPS an 8-A measurement. Figure 22 shows the UPS (He II) spectrum of the BFA film. The peaks in the spectrum are well resolved. The doublet structure of the HOMO (peak A) is common for the triphenylamine group, such as in TPD and in NPB. Its Ip was measured as 5.4 eV from UPS (He I) spectrum in this experiment. Shown in Figure 23a are the UPS (He I) spectra of different layers. These layers are built up one by one after each characterization step; meanwhile, the layers are kept thin enough without severe charging effect. Curves a–d were taken from ITO, BFA on ITO, PFO on BFA, and Ca on PFO, respectively. The enlarged views of the cutoff at the high binding energy side and the HOS at the low binding energy side are shown in Figures 23b and c, respectively. When we look at the cutoff portion in Figure 23b, we find that the BFA has almost the same vacuum energy level (Evac) as that of the ITO substrate after it was spin-coated on the substrate; the Evac of PFO film was

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Figure 22 UPS spectrum of a BFA film on ITO substrate obtained with He II excitation. The inset is the chemical structure of BFA.

0.4 eV lower than that of BFA after PFO film was formed on BFA layer; ˚ -thick Ca was deposited on and the Evac of PFO was 1.6 eV lower after 8-A it. This means that dipoles were formed at the interfaces of PFO/BFA and Ca/PFO. In Figure 23c, we show that the band offsets of BFA/ITO and PFO/ BFA measured 0.7 and 0.6 eV, respectively. Since there are carboxyl groups in the repeat unit of the BFA molecule, and the O-C — —O peak in XPS core level superposed on the shake-up peaks in the C1s core level, it is difficult to determine the ␲ – ␲* transition via the shake-up peaks in this situation. Therefore, we used PLE to estimate the Eg of BFA film. Shown in Figure 24 are the PL and PLE spectra of BFA film coated on a clean quartz substrate. The PL peak is at 2.61 eV (␭ = 475 nm). The PLE peaks are at 4.28 eV (␭ = 290 nm), 3.65 eV (␭ = 340 nm), and 3.13 eV (␭ = 396 nm), respectively. We chose 3.13 eV as the Eg of the BFA film [27]. Table 2 shows the energy band parameters extracted from Figures 23 and 24. Figure 25a shows the energy level alignment before Ca deposition. The hole injection barriers from ITO to BFA and from BFA to PFO are 0.7

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Figure 23 UPS spectra of different layers on ITO substrate. (a) The full-range spectrum. (b) Enlarged view of the high binding energy portion. (c) Enlarged view of the low binding energy portion. Curve a: As-prepared ITO substrate. Curve b: ˚ -thick Ca deposited BFA film on ITO. Curve c: PFO film on BFA film. Curve d: 8-A on PFO film.

and 0.6 eV, respectively. Figure 25b shows the energy level alignment after Ca deposition. In order to align the EF between the ITO and Ca layer, the energy levels of BFA and PFO have to be decreased. As a result, the electron injection barrier from the Ca layer to PFO is estimated at about 0.2 eV. Figure 25b indicates that, in PFO-based LEDs, the electron injection barrier could be much lower than the hole injection barrier by using low work function metals as the cathode. Unlike OLEDs, the carrier recombination mechanism in the PFO-based LEDs may be limited by holes. Therefore, reducing the hole injection barrier should be very important for the improvement of device performance in PFO-base LEDs. Further investigation is in progress.

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Figure 24 The photoluminescence (PL) spectrum (curve a) and photoluminescence excitation (PLE) spectrum of a BFA film on a quartz substrate.

Table 2

Energy band parameters extracted from Figures 23 and 24 (eV) ITO

a

Ip ⌽b Eg LUMO-EF EF EF-HOS a b

4.7

0

BFA on ITO

PFO on BFA

5.4 ⬃4.7 3.1 2.4 0 0.7

5.6 4.3 3.1 1.8 0 1.3

Ca on PFO 2.7

0

Ip = photon energy (21.22 eV)–(HOS-cutoff). ⌽ (work function) = Ip ⫺ (EF ⫺ HOS). We use cutoff position (edge) at the high binding energy side of UPS spectrum to calcuate ⌽. If the midpoint value on the slope of the high binding energy side would be used for the calculation, the ⌽ then would be increased by ⬃0.2–0.4 eV.

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Figure 25 Energy level alignment based on the data in Table 2. (a) PFO/BFA/ ITO and (b) Ca/PFO/BFA/ITO. The energy unit is eV.

G.

Stability Study of PFO film—Ambient Effect

As for all OLEDs, stability is one of the essential requirements for practical applications. And the stability of light-emitting films is strongly related to that of devices. The objective of this section is to investigate the effect of ambient storage on the electronic structure of PFO film. UPS and XPS spectra were taken from the as-prepared PFO sample and those stored in the open air for 1, 2, 3, and 4 weeks (at ⬃60–70% RH). Figure 26 shows the evolution of the UPS (He II) spectrum of the PFO film via the ambient storage time. As the storage time increased, the intensity of peaks A, B, and C gradually decreased. However, all of the peaks in the spectrum were still resolved after storage in air. Moreover, according to the UPS (He I) spectrum (not shown here), the ionization potential of the PFO film changed only 0.2 eV after storage for 4 weeks in air. This means that the surface of the PFO film cannot be easily contaminated or attacked by

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Figure 26 Evolution of UPS spectrum (He II) of PFO film with increasing storage time in air.

ambient species due to its strong bonding. However, for other kinds of materials, such as tris(8-hydroxyquinoline)aluminum (Alq), their electronic structure obviously changed, even after the film was exposed to air for only 1 h. Figure 27 shows the evolution of the XPS spectrum of the PFO film with the ambient storage time. The FWHM increased slightly and the intensity of the shake-up peaks decreased slightly with increased storage time. However, these changes are very small in comparison to other light-emitting materials. Shown in Figure 28 are the surface morphologies of the PFO film on Si substrate taken with atomic force microscopy. The as-prepared film had a root-mean-square roughness of 1.8 nm in a 10 ⫻ 10 ␮m2. After storage in air for 2 weeks, the morphology did not show obvious changes and the roughness was almost the same as before. After storage in air for 4 weeks, the morphology showed some changes and the roughness increased to about 2.1 nm. However, these are small changes compared to that of Alq. We found that the surface roughness of Alq film can be increased from 3 nm to 30 nm after storage in air for 4 weeks. Therefore, PFO film is more stable under ambient conditions. This is very helpful during device processing.

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Figure 27 Evolution of the XPS C1s core level of the PFO film with increasing storage time in air. The inset is the enlarged shake-up portion.

IV.

SUMMARY

The electronic structure of PFO film was investigated. From the UPS measurement, we obtained the Ip of the PFO film: Ip = 5.60 ⫾ 0.05 eV. From the XPS measurement of the shake-up peaks of C1s core level, we estimated the Eg of the film: Eg = 3.10 ⫾ 0.10 eV. The interactions between Ca and PFO film and between Ag and PFO film were studied. As the Ca coverage increases, the vacuum energy level (Evac) of PFO decreases gradually in order to match the Fermi level (EF) of

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Figure 28 Surface morphologies of the PFO film on a Si substrate taken with the atomic force microscopy in a 10 ⫻ 10 ␮m2. The root-mean-square roughness (rmr) is indicated in the figure.

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Ca. The original highest occupied molecular orbital of PFO moves away from the EF and vanishes eventually. Bipolaron states separated by 1.8 eV are formed in the former energy gap. The original lowest occupied molecular orbitals of PFO are broadened, resulting in a featureless shake-up in the XPS C1s core level. The interface region of Ca/PFO with bipolaron states is detrimental to the performance of the device. Therefore, modifying the interface and stopping the formation of interface region filling with bipolaron states is demonstrably useful for the improvement of device performance. However, because Ag atoms have little reaction with PFO film and are more diffusible in PFO film than Ca atoms, it is obvious that the interface formation between metal and PFO is metal dependent. We have investigated the interface formation in the Ca-doped interfacial region in PFO film upon oxygen exposure using UPS and XPS. The UPS results indicated that the deformed UPS spectrum after Ca deposition can be partially recovered and the Ca doping–induced bipolaron states in the former forbidden energy gap can be removed upon O2 exposure. The XPS results indicated that the interaction between Ca and C can be reduced and the LUMO can be recovered after O2 exposure. However, too much O2 exposure will break the original LUMO and result in a very wide band gap, which might be due to the change in the chemical structures of PFO in the Ca-doped region after extra O2 exposure. By combining both XPS and UPS data, we can explain the evolution of oxygen effect on the carrier injection and recombination in the Ca-doped region. PFO surface was modified by 1.0 keV Ar⫹ ions with a dose of about 6.5 ⫻ 1014 ions/cm2 before Ca deposition. This modified surface can effectively stop Ca diffusion into PFO film and form a sharper metal contact on the surface of the PFO film upon Ca deposition by comparing with the unirradiated surface. Therefore, suitable surface modification of polymer surface may be useful to the improvement of the devices. The energy level alignment of the PFO-based PLEDs with a device configuration of Ca/PFO/BFA/ITO was measured. We found that the electron injection barrier could be much lower than the hole injection barrier by using low work function metals as the cathode in the devices. Unlike OLEDs, balanced carrier recombination in the PFO-based LEDs may be limited by holes. Therefore, reduction of the hole injection barrier should be an important consideration for the improvement of device performance in PFObased LEDs. The stability of the electronic structure of PFO film was investigated. When the PFO film was exposed to ambient conditions for several weeks, its electronic structure did not show obvious changes. Therefore, we believe that PFO is a stable polymer material useful for PLEDs.

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molecular orbitals (HOMOs) of the donor and the acceptor and will be referred as HOMO-HOMO transfer in this chapter. The second class includes the transfer reactions between the excited state manifolds or the lowest unoccupied molecular orbitals (LUMOs) of the donor and the acceptor centers, and will be referred to as LUMO-LUMO transfer. Interestingly, the same molecular species of a particular heterostructure can display both types of CT depending on the nature of the excited states and the excitation energies involved. In the following, we shall focus on the LUMO-LUMO type of electron transfer, although the organic heterostructure studied shows both types of CT. In the case of LUMO-LUMO transfer, the key factors are the energy level alignment between the LUMO levels of the acceptor and the donor and the strength of the electronic coupling between the excited state manifolds. The first factor is critical in determining whether the reaction is energetically feasible, i.e., whether a barrierless ET can be realized. The second factor determines the overall rate and the efficiency of the reaction, and it is subject not only to the photochemistry of the donor and the acceptor but also to the film and the interface properties, e.g., morphology, orientation effects, etc. For a barrierless transition these factors can be incorporated in a simple rate equation [3]: k ET =

2␲ 2 V ␳ ប

(1)

Equation (1) defines the effective rate as the product of the electronic coupling V 2 between the two manifolds and the combined density of states ␳ of the donor–acceptor complex. These kinds of barrierless ET reactions in solid systems have been found to be extremely rapid, ranging from a few tens of femtoseconds to a picosecond, with the fast transfer rate being an important manifestation of the high CT efficiencies achievable in molecular systems. Earlier observations of ultrafast ET in solid systems have been based almost solely on the measurements of fully optical transients. In time-resolved fluorescence and transient absorption studies of oxazine-covered SnS2 surfaces, the time constant for the ET from the dye to the semiconductor was determined to be 40 fs [4,5]. Similar transient absorption measurements on ruthenium dye and perylene chromophores anchored onto a TiO2 film have deduced reaction times between 25 fs (ruthenium) and 190 fs (perylene) [6,7] for ET from the LUMO of the dye to the conduction band of the semiconductor. In all of these experiments, the LUMO of the dye molecule was found to be well above the conduction band minimum of the semiconductor and therefore the ET was considered barrierless. More recently, ultrafast ET has also been reported for a fully organic solid heterostructure of

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polymer/C60 composites [8,9]. In these studies, the transfer rates for ET from the polymer to C60 were found to be a few hundred femtoseconds. In this chapter, we will present and discuss the time-dependent photoemission spectroscopy (TR-PES) results on an organic heterostructure of N,N⬘-diphenethyl-3,4,9,10-perylenetetracarboxylic diimide (DPEP) and N,N⬘-diphenyl-N, N⬘-(3-methylphenyl)-1,1⬘-biphenyl-4,4⬘-diamine (TPD). DPEP is an organic photoconductor material that is well known for its charge generator properties, and TPD is a widely used organic hole transport material in many photoreceptor applications. Together these two materials form a photoreceptor structure whereby either of the molecules can act as a donor depending on the excitation energy. We found that the TR-PES spectra of the heterostructure were dominated by a two-exponential decay. The two components were separated by both intensity and duration so that the short component was typically less than 100 fs and the long component was a few hundreds femtoseconds in length. The time constants obtained for the ultrafast ET between TPD and DPEP compare very nicely with these previous experimental findings. The significant new property of the system studied here is that the ET is between two organic molecules in contrast to the organic–inorganic interface of the earlier experiments. The observation of the ultrafast ET process in a fully organic system is noteworthy in the sense that the strong coupling necessary for the fast reaction can exist between the two manifolds of a molecular acceptor–donor system.

II.

EXPERIMENT

The organic films were made in situ by vacuum evaporation (base pressure 1 ⫻ 10⫺9 Torr) of DPEP and TPD. The substrate materials used were indiumtin-oxide (ITO)–covered glass and vacuum-cleaved graphite. The morphology of the film was controlled by varying the substrate temperature and the deposition rate. For the experiments described in this chapter, we generated two DPEP film structures characterized by very different absorption spectra and film colors. The DPEP film with an absorption maximum at 500 nm (DPEP-I) was evaporated at the substrate temperature of ⫺60⬚C and with ˚ /s. The film with absorption maximum at 630 nm evaporation rate of 8 A (DPEP-II) was evaporated at an elevated substrate temperature of 110⬚C with ˚ /s. an evaporation rate of 1 A The TR-PES measurements were performed with an experimental setup that has been described in detail elsewhere [10]. Briefly, the laser system consisted of an active mode–locked Ti/sapphire laser pumped by an Ar⫹ laser. The frequency-doubled beam was split into two with a nonpolarizing beam splitter. After passing of one of the beam arms through a com-

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puter-controlled delay stage, the beams were recombined to create coaxial beams to ensure maximum beam overlap at the sample. The TR-PES experiment was carried out by measuring the photoemission intensity as a function of the delay time between the pump and the probe pulses in a monochromatic mode with cross-polarized pump and probe beams. The photon energy of the frequency-doubled light was tunable from 3.1 to 3.3 eV (375–400 nm), and the full width of half maximum (FWHM) of the Gaussian-shaped two-pulse instrument response function was typically 60–65 fs. The photoelectrons were detected by an angle-integrated hemispheric energy analyzer in an ultrahigh vacuum (UHV) chamber with the base pressure of 7 ⫻ 10⫺11 Torr. The resolution of the analyzer was set to 100 meV. In order to improve the detection efficiency, the sample was biased with ⫺10.0 V. TR-PES spectra were recorded both on pristine DPEP films and TPD/ DPEP bilayer structures. After the transient spectroscopy, ultraviolet photoemission spectroscopy (UPS) measurements were performed on the film structures in order to determine the ionization potentials and the position of the HOMO levels of the films. Inverse photoemission spectra (IPES) were collected in a separate apparatus of similar attributes as the TRPES system. In addition to the in situ measurements, the absorption spectra were recorded in the ambient conditions for a separate set of pristine film samples.

III.

RESULTS AND DISCUSSIONS

A.

Relaxation in Pristine DPEP Films

The TR-PES spectra were recorded with cross-polarized beams, i.e., a ppolarized pump beam and an s-polarized probe beam. In this case, the pumpprobe signal I(t) is a convolution of the two-pulse instrument response function R(t) and a single exponential decay term [11]: I(t) = R(t) 嘸 e⫺兩t兩/␶

(2)

where ␶ is the relaxation lifetime of an intermediate state electron. The lifetime was extracted by fitting the convoluted function to the experimental data. The pump-probe spectra for DPEP-I and DPEP-II at E ⫺ EHOMO = 2.2 eV with fits to the convolution functions are shown in Figure 1. Also shown in Figure 1 are the instrument response functions measured at E ⫺ EHOMO = 3.1 eV with fits to the convolution function. The absorption and emission spectra of the two films evaporated on glass substrates are shown in Figure 2. The film with an adsorption maximum at 500 nm (DPEP-I) was evaporated at the substrate temperature of ˚ /s. The film with absorption maxi⫺60⬚C and with evaporation rate of 8 A mum at 630 nm (DPEP-II) was evaporated at an elevated substrate temper-

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Figure 1 Upper: Pump-probe spectra for the DPEP-I (solid squares) and DPEPII (open circles) at E ⫺ E HOMO = 2.2 eV with fits to the convolution functions (solid lines). Lower: Pump-probe spectrum for the instrument response function (triangles) and for the DPEP-II film (circles) measured at E ⫺ E HOMO = 3.1 eV with fits to the convolution function (solid lines) given in the text.

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Figure 2 Upper: The absorption and emission spectra of DPEP in DMSO (10⫺6 mol). Lower: The absorption of DPEP-I film (dashed line) DPEP-II film (solid line) and the emission spectrum of both films.

˚ /s. In addition to the solidature of 110⬚C with an evaporation rate of 1 A state spectra, absorption and emission spectra were taken for a dilute solution of DPEP in dimethyl Sulfoxide (DMSO). It is clear from the lower panel of Figure 2 that the film morphology has a pronounced effect on the absorption properties of DPEP films. The absorption spectrum of DPEP-I shows the vibrational levels of the S1 state, which are both broadened and red-shifted

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compared with the solution spectrum of DPEP shown in the upper panel of Figure 2. The broadening and the shift of the spectral peaks are attributed to the polarization effects present in the solid state. The polymorph of the DPEP-I film has previously been described in the literature as an amorphous phase [12,13] where molecules are more or less randomly oriented. However, more recently, it has been shown that the composition of the DPEP-I film is small crystalline grains embedded in amorphous material [14]. The absorption spectrum of DPEP-II shows a further broadened S1 band whose long-wavelength part is red-shifted. The spectrum compares well with similar spectra of solvent-treated DPEP films [15] and polycrystalline films created directly through vacuum evaporation [13] reported in the literature. The extended states forming in the long-wavelength region are attributed to intermolecular interactions found in more ordered DPEP-II films [16], which is also used to explain the greater photocharge generation efficiency of these films [17]. Interestingly, the observed emission spectra are very similar for both polymorphs. It is centered at 690 nm (1.80 eV) and has broad shapeless features. A more careful analysis [14] that takes account for the reabsorption of some of the emitted radiation shows that the emission peak of the DPEPII film is blue-shifted by about 4 nm with respect to the emission peak of the DPEP-I film. The absorption and emission spectra of both forms of DPEP films are shown in energy scale in Figure 2. Following the assignment of the absorption peaks for pure perylene [18] the peaks at 2.25 eV (550 nm), 2.47 eV (500 nm), and 2.63 eV (470 nm) of the DPEP-I spectrum are assigned to the vibrational levels of the S1 state, and the low-intensity peak at 3.30 eV (375 nm) to the S2 state. Pulsed excitation produced no observable difference in the emission spectra when compared with the spectrum produced by cw excitation. The variation of the TR-PES lifetimes of both DPEP films as a function of energy is shown in Figure 3. The energy scale used in Figure 3 represents the energy of the intermediate state and was calculated by using E ⫺ E HOMO = IP ⫹ E K ⫺ ប␻

(3)

where E K is the kinetic energy of the electron detected by the analyzer, IP is the ionization potential of the molecular film, and ប␻ is the one-photon excitation energy. At high energies, both forms of the film show very fast lifetimes, about 30 fs, which decrease linearly with increasing energy. However, the DPEPII film shows significantly shorter lifetimes (32 fs) than the DPEP-I film (55 fs) at the lower excitation energies (E ⫺ E HOMO = 2.1 eV). The samples were also cooled down to 77 K, and the relaxation times were measured as

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Figure 3 Lifetimes measured with TR-PES for photoexcitations in a DPEP-I (squares) and in a DPEP-II (circles) film. Also shown is the position of the highest peak of S1 manifold (dashed line).

described above. No significant changes were observed in the lifetimes compared with the room temperature values. By combining the UPS and the IPES spectra of a molecular film it is possible to obtain the energy separation between the HOMO of the cation and the LUMO of the anion which are ionized molecules with a hole and an excess electron, respectively. This is possible because UPS and IPES experiments share a common Fermi level from which we can measure the position of spectral features in the energy axis. The Fermi level was determined from the UPS and IPES spectra of the gold substrate prior to the evaporation of DPEP films. The features of the combined UPS-IPES spectra of the two DPEP polymorphs appear to be very similar but not identical, as

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can be seen in Figure 4. Both the HOMO and LUMO peaks are relatively broad with FWHM = 1.0 eV. There are a number of factors contributing to peak widths in molecular solids, including variations in bulk and surface polarization energies, intramolecular vibrations, intrinsic linewidth of the HOMO and LUMO levels, and disorder effects [19,20]. The UPS and IPES spectra have a background component which needs to be subtracted to determine the HOMO and LUMO peak positions and widths accurately. The UPS background is due to low-energy secondary electrons, and there are standard procedures for doing the background subtraction [21,22]. In the case of IPES, the background is not as well understood, and it was found to vary from sample to sample in the experiments discussed here. This is consistent with other observations in IPES experiments on similar molecular

˚ -thick DPEP films with fits (solid Figure 4 Combined UPS-IPES spectra of 400-A lines) to sums of Gaussians (dashed lines) and background functions (dotted lines).

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films [23]. Therefore, the background was chosen to be generated by an arbitrary smooth function, which was a third-order polynomial in this case. The UPS spectra were fitted to a sum of three Gaussian functions (only the one for HOMO is shown). The position and the width of the HOMO peak were the same for both the DPEP-I and the DPEP-II films. The peak was found to be at 2.4 eV below the Fermi level with width of 1.1 eV. The IPES spectra of the two polymorphs look equally similar with almost identical peak positions. The lowest LUMO peaks are located at 1.2 eV above the Fermi level in both the films. However, the higher peaks have slightly different location in the two films: 2.6 eV and 4.1 eV in DPEP-I and 2.5 eV and 4.0 eV in DPEP-II. The widths of the two lowest LUMO levels are also different: the FWHM values are 1.1, 1.0, and 1.9 eV in DPEP-I and 1.0, 1.2, and 1.9 eV in DPEP-II. The HOMO-LUMO separation is 3.6 eV in both the films and is indicated in Figure 4. The energy gap indicated as the energy separation between the HOMO and the LUMO cutoffs in Figure 4 is 2.0 eV in the DPEP-I spectrum and 2.1 eV in the DPEP-II spectrum. From the experimental results presented above, we can divide the excited-state relaxation process into several stages distinguished by the duration and the interaction energies involved. The first stage is extremely fast and results from the interaction of the electron and hole with polarization fluctuations of the medium. In the second stage, the excited molecule undergoes a rapid internal conversion process from the highly excited S2 state to the lower S1 and S0 levels. This is then followed by a fast vibrational relaxation of S0 and S1 levels to the ground state and the lowest excited S1 state, respectively. It is from the low-lying S1 state that the 690-nm (1.80eV) fluorescence is observed. We explain the short lifetimes measured in the TR-PES experiments by a rapid internal conversion (IC) process from S2 and S1 levels to the vibrational manifolds of S1 and S0 states. We rule out the relaxation due to the valence electron polarization effects of the medium because it is simply too fast (0.4–0.5 fs) [24] to be observable with the given time resolution of the experiment (10 fs). Conversely, the duration of radiative decay in thin DPEP films is several hundred picoseconds or a few nanoseconds as found in the single photon counting experiment. These vales are several orders of magnitude longer than the lifetimes observed in the TR-PES experiment. The proposed model is illustrated in Figure 5. The 3.3 eV pump photon excites the molecule to the S2 level, which is isoenergetic with the highly excited vibrational levels of S1 and S0 states, and consequently the IC process can be extremely fast [25,26]. The IC is then followed by vibrational relaxation within the S1 and S0 manifolds and IC from S1 levels to the S0 manifold. The model depicted here is similar to the one used to explain the fast relaxation of highly excited benzene monomers and dimers observed in recent

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Figure 5 Two-photon photoemission process and the population relaxation mechanism in DPEP films. A 3.3-eV pump photon excites the electron to the S2 level from which rapid internal conversion (IC) to the S1 and S0 manifolds proceeds.

photoemission experiments [27,28]. In these experiments, the S2 states of benzene monomer and dimer were found to undergo a rapid IC process from S2 to S1 and S0 manifolds with a time constant of 50 fs [27]. This value compares nicely with the 20 to 50 fs lifetimes measured in our TR-PES experiments. In the proposed model, it is assumed that the relaxation time of the subsequent intramolecular vibrational energy redistribution (IVR) is comparable to the initial IC. The time scales for ultrafast IVR can be less than 100 fs [29]. The effectiveness of this dissipation mechanism can be further illustrated by considering the energy redistribution process that occurs during the IVR. The lowest lying level of the S2 manifold has very little vibrational energy and therefore most of the energy of the state is included in the electronic part of the wavefunction. As the excited molecule undergoes IC, the energy of the S2 state E S2 is split between the electronic part E S1 (E S0) and the vibrational part of the S1 (S0) state, Thus, the energy conservation in the IC process can be expressed approximately as: E S2 = E S1 ⫹

冘

n k ប␻ k

(4)

k

where the sum is over the vibrational quanta ប␻ k of the S1 state. The IC process is typically dominated by a single or a few promoting modes [30] and hence the summation in Eq. (4) contains only a few terms. This also means that any one of these promoting modes carries a significant portion of the total vibrational energy present. As the energy of this mode is split between a few low-energy modes, the energy level of the excitation drops dramatically and the electron transition to vacuum is no longer energetically

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feasible. We can estimate the energy of the active modes from the solution spectrum of DPEP shown in Figure 2. The absorption lines are separated by approximately 0.17 eV, which gives an idea of the magnitude of energies involved in the IVR. If all the energy was located on this single mode it would require approximately 8 quanta to account for the energy difference between the S2 and S0 levels. It is therefore evident that because of large vibrational energies involved, the rapid reduction of active-mode energies can be realized in the DPEP molecule. The probability for IC from S2 to S1 in benzene monomers and dimers has been found to be 1–2% [28]. This observation explains the absence of a long component in the pump-probe spectra of DPEP, which would otherwise be expected for the slower relaxation of the low-lying vibrational levels of S1 [31]. Instead, most of the excitation energy of the DPEP molecule is channeled to the S0 manifold, which then rapidly decays nonradioactively. It has also been noted that the transition probability from the intermediate state of a highly excited vibrational level of the S1 and S0 states to the final state may be reduced when exciting with small photon energies [27]. This is because of small Franck-Condon factors between the intermediate and final states, i.e., the vibrational content of the intermediate and final states is significantly different and subsequently the overlap integrals of the wavefunctions are small, resulting in substantial reduction of the transition probability. This may further contribute to the fast disappearance of the pump-probe signal. The linear increase of the lifetimes with decreasing energy indicates that after the isoenergetic IC process from the S2 level, the resulting S1 and S0 levels relax more slowly. In other words, the combined rate of IC from S1 to S0 and the vibrational relaxation of the S1 and S0 manifolds decreases with decreasing energy. This is consistent with a picture in which the density of states decreases with decreasing energy. This picture is, however, complicated by the static disorder that is present in an organic solid. The static disorder can be viewed as fluctuations of site energies, which result from the relatively weak interactions between the molecules in their ground state. The static fluctuations are in the range of 0.4–0.6 eV, and they cause the photoemission lines to be broadened [19]. Hence, we expect IC from the S2 level to contribute to the observed total relaxation rate over a similar range of energies below the 3.3-eV excitation energy. The effect of disorder is also observed in the UPS spectra of DPEP films where the HOMO levels are found to be very broad. This makes the exact determination of the ionization potential difficult and explains the relatively large error margin (⬃0.2 eV) obtained for IP. The measured lifetimes in DPEP-I and DPEP-II are very similar at

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high excitation energies, but starting at 2.6–2.7 eV above E HOMO, the lifetimes in DPEP-I are significantly longer than in DPEP-II. This ‘‘onset’’ is located about 0.65 eV below the S2 level, which matches the energy separation between the highest S1 peak at 2.65 eV and S2 level at 3.3 eV above E HOMO (see Figure 1). This suggests that the additional fast relaxation process observed in DPEP-II films originates from the S1 level. The absorption spectrum of DPEP-II shows the formation of extended states arising from S1 and having energies ranging from about 2.9 eV to below 2.1 eV. Therefore, the most obvious explanation for the reduced lifetimes in DPEP-II is a fast relaxation within this manifold of extended states. Such fast transitions have been predicted by theoretical calculations for J-aggregate systems with a similarly well-ordered stacking structure where the main energy loss mechanism is through emission of phonons [32]. Indirect evidence for the nature of the excited states in the DPEP films can be gathered from the IPES spectra in Figure 4. The features in the IPES spectra of the two polymorphs are strikingly similar, and the energy positions of the peaks are practically the same. Since the density of states of the unoccupied levels measured in IPES corresponds to the anionic states of the molecule, it means that differences observed in optical absorption and relaxation dynamics are not likely to be related to the anionic properties of the polymorphs. Rather, the excited state appears to be excitonic, i.e., a correlated electron-hole pair, which gives rise to the delocalization effects observed in the DPEP-II film. Other possible explanations for different relaxation rates include transport of excitations to the surface and defect sites, resulting in direct recombination [33], and spectral diffusion. These are both enhanced by the increased mobility of more delocalized excitations in the DPEP-II film. Spectral diffusion with extremely short exciton hopping times (10 fs) between different molecular sites has been observed for DPEP-II films [34]. However, we expect these last two processes to be more prominent at energies closer to the emission energy, 1.8 eV. Also, where the spectral diffusion is found to be prominent, the time regime is typically from hundreds of picoseconds to a few nanoseconds, and therefore if present this process is unlikely to be observed in the pump-probe spectra. In general, it is very interesting to see such a clear dependence of the relaxation dynamics on the film morphology. In particular, according to this result, the intermolecular interactions have a significant role in the initial relaxation of the photoexcitation in DPEP films. This may have important implications from a materials design point of view, since the photocarrier generation efficiency has been shown to depend on the nature of exciton state formed and hence on intermolecular interaction in organic photoconductors [35].

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Charge Transfer Dynamics in TPD/DPEP Bilayer Structure

As another material is deposited onto a surface one might expect to observe changes in the energetics and in the level densities of the occupied states. Since it is the position of these levels in energy that largely determines the observed kinetic energy spectrum in one-photon (1PPE) and two-photon (2PPE) photoemission experiments, it is critical to observe and understand any changes induced by the adlayer. In conjunction with the TR-PES measurements of the bilayer structures, we measured the UPS spectra on the films after each TPD evaporation. The progression of the UPS spectra of a TPD/DPEP-II film is shown in Figure 6 as a function of the kinetic energy observed in the energy analyzer. We have chosen to show here only the spectra of DPEP-II surfaces since it was found that TPD deposition onto the two different polymorphs produced practically identical UPS spectra just as in the case of pristine DPEP films. The effect of the added TPD layer can be observed as a shift of the ˚ coverages. The HOMO cutoff of the pristine HOMO level starting at 4-A ˚ the level has shifted to 29.60 DPEP-II film is at 29.45 eV, whereas at 24 A ˚ TPD eV. At the same time, the vacuum level cutoff of the film with 24 A is shifted down from the pristine value of 13.85 eV to 13.70 eV. This means that the bilayer film has a PE spectrum that is 0.3 eV wider than the value for the pristine DPEP film. Consequently, the ionization potential of the bilayer film is 5.3 eV, which is 0.3 eV lower than the ionization potential of the pristine film, 5.6 eV. Furthermore, the shift of the vacuum levels in the bilayer films indicates the presence of vacuum level offset between the two constituent films. This leads to level shifting at the film interface (see the diagram in Figure 6), which is a consequence of unequal work functions (ionization potentials) of the two materials that results in a built-in potential at the interface. The exact origin of the level shifting is under debate, and a number of different mechanisms have been proposed, including trapping in Scottky-type defects [36] and formation of a dipole layer [37] at the interfaces. The observed vacuum level shift of 0.15 eV can be considered small, and it compares well with shifts reported for other organic–organic interfaces where one of the composite layers is TPD [38]. Based on the vacuum and HOMO levels determined from the UPS spectra, the energy level diagram shown in the inset of Figure 6 was constructed. The HOMO level of the TPD/DPEP structure is 0.15 eV higher in energy than the DPEP HOMO level. This affects the kinetic energy spectrum measured in the TRPES and will be discussed later. Figure 6 also includes a UPS spectrum of a pure TPD film. The ion˚ film was found to be 5.2 eV, which compares ization potential of the 400-A

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trum of the composite films. Therefore, the fits shown for the composite films in Figure 7 include two exponential decay terms. I(t) = R(t) 嘸 [A1 exp(⫺兩t兩/␶1) ⫹ A2 exp(⫺兩t兩/␶2)]

(5)

where the amplitudes A1 and A2 were allowed to vary during the fitting process. This produced excellent fits, which are shown as solid lines in Figure 7. Interestingly, the two lifetime components included in the fits are separated by roughly an order of magnitude at high TPD coverages, as can be seen in Figures 8 and 9, where the deconvolved lifetimes are plotted as a function of energy above the DPEP HOMO level. The fast components range from 40 to 80 fs in both films, and the long components are from 200 to 1.6 ps in the TPD/DPEP-I film and from 200 to 800 fs in the TPD/DPEPII film. The short components in bilayers are also consistently longer than the corresponding single exponential lifetimes in pristine DPEP films. The lifetime vs. energy plots reveal another remarkable effect of the added TPD layer. The energy dependence of the lifetimes is different in bilayers when compared with the trend found in pristine DPEP films where lifetimes decrease monotonically with increasing energy. In bilayers, the short components seem to be almost energy-independent up to 3.0 eV above the DPEP HOMO. However, above 3.0 eV the short components increase with increasing energy, which is the opposite of that observed in pristine DPEP films. The quality of data acquired for the DPEP-II films appears to be somewhat better than the data for DPEP-I films. In other words, the short components of DPEP-I show more ‘‘scatter’’ than the same components of DPEP-II. This is most likely due to the lower signal-to-background ratio present in the DPEP-I-based bilayer data. Also, the same TPD coverage seems to have greater effect on the observed lifetime components in the DPEP-I film than in the DPEP-II film. At 2.5 eV above DPEP HOMO, with ˚ , the short and long components in the DPEP-II film a TPD coverage of 4 A are 44 fs and 190 fs, respectively, whereas the corresponding components in the DPEP-I film are 82 fs and 1.19 ps. This difference between the two polymorphs can be understood by the effective surface areas of the two pristine films. Because of the higher average roughness, the total surface area of the DPEP-II film is approximately 20% higher than the surface area of DPEP-I film [14]. Also, the growth of the TPD film may be very different on the two films. For instance, if the growth is dominated by clusters in one film but the TPD molecules assemble in a layer-by-layer fashion in the other, the effective TPD film thickness may vary greatly from film to film. The most dramatic change observed as the thickness of the TPD layer is increased is the sudden growth in the absolute signal intensity. The measured intensity for the short component at E ⫺ E HOMO = 2.55 eV in Figure

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Figure 7 Upper: A comparison of the single (dotted line) and two exponential ˚ TPD ⫹ DPEP(solid line) fits of the TR-PES spectrum of a bilayer structure, 16 A II, at E ⫺ E HOMO = 2.55 eV. Lower: The short and long components (dashed line) of the two-exponential fit (solid line).

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Figure 9

461

Long- and short-lifetime components in TPD/DPEP-I film.

The appearance of two exponential components in the bilayer decay dynamics indicates the presence of two distinct electron populations with different decay mechanisms. It is clear from the pronounced dependence of the lifetime components on the TPD deposition thickness that the decay dynamics observed in TR-PES experiments are interface and adlayer driven. Therefore, the first step in understanding the observed decay dynamics is to identify the possible electron populations formed in the bilayer structure upon photoexcitation. An obvious question is whether the two lifetimes components are just intrinsic decay times of photoexcitations in the two materials, DPEP and

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show very different energy dependence. Perhaps the most dramatic change observed in the short component is the three orders of magnitude increase in the intensity of the component as the thickness of the TPD layer is increased. These three observations about the fast-decay components in the bilayers indicate that the fast transients are not only due to electron relaxation within the DPEP film but that an additional mechanism is likely to influence the rapid relaxation dynamics in bilayer structures. In the model depicted in Figure 11, after photoexciting a TPD molecule the electron is quickly transferred from the TPD molecule across the interface into the DPEP film. Once in the DPEP the electron rapidly loses its energy, relaxing to the lowest level of the excited anion. Therefore, the thickness dependence of the lifetime components seems to indicate that the excitation decay measured in the TR-PES experiments is injection-limited, i.e., the lifetime components actually reflect the injection rates of electrons into DPEP. The latter statement can be clarified by considering a population model where two upper level populations n1(t) and n2(t) undergo a decay process at rates k1 and k 2, producing the population n0(t) with its own decay rate k 0. The kinetics of the three populations can be described by three coupled first-order differential equations [39]: dn1(t)/dt = ⫺k1n1(t) dn2 (t)/dt = ⫺k 2 n2 (t) dn0 (t)/dt = ⫺k 0 n0 (t) ⫹ k1n1(t) ⫹ k 2 n2 (t)

(6)

The set of first-order differential equations can be solved by standard methods, e.g., by the variation of the constant, and the solutions for the three populations are:

Figure 11

Electron transfer from TPD to DPEP.

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n1(t) = n10e⫺k1t n2 (t) = n20 e⫺k 2t n0 (t) =

k1n10 k 2 n02 [e⫺k1t ⫺ e⫺k 0 t ] ⫹ [e⫺k 2t ⫺ e⫺k 0 t ] k 0 ⫺ k1 k0 ⫺ k2 ⫹ n00 e⫺k 0t

(7)

where n0i = ni (0). The solutions for the parent populations n1(t) and n2(t) are single exponential decays whereas the expression for the daughter population n0(t) includes three exponential terms that account for the intrinsic decay of the population and for the injection of particles produced by the parent populations. Depending on the relative magnitudes of the decay constants k 0 , k1 , and k 2 , the decay of the population can appear to be single- or multiexponential. If k 0 > k1 , k 2 , then the exponential terms containing k 0 will vanish faster than the other two terms, including k1 and k 2 , and these two decay constants will therefore determine the observed kinetics of the three-population system. The above considerations of the three-population kinetics can be related to bilayer kinetics by assigning the decay rates and populations to the ones observed in the TR-PES experiments. We believe that there are two parent populations n1(t) and n2(t) with fast k1 and slow k 2 injection rates, respectively, created in the TPD layer upon photoexcitation. These rates are slower than the inherent relaxation of the electron in DPEP described by rate k 0 after its injection from TPD. Therefore, the two exponential terms describing the decay of the pump-probe spectra correspond to the fast and slow injection of electrons from the TPD layer into DPEP. The critical assumption included in the injection model is the fast relaxation of the injected electron in DPEP. Since the relaxation rate cannot be measured directly in the experiment where the dynamic is dominated by the charge injection, we had to estimate the rate indirectly. This was done by fitting Eq. (7) to the pump-probe curves shown in Figure 7 and by keeping the previously found two components k1 and k 2 fixed. The additional component representing the decay in DPEP turned out to be 20–30 fs (=1/ k 0) in each case. The resulting fits could not be distinguished from these shown in Figure 7. Hence, we estimate the relaxation time for the electron in the DPEP after injection to be 20–30 fs or less. These values compare relatively well with the relaxation times measured in pristine DPEP films. However, we need to note that the excited states and possibly the relaxation mechanisms in the two systems are different. In the case of the pristine film, the photon absorption results in an excited state of a neutral DPEP molecule, whereas in the bilayer structures the charge injection is followed by the creation of an excited state of an anion (= neutral molecule ⫹ electron).

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Figure 12 Fits of the ET model to the fast components in the TPD/DPEP-II structure with the deduced cutoff for the ET at 3.1 eV (vertical line).

at lower energies reflect the most efficient injection rate near the TPD LUMO. The prefactor K 0 has an almost linear dependence on the TPD thickness, as shown in the inset of Figure 12. K 0 is proportional to 2␲ V 2/ប␭1/2 where V 2 represents the electronic coupling between the donor and the acceptor levels. Therefore, the thickness dependence of K 0 is an indication of how the orbital overlap integral depends on the separation distance and orientation between the acceptor and the donor, whereas the reorganization energy [42] ␭ is assumed to be constant. Here we need to stress the fact that we expect the effective donor–acceptor distance to be much less than the overall TPD coverage. We believe that the main contribution to the short component comes from the ET in the near-interface regions, and thickness dependence of the prefactor K 0 results from the small increase of the average

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distance of the near-interface region of the TPD layer as the amount of TPD is increased. The slow components of the charge injection are separated from the fast components not only by their magnitude but by the fact that they are practically independent of energy just below and above the 3.1 eV cutoff energy for ET. This observation leads one to believe that the electrons contributing to the slow injection rate are likely to originate from a different population of excitations within the TPD layer. What makes identifying the nature of the second population and the charge injection process difficult is the fact that in the time scale in which the slower injection takes place, there could be both charge transfer and charge transport processes involved. Further, the dominating process might be different at different TPD coverages. We first considered a simple charge transfer mechanism with a typical donor–acceptor separation dependence. The separation dependence of the transfer rate is usually thought to be exponential [42]: k(r) = k 0 e⫺␤ (r⫺r0 )

(9)

where r0 can be treated as the average separation distance between the donor and acceptor centers, ␤ is the attenuation constant, and k 0 is the prefactor equal to the ET rate at r = r0. The exponential form for the separation dependence comes from the tunneling probability between the two reaction centers. The tunneling is expected to dominate the charge transfer at low temperatures near the zero-point vibrational energy of the reactant pair. Therefore, describing the slower charge transfer with Eq. (9) involves a postulation for a state that is less strongly coupled to the interface. Such a state could be a disorder-induced trap state that has less energy available for ET reaction or has such an orientation with respect to the interface that the coupling to the acceptor surface is weaker. To account for the distribution of the separation distance r0 in a film of thickness d, we can calculate the average value for the transfer rate:

冕 冕 d

drk(r)

具k(r)典 =

0

= k 0 e␤r0

d

dr

1 ⫺ e⫺␤d ␤d

(10)

0

Equation (10) was fit to the reciprocals of the long components, i.e., k 2 = 1/␶2 shown in Figure 13. The value of ␤ deduced from the fit is ap˚ , and the value for r0 is 12 A ˚ . In recent experiments, the proximately 15 A ˚ in effective transfer length has been found to be even longer, about 80 A

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Let us briefly consider a possible transport mechanism. For all the coverage of the TPD/DPEP-II and for all but the highest coverage of the TPD/DPEP-I, the long transients are all below 1 ps. A typical thermalization time for an electron in simple metals is 1 ps or longer [44,45], which is comparable to the vibrational relaxation time in organic molecules. This implies that the excited charge distribution responsible for the slower injection does not have time to equilibrate before the charge is moved across the interface. A simple diffusion process does not provide an adequate description of the transport of a hot particle distribution where the nature of particle motion can be characterized more accurately as ballistic than as a random walk. The ballistic transport model has previously been employed in the analysis of the lifetimes measured with TR-PES at simple metal surfaces [46,47], and it was found that the ballistic transport would limit the observable lifetime below 100 fs. Since the effective probing depth in the metal ˚ , we would not expect to see any slow comsamples is thought to be 50 A ˚ -thick TPD films if the injection were due to ballistic ponents from the 24-A transport. The transport analysis is further complicated by clustering and incomplete coverage of the TPD film discussed earlier. In the absence of a well-defined film thickness at low coverages, it is somewhat difficult to do any quantitative analysis based on transport effects. The semiquantitative analysis more closely fits the second scenario wherein the ultimate challenge is estimating the effective transfer distance and understanding the nature of the transfer states. There exists at least a few experimental approaches that might help to elucidate the listed processes. One is measuring the temperature dependence of the slow component, which ought to help to distinguish between the transport and the tunneling processes. Namely, the latter is temperature independent at low temperatures but shows temperature dependence at high temperatures [42]. The other approach is creating overlayers whereby the donor–acceptor distance is strictly controlled. This is possible by attaching anchor groups to the donor molecule, which are then kept at the distance of the anchor length from the acceptor surface. These two approaches were not available in the experiments reported here and therefore need be considered as proposals for future investigations.

IV.

CONCLUSIONS

We have performed femtosecond time–resolved photoemission studies on the relaxation and interface charge transfer dynamics of two representative organic materials for photoreceptor applications, DPEP and TPD. Careful data analysis and combination of results obtained with TR-PES, UPS, IPES,

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and optical spectroscopy reveal intriguing dynamic behavior of these systems and shed light on the understanding of charge transfer processes across interfaces in organic semiconductors. The TR-PES results show that the initial relaxation of a photoexcitation in thin DPEP films is extremely rapid. A typical relaxation time is measured to be about 30–35 fs at 2.5 eV above the HOMO level. We attribute the fast relaxation to a rapid internal conversion process from S2 and S1 levels to the vibrational manifolds of S1 and S0 states and the vibrational relaxation of the latter. An interesting difference is observed between the energy dependence of the relaxation time in polycrystalline and amorphous films. The relaxation rate is almost twice as fast at lower electron energies in the polycrystalline films than it is in the amorphous films. We attribute this difference to fast transitions within the manifold of extended states that are shown to form in ordered polycrystalline films. This leads us to conclude that intermolecular interactions are important for the relaxation of excitations in DPEP thin films and are likely to affect the photocharge generation efficiency of the films. The TR-PES spectra of the bilayer structures reveal a fast two-component electron injection from TPD to DPEP. The short component, with duration less than 100 fs, is assigned to ultrafast ET from the excited TPD manifold to the unoccupied DPEP manifold. The ET rate is dominated by the excess energy of the electron on TPD, and we were able to determine the cutoff energy for the ET process at 3.1 eV above the DPEP HOMO using the formalism based on the classical description of average FranckCondon densities. The long component, with a pronounced dependence on the TPD overlayer thickness, is attributed to disorder-limited injection with ˚. the effective donor–acceptor distance of 12 A

ACKNOWLEDGMENTS The authors gratefully acknowledge the support from the National Science Foundation for Science and Technology Center Grant (No. CHE-9120001), DMR-9982988, and the Xerox Foundation. A. J. M. acknowledges support from the Va¨isa¨la¨ Foundation of the Finnish Science Academy.

REFERENCES 1. 2. 3.

B. O’Regan, M. Gra¨tzel, Nature 353:737, 1991. A. Hagfeldt, M. Gra¨tzel, Chem. Rev. 95:49, 1995. J. M. Lanzafame, R. J. D. Miller, A. A. Muenter, B. A. Parkinson, J. Phys. Chem 96:2820, 1992.

Charge Transfer Dynamics in Organic Photoreceptors 4. 5. 6. 7. 8. 9. 10. 11.

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14 STM-Excited Electroluminescence and Spectroscopy of Conjugated Polymers Santos F. Alvarado Zurich Research Laboratory, IBM Research, Ru¨schlikon, Switzerland

I.

INTRODUCTION

The discovery of conjugated polymers has recently led to the development of high-efficiency organic light-emitting devices (OLEDs) [1,2], and many experimental and theoretical investigations have been undertaken to elucidate the underlying fundamental physical processes. Knowledge of the relative alignment of the energy levels at interfaces between organic materials and an injection electrode is crucial to an understanding of the device operation and, in particular, the physics of charge-carrier injection, transport, and radiative recombination, with the ultimate goal of improving electroluminescence (EL) efficiency. The alignment of the highest occupied molecular orbital (HOMO) of the organic materials of an OLED is usually estimated from the difference between the ionization potential (IP) of the polymer and the work function of the electrode, i.e., by using the SchottkyMott (SM) rule [3], which assumes that the energy levels correspond to a common vacuum level. This approximation thus neglects surface, interface, and charge transfer effects. The alignment of the lowest unoccupied molecular orbital (LUMO) is then estimated by adding the optical absorption edge energy to the HOMO energy, a procedure that neglects the exciton binding energy. This method for constructing OLED energy diagrams is a useful first approximation but can lead to significant errors. The applicability of the SM 473

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rule is a topic of current debate. For instance, studies carried out on various conjugated polymers and small-molecule films using techniques such as ultraviolet photoemission spectroscopy (UPS) [4–9], internal photoemission [10,11], and scanning tunneling microscope (STM) spectroscopy [12–15] show clear evidence of significant deviations from the SM rule for organic– metal as well as for organic–organic interfaces. However, other reports of UPS studies [16,17] on thin films of precursor poly( p-phenylenevinylene) (PPV) or of poly(9,9-dioctylfluorene) (PFO) deposited on substrates with different work functions show that, at least in certain cases, the SM rule holds quite well and that the vacuum level misalignment at the polymer– electrode interface is only a few hundred meV. Furthermore, studies show that the actual gap for charge-carrier injection into organic materials can be significantly large than the optical band gap [10,11,15,18; see also Ref. 19 and references therein]. The excess energy is attributed to the binding energy of the singlet exciton [20]. Ba¨ssler and coworkers [21] have proposed that the absorption and luminescence spectra of materials like PPV and its derivatives can basically be described by a molecular exciton model wherein charge carriers are single-charged polaronic states instead of free charge carriers [22,23]. There is much experimental support for the exciton model; see, e.g., Ref. 24 and references therein. The controversial and often debated issue is whether the emission of light from PPV is from the decay of a weakly bound exciton, of binding energy below a tenth of an eV, or from a tightly bound exciton of binding energy as high as 1 eV. Note that in conjugated polymers the exciton diffusion length and luminescence efficiency are strongly influenced by the exciton’s lifetime and binding energy. Here an overview is given on how STM-based techniques can be employed to probe the electronic and transport properties of organic materials. These studies allow the above controversies to be resolved. One advantage of using STM spectroscopy techniques, besides their high spatial resolution, is that they can be used to probe directly not only the occupied but also the empty states of organic materials. A technique that allows the alignment of the empty levels to be determined in a direct way is particularly welcome because photoemission spectroscopy techniques have so far proved to be useful only to probe the occupied states of organic materials. In this chapter, an in-depth description of STM-based spectroscopy techniques is given, along with results obtained with thin films of PPV and some of its derivatives as well as with thin films of polyfluorene (PFO) are presented. The materials investigated are spin-coated on Au or indium-tin-oxide (ITO) substrates.

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CHARGE-CARRIER INJECTION FROM THE TIP OF AN STM INTO ORGANIC MATERIALS

A basic condition in standard STM experiments is that the tunneling resistance RT = VT/iT has to be much higher than the resistive losses associated with current flow through the sample (spreading resistance). Here the charge injection takes place by means of tunneling through a vacuum potential barrier (Figure 1). As the applied bias voltage, VT, drops at the tunneling barrier, the charge-carrier injection energy can be tuned through the polarity and magnitude of VT. This is the basis of the spectroscopy technique for the determination of the density of states in the surface region by means of i– V curves collected with the STM feedback loop disconnected [25]. However, if the resistive losses within the organic material are allowed to become comparable to RT, the surfaces of the tip and the sample can come into close proximity and make physical contact, causing the vacuum barrier to collapse. In the following charge-carrier injection as well as the electric field distribution for this particular case are discussed, whereas subsequent sections show how STM in contact mode can be used to probe the alignment of molecular energy levels at organic interfaces as well as electronic excitations. We begin by noting that when the tip makes contact with the organic surface the Fermi level (EF) of the former is pinned at an energy within the forbidden energy gap. This gives rise to the formation of a Schottky barrier, and the bias voltage drops fully within the organic material. Figure 2 shows the geometry and energy diagrams for negative tip polarity (tip negative, VT < 0, relative to the anode’s Fermi level, E⬘F ) and for a magnitude of VT such that the Fermi level of the tip is above the threshold for injection into the lowest electron polaron state of the organic material. Note that the energy diagram corresponds to that of a single-layer planar OLED but with one significant difference: In STM experiments the strength of the electric field is much higher at the highly curved apex than at the planar substrate interface. In this configuration, injection of charge carriers takes place by tunneling through the Schottky barrier into polaron states of the organic material [26]. Charge injection from Schottky barriers into organic material has been the topic of several theoretical studies; see, for instance, Refs. 26 and 27. The ratio between the field at the apex and the fields at the substrate, ET and ES, is approximately [28] ET/ES ⯝ [(␬R ⫹ d)/␬R]n

(1)

where R is the radius of curvature of the tip apex, typically 5 nm up to a few 10 nm, ␬ is a modifying factor to account for the geometry of the real system, and d is the distance between the tip apex and the substrate (Figure 2).

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Figure 1 Schematics of (a) the geometry and (b) the energy diagram for a typical STM configuration, where charge injection into the organic material occurs by tunneling through a vacuum barrier.

For instance, ␬ > 1 for R << d due to the effect of the tip’s shaft, whereas ␬ < 1 for R ⯝ d due to the proximity of the planar substrate [28]. In the spherical-tip approximation n = 2. The average field at the injection spot is Em = (VT ⫺ Vbi)/d, where Vbi is the contact potential. It is clear that if the magnitude of VT is decreased while iT is kept constant, the STM feedback loop causes the tip to move into the sample, decreasing d to compensate for the decrease in RT. Here the requirement is that the tip motion compensate for the electric field reduction associated with the change in VT to maintain

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Figure 2 Schematics of (a) the geometry and (b) the energy diagram for tunneling with the tip in point-contact mode. Injection into the organic material occurs by tunneling across a Schottky barrier of thickness DB.

constant tunneling current flow across the barrier. Measurements of the relative tip displacement vs. the applied voltage (z–V curves) can be useful as a local probe of the electronic properties, as are the barrier for charge-carrier injection and the exciton binding energy (see below). In addition, a local qualitative characterization of the charge-transport properties of organic materials can be done by means of z–V curves [28]. The question arises as to the current balance for the point-contact injection geometry compared to the planar OLED case. At the tip apex the

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electric field strength is significantly higher than at the substrate interface, which translates exponentially into a much higher injection probability of charge carriers across the corresponding Schottky barrier from the tip than from the substrate. Consequently, the current passing through the tip–organic material–substrate system is expected to be predominantly monopolar. There are, however, some requirements to be met for this to hold, such as that R be small and that the barrier height for oppositely charged carrier injection from the substrate not be much smaller than that of the tip [28].

III.

ELECTROLUMINESCENCE GENERATION BY CHARGE-CARRIER INJECTION FROM THE TIP OF THE STM

Photon emission generated via injection of electrons from the tip of an STM can be used to obtain spectroscopic information about the physical and electronic structure of surfaces at nanometer scales. This powerful technique was used initially to image the light emission and to probe the electronic properties of single III–V quantum wells [29–32], quantum wires [33,34], and quantum dots [35]. Recently it has also been used in the study of various organic materials [12,13,28,36–39]. By applying a bias to the STM tip, charge carriers may tunnel directly into polaron states of the material. Upon combination with carriers of opposite polarity injected from the substrate, excitons are formed that can decay radiatively with the emission of light. The STM-tip/polymer/substrate structure mimics a conventional OLED but in a very localized fashion. Therefore, this type of experiment allows the EL efficiency to be mapped for a device-equivalent structure at near-nanometer dimensions. In the following, four kinds of STM-excited electroluminescence (STL) experiments performed on thin films of conjugated polymers are described: (1) simultaneous collection of topography and wavelength-integrated EL intensity (iL) maps; (2) measurements of iL vs. VT curves; (3) EL spectroscopy, in which wavelength-resolved spectra are collected at different spots of the sample; and (4) EL intensity decay with time for charge-carrier injection at a fixed point of the sample. All experiments are performed in constant-current tunneling mode. The experimental setup is shown schematically in Figure 3. The STM system is mounted inside an ultrahigh vacuum (UHV) chamber, and all experiments are performed at 10⫺10 mbar and at room temperature. EL is excited either by charge injection from the tip of the STM or in field emission mode for cathodoluminescence (CL). The CL measurements, performed at various spots of the thin film, are taken by raising the tip ⯝ 1 ␮m above the surface and applying a negative bias of a

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Figure 3 Experimental setup used for the EL excitation by charge injection with the tip of an STM. For wavelength-integrated intensity measurements, the optical fiber and the optical multichannel analyzer (OMA) system is replaced by an avalanche photodiode. (From Ref. 12.)

few hundred volts to the tip. Typically, the CL spectra are found to be very similar to the photoluminescence (PL) spectra of the material. EL is collected using a lens system and transmitted outside the chamber via a glass window. For the generation of wavelength-integrated emission maps, an avalanche photodiode is used to measure the intensity. To measure EL emission spectra an optical multichannel analyzer system is used. For the latter measurements the STM serves primarily to select and define the location of the experiment and the tip is held at a fixed spot during the collection of the spectra (see Ref. 12 for further details). By scanning the STM tip over the surface, photon emission can be recorded as a function of the injection point and compared with the conventional STM topographic image. The photon maps represent the local efficiency of electron vs. hole injection (charge-carrier balance) convoluted with the local EL efficiency of the material. Owing to the migration of injected charge carriers and of excitons, the radiative recombination process giving rise to light emission might occur at a different position from the injection point. The photon map corresponds to a map of local EL efficiency only if the charge-carrier balance is uniform over the surface, and the chargecarrier and exciton diffusion lengths are constant.

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Figure 4 shows a typical result comparing topographic and EL maps collected on a PPV thin film (a few 10 nm thick) deposited on a Au(111) substrate (cf. Ref. 12). Typically the polymer forms smooth structures without pin holes. The tunneling parameters were VT = ⫺4.6 V and iT = 340 pA. Under these conditions the EL intensity shows no significant degradation with time for the duration of the experiments (less than a factor of 2 decrease in intensity over several hours). The time required to collect an image is

Figure 4 (a) Topography and (b) EL maps obtained simultaneously on a PPV thin film deposited on a Au substrate. Note the correlation between topographic and EL features. Tunneling parameters: VT = ⫺4.6 V (tip negative with respect to the Au substrate), iT = 340 pA. (c) Profile of the EL intensity revealing large local variations of the EL efficiency over scales of a few nanometers. (From Ref. 12.)

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approximately 30 min. The luminescence intensity is neither homogeneous nor completely random with respect to the surface topography. As the tunneling current is constant, this image indicates that very large variations in the observed EL efficiency occurs over length scales of a few nanometers. In some regions the intensity map reveals a clear correlation with topographic features, whereas in others the luminescence is anticorrelated. Some topographic features, such as hillocks, appear dark (region A in Figure 4), whereas others appear bright (region B in Figure 4). The peak-to-peak fluctuations of the luminescence intensity is about two orders of magnitude. Although this luminescence map is quite typical, we note that the size of the bright and dark areas varies over the sample and that 100-nm-wide regions of homogeneous luminescence can be found. Experiments on PPV thin films (about 20 nm thick) deposited on ITO exhibit inhomogeneities in the EL intensity maps very similar to those shown in Figure 4. Figure 5 shows topographic and EL maps collected on a thin film of the soluble PPV derivative poly(1,3-phenylenevinylene-co-dioctyloxy-1,4phenylenevinylene) (PmPV-co-DOctOPV) spin-coat deposited on a Au(111) substrate [38]. Here VT = ⫺2.5 V and iT = 100 pA. Similar to the PPV sample, the thin film of this material exhibits smooth features and is free of pin holes. In this PPV derivative we again find that the EL intensity correlates with the topographic features. In addition, no significant degradation of the EL signal occurs during the measurements. It is apparent that the triangular region of the film emits much less light than the rest of the surface. Other regions in the photon map appear to be highly emissive but do not seem to correlate obviously with topographic features. The boundary between the light and dark regions is very sharp, with a change in EL intensity of more than two orders of magnitude occurring on a length scale below 2 nm. These results are very similar to those of the PPV sample above, indicating that this behavior may be common to both solution-processed as well as precursor-type polymers. The inhomogeneity of the local EL efficiency is primarily related to variations of the injection threshold as well as of the local electronic properties of the conjugated polymers (see below). The local electronic properties of semiconducting materials and their interfaces can be probed with STL by means of measurements of the wavelength-integrated EL intensity vs. VT [29,30,40] or of the EL spectrum [31,32,34]. With the former technique, measurements of the position of the energy of the conduction band edge of the bulk, just behind the depletion zone at the injection interface, of III–V semiconductors can be determined [30]. This technique has also been applied to probe the injection threshold of electrons [13] as well as of holes [28] for organic materials. However, in the contact mode experiments discussed here, the charge injection threshold, and thus the luminescence threshold, corresponds to the injection barrier of

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Figure 5 (a) Topography and (b) EL maps obtained simultaneously on a thin film of PmPV-co-DOctOPV deposited on a Au substrate. Note that, similar to the results for PPV/Au shown in Figure 4, the EL intensity is inhomogeneous over scales of a few nanometers. Tunneling parameters are VT = ⫺2.5 V, iT = 100 pA. The regions marked A and B are two positions where EL spectra were recorded; see Figure 6. (Reprinted with permission from Ref. 38. Copyright 1997 American Institute of Physics.)

the conjugated polymer at the interface with the substrate (see Figure 2 and Section IV below). In order to elucidate the origin of the strong variations of the EL efficiency, iL vs. VT curves are collected on different spots of a PPV/ITO sample (Figure 6). Near EL threshold the data follow the relationship: iL = k(eVT ⫺ EP⫺)l,

(2)

where k is a proportionality constant, EP⫺ is the energy threshold for electron

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Au(111) substrates (see below). It is also evident that the shape of the iL vs. VT curves can be different, even for curves exhibiting very similar injection thresholds. The latter result suggests that the large variations of the electronic properties of PPV can occur over the distance of a few nanometers and that these variations can partially account for changes in the EL efficiency. EL spectra recorded on spots with different EL efficiencies of the PmPV-co-DOctOPV show that the spectral signature of the EL, and thus the electronic properties of the polymer, exhibit local variations. With the STM tip positioned over two different regions of the surface marked A and B in Figure 5, the EL spectra shown in Figure 7a and b were recorded [38]. The spectrum from point A is very similar to the CL spectrum, but the spectrum taken at point B appears broad and featureless, and the 0–0 and 0–1 transitions are relatively suppressed and show significant changes in the local Franck-Condon (FC) distribution. A correlation between the EL efficiency and the FC distribution is found in CL experiments on PPV cured at different temperatures in UHV [12], which provides indirect evidence of the influence of the conformation of the molecular structure of the organic material and the luminescence efficiency. Such effects observed in both PL and EL from bulk films have previously been ascribed to the formation of interchain excimer states [42]. Alternatively, spectral variations may be due to other aspects of the local order of the polymer chains [43]. The observation of such variations in local spectra highlight the power of the STM-excited EL technique to act as a probe of the molecular environment, as the emission occurs from a limited number of molecules and is consequently less inhomogeneously broadened. It is necessary to consider also other phenomena that could induce variations in EL efficiency, such as emission quenching due to the presence of nearby metallic surfaces, which has been well documented [44–46]. This effect occurs at distances of about 20 nm, which is greater than the thickness of most of the thin films used in the STL experiments. But although a certain contribution of this effect might be present, it cannot be the origin of the sharp local variations of the EL intensity or of the anticorrelation observed between some topographic features and the EL intensity discussed above. Electric fields have been shown to quench EL by up to 30% for fields up to 2 ⫻ 106 V/cm due to field-induced exciton breaking [47]. For the present experimental parameters, electron injection is in contact mode, at least for low bias voltages. This means that, as discussed above, most of the applied potential drops near the apex, at the Schottky barrier. Thus, even though the mean field between the tip and the substrate is ⯝7 ⫻ 106 V/cm, the actual field in the zone where the recombination takes place, which is behind the tunneling barrier, is significantly smaller, as can be estimated from Eq. (1). Furthermore, there is no clear correlation between feature heights in the

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Figure 7 (a, b) EL spectra collected with the tip placed over points (marked A and B, respectively) of the topography/EL maps shown in (c). Cathodoluminescence spectrum generated by using the STM tip as a field emission source of electrons. (Reprinted with permission from Ref. 38. Copyright 1997 American Institute of Physics.)

topographs and in the luminescence intensity. Hence, it appears that variations in the local electric field are of only secondary importance for the variations in the EL maps. For a homogeneous material, charge carrier and exciton diffusion determines the volume from which EL will occur. In PPV this distance is about 6–8 nm [44]. It is possible that the low EL efficiency regions of the polymer arise from the particular local molecular packing favoring nonradiative decay channels. If the local nonradiative rates were higher, the lifetime of the excitons (and hence their diffusion length) would be reduced, effectively

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Figure 8 EL intensity decay with time. The short-term intensity fluctuations are typical of the STM-excited technique and are related to fluctuations of the tunneling barrier, due to tip–sample interactions.

pA. From Eq. (3), we extrapolate that for an OLED device operated at 10 mA/cm2, the half-lifetime would exceed tens of thousands of hours even for the typically claimed Coulomb aging case where x = 1.

IV.

MEASUREMENT OF THE BARRIER HEIGHT FOR CHARGE-CARRIER INJECTION

The measurement of z–V curves can be used as a spectroscopy technique to probe the density of states of the material in a manner that differs from the standard i–V spectroscopy technique [25]. More specifically, as the magnitude of the bias voltage is decreased, the tip penetrates the organic material and can inject charge carriers in the neighborhood of the substrate interface (Figure 2b). Obviously, charge-carrier injection into the organic material is possible up to a characteristic threshold voltage at which the Fermi level of the tip moves into the forbidden energy gap at the interface with the counterelectrode. At this point the tip penetrates deeply into the organic material, reaching the interface region where charge carriers can be injected into the substrate either by directly tunneling into it or via nonresonant tunneling through molecules in contact with it. This threshold voltage (Vth), typically identified by a sharp decrease in the slope (dz/dV) of the z–V curve, marks the onset of injection into the lowest electron polaron state for negative tip

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polarity, and is thus a measure of the potential barrier for electron injection at the buried interface of the organic material (Figure 2b). Repeating this gedanken experiment for positive tip polarity we can convince ourselves that the barrier for hole injection into polaronic states can be measured using the same procedure. Note that this technique works only for soft materials that yield to the pressure exerted by the tip. Figures 9–11 show results obtained on various conjugated polymers chosen for having different optical absorption gaps, namely, the soluble PPV derivative poly[(2-methoxy-5-dodecyloxy)-p-phenylenevinylene-co-p-phenylenevinylene)] [poly(1,12)AOPV-co-PPV], Ea ⯝ 2.34 eV; PPV, Ea ⯝ 2.4 eV; and PFO, Ea ⯝ 2.9 eV. The optical absorption spectrum of each material is shown in Figures 9–11. The PPV sample was prepared in situ under UHV conditions by thermal conversion of a nonionic n-octylsulfinyl precursor polymer soluble in toluene [49]. All soluble materials were deposited on Au(111) substrates under ambient conditions. The z–V curves were collected under UHV conditions; see Refs. 14 and 15 for details. Control measurements with uncoated gold surfaces show that the tip–surface separation increases gradually with increasing voltage, with a maximum increase of approximately 0.6 nm over the bias range of 0.3–3 V, shown as dotted lines in Figures 9–11. Note that the difference between the two curves is a measure of the distance between the tip and the substrate, and, at high voltages, of the approximate local thickness of the polymer film. The point at which the z–V curves diverge, which we determine by linear extrapolation, corresponds to the threshold voltage for charge-carrier injection into polaronic states, P⫺ and P⫹ of the conjugated polymers. Typically, the transition can be determined to within approximately 100 meV. In addition, we observe a corresponding increase in the junction noise for injection into the polymers relative to direct injection into the substrate. This increase in junction noise is related to fluctuations of the tunneling barrier due to tip–sample interactions, which are larger for soft materials. Table 1 summarizes the values of the injection barriers for electrons and holes, EP⫹, EP⫺, determined from a series of measurements on the various samples. The STL result for PPV on ITO is also included in Table 1. Note that because the injection barrier for holes and electrons is measured directly in the STM experiments, no knowledge is necessary of the IP, electron affinity, or the work function of the substrate. We now compare the STM results for the barrier for hole injection with the values obtained using the SM approximation: EP⫹(SM) = IP ⫺ ␾, where ␾ is the work function of the substrate. The difference ⌬ = EP⫹(SM) ⫺ EP⫹ can be taken as a measure of the ‘‘shift’’ between the vacuum levels of the two materials, where ⌬ = 0 in the SM approximation. It should be noted, however, that in order to establish whether deviations from the SM

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Figure 9 (a) z–V curves collected on a thin film of poly(1,12)AOPV-co-PPV deposited on a Au(111) substrate. The dashed curve represents the typical tip displacement measured for the bare, clean Au(111) substrate. The threshold for charge-carrier injection into electron (P⫺) and hole polaron (P⫹) states is marked by the vertical dashed-line segments. (b) Optical absorption spectrum of the organic material. (Adapted from Ref. 15.)

rule occur for a given interface, it is necessary to know exactly the IP of the conjugated polymer. This is important because the optimization of OLED devices requires a good theoretical description of metal–polymer contacts. The IP of free organic surfaces can be measured by UPS. However, despite the extremely clean UHV working conditions under which these experiments are usually conducted, the results of different laboratories diverge substantially. For PPV, for example, values in the range of 5.2 eV [50] to 6.2 eV

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Figure 10 Same as Figure 9 but for PPV/Au(111). (Reprinted from Ref. 14. Copyright 2000, with permission from Elsevier Science.)

[51] have been reported. For the modeling of OLEDs this translates into an uncertainty of 1 eV in the height of the injection barrier at PPV interfaces. These apparent discrepancies can be related in part to the well-known variation of material properties found for a given substance even when prepared by the very same research group. With this caveat in mind ⌬ is calculated (see Table 1) by taking the values of IP reported by various research groups [17,18,50,52,53]. For a comparative review of the different techniques used to determine IP, see, for example, Ref. 54. The work function of Au(111), ␾(Au) = 5.31 eV [55], is taken although we may not have examined an ideal interface, as the substrates have to be exposed to ambient conditions for spin-coating. The results show that even though there might be substan-

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Same as Figure 9 but for PFO/Au(111). (Adapted from Ref. 15.)

tial uncertainty regarding the values of IP for the various conjugated polymers, the values of ⌬ show a significant tendency toward negative values, i.e., the molecular orbital of the polymers are clearly shifted to lower binding energies with respect to the predictions of the SM rule. Similar results have been found using other experimental techniques for various organic compounds in contact with metallic surfaces [4–10]. This shift indicates the existence of image forces [56] or of interface charge transfer effects that lead to the formation of dipole layers. Note that the effect of image forces is expected to be important in materials with internal charge separation. This appears to be the case for the alkoxy-PPV derivatives for which the shift is largest. Thus, the actual alignment is determined by physical and chemical interactions at the substrate–polymer interface, the nature of which is additionally influenced by atmosphere and solvent exposure. All these effects

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are neglected in the common-vacuum-level model. Thus, although we may not have examined an ideal interface, our results provide insight into electrode conditions generated in actual OLED fabrication. In addition, the results in Table 1 show that the hole injection barrier of MEH–PPV is lower by about 100 meV than that of PPV, whereas for poly(1,12)AOPV-co-PPV any difference would be within the experimental error. The trend observed between PPV and MEH–PPV is in the direction predicted by calculations of the influence of the substituents on the IP of the material [57]. However, this effect is smaller in magnitude: the theoretical prediction is a difference of 300 meV between the IPs of these two polymers. V.

PROBING THE EXCITON BINDING ENERGY

The combination of a hole polaron and an electron polaron, of energy EP⫹ and EP⫺, results in the formation of an exciton. The exciton binding energy has been defined [58] as Eb = Egsp ⫺ Ea, where Egsp, the difference between the electron and hole polaron energies, is called the single-particle energy gap [11] and Ea is the energy required to create a molecular exciton, which can be determined from optical absorption spectra. Figure 12 illustrates this definition. Here HOMO and LUMO represent the one-electron band picture, which neglects Coulomb and exchange interactions as well as molecular relaxation effects. Supplying the exciton with energy Eb gives rise to the creation of a pair of oppositely charged polarons. The radiative decay of a singlet exciton results in emission of a photon. It is important to note that Eb is the sum of the Coulomb interaction and the energy difference between the molecular configurational relaxation of the exciton and of the indepen-

Figure 12 Energy diagram illustrating the definition used for determining the exciton binding energy, Eb = Egsp ⫺ Ea. The quantity Egsp is the single-particle band gap determined from z–V curves, see Figure 3; Ea is the optical absorption threshold. (From Ref. 15.)

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dent polarons. It can be seen that there is no intrinsic dependence of Eb on the band gap. Note that the z–V curves allow one to measure the threshold for charge-carrier injection into the P⫹ and P⫺ levels. Therefore, in the present analysis we use the optical threshold, Ea, to determine Eb, which minimizes errors due to inhomogeneous broadening and is thus more accurate than taking the energy of the first absorption transition, as done in Ref. 11; cf. also Ref. 58. Combining the STM data with the optical absorption data, we obtain the results given in Table 2. Figure 13 illustrates the correlation between Eb and Ea. For PPV and its derivatives, Eb has been determined by a number of other different experimental techniques, and values between 200 and 800 meV have been obtained [10,11,19,58–69]. However, extreme values as low as 25 meV [23] or as high as 1.1 eV [64,70] have been reported. There have also been several theoretical efforts to estimate Eb. Chandross et al. [64] predict that Eb = 0.95 ⫾ 0.15 eV and Rohlfing and Louie report a value of 0.9 eV for an isolated PPV chain [71], whereas Lavrentiev et al. [72] give a value of 0.6 eV for an infinite chain. However, as in those calculations the system is quasi-one-dimensional, interaction is very strong, and the exciton energy is expected to be significantly higher than in bulk material. Recent calculations by van der Horst et al. [73], where long-range (bulk) screening is introduced by use of the bulk dielectric constant, yield a value of 0.54 eV. On the other hand, Monte Carlo calculations reported by Scheidler et al. [63] yield a value of 0.3 eV for PPV. The challenge for the calculations of the electronic properties of conjugated polymers is the inclusion of geometrical relaxation effects. The STM spectroscopy results obtained on PPV and alkoxy-PPV derivatives clearly favor the intermediate range of values, i.e., of 300–600 meV (Table 2). The results on MEH–PPV agree well with the exciton binding energy of 0.35 eV deduced in Refs. 15 and 58 from internal photoemission results [11]. Note also that the value of Eb for the PPV is found to be significantly grater than that of the alkoxy-PPV compounds. Moreover,

Table 2

Exciton binding energies, Eb, of various conjugated polymers

Polymer

Egsp = EP⫺ ⫺ EP⫹ (eV)

Ea (eV)

⫾ ⫾ ⫾ ⫾

2.1 2.34 2.4 2.9

MEH-PPV Poly(1,12)ADPV-co-PPV PPV PFO

2.38 2.65 2.88 3.2

0.13 0.1 0.14 0.1

Egsp, single-particle gap; Ea, optical absorption edge.

Eb (eV)

Ref.

⫾ ⫾ ⫾ ⫾

14 15 14 15

0.28 0.36 0.48 0.3

0.13 0.1 0.14 0.1

STM-Excited Electroluminescence and Spectroscopy

Figure 13 polymers.

495

Exciton binding energy vs. optical absorption gap of various conjugated

measurements on methyl-substituted PPV (Me-PPV) yield Eb ⯝ 0.6 eV [74]. It has been proposed [14] that the lowering of Eb in the alkoxy-substituted compounds is induced by the effect of the substituents on (1) the effective polaron-exciton Coulomb contribution to the confinement potential and (2) the molecular conformational relaxation energy. The latter contribution to the exciton binding energy can be nonnegligible owing to electron–phonon coupling of relatively high vibrational energies. The analysis of luminescence spectra shows that the energy of some of the vibrational modes, particularly those that are active in fluorescence, shift to lower energies upon alkoxy substitution, suggesting some influence on the exciton binding energy in the polymer [14; see also Ref. 75]. This is plausible in terms of theoretical models where the structural relaxation energy contribution to exciton binding energy is nonnegligible [76]. Moreover, the Coulomb interaction contribution to Eb is significant owing to the low dielectric constant of the materials, ␧ = 3–4 [47]. Thus, the lower Eb of alkoxy-PPVs relative to PPV might also indicate a weaker Coulomb contribution to Eb in the former materials. This is reasonable because the substitution of hydrogen atoms in pristine PPV by alkoxy radicals will increase the polarizability in the immediate neighborhood of the exciton, leading to a reduction of the Coulomb interaction. We note that the reduced binding energies in alkoxy PPV compounds relative to PPV or Me-PPV cannot be explained by the smaller conjugation length in the alkoxy compounds [75]. The reason is that the binding energy

496

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of the exciton would actually be expected to increase with decreasing number of phenylene units [19,72]. Regarding PFO, the blue-emitting polymer that has recently attracted much interest [77], less research work on its electronic properties has been reported than for PPV and its derivatives. In a recent cyclic voltammetry study, Janietz et al. [18] find that the energy gap between the IP and the electron affinity is about 0.76 eV greater than the optical absorption edge. However, owing to solvation effects, it is not straightforward to deduce the exciton binding energy from that result for comparison with the value of 0.30 ⫾ 0.10 eV deduced from the z–V measurements. Lately, the electronic structure and the electron energy gap of PFO deposited on Au were studied by UPS and X-ray photoemission spectroscopy [78]. There it is reported that the electron energy gap determined from the shake-up peak of the Cls core level is at least 0.2 eV greater than the optical absorption edge of this material.

VI.

CONCLUDING REMARKS

STM-excited electroluminescence and z–V spectroscopy techniques have been employed to probe charge injection and electronic properties of conjugated polymers, which are of importance for OLED applications as are PPV, PPV derivatives, and PFO. STM imaging and electroluminescence maps are used primarily to select and define the location of the investigated region of the sample with nanometer resolution. These techniques allow one to probe (1) the barrier heights for injection of positive and negative charge carriers across interfaces, (2) the energy gap for single charge-carrier injection, and (3) the exciton binding energy. The results show that the single-particle band gap is significantly larger than the optical band gap. This is at odds with the commonly held practice of determining the threshold energy for electron injection by adding Ea to the IP energy. Relative to Ea the excess energy for injection of charge carriers of both polarities into the organic materials is given by the exciton binding energy. It is found that it lies in the intermediate range of Eb = 300–600 meV, and, as shown for PPV and its derivatives, it depends on the nature of the substituent. Furthermore the height of the barriers for electron and hole injection, respectively, into the polymers reveals significant discrepancies with the predictions of the commonly used SM rule, indicating the effect of image forces and the formation of dipole layers at the organic–metal interface. In addition, it is found that the EL efficiency of the polymers can exhibit variations over length scales of a few nanometers and that these

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Figure 1 Schematic band diagram of a polymer light-emitting diode (PLED) under forward bias. Indicated are the three basic processes that govern the performance of the PLED: injection, transport, and recombination. The conjugated polymer used in this study is poly(dialkoxy p-phenylenevinylene), whereas ITO and Ca serve as the anode and cathode, respectively.

which match up with lowest unoccupied molecular orbitals of PPV, must be used. It is evident that control of the electrode–polymer interfaces is very important with regard to the device performance of a PLED. The nature of metal–polymer interfaces and their formation have been studied both experimentally and theoretically [2]. The understanding of the metal–polymer interface is complicated by the occurrence of chemical interactions, which depend on both the metal and the polymer involved; the cleanliness of the materials; and the vacuum system used in the evaporation process [2]. However, in spite of these chemical interactions, measurements of the built-in electrical field have demonstrated that the energy barriers at the electrode– polymer interfaces still scale with the work functions of the electrodes [3]. Thus, the contact barrier for a given electrode–polymer combination is still a relevant parameter for estimating the role of charge injection. It is crucial for the understanding and optimization of polymer LEDs to obtain an answer to the question of whether the device characteristics of polymer LEDs are controlled by injection or transport. Experimentally, the discrimination between charge injection and charge transport on the basis of the device characteristics of a PLED has proven to be difficult. Initially, from the absence of a thermally activated behavior as well as a highly nonlinear J–V characteristic at high electric fields, it has been suggested [4,5] that the dependence of J on V resembles that of Fowler-Nordheim (F-N)

Contact-Limited Hole Current in PPV

507

tunneling through a barrier. From the slope of the J–V characteristics in an F-N plot [ln(J/E2) vs. E⫺1)], energy barriers of 0.1 and 0.2 eV were deduced for the Ca/PPV and ITO/PPV interfaces, respectively [5]. However, quantitatively the currents predicted by the F-N theory for these low-energy barriers exceed the experimentally observed currents by several orders of magnitude. Furthermore, it has been demonstrated [6] that the hole currents in our PPV-based PLEDs are determined by the bulk transport properties of the PPV and not by charge injection.

II.

MODEL CALCULATION

As a first estimate of whether the current for a given polymer–electrode combination will be injection or transport limited, the effect of an injection barrier ␸b on the J–V characteristics of a PLED has been modeled [6]. However, a major problem for setting up such a model calculation is the fact that the knowledge of the mechanisms of charge injection into conjugated polymers is not nearly as comprehensive as for inorganic semiconductors [7]. Therefore, let us first review the conventional charge injection mechanisms and then discuss their application in case of a conjugated polymer. First, we apply a bias V across a device with thickness L, which results in a current density J. As a result of the injection of charges at x = 0, the dependence of the electric field on the position x (between 0 and L) is given by [8]: E(x) =

冑

E(0)2 ⫹

2Jx ␧0␧r␮

(1)

with E(0) is the electric field at the contact, ␮ the charge carrier mobility, and ␧0␧r the permittivity of the semiconductor. The second term in the righthand side of Eq. (1) takes into account the effect of space charge in the device. For a device that is strongly contact limited, only a small number of charge carriers are injected, which results in a low current. As a result, E(0) >> 2JL/␧0␧r␮ and the electric field E(x) = E(0) = V/L is constant over the device. For efficient charge injection E(0) ⱕ JL/␧0␧r␮ and space-charge effects have to be taken into account.

A.

Ohmic Contact

For an Ohmic contact, E(0) = 0 and the J–V behavior is directly given by [9]:

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J=

9 V2 ␧0␧r␮ 3 8 L

(2)

This quadratic dependence of J on V is characteristic for the so-called spacecharge limited current (SCLC). In Figure 2 the calculated bulk SCLC is shown (dashed line) for a device with L = 100 nm and ␮ = 10⫺10 m2/Vs, which is typical for PPV. B.

Contact Limited Current

The current density–voltage (J–V ) characteristics for thermionic emission of charge from a metallic electrode into an insulator are described by the Richardson–Schottky equation [10]: J = A*T 2 exp

冉 冊 ⫺q␾b kT

(3)

where A* is the Richardson constant (4␲qm*k 2/h3), T the temperature, and ␸b the barrier height arising from the band offset between the insulator and the electrode. The barrier height ␸b is lowered by the image force effect:

Figure 2 Calculated J–V characteristics (solid line) for a device with thickness L = 100 nm, an injection barrier height ␸b0 = 0.35 V, a mobility ␮ = 10⫺10 m2/Vs, and a dielectric constant ␧r = 3. For comparison, the bulk-limited SCLC [dashed line, Eq. (5)], the diffusion-limited injection [dash-dotted line, Eqs. (2) and (3)], and the Fowler-Nordheim current [dotted line], assuming F(x) = F(0) = V/L are included as well.

Contact-Limited Hole Current in PPV

␾b = ␾b0 ⫺ ⌬␾ = ␾b0 ⫺

509

冑

qE(0) 4␲␧0␧r

(4)

where E(0) is the electric field at the contact. It has been pointed out by Simmons that this expression is invalid in case of an insulator with a low carrier mobility [11]. For this special case, a large amount of charge will build up at the contact such that back-diffusion from the insulator to the metal will occur. In this diffusion limited regime the J–V characteristics are given by [11]: J = qp(0)␮E(0) = qNc exp

冉 冊 ⫺q␾b kT

␮E(0)

(5)

where p(0) is the charge carrier density at the contact, ␮ the charge carrier mobility, and Nc the effective density of states. For a given current J, Eqs. (4) and (5) directly provide the boundary condition for the electric field E(0) at the injecting contact. By integration over x using Eq. (1), the J–V characteristic can be obtained for arbitrary ␸b0. However, the main result of Eq. (5) is that in case of a low-mobility material it is very difficult to disentangle contact-limited and bulk-limited processes. The charge-carrier density p(0) depends on the injection process, whereas the mobility, which determines the velocity of the charge carriers away from the contact, is a bulk property of the semiconductor. Therefore, for conjugated polymers, which exhibit extremely low mobilities as compared to conventional semiconductors [6], it is expected that besides the injection process the mobility still plays an important role in contact-limited devices. In order to investigate the role of an injection barrier ␸b0 in a low mobility material we have calculated the J–V characteristics using Eqs. (1), (4), and (5) for a device with L = 100 nm, ␮ = 10⫺10 m2/Vs, ␸b0 = 0.35 V, and ␧r = 3, as shown in Figure 2 (solid line). For comparison, the limiting cases for an Ohmic contact [E(0) = 0], implying SCLC according to Eq. (2), and diffusion-limited injection, Eqs. (4) and (5) with E(x) = E(0) = V/L, are also given. At low bias (V < 5 V) the calculated current is approximately equal to the diffusion-limited injection current, and at high bias (V > 15 V) the current approaches the bulk SCLC current. This transition from contact-limited current to bulk-limited current at high fields results from barrier lowering [Eq. (4)], which then eliminates the role of the injection barrier. Also plotted in Figure 2 is the result for F-N tunneling with E(0) = V/L. It is clear that at low bias this injection mechanism is negligible with regard to diffusion-limited injection, whereas at high bias space-charge effects dominate the conduction. In general, it will depend on both ␸b0 and L whether the conduction is injection or space-charge limited. By repeating

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these calculations for various injection barriers ␸b0 , it appears that for ␸b0 ⱕ 0.2 V the J–V characteristics are given by Eq. (2) in the whole bias regime. Thus, from this simple estimate we have obtained that an electrode with an injection barrier of ⱕ0.2 V at the electrode–PPV interface is able to supply the bulk SCLC. Similar conclusions have been drawn from model calculations by Malliaras and Scott [12] and Campbell et al. [13] who reported that bulk-dominated behavior in PLEDs is expected for contact barriers lower than 0.3 eV or 0.4 eV, respectively. From the first simple estimates described above it became clear that the low charge carrier mobility of the PPV-based materials enhances the occurrence of space-charge effects in a PLED. Consequently, knowledge about the mobility is indispensable in disentangling the contributions from the bulk (mobility) and the electrode (carrier density) to the contact-limited currents in low-mobility materials. Thus, a systematic experimental study of contact-limited currents in conjugated polymers requires the use of both Ohmic contacts (for determination of the mobility) and blocking contacts (for evaluation of the injection mechanism). It should be noted that the equations used in the calculations for charge injection so far have been derived for inorganic semiconductors with sharply defined band edges. For organic semiconductors with an energy distribution of localized states the injection mechanism is expected to be more complicated [14,15]. Alternatively, a model of thermally assisted tunneling of carriers from the contact to localized states of the polymer has been formulated [16]. This model has been further investigated by including energetic disorder and image force effects in a Monte Carlo simulation [14,17]. These simulations indicate that in conjugated polymers an increase of J with V is mainly due to the field dependence of the mobility as well as an additional increase of the carrier density at the contact due to image force effects [17]. In a recent study by Arkhipov et al. [18], charge injection from metallic electrodes into a random hopping system has been described. The mechanism consists of injection from the Fermi level of the electrode into the tail states of the distribution of hopping states of the disordered conductor, followed by either a diffusive escape from the attractive image potential or return to the electrode. In the present study we use the conventional injection model [Eqs. (4) and (5)] to analyze the experimental injection-limited currents in PPV. In this approach the presence of localized states in the PPV is taken into account by the use of an empirical charge carrier mobility, which represents the charge transport in a disordered hopping system. From the experiments we then determine the field and temperature dependence of the metal–polymer interface charge carrier density p(0). The evaluation of p(0) for various

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electrodes, fields, and temperatures then provides direct information about the charge injection mechanism in PPV, which is the subject of this study.

III.

EXPERIMENTAL

The J–V characteristics of holes in PPV are investigated in both the bulk and contact-limited regime using various hole-injecting electrodes. The devices under investigation consist of a single polymer layer sandwiched between two electrodes on top of a glass substrate. The polymer is a soluble poly(dialkoxy p-phenylenevinylene) [6] that is spin coated on top of a patterned ITO electrode. As a top electrode we use evaporated Au, Ag, Al, and Cu contacts. From cyclovoltametric experiments it has been shown that the hole transport states of our PPV are located approximately 5.3 eV below the vacuum level (0 eV). From electroabsorption measurements on poly[2methoxy,5-(2⬘-ethylhexyloxy)-1,4-phenylenevinylene) (MEH-PPV), which is very similar to our PPV, an energy of ⫺5.35 eV has been obtained [19]. The work function of our ITO amounts to 4.8 eV, from which a contact barrier for hole injection of about 0.5 eV is expected. Figure 3 schematically shows a band diagram indicating the work functions of the various electrodes

Figure 3 Band diagram indicating the work functions of the electrodes used in this study. Under forward (⫹) bias, holes are injected from the ITO and ITO* electrode; for negative bias (⫺), holes are injected from the evaporated Al, Ag, Cu, and Au electrodes, respectively. As a result of the large energy barriers with regard to the conduction band, electron injected can be neglected.

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Figure 4 Contact potential difference (CPD) between the ITO and an Au reference electrode vs. time without (dashed line) and after UV/ozone cleaning of the ITO (solid line).

used in this study with regard to the transport states of PPV. By using a UV/O3 or an O2 plasma cleaning treatment the work function of ITO can be increased [20,21] by up to 0.8–0.9 eV, as shown in Figure 4. These measurements have been performed using a Kelvin probe in a clean-room environment. The oxidation of the ITO surface is responsible for the observed work function increase [20] and such a well-cleaned ITO contact is hereafter denoted as ITO*. An additional difficulty when studying hole injection from evaporated metallic contacts into conjugated polymers is that some metals, such as Ca and Al, chemically interact with the polymer [22–24], which may complicate the charge injection process. However, noble metals deposited on organic conductors form abrupt unreacted interfaces [25]. The J–V measurements are performed in a nitrogen atmosphere in the temperature range 200–300 K. For a positive bias the holes are injected from the bottom ITO or ITO* electrode, whereas a negative bias implies hole injection from the evaporated metal electrode. Furthermore, in the measured voltage range there is no light emission from the devices, indicating that the devices can be regarded as hole-only devices. A.

Ohmic Contact

For a well-cleaned ITO* contact it has been demonstrated that the J–V characteristics are (bulk) limited by space-charge effects [6]. This observation demonstrates that for an ITO* contact on PPV the hole current is completely governed by the bulk conduction and not by charge injection. In

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Figure 5 the J–V characteristics for holes injected from the ITO* contact into the PPV are shown at various temperatures. The occurrence of SCLC in our ITO*/PPV devices enables us to directly obtain the E and T dependence of the hole mobility in PPV [26], given by:

␮p (E, T) = ␮0 (T)exp(␥兹E) with

冉 冊

␮0 (T) = ␮0 exp ⫺ and

␥=B

冉

⌬ kBT

冊

1 1 ⫺ k0T k0T0

(6)

(7)

(8)

with ⌬ = 0.50 eV, B = 3.1 ⫻ 10⫺5 eV (V/m)⫺1/2, T0 = 420 K, and ␮0 = 1.0 ⫻ 10⫺2 m2/Vs. The corresponding calculated SCL J–V characteristics are also shown in Figure 5. This functional form of the field E and temperature T dependence of the carrier mobility [Eq. (6)] is an intriguing feature of disordered organic semiconductors. The stretched exponential form has first been observed for poly(N-vinylcarbazole) by Gill in 1972 [27]. Numerous

Figure 5 Forward hole current density J vs. voltage V of an Ohmic ITO* contact using a ITO*/PPV/Cu hole-only device with thickness L = 300 nm at various temperatures. The calculated J–V characteristics for positive bias following the SCLC model, using the field-dependent mobility [Eqs. (6)–(8)] are plotted as solid lines.

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experimental studies on molecularly doped polymers, pendant group polymers, and amorphous molecular glasses have revealed a similar behavior [28–30]. A quantitative microscopic interpretation of this ubiquitous mobility requires a model for the charge transport in these materials. Charge transport in disordered organic conductors is thought to proceed by means of hopping in a Gaussian site–energy distribution. This density of states (DOS) reflects the energetic spread in the charge transporting levels of chain segments due to fluctuation in conjugation lengths and structural disorder. Ba¨ssler and coworkers [31,32] have performed numerical simulations of charge transport in a regular array of hopping sites with a Gaussian distribution of site energies. In this Gaussian disorder model (GDM), the following functional dependence of ␮ has been proposed [31]:

冋 冉 冊 冉冉 冊

␮GDM = ␮⬁ exp ⫺

2␴ 3kBT

2

⫹C

␴ kBT

2

冊 册

⫺ 2.25 兹E

(9)

with ␮⬁ the mobility in the limit T → ⬁, ␴ the width of the Gaussian DOS, and C a constant (depending on, say, the site spacing). Thus, the phenomenological parameters ⌬ and ␥ [Eq. (6)] may be related to the microscopic material parameter ␴. Applying Eq. (9) to the zero-field mobility of our PPV, as shown in Figure 6, yields for the width of the Gaussian DOS a

Figure 6 Zero-field mobility ␮0(T) vs. 1/T 2 (symbols) as obtained from the J–V characteristics of the ITO* contact. From a comparison with the Gaussian disorder model [Eq. (9)], a width of the DOS of 0.1 eV is obtained (solid line).

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value of ␴ = 0.11 eV. The GDM simulations reproduced the ␮ ⬀ exp(兹E) law only at relatively high fields (108 V/m). At low fields, agreement with experiments could be improved by taking into account spatial correlations between the energies of neighboring sites [33]. In an analytical one-dimensional model the form of the empirical mobility, as defined by Eqs. (6)–(8), has been attributed to energy correlations associated with charge–dipole interactions [34], which has been confirmed by simulations for three dimensions [35]. The modification of Eq. (9) by incorporation of correlation effects only slightly (<10%) affects the value of ␴ as