Physics and Technology of Thin Films IWTF 2003: Proceedings of the International Workshop, Tehran, Iran 22 February - 6 March 2003

Physics and Technology of Thin Films Editors A.Z. Moshfegh, H.v. Kanel S.C. Kashyap, M. Wuttig World Scientific Pro...

Author: A. Z. Moshfegh | H. V. Kanel | Subhash C. Kashyap | M. Wuttig

11 downloads 817 Views 32MB Size Report

This content was uploaded by our users and we assume good faith they have the permission to share this book. If you own the copyright to this book and it is wrongfully on our website, we offer a simple DMCA procedure to remove your content from our site. Start by pressing the button below!
Report copyright / DMCA form

DOWNLOAD PDF

Physics and Technology of

Thin Films

Editors

A.Z. Moshfegh, H.v. Kanel S.C. Kashyap, M. Wuttig World Scientific

Proceedings of the International Workshop on


Thin Films

This page is intentionally left blank

Proceedings of the International Workshop on


Thin Films I W T F

Tehran, Iran

2 0 0 3

22 February-6 March 2003

Editors A.Z. Moshfegh Sharif University of Technology, Iran

H.v. Kanel Politecnico di Milano, Italy

S.C. Kashyap Indian Institute of Technology-New Delhi, India

M. Wuttig I. Physikalisches Institut der RWTH Aachen, Germany

\[p World Scientific NEWJERSEY

• LONDON • SINGAPORE • BEIJING • SHANGHAI • HONG KONG • TAIPEI • CHENNAI

Published by World Scientific Publishing Co. Pte. Ltd. 5 Toh Tuck Link, Singapore 596224 USA office: Suite 202, 1060 Main Street, River Edge, NJ 07661 UK office: 57 Shelton Street, Covent Garden, London WC2H 9HE

British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library.

THE PHYSICS AND TECHNOLOGY OF THIN FILMS Copyright © 2004 by World Scientific Publishing Co. Pte. Ltd. All rights reserved. This book, or parts thereof, may not be reproduced in any form or by any means, electronic or mechanical, including photocopying, recording or any information storage and retrieval system now known or to be invented, without written permission from the Publisher.

For photocopying of material in this volume, please pay a copying fee through the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, USA. In this case permission to photocopy is not required from the publisher.

ISBN 981-238-770-6

Printed by Fulsland Offset Printing (S) Pte Ltd, Singapore

International Workshop on Physics and Technology of Thin Films (IWTF 2003) International Scientific Organizing Committee S.C. Kashyap, Indian Inst, of Technology-New Delhi, India A.Z. Moshfegh, Sharif Univ. of Technology, Iran (Organizer) M. Ohring, Stevens Inst, of Technology, USA G. Ottaviani, Univ. Degli Studi di Modena, Italy A. Zvezdin, General Physics Inst., Russian Academy of Science, Russia

Invited Speakers M. Farle, Gerhard-Mercator-Universitaet Duisburg, Germany A. Iraji-zad, Sharif Univ. of Technology, Iran H.v. Kanel, Politecnico di Milano, Italy S.C. Kashyap, Indian Inst, of Technology -New Delhi, India S.H. Keshmiri, Ferdowsi Univ., Iran S.K. Kulkarni, Univ. of Pune, India J.G. Lin, National Taiwan Univ., Taiwan M. Mirsalehi, Ferdowsi Univ., Iran A.Z. Moshfegh, Sharif Univ. of Technology, Iran (Organizer) M. Ohring, Stevens Inst, of Technology, USA A.I. Popov, Moscow Univ. of Electronics-MIET, Russia N. Radic, Ruder BoSkovic Institute, Croatia B. Rashidian, Sharif Univ. of Technology, Iran D. Rassi, Univ. of Wales, Swansea, UK H. Salamati, Isfahan Univ. of Technology, Iran P. G. Soukiassian, CEA and Univ. of Paris-Sud, France M. Vesaghi, Sharif Univ. of Technology, Iran M. Wuttig, Physikalisches Institut der RWTH Aachen, Germany Y. Zhuravlev, Univ. of Wales Swansea, UK

v

Local Organizing Committee (Sharif University of Technology) M. Akhavan, Department of Physics A. Amjadi, Department of Physics A. Arfaei, Department of Physics A. Barzegar, Public Relations Office G.H. Farrahi, Department of Mechanical Engineering M. Ghorbani, Department of Material Science and Engineering A. Ghorbanzadeh, Department of Physics A. Iraji-zad, Department of Physics F. Kaymaram, Department of Management and Economics S.M. Mahdavi, Department of Physics A.Z. Moshfegh, Department of Physics (Organizer) B. Rashidian, Department of Electrical Engineering M.A.Vesaghi, Department of Physics R. Zamani, Research & Technology Administration WORKSHOP SUPPORTERS Organizers Sharif University of Technology, Tehran, Iran Ministry of Science, Research and Technology, Iran Sponsors United Nations Educational, Scientific and Cultural Org. (UNESCO) The Abdus Salam International Center for Theoretical Physics (ICTP) Sharif University of Technology Ministry of Science, Research and Technology (MSRT) United Nations Development Projects. (UNDP) Ministry of Industries and Mines (MIM) Ministry of Communication and Information Technology (MCIT) Iran Air Center for International Research and Collaboration (CIRC) International Scientific Meetings Office (ISMO) High Technologies Organization (HTO) Technology Cooperation Office (TCO) Telecommunication Company of Iran (TCI) Electronic Components Industries (ECI) YarSanat Co. Ltd.

VII

Co-sponsors The Intl. Union for Vacuum Science, Technique and Appl. (IUVSTA) Iranian National Commission for UNESCO Advanced Manufacturing Research Center (AMRC) Industrial Development and Renovation Organization of Iran (IDRO) Information System of Iran (ISIRAN) Emad Semicon Co. (ESC) Pajouheshi Electron Co. Ltd. Iran Cutting Tools Mfg. Co. (TABA) Jam Ara Co. (JAC) Peres Sanco Co. Ltd. Contributors Ministry of Foreign Affairs, Iran Ministry of Culture and Islamic Guidance, Iran Isfahan University of Technology Isfahan Optical Industry (IOI) Iranian Academic Center for Educations Culture and Research (ACECR)-Sharif Branch

Tehran Sakkoo Co. Almasehsaz Company Farapajouhesh Co. Bimeh Iran Co. Laleh Hotel Inn Academy of Persian Language Literature Workshop Staff O. Akhavan M. Alempour A. Azarm R. Azimirad E. Bagheri M. Bashlideh H. Bayat L. Chahoshizadeh M. Dashti M. Ebadi

F. Falahi M. Kianpisheh S. Mohajer F. Nasiripour K. Ourami P. Rajai P. Sangpour S. Sanjabi F. Shahbandi

.2

e B en

o o

H

so

PREFACE Thin film technology has been developed primarily for the need of the integrated circuit industry. The demand for smaller and smaller devices with higher speed especially in new generation of integrated circuits requires advanced materials and new processing techniques suitable for future giga scale integration (GSI) technology. In this regard, physics and technology of ultra-thin and nanostructural thin films can play an important role to achieve this goal. Thin film technology is based on three foundations: fabrication, characterization and applications. The fabrication of thin films is carried out by employing conventional physical and chemical vapor deposition techniques and their modifications viz. ion-assistance and laser ablation. Thin film application categories include 1) electronic components, 2) electronic displays, 3) optical coatings, 4) magnetic films for data storage, 5) optical data storage devices 6) antistatic coatings and 7) hard surface coatings. Characterization of thin films can be investigated based on film thickness, structure and their chemical composition. The characteristic of a thin film can be quite different from those of bulk material because thin films as a two dimensional systems have a large surface to volume ratio (A/V). In addition, the morphology, physical structure and chemical nature of thin films are differed from the corresponding bulk materials. Another point to be considered is that the surface and/or interface properties of the substrate can drastically influence thin film characteristics due to surface contamination, nucleation effects, surface chemical reactions, surface mobility, stress effects due to thermal expansion mismatch and others. Therefore, There are specialized techniques for the analysis of crystallographic and electronic structure at nano-level and for surface/interface morphology and composition. The international workshop on physics and technology of thin films (IWTF 2003) was the first series of workshop planned to be held triennially in the developing countries. The aims of this workshop series are three-folds: 1) to promote and propagate the recent trends in physics and technology of thin films, among international groups of researchers and technologists, especially young ones, 2) to provide a forum for the young scientists to present their recent results, and to discuss new and current problems in this important field with the experts. 3) to increase scientific collaboration among the participants. IWTF2003 was organized by Sharif University of Technology with the support extended by more than 35 domestic and international institutions. It was held in Tehran from 22 February to 6 March 2003. During this 12 days event, IX

X

about 70 papers from 22 countries were presented as invited talks, contributive seminars or as a poster. In addition to these presentations, there were scientific visits to seven thin film/surface laboratories on a rotational basis. Further, in conjunction with the workshop, there was an exhibition of laboratory equipments and materials as well as internationally published recent books. The present volume comprises of 44 selected manuscripts out of the total contributions presented at the workshop. These proceedings were reviewed by editors and some other experts and accepted on the basis of technical merit and timing. This collection covers various aspects of the broad field of physics and technology of thin films. It is hoped that the proceedings should be useful for both graduate students and professional scientists and engineers. Throughout this collection, the emphasis is on practical application of the basic principles of thin film materials. The editors would like to thank the authors, reviewers, members of international scientific organizing committee and local committees as well as organizers, sponsors, co-sponsors and contributors, for their support through these proceedings. Finally, we are especially grateful to Mr. Azimirad and Ms. Bashlideh for their patient and careful collaboration in the preparation of the proceedings. It is believed that the above aims were fulfilled during this workshop. We look forward to active participation of both the senior and young scientists and technologists in the next IWTF workshop, which will be held in Prague, Czech Republic in 2006.

A.Z. Moshfegh H.v. Kanel S.C. Kashyap M. Wuttig February 2004

CONTENTS I. GENERAL

Welcome Address by: A.Z. Moshfegh Inaugural Address by: S. Sohrabpour Closing Address by: A.Z. Moshfegh Panel Discussion

3 5 6 7

II. DEPOSITION PROCESSES Vacuum Technology: Principles and Applications (Invited) A. Z. Moshfegh

11

PVD Growth Method: Physics and Technology (Invited) A. Z. Moshfegh

28

Introduction to Semiconductor Epitaxy (Invited) H. von Kdnel

54

Semiconductor Superlattices (Invited) H. von Kdnel

70

Oxide Thin Film Growth on Silicon Carbide Surfaces (Invited) P. G. Soukiassian

85

Al-W Amorphous Thin Films (Invited) N. Radic, T. Car, A.Tonejc, J. Ivkov, M. Stubiear and M. Metikos-Hukovic

101

Heat and Mass Transfer During ZnSe CVD Deposition Process V. G. Minkina

119

III. CHARACTERIZATION TECHNIQUES Thin Films Analysis Using Photoelectron Spectroscopy (Invited) S. K. Kulkarni

XI

129

XII

Passivation Investigations of GaAs (100) Surface R. Purandare, B. A. Kuruvilla, S. M. Chaudhari, D. M. Phase and S. K. Kulkarni

149

Correlation Between Microscopic and Macroscopic Properties of Yttria-Stabilized Zirconia Thin Films M. Hartmanova, M. Jergel, V. Navrdtil, K. Navrdtil, K. Gmucovd, F. C. Gandarilla, J. Zemek, S. Chromik and F. Kundracik

158

IV. SURFACE PROCESSES Reliability and Failure of Electronic Materials and Devices (Invited) M. Ohring

171

Diffusion in Multilayers S. Luby, E. Majkova and A. Luches

180

Dynamics of Interacting Adatoms on Complex Surfaces Z Chvoj

187

Copper Surface Segregation During V 2 0 5 Thin Film Deposition M. M. Ahadian, A. Iraji-zad, M. Ghoranneviss and M. Hantizadeh

198

The Preparation and Surface Studies of Fe/Pt Thin Films G. Varghese

205

V. NANOMATERIALS lD-Nanostructures on Silicon Carbide Thin Films (Invited) P. G. Soukiassian

213

Giant Magnetoresistance in Electrodesposited Nanogranular Thin Films (Invited) S. C. Kashyap

228

Self-Assembled Quantum Dots: Structural and Optical Properties, and Device Applications M. Henini

244

XIII

Preparation and Characterization of Ultrathin Films and Film Coatings for Microelectronics Y. A. Pogoryelov

256

Nanocrystalline Films in the Ag-Ni System /. K. Bdikin, G. K. Strukova, D. V. Matveev, S. A. Zver'kov, V. V. Kedrov and G. V. Strukov

265

Fragmentation of Positively Charged Metal Clusters in Stabilized Jellium Model With Self-Compression M. Payami

271

VI. OPTICAL MATERIALS Organic Films for Optoelectronic Applications (Invited) X. Liu, T. Michely and M. Wuttig

285

Development of Highly Reactive Photo-Catalytic Ti0 2 Films (Invited) S. H. Mohamed, R. Drese, M. M. Wakkad and M. Wuttig

297

Multilayer Thin-Film Optical Filters: Design, Fabrication, and Applications (Invited) S. H. Keshmiri and M. M. Mirsalehi

306

Thin Films for Optical Recording A. Kikineshi

318

Diffusion of Atomic Hydrogen and Passivation of Structural Defects in Silicon and in Transparent-Conducting Thin Films (Invited) S. H. Keshmiri

324

The Effect of Particle Size on Optical Properties of CdS Films Formed by Photochemical Technique S. M. Mahdavi, A. Iraji-zad, F. Razi and M. Rezaesmaeili

337

Enhancement in Physical Properties of ZnO Transparent Conducting Coating by Al Incorporation B. N. Pawar, S. R. Jadkar, K. C. Mohite andM. G. Takwale

344

XIV

Optical Energy Gap of Magnetically Confined Arc Discharge D.C. Sputtered Hydrogenated Amorphous Silicon M. C. Abdulrida, H. A. Hamed and B. A. Hassan

355

Photocatalytic Study of Ti0 2 Thin Films Deposited by DC Reactive Magnetron Sputtering and Spray Pyrolysis Methods A. I. Martinez, D. Acosta and A. Lopez

363

Novel Transparent and Highly Conductive ZnO-Based Coatings B. M. Ataev, A. M. Bagamadova, I. K. Kamilov, V. V. Mamedov, A. K. Omaev and S. Sh. Makhmudov

374

Low-Temperature CVD Growth of ZnO Films Stimulated by RF-Discharge Plasma B. M. Ataev, A. M. Bagamadova, I. K. Kamilov, V. V. Mamedov, A. K. Omaev and S. Sh. Makhmudov

380

VII. SUPERCONDUCTIVITY Physics and Applications of YBa2Cu307/Lao.7Sro.3Mn03 Heterostructures (Invited) /. G. Lin

387

Domain Structure of YBa2Cu3Ox Films on NdGa0 3 Substrates /. K. Bdikin, P. B. Mozhaev, G. A. Ovsyannikov, P. V. Komissinski and I. M. Kotelyanskii

405

Raman Active Apical Oxygen Modes in Cu1.xTlxBa2Ca3Cu4012_5 Superconductor Thin Films N. A. Khan and H. Ihara

414

Generation and Amplification of Electromagnetic Radiation by Superconducting Films — a Superconductor Maser A. N. Lykov

420

Fabrication of YBCO and BSCCO Thin Films H. Salamati, P. Kameli and M. Akhavan

428

XV

VIII. MAGNETIC THIN FILMS Some Aspects in Thin Film Magnetism (Invited) M. Fade

441

Microscopic Mechanisms of Magnetooptical Activity in Epitaxial Garnet Films (Invited) A. I. Popov

454

Design, Fabrication and Applications of Multilayer Thin-Film SQUID Sensors (Invited) D. Rassi and Y. E. Zhuravlev

469

Deposition and Characterization of Fe/Si Multilayers S. Kharrazi, S. Ashtaputre, S. Kulkarni, R. Choudhary, S. Shinde andS. Ogale

479

Surface Nonlinear Magneto-Optical Effect in Antiferromagnetics A. K. Zvezdin, A. R Pyatakov, V. I. Belotelov and V. A. Kotov

489

Fabrication and Characterization of the Co/Cu/Co/NiO/Si( 100) Magnetic Multilayer A. Z Moshfegh, P. Sangpour, O. Akhavan, G. Kavei and A. Iraji-zad

499

Optimisation of Thin Film Multi-Layers by Micromagnetic Simulations for MR Applications P. Gornert, D. V. Berkov and N. L. Gorn

507

Conference Photos

515

Author Index

529


WELCOME ADDRESS A. Z. Moshfegh Workshop Organizer

Mr. President Distinguished Participants Ladies and Gentlemen Good morning on behalf of all the Organizing Committees and as the organizer of the international workshop on physics and technology of thin films, it is my great pleasure to welcome you all in here. I would like to thank professor Sohrabpour the president of Sharif University of Technology, for his financial and moral support as well as encouragement. Now, we will ask him to inaugurate this workshop. First, I would like to present a brief introduction about the importance and use of thin films in various scientific and technological areas. Thin films are playing the key role in many technological and sophisticated industries including microelectronics, optoelectronics and sensors. In addition, thin films are performing the similar function in data storage devices industries. They are used in magnetic memory such as hard and floppy disks, and in optical CD memories. Thin films also use in electronic devices including flat panel displays and data storage. Rapid progress and advancement in thin film materials, growth, characterization and application especially in the last two decades, is enormous. Thus due to importance of this field and our potential resources and great interests in this university, it leads us to promote and develop this key technology to national and international scientific communities. In this regard, we have organized such an international event. The initiation and preparation of this workshop has been begun more than two years ago. The objective and main goal of this workshop is to present the latest research work and an overview of the advancement in physics and technology of thin films as well as transfer of knowledge and experience to other interested national and international scientists especially young researchers who work on this important field. In addition to this, we are also seeking the following aims: 1. To increase international scientific collaboration. 2. To promote and enhance higher education and cooperation. 3. To inform and attract attention of government authorities and industries to invest on this key and high technology especially in this information era in 21 century. 3

4

The scientific program of the workshop is including the following activities: 1. Invited speakers presentations that are 90 or 60 min. talk with discussion. 2. Contributive seminar presentations that are 20 minutes talks or poster presentation. 3. Visiting thin films/surface laboratories in a rotational based program. This third activity is organized and arranged to visit seven thin film operative laboratories for experimental work in order to balance theoretical and fundamental sessions. 4. Exhibition of laboratory equipments and components presented by ten different local companies and research institute, and research institutes. At this moment, I would like to thank to all students, faculty members of different committees as well as my colleagues who work in various departments and divisions of this university for their support and assistance. It is necessary to express my sincere appreciation to the international scientific organizing committee (SOC) members as well as to all invited speakers and seminar contributors for their scientific contributions. Concerning our international and national financial supports, I would like to appreciate all of our supporters. Some of our major supporters are including UNESCO, International Center for Theoretical Physics (ICTP) Trieste, Italy, Ministry of Science Research and Technology, Ministry of Industries and Mines, Ministry of Communication and Information Technology and Iran Air. As the organizer, I would like to express my apology for any kind of deficiency and insufficiency in our services that will be presented during the workshop period. At the end, I sincerely look forward to benefiting from this gathering through presentation, exhibition and group discussions, and I hope the workshop be useful and fruitful for all of you especially young researchers working in this important field. You are all very welcome and thank you for your attention.

INAUGURAL ADDRESS Professor S. Sohrabpour President, Sharif University of Technology

Thin films science and technology plays an important role in the high-tech industries. Production of thin films for device purposes is a development of the past 40 years. Thin films as a two dimensional system are of great importance to many real-world problems. Their material costs are very little as compared to the corresponding bulk material and they perform the same function when it comes to surface processes. Thus, knowledge and determination of the nature, functions and new properties of thin films can be used for the development of new technologies for future applications. Some of the important applications of thin films in technology and industries are including: microelectronics, optoelectronics, communication, all types of sensors, catalysis, coating of all kinds (for examples mirrors in lasers and in telescopes) as well as in energy generation and conservation strategies. Therefore, the impact of thin film science and technology on our modern life is enormous. In the other word, thin films are currently used in various aspects of both daily life and sophisticated and hi-tech applications. Unfortunately, the countries in this region suffer from weak scientific relations among themselves and with the advanced world scientific community. As a result, the region is behind in the frontier of science and technology, especially in physics and technology of thin films. Therefore, the objective of the workshop is to emphasize the importance of Thin Films and Sensors in the new technologies and to increase cooperation in the region and with the international community workshop in physics and technology of thin films. Thus, it is believed that this event will promote and enhance the international scientific cooperation in the area of thin films research. I have studied the workshop program thoroughly and I found that it contains high degree of scientific values as well as diversity in topics. I believe, that we will have a successful and fruitful workshop. I am proud to inform you that this university has a great potential on the important fields of thin films and related subjects including surface, interfaces and sensors from different view points including human resources, instrumentations as well as journals and books. Finally, I would like to thank you all once again especially distinguished foreign delegates and hope that you all have a pleasant stay in this beautiful and ancient country. I am certain that presentation of valuable scientific papers and the interaction between scientists will advance the physics and technology of thin films in scientific and industrial communities. 5

CLOSING ADDRESS A. Z. Moshfegh Workshop Organizer

Ladies and Gentlemen, today is the last day of the International Workshop on Physics and Technology of Thin Films (IWTF2003). On behalf of all the organizing committees, I would like to thank all of our participants and contributors for their efforts and cooperation. In this workshop, we have had 30 invited speakers with each talk was either 90 or 60 minutes and about 40 contributive seminar presentations. In addition to those presentations, we have had 5 days scientific visits from seven thin films/surface laboratories. Moreover, we have arranged an exhibition about newly international published books as well as laboratory equipments and components. At this time, I would like to thank professor Sohrabpour the president of Sharif University of Technology for his financial and moral supports. In addition, I would like to thank all personnel of various divisions and departments of the university including international relations, finance and administration, research and technology, academic affairs, public relations and all other staff. It is time to express my sincere thank and appreciation to all students, faculty and staff of the Physics Deportment for their warm and effective supports. It is also necessary to appreciate all of our international and national supporters including UNESCO in particular Paris, and Jakarta offices. Iranian National Commission for UNESCO specially Dr Tavakoli and Dr Gazeni, International Center for Theoretical Physics (ICTP) Trieste, Italy. Ministry of Science Research and Technology, Ministry of Industries and Mines, Ministry of Communication and Information Technology, Iran Air, as well as Center for International Research and Collaboration. It is to note that more than 35 international and national institutions have supported the event that I would like to express my deep appreciation to all of them. At this moment, I would like to state my special thank to our Czech colleagues especially Dr. Chvoj for accepting the host of the 2nd International Workshop on Physics of Technology Thin Films that will be held in Prague, Czech Republic, in 2006. I hope that you have enjoyed and had a good time during your stay in Iran, and I hope that all of you go back safely to your home country with good memory. The workshop ends with a three days post workshop tour to the historical cities of Shiraz (Persepolis) and Isfehan. Thanks to all of you 6

PANEL DISCUSSION On the last day of the event, the workshop organizer A. Z. Moshfegh invited some of the major speakers to present their views on scientific quality and the impact of the workshop on future scientific collaboration between the participants. The name of these contributors are listed below: M. Ohring (U S A) H.v. Kanel (Italy) P.G. Soukiassian (France) A.I. Popov (Russia) D. Rassi (UK) J.G. Lin (Taiwan) S.K. Kulkarni (India) The main goal of panel discussion was to strengthen cooperation based on common interests. After the panellists expressed their views, some of the participants presented their opinions through lively open discussion. Some of the important results and outcome of the discussion is briefly reported. First most of the participants asked for continuation of this activity in one of developing countries holding with two or three years period in future. For the advancement of thin film science and technology in the region and other developing countries, it was also suggested encouragement for exchange program for scientists, especially young researchers in order to reach that important aim. Furthermore, collaboration based on common interests is emphasized and it was proposed for initiation and support soon. Another suggestion that was expressed on future collaboration among the participants, was through the use of international facilities such as Synchrotron Light for Experimental Science and Application in the Middle East (SESAME). This facility can provide characterization services for physicists, chemists, material scientists, as well as biologists, especially for scientists who live in the region or other developing countries. At the end of this session, Dr. Chovj and his colleague from Institute of Physics Czech Academy of Science, have accepted to host the second international workshop on physics and technology of thinfilmsin Prague, Czech Republic in 2006.

7


II. DEPOSITION PROCESSES


VACUUM TECHNOLOGY: PRINCIPLES AND APPLICATIONS A.Z. MOSHFEGH Department of Physics, Sharif University of Technology, P.O. Box 11365-9161, Tehran, Iran E-mail: [email protected] This work is devoted on principles and applications of vacuum technology. Classification and properties of vacuum are discussed. Various pumping mechanisms as well as three basic flow regimes namely viscous, intermediate and molecular are briefly presented. Gas-surface interaction concepts including physisorption and chemisorption states with their distinctive character as well as desorption phenomenon are considered. Two types of surface reaction mechanisms, Langmuir-Hinshelwood and Eley-Rideal are introduced. Applications of vacuum technology in the field of surface science, microfabrication, particle accelerators and analytical techniques are described. Finally, the use of vacuum in different industries with their corresponding applications is briefly reviewed.

1. Introduction Generally, Term "vacuum" refers to a given space filled with gas(es) below atmospheric pressure (P„). The degree of vacuum increases as the pressure exerted by the residual gas decreased below Pa. In the other words, a vacuum is the absence of material including the gases, moisture and particles, which fill our environment. Evacuated spaces or reduced pressure environments can be used for many processing methods or techniques (see sections 5 for details). This type of environment is usually made up of neutral (uncharged) atoms and molecules. However, in some cases electrons and ions are present in plasma. There are four basic concepts and definitions which are related to any vacuum environment namely a) molecular density, b) mean free path of colliding gas, c) the time to form a monolayer (ML), and d) impingement rate (I). By definition, one ML is about ~1015 atoms/cm2. The quantity I is defined as number of particles strike a surface per unit area per unit time expressed by: I = PI(2TTMRT)112

(1)

where P is gas pressure, M molecular weight, R universal gas constant and T temperature in Kelvin. Therefore, these characteristic parameters can express any vacuum system. The degree of vacuum depends on pressure of a given vessel. Thus, there are some regions in vacuum each with different properties. The classification of vacuum is usually described in six different pressure ranges: 1) low, 2) medium, 11

12

3) high, 4) very high, 5) ultra high vacuum (UHV) and 6) extreme high vacuum (XHV). These divisions and corresponding pressure ranges are listed in Table 1. A detailed descriptions and properties of these regions as well as type of materials that can be used in vacuum are discussed in [1,2]. Vacuum systems have been widely used in laboratories and industries (e.g. lamps and vacuum tubes) for many years. They have evolved with improvement in their pumping speed, materials performance, purity and ultimate pressure. A most recent review on progress and advancement of vacuum science and technology during last fifty years of AVS activities (1953-2003), specially from the development of the Bayard-Alpert gauge in 50's as well as history and its future advances is described in [3,4]. Table 1. Various vacuum regions with their pressure ranges.

Region

Pressure Range (torr)

Low Vacuum

~103-1

Medium Vacuum

1-10"3

High Vacuum

10"3 - 10"6

Very High Vacuum

10"6-10"9

Ultra High Vacuum (UHV)

10"9-10"12

Extreme High Vacuum (XHV)

Ti then, the atom usually gains some kinetic energy corresponding to its new equivalent temperature Tr such that Tj

^ s

1 1 1 Hlli

Energy (keV)

(b

0 18"*

I i 1111 ii|—i i i mill—i i 111 MI | 10" 1 10

M I inn—i i i mill—i i ii ni] 10* 10' 10"

Energy (keV)

C 10-

I I lllf| 1

I ITI'lllll 10

I I I I Uli I I I I I i 11 j 18' 10 a

1 I | | illlj 10-

Energy (keV)

1

"1 I I I llll| 10

I I I I llll| 1 I I I llll I I I I Illlj I I I I 1111| 10 a 10* 10" 10"

Energy (keV)

Figure 5. Computer simulated sputtering yield for six transition metal targets under bombardment of different ions a) He+, b) Ne+, c) Kr+ and d) Xe+.

35

In addition, a Monte Carlo code TRIM was also developed to estimate sputtering yield based on binary collisions and in recoil cascade [16]. A linear increase of sputtering yield is observed for many conditions under incident ion energy of about 2000 eV especially for relatively low mass ions. At higher energies, the incident ions penetrate too deeply into the target, and as a result the sputtering yield decreases to a lower value. This is because that the sputtering process is a surface phenomenon. The sputtering yield also depends on angle of incident of the projectile. Yamamura et al. [17] proposed an empirical expression to describe the overall dependence of the sputtering yield on the angle of incidence n. as given below: Y(0) 1 Y(rj) = y - . exp[/. cos 7jopt -(I )] (3)

(cosy n)

cos/7

where Y(0) is the sputter yield at normal incidence to the target surface and n, opt is the angle of incidence corresponding to the maximum yield. In the above equation f is defined as [17]: m f= (Eb)1/2 [0.94 -1.33xlO-3 — - ] (4) m P where Eb surface binding energy in eV and mr and mp are mass of recoil (sputtered atom) and mass of discharge gas in amu, respectively. As an approximation, the sputtering yield at other energies can be estimated by knowing the yield at 1 keV. According to a recent model [9], the sputtering yield at the energy (E) can be defined by: Y(E) = Y(lkeV)[

]0-5

(5)

IkeV this equation is valid in the energy range between 0.5 keV < E < 2 keV. In addition to sputtering yield, several models have been also developed to describe the deposition rate during sputtering process. Stutzin et al. [18] have proposed a useful model to estimate deposition rate for DC diode sputtering. For insulating films, radio frequency (RF) sputtering (in stead of direct current) must use to deposit this type of materials due to their high electrical resistivity. A detailed description on principles, process and applications of RF sputtering technique can be found elsewhere [19,20]. In addition to magnetron and RF sputtering, reactive sputtering can be also used to deposit various materials particularly oxides and nitrides. Westwood [21] has reviewed kinetics and mechanism of reactive sputtering process. Based on the above discussion, different types of sputtering techniques and processes including RF diodes, reactive, and ion beam assisted for metallic, alloys and compound films are described in a recent review article [22].

36

Moreover, various applications of sputtering for deposition of thin films along with its future directions are given in [23]. A co-deposition technique can be also used for fabrication of special compound materials. A combinative sputtering-evaporation method was applied successfully for deposition of the YBa2Cu307_5-Ag thin film (Ag evaporated) over MgO(lOO) substrate with good superconducting property and improved ductility [24]. In many deposition techniques, the depositing flux posses a thermal energies (kT) less than 0.1 eV. Some of these techniques are including conventional thermal evaporation, CVD, MBE and etc. According to Mattox investigation [25], the extra energy added to the surface in the form of energetic atomic flux improved film properties. These ion assisted deposition techniques have some advantages in modifying surface properties of growing film. During sputtering, however there is considerable randomization of travel direction, but a bias voltage can be applied to provide directionality of charged species leading to better uniformity of deposition on stepped surfaces. Thus, bias sputtering at low ion energies (E < 300 eV) can be used to improve the purity of growing film by removing loosely bonded impunity atoms. In addition to improvement in film purity, bias sputtering can effectively alter various properties of deposited films [7] including electrical resistivity, dielectric property, step coverage, hardness, film morphology, density and adhesion. Recently, bias sputtering of Co layer has been applied during the growth of the Cu/Co(Vb)/NiO/Si(100) magnetic multilayer system resulting in reduction of its sheet resistance and surface roughness at an optimum negative bias voltage of Vb=-60 V by using atomic force microscopy (AFM) technique [26]. Sputtering has some disadvantages. For example, for the growth of high temperature superconducting (HTSC) thin film materials, on-axis stoichiometry affected by ions and neutrals bombardment as a result a non-stoichimetric compound sometimes is obtained. However, off-axis sputtering of single target of YBa2Cu307_g-Ag thin films over LaAlO3(100) substrate showed an improvement in surface morphology of the deposited film having smaller grain size with Ag particles residing at the grain boundaries [27]. Other disadvantage of the sputtering technique is its slow deposition rate as compared to laser deposition. 3.2. Laser Deposition Laser evaporation is another physical process that has a much higher energy of arriving atoms and molecules at the film surface. It is also called " flash evaporation" since a powerful laser beam (usually excimer) strikes a target

37

surface (-0.1 cm2) producing a considerable vapor. The vaporized region of the target is about ~ 100 nm thick below its surface. In PLD process, a conical plume of evaporant is created along the direction normal to the target surface. The speed of evaporant particles (neutral and ions) is about 3xl0 5 cm/s corresponding to kinetic energy of about 3 eV. After absorption of laser beam energy by the target surface atoms, evaporants form a plume above the target consisting of collection of energetic neutral atoms, molecules, ions, electrons, atom clusters, micron-size particulates and molten droplets. The plume is highly directional i.e. c o s ^ n where 8r

^_ ^__

a

M — ^ H

^

"•

z

Figure 3. Schematic real space picture of a superlattice made from periodically arranged quantum wells. The minibands are shown as shaded bars.

1.0 n 0.8 0.6

c

0.4 02 0.0-

-4n/a

-n/a

n/a

4jt/a

K

Figure 4. Dispersion relation calculated for a periodic potential like that in Figure 3 with the following parameters: period a = 100 A, width of the tunnel barriers 20 A, effective mass m* = 0.067, and V0 = 0.2 eV.

For an infinite periodic system the wavevector k is well defined, and we obtain the dispersion relation shown in Figure 4, again for parameters applying to the conduction band of GaAs. This is nothing else but the solution of the Kronig-Penney Model familiar from text book physics. In the reduced zone scheme, we obtain from Figure 4 the familiar band diagram of Figure 5.

74

_ •

E

,r"

"

T

T, » T

^ , * •" ,*"

T

" " T

T

» "••••--....

-^..^

.^^^ .^^^"^

' ' " ' ' * ' • ' ' .

*•••'

-7i/a

0

7c/a

•

k

Figure 5. Dispersion relation for the Kronig-Penney model reduced to the first Brillouin zone. The parameters are the same as those in Figure 4.

Evidently, SLs can be considered as model solids, the properties of which can be tuned by varying the SL period, the width of the runnel barriers, and the height of the confining potential V0. These are all experimentally accessible parameters, determined by the GaAs thickness, and the thickness as well as composition of the AlGaAs barriers. The fabrication of semiconductor SLs was first realized by Esaki and Tsu [1]. It should be emphasized that throughout the discussion above we have only been considering particle motion in the direction parallel to the growth direction (or perpendicular to the interfaces). This is in other words a purely one dimensional model. In reality, the electrons are free to move in the direction perpendicular to the growth direction, and the kinetic energy contribution of this free motion has to be taken into account. For non-degenerate wavefunctions of the isolated QWs, as discussed so far, the free motion within the interface plane does not lead to any difficulties. We just have to add the energy for a free electron to each bound energy state, where the wavevector kj_ is twodimensional. In other words, the band structure consists of a series of 2D subbands. The result for an isolated QW is shown in Figure 6. As can be seen, the corresponding density of states consists of steps, each step belonging to one 2D subband, and beginning at one of the eigenvalues discussed above.

75

In a SL we have seen that the states of the corresponding isolated Q Ws broaden into minibands. This has the effect of broadening the steps of the density of states shown in Figure 6, such that the transition from one plateau to the next occurs within the miniband width.

-

VV/ E

\

1

*x

'

2

V

E

1

- = ! - •

1

0

L

2

u

2

z

J

0

L

P(E)

Figure 6. Bound states of a QW containing two states (a), dispersion relations of the 2D subbands (b), and density of states (c).

1.2. Energy Levels for Valence Band States For degenerate energy levels, such as are typical for the valence bands of all important semiconductors, the situation is considerably more complex. Semiconductors with the diamond and zinc-blend structure are characterized by heavy and light hole bands degenerate at the T-point. Here our previous discussion is applicable only for k± = 0. In this case, an isolated QW contains two sets of energy levels, one with smaller spacing for the heavy hole states, and another with a larger spacing for the light hole states. Put differently, the heavy hole states are more tightly bound than the light hole states, since their wavefunction penetrates less into the barrier regions. Correspondingly, SL-minibands stemming from heavy hole states are narrower than the corresponding minibands associated with the light holes. The true difficulty arises as one moves to finite kj., as now heavy and light hole states become coupled. It is hence no longer possible to solve the Schrodinger equation only for motion in the growth direction, and then adding the perpendicular free particle motion afterwards. The consequence is that miniband dispersion relations for holes may become very complex, exhibiting even electron-like behaviour in some cases (negative hole masses). An in-depth treatment of this subject can be found in Ref. [2].

76

1.3. Types of Heterostructures Before discussing specific examples, we wish to illustrate the fundamental types of interfaces or QWs characterizing different combinations of semiconductors. Some important representatives exhibiting band line-ups as shown in Figure 7 (a) are tensile strained Si quantum wells on top of relaxed SiGe buffer layers. Representatives of Figure 7 (b) are compressively strained SiGe alloys on Si or relaxed SiGe. These are all type II quantum wells. Important examples applying to the type I QW of Figure 7 (c) are GaAs QWs on Gai.xAlxAs (lattice matched) or GalnAs alloys on GaAs (mismatched). It is evident that heterostructures of this kind will have better optical properties, such as luminescence efficiencies, since holes and electrons are confined to the same layer. Ec(z) -

"Ur (a) :r-ij

(c)

Ec(z)

E„(z) Ev(z)

(b) Ec(z)

EcW

EvW

E„(z)

(d)

Figure 7. Different kinds of band line-ups which may occur when two semiconductors of different band gap are joined to form heterojunctions and QWs. In (a) the electrons are confined in the A layer, holes in the B layer (type II quantum well). Case (b) is the same except that the A layer is no longer a barrier for holes. In (c) electrons and holes are both confined in the A layer (type I quantum well), (d) is essentially the inverse of (a) in the sense that the A layer acts as a quantum well for holes and as a barrier for electrons.

2. Structural Characterization of Superlattices 2.1. X-ray Diffraction Since usual superlattices are periodic structures, we may expect diffraction effects to occur, whenever a wave with an appropriate wavelength is incident on such a structure. The most powerful experimental tool to investigate the

77

structure of SLs is X-ray diffraction (XRD). Here, as in ordinary XRD, photons with wavelengths of the order of A are incident on the crystal (Figure 8). For a lattice-matched system, such as GaAs/AlGaAs, the situation is particularly simple. We consider the important case of an (001) oriented substrate. The presence of the SL implies additional reciprocal lattice points along the direction of the surface normal. The spacing between these points amounts to 2n/D, if D is the period of the SL. Since we are considering a lattice-matched case, the period D will simply be some multiple of the distance d between the set of lattice planes of the parent materials. The length 2rc/d of the reciprocal vector q is hence a multiple of the SL Q -vector, 2n/D. Thus the SL is commensurate with the lattices of the parent materials.

Analyser

Figure 8. Scattering geometry for a symmetric Bragg reflection. There are two main possibilities: (1) in a 6-26 scan reciprocal space is probed along the surface normal. (2) In a rocking curve, only the incidence angle 6 is varied, and the detector opening is open wide, such that all the scattered intensity is recorded.

The most common way of performing the actual experiment is to acquire so-called rocking curves. Here, a high-resolution diffractometer is used in such a way that only the incidence angle of the X-rays is varied, while the detector stays fixed. In this mode the detector opening is chosen to be sufficiently large in order to detect any scattered radiation in the vicinity of a reciprocal lattice point. One then expects to find each ordinary Bragg reflection to be surrounded symmetrically by superlattice reflection peaks (Figure 9). The situation is somewhat more complicated when the components of a SL are not lattice matched to the substrate. In such a case, the average lattice constant of the SL in general differs from that of the substrate. Accordingly, the Bragg reflection of the corresponding alloy does not coincide with that of the substrate. Instead, it gives rise to the so-called 0th order superlattice peak, situated at some distance from the substrate peak.

78

2TC/D

J

-• 3 -• 2

• 1 # 0 • 1• 22n/d

reciprocal space

• 3•

Figure 9. Region of reciprocal space around some reciprocal lattice points (large dots) in the presence of a SL with period D giving rise to additional points spaced by 2;t/D in the growth direction. The example applies to a lattice matched system.

An important example of this kind are strained-layer SiGe/Si and Ge/Si SLs on Si. Figure 10 shows rocking curves obtained on two such SLs. From the angular separation between the substrate peak (S) and the 0th order SL peak, the average perpendicular lattice parameter and hence the perpendicular strain of the SL can be calculated. In order to check if the SL is fully strained, that is if all lateral lattice parameters are equal to that of Si, some asymmetric Bragg reflection has to be measured as well. In that case, the scattering vector has a component parallel to the surface, such that in combination with the symmetric Bragg reflection, both the parallel and perpendicular lattice parameters (and strain) can be deduced. Once the average lattice parameters are known, the lattice parameter a0 of a relaxed alloy can easily be evaluated by making use of elasticity theory. For the case of a Si substrate we then get by using Veygard's law: «o =

x

• aGe + 0 - x)asi

(!)

Here, x is the concentration of an alloy with the same average Ge content as that of the SL. This quantity can hence be deduced immediately from the position of the 0* order SL peak. From the measurement of the SL period, and the knowledge of the Ge content, one can then deduce the individual thicknesses of the Si and Ge (or SiGe) layers, provided that intermixing is unimportant. There is, however, computer simulation software for rocking curves, taking into account also interdiffusion and interface scattering.

79 (400)

10' 103 10* 10'

f 10° 10"' 10"2 10 3 10 J

lo-'h 10"

-0.06

jWX -0.04

-0.02

0.00

0.02

e-eB (rad) Figure 10. High-resolution X-ray rocking curves recorded in the vicinity of the Si(400) reflection for two Si/Ge superlattices grown by MBE on Si(100). The first SL (lower trace), #233, has a period d = 72.3 A and Ge thickness dc* = 8.3 A, the second (upper trace), #230, has a period d = 50 A and dGE = 5.5 A (from [3]).

2.2. Raman Scattering Raman scattering is a way in which all kinds of fundamental excitations in solids (such as phonons, magnons, electrons, etc.) can be studied. In its simplest version (first order Raman scattering) only zone-center excitations are possible because of the small wave vector of the incident photons. Hence in the case of Raman scattering from pure Si or Ge one expects one single peak, located at the zone center LO frequency of approximately 520 or 302 cm"1, respectively. A superlattice modifies the phonon dispersion relation fundamentally because of the new periodicity it introduces. The most drastic effects occur at low wave numbers, where the constituents do not exhibit any Raman scattering at all. To a first approximation, it is sufficient to take into account only the SL periodicity in the way outlined in Figure 11.

80

0

Figure 11. Folding of the longitudinal acoustic phonon branch in a Si/Ge superlattice. First the sound velocity of an alloy with the same average Ge content as the SL is calculated. This acoustic phonon branch is then folded back to the first Brillouin zone of the SL with zone boundary 7t/d, where d is the SL period.

The q -vector of the photon is also indicated in Figure 11 as qs. Obviously, the zone folding has introduced a large number of optic modes which should appear as doublets in the Raman spectrum (except for the lowest energy mode). This is indeed observed experimentally, as shown in Figure 12 for two Si/Ge SLs on Si(100) with different periods. Also indicated in the figure is a calculation of the phonon spectrum for a laminar medium composed of elastic continuum layers with different mass densities and sound velocities [5]. This more exact model gives very similar results as the simple zone folding scheme outlined in Figure 11, except that it also shows the formation of minigaps in the acoustic spectrum, in close analogy to the electronic case outlined in Figure 4 and Figure 5. It should be evident from an examination of Figure 11 that by measuring the doublet energy the SL period can immediately be deduced.

81

200 inelastic shift [cni ]

O

lOO inelastic shift [cni

200

Figure 12. Raman spectra of Si/Ge SLs on Si(OOl) for parallel and perpendicular polarization of incident and scattered light. (I) SL #233 with period d = 72.3 A and Ge thickness dGe = 8.3 A, and (II) SL #230 with period d = 50 A and dGe = 5.5 A. The SLs have been grown by MBE at substrate temperatures in the range of 460 - 480° C, by using electron beam evaporation for Si and an effusion cell for Ge (from Ref [4]).

The spectral region of the optic modes is equally interesting. First of all, a peak associated with Si-Ge vibrations is observed (Figure 13), which is an indication that the interfaces are not perfectly abrupt and/or planar. The latter is actually expected in view of the Stranski-Krastanow growth mechanism for Ge/Si. In fact the Ge layer thicknesses of these SLs are above the wetting layer thickness discussed before (see lecture on Semiconductor Epitaxy). We expect hence hut clusters and pyramids to form during Ge layer growth, but also their flattening during growth of the Si layers. Close examination of Figure 13 shows that the Ge-Ge vibration has a wave number of 315 cm"1, shifted upwards by 13 cm"1 with respect to unstrained Ge. Using uniaxial stress parameters, we would expect an upward shift of the Ge LO mode by 16 cm"1 for a fully strained pure Ge layer [6]. The lower experimental shift may be attributed either to alloying at the interfaces, in accordance with the observed Si-Ge vibration, or to a confinement shift to lower energies. In fact the mode mainly attributed to Ge-Ge vibrations is largely confined in the Ge layers [7]. In view of the resemblance of the dispersion relation of LO phonons with that of holes in a semiconductor, we expect hence a downward shift of a mode squeezed into a thin layer.

82

i

1 O

u

UUL

1

100

-

1

—

1

200 inelastic shift

400

[cm 1 ]

600

Figure 13. Overview Raman spectrum of SL #230. The spectral region corresponding to the LO modes of Si and Ge shows 3 peaks, associated with Ge-Ge, Si-Ge, and Si-Si vibrations. The Ge-Ge mode is located at 315 cm"1, i.e., shifted to higher energy by strain, whereas the Si-Si mode is at the position odf 520 cm"' for unstrained Si.

3. Applications 3.1. Short-period Si/Ge Superlattices We first of all want to emphasize that there does exist a significant potential for optoelectonic applications of epitaxial SiGe heterostructures. For historical reasons we first want to discuss one attempt which, however, turned out to be not very successful. The basic concept is nevertheless highly attractive and easy to understand after the previous discussion. It is based on the idea of Brillouin zone folding, as explained in Figure 11 for phonons. Let us start with the familiar band structure of Si. The conduction band minima are located along the directions (the so-called A-line in the BZ), at about 80% of the distance from the T-point to the X-point. It turns out that a SiGe alloy has the same conduction band structure up to Ge concentrations of about 80 %. Hence Figure 14 applies equally to an alloy. Figure 14 shows one possible way of folding the A-minimum of the alloy back into the T-point. Evidently, this is possible for a SL with a period of 10 ML, where the thickness of one ML is defined as aSioe/4. One possible experimental realization is hence a Si6Ge4 or a Si7Ge3 SL. SLs with such small periods are customarily called short period superlattices. Since the valence band is folded onto itself, because its maximum is at T, the effect of this folding action is to convert SiGe into a direct bandgap semiconductor. In principle, this should thus result in a material with much improved optical properties, compared with pure Si or a SiGe alloy.

83

1„ _ _ ^ _ __ J

A

^ J U J .

\J. •

!

Figure 14. Folding of the A-minimum of a SiGe alloy into the T-point by a superlattice with period d = 10 monolayers.

Unfortunately, the matrix elements for interband transitions across the direct gap proved to be much smaller than the corresponding ones of III/V semiconductors, such as GaAs. This makes it unlikely that short period SiGe SLs will ever be used as optoelectronic materials, certainly not as light emitters. It also has to be emphasized that epitaxial growth of such heterostructures is far from easy, since the SLs are heavily strained. One possible way to solve the strain problem is to grow the SL on a relaxed SiGe alloy of the same average composition. In this way, the Si and Ge layers of the SL are both strained, but with opposite sign, such that in principle infinitely thick heterostructures should be possible. Of course, there always remains the problem of 3D growth due to the Stranski-Krastanow growth mechanism. The fabrication of short period SLs shows, however, the level of control that has been achieved in epitaxial growth. This is demonstrated by the highresolution transmission electron microscopy imagein cross-section of Figure 15. The figure shows a Si7Ge3 short period SL grown by magnetron sputter epitaxy (MSE) on Si(100), and a strain-compensated Si6Ge4 SL grown by MBE on a relaxed Sio.6Ge0.4 alloy. The periodicity can clearly be recognized. On the other hand, the interfaces are not perfectly abrupt in view of phenomena like Ge surface segregation and the tendency to 3D islanding. These effects can only be kinetically suppressed by further lowering the substrate temperature during growth.

84

Figure 15. Cross-section transmission electron microscopy obtained on (a) a Si?Ge3 short period SL grown by MSE, and (b) a strain-compensated 8uGc4 short period SL grown by MBE (from [8]).

3.2. Qusmtmm Cmscade Lasers Finally, we want to mention one recent application of superlattices in a new device, namely the quantum cascade laser* In this approach, lasing action is not achieved by optical interband transitions, as in usual solid-state lasers, but rather by intersub- band transitions, either in the conduction or the valence band. An example of the fonner are GalnAs/AlInAs heterostructures lattice matched to InP [9]. The minibands are used to inject the electrons efficiently into the excited state, and to remove them equally fast from the ground state after the radiative transition. References t 2. 3. 4. 5. 6. 7. 8. 9.

L. Esaki, R. Tsu9 IBM J. Res. Dev. 14, 61 (1970). G. Bastard, in Wave mechanics applied to semiconductor heterostructures, (les Editions de physique, Paris, 1992). E. Carlino, C. Giannini, L. Tapfer, K.A. Mlder, H. von Klnel, Microsc. Microanal Microstruct. 6, 473 (1995). W. Bacsa, H. von Klnel, K.A. MSder, M. Ospelt, P. Wachter, Superlattices and Microstructures 4, 717 (1988). M. Rytov, Akust. Zk 2, 71 (1956). F. Cerdeira, A. Pinczuk, J.C. Bean, B. Batlogg, B.A. Wilson, Appl Phys. Lett 45, 1138(1984). A. Fasolino, E. Molinari, Journal de Physique C5, 569 (1987). P. Sutter, C. Schwarz, E. Miller, V. Zelezny, S. Goncalves-Conto, H. von Klnel, Appl. Phys. Lett 65, 2220 (1994). J. Faist, F. Capasso, C. Sirtori, D.L. Sivco, J.N. Baillargeon, A.L. Hutchinson, S-N G. Chu, A.Y. Cho, Appl. Phys. Lett 68, 3680 (1996).

OXIDE THIN FILM GROWTH ON SILICON CARBIDE SURFACES P.G. SOUKIASSIAN Departement

de Physique, Universite de Paris-Sud, 91405 Orsay Cedex, E-mail: [email protected]

France

The most recent investigations into the atomic scale understanding of silicon carbide surface oxidation and subsequent initial oxide/SiC interface formation are reviewed for the 6H and 4H hexagonal polytypes. These studies are conducted using advanced experimental techniques primarily based on core level photoemission spectroscopy using synchrotron radiation at NSRRC in Hsinchu. The results indicate a very high reactivity to oxygen of the Si-rich 6H- 4H-SiC(0001) 3x3 surface reconstruction (« 3 orders of magnitude larger than for Si surfaces). Oxygen atom insertion is taking place below the surface close to the first carbon atomic plane, leaving the Si ad-atoms unaffected. By low temperature (500°C) oxidation of a predeposited Si overlayer onto the 6H-SiC(0001) 3x3 surface, a carbon free Si02 ultrathin film (» 10 A) could be grown, leading to the formation of an abrupt SiC>2/6H-SiC interface. However, the two 6H and 4H polytypes have significantly different behaviors with larger amounts of oxide products having higher oxidation states for the 6H-SiC(0001) 3x3 surface, while mixed oxides including carbon species (Si-O-C) are the dominant oxidation products for the 4H polytype surface. The oxidation rate is improved at increased surface temperatures. In all cases, the oxygen uptake remains significantly larger for the 6H polytype when compared to the 4H one. The very different behavior of the 6H and 4H polytypes seems to originate, at least in part, from the presence of two domains in the bulk for the 4H polytype (as evidenced by two bulk components in the Si 2p core level spectrum) which limits the oxygen insertion into the 4H-SiC lattice. These findings show that a "gentle" oxidation could be a promising approach to SiC surface passivation.

1. Introduction Since several decades, the general trend in microelectronics is toward higher integration densities which, according to the Moore law, are supposed to double every 18 months [1]. However, such an approach is rapidly reaching its fundamental limits with basically no known technological solution indicating that following the Moore law would still be possible in 5 to 7 years from now. In this view, one of the important issues is semiconductor surface passivation. Primarily due to the excellent properties of its native oxide (Si02) and to the low defect density of the Si02/Si interfaces, silicon is by far the most commonly used semiconductor in device technology [1,2]. The Si02/Si interface is generally grown by thermal oxidation which results in a non abrupt interface having a Si02 to Si transition layer of about 30 A to 40 A thick including non stoichiometric oxidation products having lower oxidation states (Si^ + ,Si^ + ,Si + ). 85

86

The present downsizing approach in device technology now requires ultra thin oxide layers having thicknesses below » 50 A (Fig. 1). It is therefore not anymore possible to have a transition layer at the oxide/semiconductor interface having a thickness in the same order of magnitude.

Figure 1. Schematic of the downsizing approach in oxide thin films on Si or SiC surfaces. The transition layer including sub-stoichiometric oxides is represented in gray at the interface between SiOl and Si (SiC).

Silicon carbide (SiC) is not a new material since it is actually older than the solar system. It has been discovered at the end of the 19th century (in 1895) by Henri Moissan (1904 Chemistry Nobel Prize laureate), on a meteorite located in the Diablo canyon (Arizona) [3]. However, SiC is definitively an advanced material with many existing or potential technological applications in microelectronics, in mechanical structures as ceramics in matrix composites and in. biocompatibility [4-8]. SiC is a IV-IV compound wide band gap (from 2.4 eV to 3.3 eV) semiconductor especially suitable in high temperature, high power, highfrequencyand high voltage microelectronics devices and sensors [5-8]. In addition, SiC is rather inert chemically and is resistant to radiation damages which make it very useful for devices and sensors working in harsh environments [5-8]. Furthermore, due to a small mismatch between lattice parameters, SiC is an excellent substrate for the growth of the new class of III-V nitride semiconductors in both hexagonal and cubic phases. According to Keyes and Johnson, the performances of SiC devices are expected to be better by factors ranging from 50 to 1000 when compared to conventional elemental (Si) or III-V compound semiconductors, SiC being surpassed only by diamond [5-9]. SiC exits in various crystallographic phases (including cubic (P), rhombohedric and hexagonal (a)) with more than 170 polytypes [5,6]. Recently, thanks to'advanced experimental techniques, high quality hexagonal and cubic

87

SiC surfaces could be achieved at the atomic scale [11-24]. Despite lower carrier mobilities compared to (J-SiC, hexagonal 6H- and 4H-SiC polytypes are the most commonly used for devices because of good quality crystal wafer availability and larger bandgap [5-8]. Surface passivation is a crucial issue in successful SiC device technology [5-8]. Unlike other semiconductors except silicon, Si02 is the native oxide of SiC. However, SiC oxidation could result in mixed oxides products containing C species, likely to be one of the parameter at the origin of high density of interfaces states [25]. The SiC oxidation has been investigated for various cubic and hexagonal polytypes surfaces using different experimental probes and ab-initio theoretical calculations [26-41]. In order to achieve high quality SiC>2/SiC interfaces, it is necessary to find alternative approaches to SiC oxidation. In particular, to avoid C intermixing into the oxide products, it is challenging to investigate the passivation process using a "gentle" surface oxidation controlled at the atomic scale. In this report, we describe the most recent results obtained in atomic scale SiC surface oxidation and initial Si02/SiC interfaces formation, focusing primarily on Si-rich hexagonal 6H-SiC(0001) 3x3 surface. The investigations are performed using synchrotron radiation based core level photoemission spectroscopy at NSRRC (Hsinchu). The 6H-SiC(0001) 3x3 surface is found to be highly reactive to oxygen, even at extremely low exposures with Si02 as the dominant oxidation product. Initial oxygen deposition takes place below the surface, away from Si dangling bonds, close from the carbon plane with oxygen atom in Si-O-Si bridge bonded sites. The oxidation of a pre-deposited thin Si overlayer onto the 6H-SiC(0001) 3x3 surface leads to an abrupt SiC>2/SiC interface formation. In contrast, 4H-SiC(0001) 3x3 oxidation results in dominant mixed oxide products. 2. Experimental Details Core level photoemission spectroscopy (CLPS) experiments are performed at the Synchrotron Radiation Research Center (SRRC, Hsinchu, Taiwan). The light emitted by the 3rd generation 1.5 GeV storage ring is dispersed by the high energy spherical grating (HSGM) monochromator beam line. The photoelectron energy is analyzed using a hemispherical angle integrating (± 15° solid angle) electrostatic analyzer having a 150 mm radius (VSW). The overall (monochromator and analyzer) energy resolution is 180 meV at Si 2p and 360 meV at C Is core levels. The pressure in the vacuum chamber remains always in the mid 10"" Torr range. The n-doped 6H-SiC(0001) 3x3 surface preparation is achieved by a 10 minutes Si deposition at 650°C followed by a 2 minutes annealing at 750°C. Oxygen is deposited onto the surface using research grade oxygen exposures. The surface quality is checked by low energy electron

88

diffraction (LEED) to have sharp 3x3 patterns. Additional details about high quality 6H-SiC(0001) 3x3 surface preparation and Si 2p core level curve fitting procedures are available elsewhere [14,15,20,29,36,37,42]. 3. Initial Oxidation of The 6H-SiC(0001) 3x3 Surface The Si-rich 6H-SiC(0001) 3x3 surface exhibits a rather complex structure that has been determined only recently using various advanced experimental techniques such as dynamical LEED, STM, electron holography and also synchrotron radiation-based CLPS, and by state-of-the-art ab-initio total energy theoretical calculations using the local density functional approach [14,15,20,21], The model of this 6H-SiC(0001)3x3 surface reconstruction having 4 Si atomic planes lying on a C plane has been established on the ground of these combined experimental results and theoretical calculations [14,15,20,21]. Its structure includes Si ad-atoms having one dangling bond connected to Si trimers. The latter are twisted by 30° to allow hybridization with the Si atoms belonging to the ad-layer thereby leading to a significant surface strain [14,15]. Below the Si adlayer, we find the first Si and C bulk atomic planes [14,15,20,21]. We use synchrotron radiation based core level photoemission spectroscopy performed at the Si 2p and C Is core levels for clean and oxygen exposed 6HSiC(0001) 3x3 surface [36]. Indeed, with this technique, one can probe not only the surface but also the sub-surface regions using the photon energy tunability of the synchrotron radiation. In addition, it is also possible to probe the Si and C atoms chemical environments and to trace the formation of oxide products. Figure 2 displays a set of representative Si 2p and C Is core levels for clean and oxygen exposed (0.25 L and 0.5 L) 6H-SiC(0001) 3x3 surfaces. The Si 2p core level for the clean 6H-SiC(0001) 3x3 surface reconstruction exhibits two surface shifted components SSI and SS2 and one bulk component [20]. SSI is related to the Si ad-atom while SS2 corresponds to the Si trimer + Si ad-layer [20]. Generally, oxygen interaction with a surface significantly affects the surface shifted components with binding energy and/or intensity changes. Upon 0.25 L and 0.5 L oxygen exposures of the 6H-SiC(0001) 3x3 surface, one can see here that SSI and SS2 are not at all affected which indicates no oxygen interaction with the Si surface atoms, in excellent agreement with the picture derived from STM measurements [36]. Instead, one can observe the growth of a chemically shifted component (Si + ) already at a very low 0.25 L O2 exposure and high oxidation states formation (Si ^ + and Si 3+) at only a slightly higher O2 exposure of 0.5 L. This indicates the formation of oxidation products already at room temperature and at oxygen exposures 3 orders of magnitude smaller that

89

e.g. for Si surfaces [36,44,45], therefore stressing the exceptionally high reactivity to oxygen of the Si-rich 6H-SiC(0001) 3x3 surface [36]. Further confirmation of oxygen atom adsorption below the sub-surface region comes by looking at Figures 5d, 5e and 5f at the C Is core level which shows slight binding energy changes upon oxygen exposures [36]. While this binding energy change is too small to indicate the establishment of C-0 bonds, it clearly shows that Si atoms in the neighboring of the C plane, i.e. well below the surface, are affected by oxygen atoms which results in fine electronic charge redistribution as traced by the small C Is binding energy change [36],

d)

Cls Clean ./

H

A /si 4+ \

1

^si2Ay.; Si-O-C.

\SJT A

8

7

6

284 283 282 Binding Energy (eV)

5

4

3

2

.-Clean!

1

0-1

Relative Binding Energy (eV) Figure 4. 1000 L of O2/Si/6H-SiC(0001) at 500°C: a) Si 2p core level recorded at a photon energy of hv = 150 eV and grazing emission angle (surface sensitive mode). The binding energy scale is relative to Si 2p binding energy (99.2 eV) for a silicon surface. The Si 2p for SiC (dotted line) and the chemically shifted components (continuous line) are also displayed. b) Si 2p core level recorded at a photon energy of hv = 300 eV and normal emission angle (bulk sensitive mode). The binding energy scale is relative to Si 2p binding energy (99.2 eV) for a silicon surface. The Si 2p for SiC (dotted line) and the chemically shifted components (continuous line) are also displayed. c) C Is core level shift for clean and oxidized (as above) surfaces recorded at a photon energy of hv = 330 eVand grazing emission angle (surface sensitive mode).

93

stoichiometric Si02 is the dominant oxide product in the bulk sensitive mode, as can be seen in Figure 4b while the C Is core level (Fig. 4c) for the clean and oxygen exposed at 500°C Si/6H-SiC(0001) surface keeps the same full width at half maximum (FWHM = 0.68 eV) indicating that the C atoms are not significantly affected by the oxidation process [29]. Overall, the results indicate the formation of carbon free Si02. In addition, the oxidation of a Si ultra thin film (« 10 A) pre-deposited onto the 6H-SiC(0001) 3x3 surface leads to an abrupt SiC>2/6H-SiC(0001) interface formation, with C atoms bonded to both Si and C species at the interface [29]. This situation significantly differs from the direct surface oxidation as seen above. The resulting oxide layer thus grown is thicker than one could expect from the oxidation of the silicon thin film only. This indicates that the latter also plays an active role into the oxidation of the underlying silicon carbide surface. The structure of this silicon thin film grown on hexagonal SiC is indeed of special relevance into the oxidation process of the SiC surface since it as an unexpected cubic 4x3 surface array and is highly sensitive to oxygen [43]. 5. Si02/4H- and 6H-SiC Interfaces : The Polytype Effect As mentioned above, the Si-rich 6H- 4H-SiC(0001) 3x3 surface reconstruction is found to be highly reactive to oxygen with an initial oxidation rate of about 3 to 4 orders of magnitude larger than for silicon surfaces. Initial oxygen atom interaction is taking place below the surface, just above the C-plane. For all polytypes, direct Si02/SiC interface formation is achieved already at room temperature and extremely low oxygen exposures. However, the two 6H- and 4H-polytypes have significantly different behaviors with larger amounts of oxide products having higher oxidation states for the 6H-SiC(0001) 3x3 surface, while mixed oxides including carbon species (Si-O-C) are the dominant oxide products for the 4H polytype surface - see Figures 5, 6 and 7 - [42]. The oxidation rate is improved at increased surface temperatures [42]. In all cases, the oxygen uptake remains significantly larger for the 6H polytype when compared to the 4H one as can be derived from Ols core level in figure 7. The very different behavior of the 6H and 4H polytypes seems to originate, at least in part, from the presence of two domains in the bulk for the 4H polytype (as evidenced by two bulk components in the Si 2p core level spectrum - see Figure 8) which limits the oxygen insertion into the 4H-SiC lattice. Instead the 6H polytype has only one bulk domain with a single Si 2p bulk component. Abrupt Si02/6H-SiC interfaces could be achieved by thermal oxidation of a pre-deposited Si overlayer onto the surface leading to have oxide thicknesses ranging from 10 A up to 80 A after post-oxidation with a transition layer of less than 5 A - see Figure 9 - [42]. These studies show that, using a

94

"gentle" initial oxidation approach at low temperatures and low oxygen exposures allow high quality Si02/SiC interfaces with oxides resistant to radiation damages as the result of a low temperature process. It also brings interesting novel insights into the understanding of polytype crucial effect in SiC surface oxidation.

1000 L O2/SiC(0001)-(3x3) at 25°C : Si 2p Surface hv=150eV,e e = 30°

a)

Si 2p Bulk hv = 300eV,9 = 6 0 °

b)

4H-SiC 4H-SiC

6H-SiC 7 6 5 4 3 2 1 0 - 1 Relative Binding Energy (eV)

7 6 5 4 3 2 1 0 - 1 Relative Binding Energy (eV)

Figure 5. Comparison between the Si 2p core level spectra for the 4H-SiC(0001) 3x3 and the 6HSiC(0001) 3x3 surfaces exposed to 1000 L Oj at 25°C recorded in a) surface and b) bulk sensitive modes (hv= 300 eV and emission angle 0 e = 60°). The peak decomposition for the 1000 L of O2/6H-SiC(0001) 3x3 surface show the 4 oxidation states Si 4 + , Si 3 + , Si 2 + and Si + (straight line). For the shake of clarity, the Si 2p contribution of the clean SiC surface is represented by a single peak (dot line) that includes the 2 surface shifted components (SSI and SS2) and one bulk component B.

95

1000 L 0 2 / 6H- 4H-SiC(0001)-(3x3) at 500°C oxidation Si 2p Surface

a)

Si 2p Bulk

b)

hv = 300 eV, 9 e = 60°

hv=150eV,6 =30°

A >\ 4H-SiC

' 1 )

J J

)

s "*•

/Si",

Si2+

i 6H-SiC

/si 4 + )


\4H-SiC

\ 6H-SiC

Mean V V. 1

i

•

i


Figure 6. Comparison between the Si 2p core level spectra for the 4H-SiC(0001) 3x3 and 6HSiC(0001) 3x3 surfaces exposed to 1000 L O2 at 500°C in a) surface and b) bulk sensitive mode. The peak decomposition for the 1000 L of C>2/6H-SiC(0001) 3x3 surface show the 4 oxidation products Si 4 + , Si 3 + , Si 2 + and Si + (straight line). For the shake of clarity, the Si 2p contribution of the clean SiC surface is represented by a single peak (dot line) that includes the 2 surface shifted components (SSI and SS2) and one (or two) bulk component B ( B * ) .

96

O 2 /SiC(0001)-(3x3) 12 •I 10 8

6H-SiC at 500°C 4H-SiC at 500' 6H-SiC at 25 4H-SiC at 25

6 4 2 10 100 0 2 Exposure (L)

1000

Figure 7. Comparison of the oxygen uptake on the 4H-SiC(0001) 3x3 and 6H-SiC(0001) 3x3 surfaces at 25°C and 500°C. O Is peak intensity is displayed using arbitrary units and correspond to the integrated surface of the core level peak.

Clean 4H- and 6H-SiC(0001)-3x3 Si 2p Surface hv=150eV,e,. = 30 c

a)

Si 2p Bulk

c)

hv = 300eV,6 e = 60°

J/KBISSA

4H

4H

/ y\ i

/7B* A \sspk

b)

A i n /

6H

6H

jJ 104

d)

1

\\ Wf\

102 100 98 96 104 102 100 98 96 Binding Energy (eV) Binding Energy (eV)

Figure 8. Si 2p core levels for clean 4H-SiC(0001) 3x3 and 6H-SiC(0001) 3x3 surfaces: a) 4HSiC(0001) 3x3 and b) 6H-SiC(0001) 3x3 recorded in the surface sensitive mode; c) 4H-SiC(0001) 3x3 and d) for 6H-SiC(0001) 3x3 recorded in the bulk sensitive mode. Peak decomposition shows two surfaces shifted components SSI and SS2 and one B (6H) or two B and B* (4H) bulk components.

97

O 2 /Si/6H-SiC(0001)-(3x3): post-oxidation Si 2p Surface hv=150eV 9 =30°

~

3000 L 0 2 at 650°C

Si 2p Bulk.

3000 L 0 2 at 650°C

hv = 300eV 9 =60°

. 2000 L 0 2 /> at500°C

2000 L 0 2 at 500°C v..v~--..

Si4+

4

\Si 3+ Si2+

•9A/ \\4 Si -°- C

1000 L O , a t500°C

106 06 104 102 100 98 Rinrlirirr Ttr\(*rcT\r (eV) (t Binding Energy

1000 L 0 2 at 500°C

106 98 106 104 104 102 102 100 100 91 R i n H i n r r T?n

1)

Figure 5. The XRD patterns of the AI78W22 deposited onto alumina ceramic at different stages of isochronal heating: a) as-deposited, b) after heating up to 530°C, c) after heating up to 730°C. The XRD lines of substrate and SS sample holder are marked (s) and (h), respectively. The unmarked lines in the uppermost pattern belong to the AI4W intermetallic compound.

109

Thus, the amorphous Al-W films exhibit structural relaxation upon first heating, which irreversibly change their electric resistivity, before they transform into crystalline phase. The results for the crystallization process itself, isothermal annealing at high subcrystallization temperatures, and structural relaxation will be presented in some more details below: 3.3.1. Crystallization of The Amorphous Al-W Thin Films The crystallization temperatures were determined by the steep change (increase) in the electrical resistance during isochronal heating of the samples. Various heating rates were applied in order to allow for the transformation kinetics analysis. For all Al-W amorphous films, a dominant end-product of crystallization was A14W intermetallic compound, which apparently exhibits higher electric resistivity than the original Al-W amorphous alloys. In tungstenrich samples, a pure tungsten separates alongside A14W. The temperatures of crystallization are shown in Fig. 6, for the whole Al-W amorphicity range. 700

"

Al

650

M,

:

600

l lin

W-> tilrn deportee* umo sapphire substrate held at 400°C, at the fracture.

However, the observed interplanar distance 0.383 nni, does not match any of the three established intermetallic compounds (Al^W, A15W, A14W) or aluminium, a- or p-tungsten. That leaves three metastable interaietallics A13W? A17W3> A12W (with uncertain crystallographic data) as a possible candidate for

113

whisker structure. This conjecture is supported by a disappearence of whiskerrelated XRD pattern upon heating up to 800°C - presumably due to the transformation into a more stable A14W phase. 3.4. Microhardness Variation of the microhardness for the completely amorphous Al-W films (3-4 um thick) is shown in Fig. 11 [22].

:

i 200

•

1 400

1

1 600

•

1 800

1

1 1000

•

1 1200

1

1 1400

T[K] Figure 11. Variation of the microhardness of the AI80W20 (O), AI75W25, (V), and AI67W33 (•) films with annealing temperature.

Microhardness of the as-deposited Al-W amorphous films is about 11 GPa and only slightly depends upon composition. It is considerably higher than the microhardness of pure constituents (5-6 GPa for tungsten, 0.2-0.5 GPa for aluminium), which is common hardening effect caused by alloying. In the temperature interval corresponding to the structural relaxation of amorphous films, the microhardness only marginally increased upon annealing in vacuum. However, crystallization of the amorphous films brought about increase of the microhardness to about 15-17 GPa, which upon further annealing up to 1000°C decreased to the value similar to pure tungsten. 3.5. Corrosion Properties Corrosion resistance of the amorphous and crystallized Al-W alloys was investigated in 1 M HC1 solution. Due to its possible applications in dental

114

prostethics, the Al-W amorphous films corrosion resistance in artificial saliva solution (0.15 M CI') was also examined. The summarized results obtained from the polarization measurements for dependence of corrosion current density upon film composition are shown in Fig. 12 [13, 34]. 1

1000 E

0

20

40

60

80

100

W/%

Figure 12. Corrosion current dependence upon the Al-W film composition: • - as-deposited Al-W films in 1 M HC1, • - crystallized Al-W films in 1 M HCI and D - amorphous Al-W films in artificial saliva solution.

The corrosion current density, j c o n T is a direct measure of the corrosion rate. As seen, the Al-W thin films are inherently passive materials, and the increase of tungsten content in the alloy lowers the overall rate of corrosion and increases the resistance against pitting corrosion [13,34]. Thermally crystallized Al-W films exhibit somewhat lower corrosion resistance in comparison with the amorphous films. 4. Discussion The aluminium-tungsten system has recently been investigated for the purposes of the aluminium-tungsten technology used in ULSI and VLSI microelectronics [3-8]. The Al-W alloys with a small content of tungsten were considered for the electrodes in RF-band surface acoustic wave filters [23]. The amorphous alloys of aluminium and tungsten have been investigated mostly for their excellent corrosion properties at high temperatures [9-14]. However, no systematic investigation of the electrical, thermal and mechanical properties of the Al-W amorphous alloys has been reported. Due to a large difference between the boiling point of aluminium (2740 K)

115

and melting point of tungsten (3483 K) it is difficult to prepare amorphous Al-W alloys in a wide composition range by the rapid quenching from the melt. According to Liu's phenomenological approach [24], the Al-W system belongs to a readily glass formers, and should be obtainable in amorphous phase. Furher, a large negative heat of mixing indicates that Al-W amorphous alloys might be rather thermally stable. Mechanical alloying yielded a great increase of solid solubility of aluminium in tungsten, but did not produce the amorphous phases [25]. Thus, most appropriate technique for the production of Al-W amorphous phases proved to be a codeposition of pure constituents. The amorphous Al-W alloys investigated in refs. 9-11 was produced by sputtering of composite targets, in a composition range from AU5W15 to AI55W45. We employed a codeposition of separately sputtered pure constituents, which allows a simple control of film composition and a great range of composition variation. The amorphicity of prepared Al-W thin films was checked by the XRD method and by the measurement of temperature coefficient of electric resistivity at room temperature. A well-established Mooij correlation [26] characterizes the amorphous phase of metallic alloys by high electric resistivity and negative temperature coefficient of electric resistivity at room temperature. The grain size, estimated from the broad XDR signal centered at about 20 » 40.5° diffraction angle by Scherrer formula, do not vary much with the Al-W alloy composition or the kind of the substrate. Grain size only slightly increased with substrate temperature in the LN2-400°C range, and was about 1.8-2.0 nm. A prolonged annealing at 515°C brought about the increase of grain size to about 2.5 nm. However, the temperature coefficient of electric resistivity remained the same as in the as-deposited samples, which confirmed that a modified Al-W phase retained its amorphous structure. As seen in Fig. 3, the amorphous Al-W alloys exhibit large negative values of temperature coefficient of electric resistivity. A steep increase of resistivity at crystallization allows for the fairly precise determination of crystallization temperatures. The crystallization temperatures were in the 520-620°C range. Their absolute values are in good agreement with the estimates based on kinetic model of amorphous alloys crystallization by Miedema and Buschow [27, 28], and even better agreement with the modified relationship proposed by Weeber [29]. The modified Johnson-Mehl-Avrami analysis yielded activation energy for diffusion increasing from 1.4 eV for Al80W2o to about 2.0 eV for AI75W25, the expected variation bearing in mind that the dominant product of crystallization is AI4W [20]. Finally, it is worth to note that the Al-W films deposited onto fused silica reacted with the substrate (yielding tungsten silicides) before exhibiting a phase transformation (at about 550°C). A pronounced variation of the electric resistivity due to structural relaxation

116

in Al-W amorphous alloys exceeds all other examples observed in amorphous alloys. According to the corresponding XRD patterns, no large structural change/grain growth occured during such heating up to about 500°C. Therefore, the major cause of such large effect might be related to the specific electronic structure of aluminium, which is reflected in the electric properties of its amorphous alloys. Check made by the examination of Al-Ti and Cu-Ti amorphous films prepared by the same method seems to confirm that assumption - a considerable increase of the electric resistivity is observed only in the Al-Ti amorphous alloys upon first heating. Further support of the special role of aluminium in the electronic structure of Al-W amorphous alloys is given by the results of Hall effect measurements [30]: within the Al-W amorphicity range the Hall constant exhibit a pronounced positive maximum, and is generally strongly dependent upon alloy composition. Both strong resistivity effects of structural relaxation and Hall constant variation indicate that the sp-d hybridization governs the transport properties of amorphous Al-W alloys. The details of the sp-d hybridization are apparently very sensitive to the change of short-range order introduced by thermally induced structural relaxation of the amorphous AlW alloys. An attempt to suppress the effects of structural relaxation by deposition onto heated substrates yielded formation of two-phase films. In the substrate temperature range 250-400°C, the XRD peaks corresponding to extremely textured crystalline phase were superimposed upon the signal from the amorphous Al-W matrix. The SEM investigation showed that a dense population of whiskers exists at the top of Al-W film, and it seems that the superimposed XRD peaks are their fingerprints. Whiskers are frequently observed atop aluminium-containing films [31-33], where they grow either during deposition [31], or due to the subsequent heat treatment [32, 33]. Their growth in particular deposition conditions indicate the limits of the optimalization of film properties by adjustment of the deposition process parameters. The mechanical properties of Al-W amorphous films are remarkable in a whole compositional range of amorphicity. The structural relaxation upon heat treatment below 500°C, which so strongly affects the electric resistivity, increases the microhardness only slightly. However, the increase of microhardness after crystallization of amorphous films to about 15 GPa, makes them suitable for protective coatings up to about 900°C. The corrosive properties of Al-W amorphous alloys in technical corrosive solutions are relatively well investigated [9-14]. The results presented in Fig. 12 indicate that the Al-W thin films are inherently passive materials. However, the observed effects of increase of tungsten content in the alloy are much more pronounced than reported previously [9]. The results for corrosion in simulated

117

physiological solutions (artificial saliva) make the amorphous Al-W alloys suitable for applications in demanding biological environment. Finally, a degradation of corrosion resistance of crystallized Al-W thin films for only about an order of magnitude (as measured by corrosion current density) makes them excellent corrosion protection in a wide temperature range. The observed thermal stability, mechanical and corrosive properties make the Al-W thin films, in their amorphous or crystallized phase, suitable for various demanding applications. 5. Conclusions 1) Completely XRD amorphous films in the composition range Al82Wi8 AI63W37 were prepared by magnetron codeposition. 2) Amorphous Al-W thin films were thermally rather stable, crystallization temperatures being in the range 540-620°C. However, amorphous Al-W thin films exhibited rather pronounced structural relaxation upon first heating up to crystallization temperature. 3) Amorphous Al-W films exhibit microhardness considerably greater than pure tungsten, which further increases upon crystallization. 4) Amorphous (and crystallized) Al-W thin films are stable against corrosion in hydrochloric acid and artifical saliva, which makes them suitable for protective coatings in hostile chemical environment. Acknowledgments The authors thank Mr. A. PavleSin for technical assistance. References 1. 2. 3.

4. 5. 6. 7.

Rapidly Quenched and Metastable Materials, Part I, ed. by T. Masumoto, K.Hashimoto (Elsevier Science, Tokyo 1994). N. Radie, D. Gracin, Fizika A4, 233 (1995). Y. Pauleau. Deposition Fundamentals and Properties of Metallic and Diffusion Barrier Films, in Multicomponent and Multilayered Thin Films for Advanced Micro technologies: Techniques, Fundamentals and Devices, NATO-ASI Series, Serie E: Applied Science, Ed. by O. Auciello, J. Engemann, Kluwer Academic Publishers, Dordrecht, The Netherlands, Vol. 234, 471 (1993). M. Tsukada, S.I. Ohfuji, J. Vac. Sci., Technol. A 13, 2525 (1995). D.B. Bergstrom, I. Petrov, L.H. Allen, J.F. Greene, J. Appl. Phys. 82, 201(1997). D.B. Bergstrom, I. Petrov, J.F. Greene, J. Appl. Phys. 82, 2312 (1997). M.B. Takeyama, A. Noya, Jpn. J. Appl. Phys. Part 1, 38, 2116 (1999).

118 8. 9. 10. 11.

12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34.

R. Pantel, J. Torres, P. Paniez, G. Auvert, Microelectron. Eng. 50, 284 (2000). K. Hashimoto, N. Kumagai, H. Yoshioka, H. Habazaki, A. Kawashima, K. Asami, B.-P. Zhang, Mater. Sci., Eng. A133, 22 (1991). H. Habazaki, K. Tahakira, S. Yamaguchi, K. Hashimoto, J. Dabek, S. Mrowec, M. Danielewski, Mater. Sci and Eng. A181/182, 1099 (1994) K. Hashimoto, P.-Y. Park, J.-H. Kim, H. Yoshioka, H. Mitsui, E. Akiyama, H. Habazaki, K. Asami, Z. Grzesik, S. Mrowec, Mater. Sci and Eng. A198, 1 (1995). A.Wolowik, M. Janik-Czachor, Mater. Sci. Eng. A267, 301 (1999). M. MetikoS-Hukovic, N. Radie, Z. GrubaC, A. Tonejc, Electrochimica Acta 47, 2387 (2002). L. Iglesias-Rubianes, P. Skeldon, T.E. Thompson, H. Habazaki, K. Shimizu, Corrosion Sci. 44, 751 (2002). P.A. Bancel, P.A. Heinay, Phys. Rev. B 33, 7917 (1986). H.S. Chen, J.C. Phillips, P. Villars, A.R. Kortan, A. Inoue, Phys. Rev. B 35,9326(1987). N. Radie, B. Grzeta, D. Gracin, T. Car, Thin Solid Films 228, 225 (1993). T. Car, N. Radie, Thin Solid Films 293, 78 (1997). J. Ivkov, N. Radie, Solid State Commun. 106,273 (1998). T. Car, N. Radic, J. Ivkov, E. Babic\ A. Tonejc, Appl. Phys. A 68, 69 (1999). J. Ivkov, N. Radie, A. Tonejc, T. Car, J. Non-Cryst. Solids 319, 240 (2003). M. Stubicar, A. Tonejc, N. Radic, Vacuum 61, 309 (2001). N. Kimura, M. Nakano, K. Sato, Jpn. J. Appl. Phys. Part 1 37, 1020(1998). B. X. Liu, Mater. Lett. 5, 322 (1987). Y.F. Ouyang, X.P. Zhong, W.M. Wu, Sci China Ser. A 43, 184 (2000). J.H. Mooij, Phys. Status Solidi a 17, 521 (1973). A.R. Miedema, Z. Metalkd. 70,345 (1979). K.J.H. Buschow, N.M. Beekmans, Solid State Commun. 35, 233(1980). A.W. Weeber, J. Phys. F: Met. Phys. 17, 809 (1987). J. Ivkov, N. Radie, A. Tonejc, 12th International Conference on Thin Films, (Bratislava, Slovakia, September 15-20, Book of Abstracts, 164). L. Succo, J. Esposito, M. Cleeves, S. Whitney, R.E. Lionetti, C.E. Wickersham, jr., J. Vac. Sci. Technol. A 7, 814 (1989). K. Hinode, Y. Homma, Y. Sasaki, J. Vac. Sci. Technol.A 14,2570 (1996). H. Takatsuija, K. Tsujimoto, K. Kuroda, H. Saka, Thin Solid Films 343, 461 (1999). A. Kwokal, M. MetikoS-Hukovie, N. Radie, R. Poljak-Guberina, A. Catovie, J. Mater. Sci.- Mater. Med., in print.

HEAT AND MASS TRANSFER DURING ZnSe CVD DEPOSITION PROCESS V.G. MINKINA Heat & Mass Transfer Institute National Academy of Sciences, 15 P. Brovka St. Minsk 220072, Belarus E-mail: [email protected]

The heat and mass transfer in a reactor designed for depositing large-area layers of ZnSe is discussed. The gas-phase deposition of ZnSe from the mixture of hydrogen selenide zinc diethyl in a flat channel with separate feeding of the reagents is studied numerically. Heat and mass transfer during the gas-phase deposition is described by Navier-Stokes equations with allowance made for the dependence of the physical properties of the gas mixture on pressure, temperature, and chemical composition. The influence of technological parameters (such as the flow-rate, pressure and chemical composition of the gas mixture) on the deposition of ZnSe layers is discussed.

1. Introduction The uniformity in the thickness of the deposited layer over the whole surface of deposition and the associated processes of the heat and mass transfer become very important for using large-size substrates. In this paper, we calculate the heat and mass transfer in a reactor designed for depositing large-area layers of zinc selenide. Chemical deposition of zincselenide layers from the gas phase is usually conducted in a continuous-flow reactor with heated walls and separate feeding of the initial reagents. For the initial reagents, we used the vapor-gas mixture of hydrogen selenide- zinc diethyl (ZDE). 2. Mathematical Modeling For Gas Phase Deposition-Process In this paper, we discuss the results of numerical investigation of gas-phase deposition of zinc selenide from zinc diethyl and hydrogen selenide in a flat channel, with separate feeding of the reagents, which has a square cross section and with the height-to width ratio far exceeding unity. This circumstance offers us an opportunity to consider the flow and heat and mass transfer in the reactor as two-dimensional. In our work, we take into account the separate feeding of reagents into the reactor by local streams (because ZDE and hydrogen selenide can react with each other even at room temperature) and the evolution of the profiles of velocity, temperature, and component concentrations in it. A diagram of the

119

120

reactor is presented in Fig. 1. Hydrogen selenide is fed through a central nozzle, whereas ZDE is fedfromthe periphery. 4

r

Y//////////A H^-H

1 y

from the axis to y4 = 0, = 0, i=l-3; dx dx on the rear wall (from y4 to h) T = T4+(TS-T4)^^-,

(12)

h-y4 where Ts is the temperature at the surface of deposition, and T3 and T4 are the temperatures at the points y3 and y4. At the surface of deposition (0<x> X> 5 ft oo P u < < . 3 O .5 o 'S. >> 2 2 s 2 s's I i X _ O o X -o — aj oo §8 -£?£ X MW

v s x

n 5 -3 f>

£ _— 2 g

>, 8 3 .2 o&

X-ray Photo electton (XPS) Auger Electronic (AES) y Loss; Spectro Brems ungi Spectr y(Bl & marly othisrtechni Synch:rotron Auger Elec i Spec (AES) X-ray ctron (XPS) Secomiary Ion Mass (SIMS X-ray Difft

3

•a

131

played a major role in throwing light on the basic mechanisms. Here basic principles and information contents of photoelectron spectroscopy will be discussed along with experimental techniques. 2. Photoelectron Spectroscopy 2.1. Basic Principles Electron Spectroscopy for Chemical Analysis [ESCA] or Photoelectron Spectroscopy, was developed by K. Siegbahn [22]. He showed that one could use the principle first explained by Einstein in 1905 to distinguish chemical states of atoms in different local surrounding. According to Einstein if photon of fixed energy hv is incident on an atom then an electron of binding energy EB will be ejected with kinetic energy EK according to the equation hv = EK + E B

(1)

It can be seen that knowing hv and knowing the kinetic energy of electron using an energy analyzer, binding energy of electron in atom can be determined. The equation is valid for atoms in gases, liquids or solids. In a solid additional energy (§) the work function of the solid will be required for the electron to get emitted as hv = EK + EB + t

(2)

However in practice it is not necessary to know | , the work function of the sample but energies are referred to (pSP the spectrometer work function. Work function of samples can change from material to material but that of the analyzer remains fixed. Analyzer is coated with an inert and stable coating so that its work function remains unaltered over a long period of time. Therefore one gets the measured EK which differs from the kinetic energy of photoelectron coming out from the sample. The situation is illustrated in Fig. 2. It can be easily seen that as long as Fermi levels of sample and the analyzer are aligned and analyzer work function remains stable, there would be a constant difference between measured and the kinetic energies of electrons emitted from different samples. It is therefore possible to know the binding energies of the ejected electrons from the sample. Resulting intensity versus electron energy spectrum is Electron Distribution Curve (EDC). A photoelectron in general carries the information about the material and analysis of EDC is very useful. Indeed one should remember that when electron is photo ejected from a sample a hole is created in the energy level from which it is ejected. Binding energy measured will be therefore the energy of the photoelectron in presence of the hole. All the electrons in the atom from which photoelectron is ejected

132

respond to the presence of the hole. The relaxation energy due to this many body effect depends up on the atomic number as well as the energy level for which relaxation energy is considered. The relaxation energies can be quite large, therefore measured binding energies do not represent initial state energy of photoelectron. Nonetheless measured binding energies still are characteristic binding energies, remaining constant for a given element. These energies are sensitive to their local environment. Therefore measurement of binding energies results into useful chemical information. The binding energies without consideration of presence of a hole and given by equation (2) are due to Koopman's theorem of frozen orbitals [2]. There are several review articles and books now available on photoelectron spectroscopy. Readers interested in details of basic issues and classic examples are referred to [22,23].

^ O "

Photoelectric Effect

Atom

Sample

Analyzer 1' Ev measured

Vacuum Level

"sample T analyzer

E F =0

hv

E,

hc = E e + E K +^

Figure 2. Schematic energy level diagram of sample and analyzer.

133

2.2. Why is Photoelectron Spectroscopy a Surface Technique? Photoelectron Spectroscopy can be performed using X-rays or Ultra Violet (UV) rays. Reason for using X-rays or UV rays is that they are able to remove electron from core levels, shallow core levels and valence band of solid depending upon the atom or solid and the energies involved. This is irrespective of elements and their chemical state. X-rays and UV rays can penetrate from few micrometers or fraction of a micrometer depth depending upon the energy of the radiation. However electrons have much shorter mean free path in solids. Therefore even if photoelectrons can be generated deep within a solid few electrons can escape out of solid, depending upon their kinetic energy. An empirical curve of electron escape depth as a function of their kinetic energy is showed in Fig. 3. It can be seen that if the electron kinetic energy is around 100200 eV, escape depth of the electrons is very small, less than a nanometer. One can always find for a given photon source, some photoelectrons from core or shallow energy level which will come from few Angstroms depth. Thus photoelectron spectroscopy is a surface sensitive technique.

3V

*y s

U

S.

uin

a S3

a O 1-

« -2 >- E

JL

a U

s

0)

•-

ISI a Jl

a

E » .2 )l

g> § O ^

«> «* u

B 01 1M B 01 C8 c

s o

« °,

3 Cd

4* -^.

OS

a

u

1-1

01

u
Ol

•M

3 OJC

o

E

* >* 0) N

xami RHE

135

I

i

1 !

OJ)

136

Bi2Sr2C aCu 2 0 8 /Clean

200

400

600

800

1000

1200

B.E. (eV) Figure 5. Survey scan, identifying presence of different elements in the Bi2Sr2CaCu20g sample.

ii) Spin-orbit Splitting - It can often be seen that some of the peaks in photoelectron spectra appear as doublets. These can be due to splitting of p shells [ 2p3/2 - 2p1/2, 3p3/2 - 3pi/2 etc.] as well as for d and f subshells. Spin orbit splitting for a given atom decreases with increase in the principal quantum number. It increases for a given principal quantum number with increase in atomic number. For large atomic number compounds, spin-orbit splitting can be used for finding out the oxidation state. As showed in Table 1, spin-orbit splitting differences are quite large for oxides of high Z atoms compared to pure forms. Table 1. Spin-Orbit splittings for some elements and their oxides.

z 10 (Ne) 14 (Si) 24 (Cr) 29 (Ni) 30 (Cu) 31 (Zn)

2p3/2 _ 2pi/2 eV 0.12 0.70 9.00 18.20 21.20 24.60

— NiO CuO ZnO

— — 18.1 eV 20.2 eV 22.2 eV

137

iii) Multiplet Splitting - Large magnetic moments on some of the atoms/ions arise due to unpaired electrons in their 3d, 4d or 4f shells. Photoelectron spectroscopy can detect the presence of such unpaired electrons. As illustrated in Fig. 6(a) for Fe3+ there are two paired electrons in 3 s level and 5 unpaired electrons with parallel spins in 3d level. When photoelectron is ejected from 3s level two situations can arise viz. (b) an electron left in 3 s is parallel to electrons in 3d level or (c) antiparallel to 3d electrons. The energy difference due to situation in b and c can be as large as few electron volts and depends upon number of electrons in d level. In fact b and c both have certain probability to exist. Therefore 3s photoelectron spectrum exhibits two peaks. Splitting of 3s levels can be correlated to electrons in d level and becomes a measure of magnetic moment. Ft*

-a> 3d

•

3s

Figure 6. Origin of multiplet splitting. Removal of an electron (a) from 3s level of Fe 3+ gives rise to two possibilities (b) and (c) with large energy difference in binding energy of 3s photoelectrons.

iv) Plasmons Loss - When photoelectrons with large kinetic energy try to escape to vacuum they can excite plasmons with energy ha>p = 47ine / m*

(3) 3

Where rop is plasmon frequency, n is number of electrons per cm and m* is the effective mass. Peaks accompany photoelectron peaks on their higher binding energy side. Bulk plasmon peaks even with second and higher orders can be

138 observed with intensity decreasing in higher orders. Also surface plasmon peaks at energies hropA/2 can be detected. v) Satellites - Intense or weak peaks can some times be observed along with main photoelectron peak. Origin of these peaks can be understood in terms of electron shake up, shake off or charge transfer many body interactions and collectively referred to as satellites [24]. In the process of photoionization a hole is created which changes the screening of nuclear charge. Outer or valence electrons respond to this by monopole excitation to a discrete level or continuum. The energy for excitation is derived from outcoming photoelectron. Therefore the photoelectrons losing energy, appear as a peak on the higher binding energy with respect to the ones that come without losing energy. vi) Valence Band- Photoelectron spectroscopy using X-rays or UV rays from helium are used to infer about the density of states in the vicinity of Fermi level. On single crystal surfaces by carrying out angular resolved measurements, surface bands also can be mapped in different direction. Photoemission crosssections are very sensitive to photon energy. Taking advantage of the situation that in typical commercial instrument Al Ka (1486.6 eV), Mg Ka (1253.6 eV), He I (21.2 eV) and He II (40.8 eV) sources of widely different energies are present, useful cross-section dependent analysis can be made. Valence band spectra show considerable changes when recorded with different photon energy. vii) Auger Peaks - Along with photoelectron peaks some of the Auger lines of elements in the solid also are present in a spectrum. This is a natural consequence of a hole in one of the core levels. As shown in Fig. 7, when a hole is created one of the electrons from outer level fills the vacancy. The energy difference between the two levels is either emitted as X-ray (photons) or utilized in emitting an electron from one of the outer levels. Such an electron is known as Auger electron. Energy of an Auger electron in Fig. 7 can be written as EK LII LIII = EK - ELH

(4)

where EK, E Ln and E Lm are binding energies of electrons in K, LII and LIII levels, respectively. Emissions of photons and Auger electrons are competing processes. Production of core hole for emitting Auger electron is independent of the process in which core hole is created. Photons, electrons or even ions can be used to produce core hole. Auger electrons are therefore produced along with photoelectrons. Auger electrons are characteristics of atom from which they are released and are useful part of photoelectron spectra. Auger peaks can be readily identified from photoelectron peaks in a spectrum. First of all Auger peaks are

139

broader as compared to photoelectron peaks. Secondly, the kinetic energy of photoelectrons depends upon the photon energy used as per equation 1. However according to equation 4, kinetic energy of Auger electron does not vary with energy of incident photons (electrons or ions) as long as core hole is created. As a consequence if one switches from Al K^, to Mg K„ (or vice versa) photoelectrons will not change the position on binding energy axis but Auger electron peaks will change the position with photon energy. Lm

—•—•—•—•

Lm

—•—«—a—*L— Auger electron

Figure 7. Schematic energy level diagrams showing (a) photoexcitation of an electron in K level (b) process of filling the hole in K level by release of a photon and (c) process of filling the hole by production of an Auger electron.

3. Advantages of Synchrotron Radiation for Photoelectron Spectroscopy Conventionally Al Ka, Mg Ka, He I and He II photons are used for carrying out photoelectron spectroscopy. Usually the energy and line width of some of the K„ lines increases with increasing atomic number. Table 2 gives energies and line widths of some of the X-ray targets and gaseous discharge lamps. Ideally one would prefer to tune the photon energy to probe core levels of different materials so that various layers of surface regions can be probed. However it is evident from Table 2 that there are gaps in energies of different target materials and line widths also are quite large in high Z materials. To resolve small chemical changes, high resolution would be preferred. Besides it is very difficult to keep on changing or maintaining various anodes in ultra high vacuum environment, necessary for surface analysis. However last 2-3 decades have witnessed a rapid growth in development of synchrotron radiation sources and photoelectron spectroscopy around it. Synchrotron radiation is electromagnetic radiation emitted by accelerated charged particles. It can be proved (see ref. [26] for general introduction to synchrotron radiation as well as photoemission experiments using synchrotron radiation) that a charged particle may be electron, positron, proton etc when accelerated would radiate electromagnetic waves.

140 Table 2. Various sources available for photoelectron spectroscopy.

X-rays "CuK,,,

Energy (eV)

Natural width (eV)

8048.00

2.5

"TiK^

4511.00

1.4

"AIIW

1487.00

0.9

"MgK^

1254.00

0.8

11

NaKai^

1041.00

0.7

ZrM 5

151.40

0.77

40

9

YM^

132.30

0.47

2

He II

40.80

0.1

2

Hel

21.22

0.1

10

Nel

16.85

0.1

18

Arl

11.83

0.1

Tunability of photon energy is achievable using synchrotron radiation.

Synchrotron radiation provides photons with energy from few tens to few keV. This energy range is usually a useful for photoelectron spectroscopy. In Fig 8 comparison is made of photoelectron spectra of Ga 3d and As 3d levels in GaAs recorded using Al Ka and synchrotron radiation. It can be easily seen that much higher resolution is obtainable with synchrotron radiation. In the following section photoelectron spectroscopy using synchrotron radiation is outlined. 4. Photoelectron Spectroscopy Using Synchrotron Radiation As mentioned earlier EDC involves measurements of kinetic energy of photoelectrons using fixed photon energy. Many experiments at synchrotron radiation facilities used EDC mode in order to take advantage of high surface sensitivity by choosing proper photon energy at high intensity. With synchrotron

141

radiation however different kinds of problems beyond EDC can be tackled. Here some of the possible experiments using synchrotron radiation for thin films analysis are discussed. i) Core Levels - High resolution and tunability possible using synchrotron radiation, makes core level analysis a very powerful technique for thin films. A thin 5 A Si02 film on Si(100) showed [27] the existence of silicon atom in Si1+, Si2+, Si3+ and Si4+ valence states. This was possible only due to the tunability of the radiation (hv - 130 eV) and high resolution (~ 0.2 eV). A conventional photoelectron spectrum would show only a smeared peak on the higher binding energy side of the main peak. Another example of difference between recording the core level spectra with conventional source and synchrotron radiation is illustrated in Fig. 8. Ga3d

As3d \

21

hv=1486.6eV

20

19

43 Bintlfng

41

42

Energy (eV)

-

As3d

Ga3d

AS CSOj^.^) ,0-TOeV B

21

20

19

43

Binding

*2

41

Energy (eV)

Figure 8. Tunability and high resolution enables to resolve surface and core levels.

142 ii) Valence Band - Valence band information of a solid is desired to understand nature of bonding, electronic properties etc. by finding out the distribution of occupied electron levels below the Fermi level. This is possible by scanning the kinetic energy of photoelectron at fixed photon energy. EDC is simply assumed to be displaced by photon energy. However at low photon energy the transition of electrons from occupied levels to the states at the bottom of conduction band and above, modifies the EDC substantially. In fact change in photon energy for an initial state means transition to a different final state. Spectra should therefore be recorded with different photon energies in order to reach proper conclusion. Due to high sensitivity and adequate resolution possible at low photon energies, detection of surface states is very straightforward using synchrotron radiation. Surface states are produced due to the abrupt termination of bulk extended solid. Surface states are prominently observable in case of semiconductors, as they are present in many cases in the band gap. These are therefore well separated from the bulk states. It is easy to check whether it is a surface state or bulk state due to the fact that when adsorption takes place on the surface, the surface state vanishes. Study of semiconductor heterojucnction also forms an important method using synchrotron radiation. As showed in reference [18], by depositing ZnS on cleaved Si (111), the position of the valence band edge could be determined which enabled to determine the valence band offset at the interface. Knowing the total width of the spectrum and photon energy, ionization energy can be determined. This enabled to construct the precise line up of valence band, conduction band and vacuum level. These kinds of interface studies are important in understanding the interface reactivity, transport properties etc. Certain novel experiments using the tunability of the synchrotron radiation are performed in the valence band region. Evans et al. [11] deposited ~ 53 A silver on cleaved GaAs (110) at low temperature (100 K). When annealed to a temperature above 249 K well-defined oscillating structure was observed. Thickness dependent studies demonstrated that the new peaks originate from wave vector quantization due to electron confinement. By annealing, the silver islands were formed and observed peaks were interpreted as due to quantum size effects in small metal quantum dots. These are observable in the photon energy range 37 < hv < 57 eV. For more details the original paper should be referred. iii) Cross-sectional Effect - When photoemission intensity from a particular atomic orbital is observed by varying the photon energy, it is noticed that intensity minima occur at certain energies. This is due to cancellation in the dipole matrix element and is popularly known as 'Cooper minimum' [28]. Thus it is an effect related to the probability of photoelectron emission depending on

143

incident photon energy on an atom. This is well studied in case of gaseous atoms and molecules initially and later on used in case of solids to determine how much is the influence of solid state effects or band formation on atomic levels. After all the solids constitute of atoms and is likely that in some cases few atomic characters are still left. This was well illustrated [29] in the case of Molybdenum and Molybdenum sulphide. Molybdenum has a Body Centred Cubic (BCC) structure and MoS2 is hexagonal. In case of BCC Mo, 4d bonding states dominate the photoelectron spectra and have dxy+yz+zx nature. Orbitals in BCC directions are compressed relative to those in free atoms (-1A.U. in Molybdenum crystal compared to -1.6 A in free Mo atom). Where as in MoS2 the distance between Mo-S in z direction is 4.65 A. This distance is much larger than the 4dz2 to be extension of 2.67A in z direction. Thus there is no distortion due to solid state effects. The photon energy variation for 4dz 2 atomic orbital shows a Cooper minimum behavior as expected in case of free atom. There is thus a close relation between theoretically expected atomic Mo cross section and that in MoS2. However bcc Mo does not show this behavior. iv) Resonance - Another effect associated with cross section is known as 'resonance', which shows enhancement of certain photoelectron peaks, as the photon energy is swept. This can be attributed to the fact that there can be more than one channels which give rise to the emission of an electron at the same energy. Electron Distribution Curves (EDC) obtained at different photon energies show sharp enhancement of certain features in EDC. This occurs at particular photon energies just above the photoionization energy for an inner shell excitation. This has been illustrated [30] for MnCl2. In this example EDC for valence band of MnCl2 are plotted for various energies between 30 and 100 eV photon energies. The intensity of the valence band features decreases for photon energies from 35 to 48 eV. Between 48 and 52 eV there is sharp rise in intensity. There after again the intensity falls. This can be explained as follows. The valence band features are mostly ascribed to excitation of 3d electrons of Mn due to the usual photoelectron emission mechanism as 3p53d5 - • 3p 6 3d 4

(5)

However at 48 eV the excitation of 3p electron is also possible in case of Mn as the photon energy is sufficient to emit the electron from it. In this case another way of getting the electron emission by a different route is 3p6 3d5 -> 3p5 3d6

(6)

144

As both the mechanisms produce electrons at the same energy a resonance is said to occur. The description of such a process is obtainable from Fano theory of photon absorption and can be found elsewhere [30]. v) Band Mapping Using Angle Resolved Photoelectron Spectroscopy - Band mapping of metals, semiconductors or insulators forms an important part of electronic structure studies as it not only tests various band theories but also provides basic understanding of different physical properties. Photoemission is the only direct method to map the solid state bands and is theoretically important method. However band mapping is complicated by some inherent difficulties associated with photoemission. First of all consider that the electron generated a few layers below the surface may conserve energy but not momentum. If the momentum of original photoelectron is represented as ko then it might become k on emission from the surface due to potential barrier (called inner potential) at the surface. This reduces the perpendicular component kj. of ko, although k|| is conserved. Photoelectrons suffer refraction at the surface. Additionally it might be necessary to introduce for k, a reciprocal lattice vector G for its conservation as k = ko + G due to effect of Bragg scattering inside the solid. It is also necessary to remember that surfaces undergo reconstruction or relaxation, which might complicate the band mapping. The inner potential is also unknown. In practice only the normal emission is investigated so that k =0. In that case the states along a particular direction in k space are scanned. Besides normal emission dispersion curve or mapping in other direction is also done. As kj. is not conserved at surfaces, the difficulty was faced in determining the bulk band structure. However Loly and Pendry [31] pointed out that the inherent difficulty can be overcome using quantum size effect in thin films by performing some angle resolved photoelectron spectroscopy on layer by layer grown single crystalline thin films of high quality. Details can be found elsewhere [8]. Here it can only be mentioned that even band structure is accessible by photoelectron spectroscopy. vi) Constant Initial and Constant Final State Spectroscopy - These are two new techniques, which have been added to the list of photoemission methods exclusively due to the availability of synchrotron radiation source. Both the techniques make use of tunability of the photon source at desired energy. Here one is specially looking for low energies so that one can derive the nature of electronic states in the valence band i.e. outer occupied states. Note that the conduction band states could not be investigated using conventional photoelectron spectroscopy. This, however, is possible with other spectroscopies

145 known as Bremstahlung and Inverse Photoemission spectroscopies. Using synchrotron source one obtains the information of both occupied and unoccupied states using the same equipment. One can obtain the occupied density of states using what is known as CFS mode. Here photon energy is varied and EK is fixed, i.e., intensity of electrons of a particular kinetic energy is measured. In such a situation different EB electrons would be detected. In other words, we are looking at the occupied density of states. CFS mode however is better than the usual EDC mode for valence band because as the final state is kept constant, the problem arising due to combined effect of density of occupied and unoccupied bands by transition from occupied to unoccupied band does not complicate the data. Thus obtaining the valence band density of states by CFS mode is preferred. Constant initial state spectroscopy is an additional advantage due to synchrotron radiation. Here initial state is kept constant and final state energy is scanned with photon energy. This would yield the DOS in unoccupied or conduction band. vii) Photon Polarization Technique - Synchrotron radiation is linearly polarized in the plane of orbiting electrons. Such light when incident on the sample it selectively emits the electrons i.e. there are certain selection rules for initial and final states symmetry of photoelectron which are to be satisfied. By suitably performing the angled resolved studies the adsorbate symmetries on a well defined solid can be determined. One of the very good examples for this is CO molecule adsorbates on transition metal surfaces [32]. With recent developments in insertion devices like undulators scope of synchrotron radiation techniques has considerably enhanced. One such technique that has emerged is Magnetic Circular Dichroism [33]. This is a unique technique which can unambiguously separate orbital and spin contributions to the magnetic moment. viii) Photo Electron Diffraction - All the techniques using photoelectrons described in this article reveal the electronic structure. However there is one technique viz. photoelectron diffraction which uses the electrons ejected by the photons but gives information of the surface structure. As it is well accepted by now due to many direct and indirect observations that positions of surface atoms are not same as in the bulk, it is necessary to know them. Over the past several years microscopic techniques like Field Emission Microscopy (FEM), Field Ion Microscopy (FIM), Scanning Tunneling Microscopy (STM), Atomic Force Microscopy (AFM) etc. have been utilized to investigate surface and interface structure. Some other techniques like X-ray absorption Near Edge Structure

146 (XANES), Surface Extended X-ray Absorption Fine Structure (SEXAFS), X-ray standing waves (XSW), Low Energy Electron Diffraction (LEED), Reflected High Energy Electron Diffraction (RHEED) etc. are also used extensively to determine the surface and interface structure. However each of the techniques has some problem or the other and often use of multiple technique to solve a problem is advised. Photoelectron diffraction technique offers some of the unique advantages in surface structure determination. As the technique is based on photoelectron emission using photon energy, it is atom specific. So one knows around which atom the structure is being studied. One can even know from chemical shifts the valence state of atom around which it is probed. The technique does not depend upon the existence of long range for atoms in solid. Thus local structure can be probed. Also using appropriate photon energy the technique can be made very surface sensitive. Principle of photoelectron diffraction is as follows. The wave associated with photoelectron gets scattered [34] by the near neighbors and interference of direct photoelectron - wave and scattered wave produces an intensity modulation at the detector. Thus photoelectron diffraction technique consists of observing the intensity modulation as a function of photon energy in certain direction or intensity modulation in various directions. For detailed analysis both are preferred. Analysis using Fourier transform technique yields the information of the position of atoms. Especially the technique is very useful for determining the locations and bond lengths of adsorbates on the surfaces. 5. Conclusions Photoelectron spectroscopy has become an indispensable technique in the analysis of surfaces of bulk materials, thin films, multilayers etc in last few decades. Besides the conventional thin films and other materials it is also possible to investigate nanostructured materials using high surface sensitivity of the technique. In the last few decades it is widely used to investigate metals, semiconductors, insulators, organic as well as biological materials. Surface and interface studies have become of prime importance in the present technologies. Various features like chemical shifts, spin-orbit splitting, multiplet splitting, plasmon loss, satellites, valence band and Auger peaks appearing in the photoelectron spectra using conventional x-ray, UV or synchrotron source can be interpreted to understand more details of the surfaces along with the chemical composition. However using the energy tenability and polarization property of synchrotron radiation additional effects/techniques like cross section effect, resonance, photoelectron diffraction etc are possible which enhance the scope of

147 photoelectron spectroscopy. In brief, detailed structure, composition and electronic structure of surfaces and interfaces is possible using photoelectron spectroscopy. Acknowledgement The author wishes to thank UGC, India for the financial support and her collaborators for their contributions to the work referred in this paper. References 1.

W.F. Egelhoff, Jr. Phys. Rev. B 30, 1052 (1984).

2.

S. Gokhale, N. Ahmed, S. Mahamuni, V.J. Rao, A.S. Nigavekar, S.K. Kulkarni, Surf. Sci. 210, 85 (1989). S. Gokhale, S. Mahamuni, K. Joshi, A.S. Nigavekar, S.K. Kulkarni, Surf. Sci. 257, 157(1991). G. Rossi, Surf. Sci. Rep. 7, 1 (1987). E.A. Kraut, R.W. Grant, J.R. Waldrop, S.P. Kowakzyk, Phys. Rev. B 28, 1965 (1983). B. Johansson, N. Martensson, Phys. Rev. B 21, 4427 (1980). C. Heske, U. Winkler, G. Held, R. Fink, E. Umbach, Ch. Jung, P.R. Bressler, Ch. Hellwig, Phys. Rev. B 56, 2070 (1997). T.C. Chiang, Surf. Sci. Rep. 39,181 (2000). L. Aballe, C. Rogero, S. Gokhale, S.K. Kulkarni, K. Horn, Surf. Sci. 482485,488(2001). U. Winkler, D. Eich, Z.H. Chen, R. Fink, S.K. Kulkarni, E. Umbach, Chem. Phys. Lett. 306,95 (1999). D.A. Evans, M. Alonso, R. Cimino, K. Horn, Phys. Rev. Lett. 70, 3483 (1993). A. Lobo, S. Gokhale, S.K. Kulkarni, Appl. Surf. Sci. 173, 270 (2001). B.A. Kuruvilla, A. Datta, G.S. Shukhawat, A.K. Sharma, P.D. Vyas, R.P. Gupta, S.K. Kulkarni, J. Appl. Phys. 80, 6274 (1996). Min-Gu Kang, Hyung-Ho Park, Vacuum 67, 91 (2002). D.R.T. Zahn, N. Esser, C. Muller, W. Richter, C. Stephens, R. Whittle, I.T. McGovern, S.K. Kulkarni, W. Braun, Appl. Surf. Sci. 56-58, 228 (1992). P. Perfetti, Surf. Rev. Lett. 2, 643 (1995). M. Kundu, S. Mahamuni, S. Gokhale, S.K. Kulkarni, Appl. Surf. Sci. 68, 95 (1991). Ch. Maierhofer, S.K. Kulkarni, M. Alonoso, T. Reich, K. Horn, J. Vac. Sci. Tech. B9, 2238 (1991). K. Tono, A. Terasaki, T. Ohta, T. Kondow, Phys. Rev. Lett. 90, 133402 (2003). L. Ruppel, G. Witte, Ch. Woll, T. Last, S.F. Fischer, U. Kunze, Phys. Rev. B 66, 245307 (2002).

3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

148

21. K. Hayashi, M. Sawada, A. Harasawa, A. Kimura, A. Kakizaki, Phys. Rev. 5 64,54417(2001). 22. K. Siegbahn, J. Electr. Spec. 5, 3 (1974). 23. D. Briggs, M.P. Seah (Eds,), Practical Surface Analysis, (John Wiley & Sons, 1988). 24. T.A. Carlson, Photoelectron and Auger Spectroscopy, (Plenum, New York and London, 1978). 25. CD. Wagner, W.M. Riggs, L.E. Davis, J.F. Moulder, G.E. Muilenberg (Eds), Handbook ofX-ray Photoelectron Spectroscopy, (Perkin-Elmer). 26. G. Margaritondo, Introduction to Synchrotron Radiation, (Oxford University Press, 1988). 27. F.J. Himsel, F.R. McFeely, A. Taleb-Ibrahimi, J.A. Yarmoff, G. Hollinger, Phys. Rev. B 38, 6084 (1988). 28. J.W. Cooper, Phys. Rev. Lett. 13, 762 (1964). 29. I. Abbati, L. Braicovich, C. Carbone, J. Nogami, J.J. Yeh, I. Lindau, U. del Pennino, Phys. Rev. B 32, 5459 (1985). 30. A. Kakizaki, K. Sugeno, T. Ishii, H. Sugawara, I. Nagakura, S. Shin, Phys. Rev.B 38, 1026(1983). 31. P.D. Loly, J.B. Pendry, J. Phys. C 16, 423 (1983). 32. S. Kulkarni, J. Sommers, A.W. Robinson, D. Ricken, Th. Lindner, P. Hoolins, G.J. Lapeyre, A.M. Bradshaw, Surf. Sci. 259, 70 (1992). 33. K. Starke, Magnetic Dichroism in Core-Level Photoemission, (Springer, 2000). 34. A.M. Bradshaw, D.P. Woodruff, in Applications of Synchrotron Radiation, Ed. W. Eberhard, (Springer -Verlag Berlin Heidelberg, 1995).

PASSIVATION INVESTIGATIONS OF GaAs (100) SURFACE

R. PURANDARE, B.A. KURUVILLA Department of Physics, University ofPune, Pune, India E-mail: rahul@physics. unipune.ernet. in S.M. CHAUDHARI, D.M. PHASE Inter University Consortium for Department of Atomic Energy Facilities, Indore, India S.K. KULKARNI Department of Physics, University of Pune, Pune, India

Gallium Arsenide (GaAs) is a direct band gap material useful in many high speed and optoelectronic devices. However, being a compound semiconductor, it is always observed that considerable segregation and oxidation occurs in the surface region. Thus GaAs is structurally and chemically heterogeneous leading to large defect state density. Defect states act as charge trapping centers and lead to inferior device. Attempts are therefore made to remove surface hetrogenities and defects. It has been found that chemical treatment can remove surface oxides and produce a passivating layer whose effectiveness depends upon chemicals used and processing conditions. Although some chemical like Na2S, (NH4)2S, NH4OH etc lead to some successful results, their chemical and long term stability remained questionable. It is therefore necessary to try other chemicals also. In the present case, we have used less attempted phosphorous compounds viz. PC15 and P2S5 to find out their effectiveness as passivant for GaAs (100) surface.

1. Introduction Gallium Arsenide (GaAs) compound semiconductor has been identified as a useful semiconductor material in the areas of advanced electronics and optoelectronic devices due to its high charge carrier mobility and direct band gap [1]. Although silicon technology has its own advantages, unless one goes for silicon nanoparticles or porous silicon based devices [2], GaAs based devices will continue to be very important due to their direct band gap and only 7% effective electron mass compared to that of silicon. In spite of advantageous properties of GaAs as a semiconductor material, a number of difficulties are faced in practice. Even for a clean GaAs surface there is a high density of surface states. Apart from the intrinsic GaAs surface states there is a large number of extrinsic states due to defects. The defect state density is very large in GaAs, especially due to the presence of two different elements viz. Ga and As, which vary in their chemical reactivity, vapour pressure etc.

149

150

GaAs surface exhibits around 1013/cm2/eV surface defect density along with a large surface recombination velocity of ~107 cm/sec. Large number of defects present on GaAs surface are also responsible for the Fermi level pinning and band bending. Fabrication of GaAs based integrated circuits involves the processing at a high temperature (> 300 °C). But GaAs rapidly decomposes at high temperature and due to high vapour pressure substantial outdiffusion of the constituents takes place. Therefore a suitable dielectric overlayer which would maintain the structural and stoichiometry of GaAs is essential. Deposition of various insulators is found to enhance the interface state density in GaAs. Interface between GaAs and insulators have showed considerable instabilities [3,4]. The main reason for surface deterioration in case of GaAs is the presence of surface oxides on GaAs. Interface between GaAs and its oxides is structurally and chemically heterogeneous. A large variety of suboxides of GaAs exist [5] with different conducting properties and are present along with elemental metallic layer of As [6,7]. Large amount of heat released during oxide formation gives rise to defects. Disorders at interface can create additional defects [8]. In order to achieve efficient devices based on GaAs, it is essential to reduce the surface and interface density by using a suitable passivant. A passivant is required to replace the native oxides, remove the Fermi level pinning and reduce the surface recombination velocity. It should prevent any interdiffucion of itself and any metal or metal oxides deposited on GaAs. There should also exist a sufficient barrier between GaAs and the passivating layer so that charge carriers from GaAs are not lost to the passivating layer [9-11]. Various passivating schemes have been devised using photochemical washing, forming over layers of Si0 2 [12 -14], nitirdes [15,16] fluorides [17,18], polymers and other organic materials [19,20], groups VI elements [21,22] etc. Phosphorous belonging to V group also is a very good passivant. In this communication we discuss some of our results on surface passivation of GaAs (100) using wet chemical deposition of phosphorous. Effectiveness of passivant has been tested using photoluminescence (PL); and the analysis of chemical states of elements in the surface region has been carried out by employing photoelectron spectroscopy. 2. Experimental In these experiments p-type GaAs (100) samples were used. Before passivating the samples were degreased using methanol and acetone. This was followed by chemical cleaning in a solution of F^SO^F^CÎÔ taken in 4:1:1 volume proportion. Samples were etched for different time in above solution to find out

151

optimum etching time to clean the samples. For phosphorous passivation P2S5 and PC15 solutions were prepared in NH4OH. Here too different durations of passivation treatment were used. Immediately after each chemical treatment, photoluminescence (PL) was performed to assess the improvement in the electronic properties. PL measurements were carried out using Perkin-Elmer LS-50 model. Xenon discharge lamp was used as the source of continuous radiation. The excitation wavelength used in all the experiments was 570 nm to get the band edge luminescence at 870 nm. The effect of surface treatment of the p-type GaAs (100) samples was analyzed by photoelectron spectroscopy using ESCA Lab MKII system (VG Scientific UK) in case of P2S5 treated samples and by an in-house assembled spectrometer on IUC beam line for photoelectron spectroscopy installed on INDUS-I, India synchrotron [23] in case of PC15 treated sample. In ESCA Lab system, Al Ka (1486.6 eV) served as a source of radiation and concentric hemispherical analyzer (CHA) as electron analyzer. A combined energy resolution o f - 0.8 eV was achieved. On INDUS-I, a photon energy of 134 eV was chosen, and a resolution ~ 0.5 eV was achieved. In both the situations a gold foil was pressed on to the edge of the sample so as to use its Fermi level as the reference. 3. Results and Discussion In order to investigate the effect of chemical cleaning of H 2 S04:H 2 0 2 :H 2 0 solution on GaAs (100) surface, we carried out photoluminescence studies as showed in Fig. 1, initially there is a rise in the photoluminescence intensity as compared to untreated GaAs sample. However after a treatment for just 10 seconds again the intensity of PL peak starts reducing. This is graphically illustrated in the inset of Fig. 1. Analysis of surface species was carried out using XPS analysis as showed in Fig.2. Ga 3d and As 3d shallow core levels were monitored for sample of GaAs treated for different time. It can be seen from this figure that strong oxides of Ga and As, initially present on the as received GaAs (100) sample, were removed to a large extent, by chemical etching carried out for 120 sec. However some oxide layers viz. Ga 2 0 3 and As 2 0 3 are present in each case. The 'as received1 GaAs (100) sample shows Ga 3d peak at binding energy at ~ 20.1 eV and As 3d at 41.2 and 44.2 eV. The peaks at 20.1 eV and 44.2 eV are attributed to gallium and arsenic oxides respectively. The peak at 41.2 eV is due to non-oxide arsenic (or partially covalently bonded to Ga). The peaks are broad and mostly

152

t

23.5

20- ~* 'I

10

16

JO Tim

2-

12

30 (••

-

g <MH o (U

. o 0,4tt c3

1

-8 0,200

/

l/-h 0,0-

\

i

\ •

(H+e b )/k B T

i

5

•

10

Figure 4. a) Dependence of the partial coverage of c-sublattice 0C on the chemical potential repulsive interaction. The values of parameters are the same as in Fig. 3. b) Dependence of the partial coverage 0t on the chemical potential - repulsive interaction. The values of parameters are the same as in Fig. 3.

193 But as the temperature is lowered the behavior becomes quite different. The surface coverage increases from zero to 1/3-ML. The horizontal plateau at 0 = 1/3-ML corresponds to the formation of a completely occupied sublattice of csites. All 6-sites are empty. At the critical value of the chemical potential the surface coverage jumps to 0 = 2/3-ML. All c-sites become empty, ad-particles desorb from the osublattice and all 6-sites are filled. This is quite an unusual phase transition of the first order in a lattice gas system with mutual repulsion between ad-particles. The ordered square structure of ad-particles abruptly changes its density and orientation with respect to the principal axes of the host surface. Such a phase transition is obviously impossible if ec < efi. The order of occupation of the sublattices is crucial. In this case fc-sites will be occupied first and the surface coverage will therefore be both monotonous and a continuous function of the chemical potential. The dependencies of the sublattice coverage 2/3ML, the occupation of ft-sites approaches 1 and the c sublattice is almost empty. When the site depths c and b differ from each other considerably the pair interaction cannot play its ordering role and is unable to influence significantly the adsorption of particles. Phase transitions are absent even at T = 0. Adparticles behave identically to non-interacting ad-particles. The b and csublattices are filled by ad-particles as in Langmuir's case. The adsorption isotherms are continuous and do not show any peculiarities. At first the 'deep' sublattice is occupied and then the 'shallow' sublattice is filled by the adparticles. Interaction changes only the site depth of the 'shallow' sublattice. 4. Calculation of the Diffusion Coefficient Now we focus on determining the diffusion coefficient of the model system in the quasi-chemical approximation by applying linear response theory. The square symmetry of the model lattice allows us to treat D as a scalar quantity. According to Ferrando et al. [13], D can be rewritten in the form of a product of the jump rate factor f, and the thermodynamic factor Kd

194

D=TKa

(7)

if there are no dynamical correlations between two subsequent jump events. Equation (7) is a generalized Darken formula where r is the mean value of particle jumps per unit time and Kd is a combination of the thermodynamic properties of the system as a function of coverage, and of the interactions of an adsorbed particle with the substrate and within the adsorbed layer. The thermodynamic factor is calculated as a derivative of the chemical potential with respect to coverage. Before we present the numerical results we would like briefly to show the range of the parameters k and v (in actual fact the energies AE = e c - eb and X) we use in the discussion: f

k = exp

f v = exp

AE^ kj

1

'

(8)

(9)

If we set the temperature to the experimentally reasonable value of 300 K the range of k between 1 and 10"3 corresponds to a binding energy decrease in the b site between zero and ~ 200 meV. Below a certain distance, the particles strongly repel each other and this interaction is described by the parameter v which corresponds for a value of 100 to a repulsion energy of ~ 50 meV at 300 K. In the following we will focus on the interplay between k and v with regard to D( 600°C to 800°C instead of < 150°C e.g. for silicon) [2-9]. Overall, these characteristics give to SiC many potential applications in aerospace, automotive, electronics and nuclear industries [2-9]. In addition, due to a small mismatch in lattice parameters, SiC (in both cubic and hexagonal phases) is a very suitable substrate for III-V nitride epitaxial growth [2]. SiC exits in (p) cubic, (a) hexagonal (more than 170 polytypes) or rhomboedric crystallographic phases, having band gaps ranging from 2.4 eV to 3.3 eV which could potentially allow to make homojunctions and superlattices based on the same material [11]. Its breakdown field, thermal conductance, band gap and saturated drift velocity are respectively xlO times, x3 times (same as Cu), x2 times and x2 times higher than silicon [2-4].

215

Unlike other group IV semiconductors, SiC is not a fully covalent semiconductor with a significant charge transfer between C and Si, which could give polar surfaces. With the availability of good quality samples, the understanding and control of both cubic and hexagonal SiC surfaces and interfaces has been successfully achieved only recently, contrary to conventional semiconductors [2]. Cubic SiC has the zinc blende structure with alternating Si and C planes, leading for (J-SiC(lOO) to many different surface reconstructions ranging from Si-rich 3x2, 8x2, 5x2, 7x2, 9x2, , Si-terminated c(4x2) and 2x1, C-terminated c(2x2) and C-rich lxl graphitic surfaces, as evidenced by both experimental and theoretical investigations [2,8,12-33]. Due to very large mismatches between lattice parameters when comparing p-SiC(lOO) with Si(100) (- 20%) and C(100) (+ 22%), the Si surface plane is under very large compressive stress while the C surface plane would be, in turn, under strong extensive stress [2,8,12-20,30,33]. This makes SiC as a test case to probe the effect of stress on surface organization. Indeed, these effects are dominant features in P-SiC(lOO) surface ordering such as for the c(4x2) reconstruction. Based on scanning tunneling microscopy (STM) experiments and core level photoemission spectroscopy measurements, we have shown that the (3-SiC(100) c(4x2) surface reconstruction results from Si-Si dimer rows having alternating up- and down-dimers (AUDD model) within the row [15,22]. This very particular surface ordering has not been observed for any other surface and results from a large surface stress as already indicated above [6,7,10,12,17]. The AUDD model is further supported by ab-initio total energy calculations [30,31]. We should remark that the behavior of the (3-SiC(100) surface is very different from corresponding Si(100), Ge(100) and C(100) surface reconstructions. The central issue is the control, at the atomic scale, of SiC surfaces and interfaces. In addition to high quality well defined surfaces, interesting features such as a semiconducting c(4x2) to metallic 2x1 phase transition has been discovered [24] with evidence of a non-Fermi liquid behavior [33]. Interestingly, at the phase transition between Si-rich and Si-terminated P-SiC(lOO) surfaces, the selforganized formation of highly stable Si atomic lines has been observed [8,9,13,16,19,23,33]. In addition, for the C-terminated surface [17,18,21], a very interesting temperature-induced sp to sp3 diamond-type transformation has also been discovered with the formation of sp3 carbon atomic lines [20]. Such C atomic lines could cover the all surface leading to a surface terminated by carbon atoms in a sp3 configuration [20]. This finding could potentially be very useful in providing a substrate for single crystal diamond growth [9]. In this review, I present some of these latest investigations on the control and understanding, at the atomic level, of Si atomic lines and atomic vacancies chains that are self-organized on cubic p-SiC(100) thin film surfaces. These

216

studies are primarily based on scanning tunneling microscopy (STM) experiments. Such important issues as the atomic structure, the role of stress in surface ordering and self-organized Si nanostructures are presented. These Si atomic lines have unprecedented characteristics such as unprecedented thermal stability (> 900°C) and lengths (> 1 urn) making them potentially very useful in nanotechnology. 2. Experimental Details The STM experiments are performed using room temperature and variable temperature scanning tunneling microscopes (RT-STM and VT-STM) operating in ultra high vacuum conditions. The pressure in the experimental and preparation chambers is always kept in the very low 10"" Torr range. We use single crystal, single domain 0-SiC thin films (about lum thick) prepared at LETI (Grenoble), at the Laboratoire de Multimateriaux et Interfaces, University Claude Bernard (Lyon) or at Centre de Recherche sur l'Hete>o6pitaxie, CNRS (Sophia Antipolis) by C3H8 and SiH4 chemical vapor deposition (CVD) growth on vicinal (4°) Si(100) wafers. Very high quality Si-terminated P-SiC(lOO) 3x2 and c(4x2) surface reconstructions can be routinely prepared from sequences of thermal annealing and Si deposition. This procedure is shown to result in very reproducible and clean surfaces as confirmed by sharp single domain low energy electron diffraction (LEED) patterns and specific electronic surface states in the valence band photemission spectra. The control of the various P-SiC(lOO) surface reconstructions has been achieved by core level and valence band photoemission spectroscopies using synchrotron radiation at the Synchrotron Radiation Center (SRC, Madison, Wisconsin, U.S.A.), Advanced Light Source (ALS, Berkeley, U.S.A.), Synchrotron Radiation Research Center (SRRC, Hsinchu, Taiwan) and Laboratoire d'Utilisation du Rayonnement Electromagn&ique (LURE, Orsay, France). Other experimental details about high quality SiC surface preparation could be found elsewhere [8,12-16,19,3338]. 3. Massively Parallel Atomic Si Lines and Si Dimer Chain Vacancies on the P-SiC(lOO) Surface The actual trend in microelectronics is towards much higher integration densities with a road map suggesting a doubling every 18 months (Moore law). However, some serious limitations in this downsizing approach are rising for the near future raising very fundamental questions. Another approach would be to manufacture desired patterns by assembling atoms one-by-one using e.g. STM manipulations [39,40]. However, such methods require very long processing

217

times to achieve nanostructures having the desired properties and, to limit surface diffusion, low temperatures [39,40]. This means that, as soon as the surface is warmed-up e.g. at room temperature, atom surface diffusion will destroyed .the obtained nanopattemlng. As adequately mentioned in the White House National Nanotechnology Initiative [41], there are some important questions such as i) "what new and novel properties will be enabled by naeostructures, especially at room temperature?", ii) "what are the surface reconstructions and atoms rearrangement in nanorods and nanocrystals?", iii) "can one use extensively self-assembly techniques to control nanoscale component relative arrangements?". It is interesting to correlate these questions to the recent discovery, at the phase transition between the Si-rich 3x2 and Si-terminated c(4x2) reconstructions of the P»SiC(100) surface the self-organized formation, upon temperature-induced p-SiC(lOO) 3x2 surface dismantling, of Si atomic lines having unprecedented characteristics - see Fig. 2 - [8,9,13,16,19,23,33,38]. They are I) very long with a length limited by the substrate only, ii) very stable, ill) made of Si-Si dimer lines, iv) the density/spacing of these Si atomic lines could be mediated by a single process, thermal annealing, resulting In arrangements ranging from a single isolated Si line to a superlattice of "massively parallel" Si atomic chains [8,9,13,16,19,23,33,38]. At the very beginning of the pSiC(100)3x2 surface dismantling, one can see in Fig. 3a that the Si atoms are removed dimer row by dimer row, leaving very long Si dimer vacancy leaving very long Si dimer vacancy chains on a 3x2 surface reconstruction [37]. Using a very rigorous protocol in surface preparation, we can now prepare defect free Si dimer lines as shown in a representative STM topograph (Fig. 3b) [37].

Figure 2. SI atomic lines on p-SiC(lOO) thin film surfaces. 800 A x 800 A STM topographs (filled electronic states) of p-SiC(100) surfaces after annealing: a) Si atomic lines forming a superlattice of massively parallel atomic Sines aier annealing at 1050°C; b) Si atomic lines obtained after annealing at I100°C; c) isolated single Si atomic line on the p-SiC(100) c(4x2) surface after annealing at i 150°C. The two top atomiclines are separated by 400 A.

218

Figure 3. a) Si dimer vacancy chains on the on p-SiC(IOO) 3x2 surface. 525 A x 525 A STM topographs (filled electronic states) of p-SiC(lOO) 3x2 surface reconstruction exhibiting dirtier row vacancies after a short annealing at 1050°C. b) Si dimer lines on a p-SiC(lOO) c(4x2) surface: 800 A x §00 A STMtopograph.Notice the quality of these lines that are defect free or almost defect free.

In order to identify the atom position in these lines, it is necessary to image the surface by tunneling into the empty electronic states. In order to correlate filled and empty topographs, we also perform dual scan STM-imaging. Figures 4a and 4b provide a comparison between empty and filled electronic state topographs of the same atomic lines [37]. One can clearly see in the empty state topograph that, by tunneling into Si dangling bonds, the lines are made of pairs atoms forming the Si-Si dimers observed in the filled state topograph [37]. Figure 4c displays the corresponding height profile along a dimer in the empty electronic state STM topographs. One can clearly notice that the Si-Si dimer is symmetric [37], unlike the corresponding behavior of the 3x2 surface reconstruction,, where dimer forming rows are asymmetric [8,14,19]. This indicates that, when the 3x2 surface is dismantled by thermal removal of Si atoms, the spacing between dimer rows increases thereby significantly reducing the lateral interaction [37]. Another possible interesting ordering configuration is to have 'these atomic lines self assembling by pairs in a very particular 8x2 surface array that are imaged by filled and empty STM topographs in figures 5a and 5b respectively, with a joint heigh profile in Fig. 5c [23]. A height profile also shows that the dimers are already symmetric [23]. This particular 8x2 array is taking place at the phase transition between the 3x2 (Si-rich) and the 5x2 (equidistant Si atomic lines) surface reconstructions.

219

0 S 10-15 20 26 30 35 40 9-

A Figure 4 . Identification of the Si atom positions for Si atomic lines: a) Filled electronic states 125 A x 125 A STMtopographshowing the Si-Si dimers forming atomic lines on the p-SiC(lOO) c(4x2) surface, b) 125 A x 125 A STMtopographs(empty electronic states) showing the Si atoms forming the atomic lines, c) Height profile along XX' showing the symmetric nature of the Si-Si dimers.

Since these Si atomic Ikes have their length limited by the substrate only, i.e. by the steps, it is challenging to explore if one can built extremely long atomic lines on very large terraces. Most interestingly, figure 6 shows spectacular self-assembled Si atomic lines on such very large terraces. One can see that they are forming a network of massively parallel atomic lines having a length reaching micron scale (several thousand atoms)9 and probably much longer [33]. Despite such very long lengths, these Si atomic lines still remain very straight. This achievement results in probably what are the longest atomic lines ever built on a surface [33].

220

Filled

States

IMVSJHV S c : ! f ' 1

Filled States XX1

Distance (A)

Figure 5. Pairs of Si atomic lines on p-SiC(lOO) forming a 8x2 surface reconstruction: a) 100 A x 100 A filled electronic state STM topograph. The Infra-pair distance dl represents the lateral row-to-row distance within an atomic lines pair. The inter-pair distance d2 represents the distance between the centers of two neighboring atomic line pairs. b) 100 A x 100 A empty states STM topograph with dl and d2 same as in a). Note overlap between dangling bonds from two adjacent Si atoms belonging to two different atomic lines from the same pair. c) Height profiles covering two line pairs along a) XX' (filled electronic states) and b) YY' (empty states). Notice that as for Isolated atomic lines, the Si-Si dimer is symmetric.

Figure 6. Imaging very long Si atomic lines on a large p-SiC(100) surface: two assembled 2000 A x 2000 A filled electronic state STM topographs. This gives atomic lines having lengths over 0.4 jim and much longer since the data acquisition was limited by the scanning capabilities of the AFM/STM instalment used here. These atomic lines, which form a network of "massivelly parallel" chains^ are probably the longest one's ever built on a surface.

221

4 High Temperature Dynamics and Dismantling of Si Atomic Lines In order to explore the stability of these atomic lines, to study their dynamics and to reach the threshold of their dismantling, high temperature STM experiments are performed [38].-Figure 7 exhibits a serie of STM topographs (filled electronic states) recorded at surface temperatures rangingfrom25 °C to 900 °C [38]. As can be seen from Fig. 7, these Si atomic lines are stable at 600 °C and 700 °C with none of them broken at such high temperatures [38]. At 700 °C? they are regularly spaced while the situation seems to change at 800°C: although almost all dimer lines are still not broken, one can see some gradual changes with very few vacancy segments and an apparent higher line density at the step edge.

T-600°C

T - 25 *€ ^

T - 700 X

I ~ 850 °C

•92::'C

Figure 7. 300A x 300A STM topographs of Si atomic dimer lines on the p-SiC(100) surface imaged at temperatures ranging from 25°C to 925°C. Note that some of these topographs have been recorded on different surfaces and that the difference in Si line density does not necessarily result only from the effect of the temperature. At 800°C, one can already notice the variations in line density in particular at the step edge.

222

The latter feature indicates that atomic lines are moving one by one in a perpendicularly to the line direction and probably eliminated in a collective mechanism at the step edge. When the temperature is raised to 850°C and 900°C, one can observed that the atomic lines are "sizzling" probably due to the large stress resulting form increasing temperatures, but it is also possible that such high temperatures might correspond to the STM instrumental limitation. Anyway, one can clearly notice that the atomic Si lines are still not broken. When the surface temperature is raised by 25°C at 925°C, one can see that the threshold of temperature-induced atomic line dismantling has been reached with only few lines remaining and Si island formation taking place [38], This means that at 925°C, the Si atom back bonds are broken leading to Si surface migration with island formation. This further shows that the bonding of the Si dimers with the silicon carbide substrate is very strong which, together with a strong dimerdimer interaction along the atomic line are at the origin of their unprecedented stability. Incidentally, these STM experiments represent the highest temperature atom resolved imaging. Subsequently and as far as we know, they also show what is probably the highest temperature stability ever achieved for nanostructures built on a surface [38]. Let us now look at the temperature-induced dynamics. Figure 8 displays a serie of STM topographs (filled electronic states) for the same area of Si atomic lines that are recorded during 25 minutes at a 800°C fixed temperature [38]. We follow with time the behavior of an atomic segment line (AS) and a vacancy segment (VS) indicated by an arrow in Fig. 8 which displays such a sequence. We have 8 representative STM topographs (a to h) of the same 100A x 100A area, all recorded at 800°C. As landmarks to follow the evolution of the same measurement, two defects Dl and D2 are used and keep the same position with the atomic line density remaining about the same except for one, labeled XX' which is of particular interest. The latter, located between Dl and D2, appears to be discontinued with two atomic segments labeled AS 1 (9 dimers) and AS 2 (8 dimers) separated by a vacancy segment VS (about 5 missing dimers) (Fig. 8a), the distance between two dimers along a Si line being 6.16 A [16,19]. AS 1, AS 2 and VS evolution with time is followed at a 800°C constant temperature. In Fig. 8b, one can see that, after 3 minutes, AS 1 and AS 2 exhibit the loss of one and two dimers respectively with VS becoming longer (7 missing dimers) indicating that AS 2 is also moving away from AS 1 which remains stable. 2 minutes later (Fig. 8c), AS 1 show no change while AS 2 has lost additional dimers resulting in an increased vacancy segment VS length by one dimer. At 7 minutes, AS 2 has only one dimer left with VS reaching a length corresponding to about 14 missing dimers. This suggests that the remaining AS 2 is still moving away from AS 1 (Fig. 8d). From 8 to 25 minutes, the last dimer belonging to AS

223

2 has disappeared, leading to the opening of a much longer vacancy segment VS (> 25 missing dinners). This sequence shows that the Si atomic line dismantling also resultsfroman individual mechanism with one-by-one dimer removal [38].

t = 6.5

Figure 8 . Dynamics of Si dimer lines at fixed 800°C temperature shown on a serie of 100A x iOOA STtvf topographs. We follow the dismantling with time (between 0 and 25 minutes) of the Si atomic line labeled XX' into atomic segments (As) and vacancy segments (Vs) (a to h). Two defects labeled Dl and D2 are used as landmarks to follow the evolution of the same measurement area.

224

Also we have found that at temperatures above §00°C5 the SI atomic lines are also moving laterally with a higher line density at the step edges. This suggests that the lines are removed one-by-one at the step edges. So the Si thermal elimination on the p-SiC(lOO) surface resultsfromboth individual (oneby-one dimer removal) and collective (line-by-line removal at the step edges) mechanisms [38]. These interesting features are also experimentally advantageous since they probably limit the Si evaporation onto the STM tip, therefore making atomic scale STM imaging at such extreme temperatures somewhat easier. Overall, these experiments stress once again the strong interaction between SI dimers belonging to the same line, this interaction possibly taking place through the SiC surface. 5. New Developments and Perspectives We have shown that It is possible, to control at the atomic scale, surfaces and nanostructures on silicon carbide. The SI atomic lines that are self»organized on the SIC surface have unprecedented characteristics since they probably have the highest thermal stability (900°C) and the longest lengths ((im range) ever observed for an atomic line built on a surface. It is also possible to monitor the line density/spacing in a single step process, thermal annealing, with arrangements rangingfroma single isolated SI atomic line to a large super-lattice of massively parallel atomic lines. If one compares with a line network of an Integrated circuitfromthe late 80's/early 90's (Fig. 9), one can notice that the line density that can be achieved with the SI atomic lines are several orders of magnitude larger. All things being equal, the surface covered by SI atomic lines is Iff8 smaller than those covered by Cu or Al lines.

umoouA

too A "

Figure % Size comparison between a late SO's/early 90's integrated circuit (40 \un x 28 \un) and a super lattice (250A x 17SA) of Si atomic lines. The latter has a surface nearly 8 orders of magnitude smaller.

225

We have also recently found that, by selective adsorbate deposition, the reactivity of these lines with molecules or metal atoms could be very different from that of the underlying surface. This feature open-up many possibilities to built nanostructures having very versatile properties. Applications are therefore possible in nano-electronics, the nanometer scale being recently reached for devices such as a 1.5 run transistor as already successfully achieved at IBM [42], but also in catalysis or in nano-chemistry, since such Si atomic lines could be used as a template e.g. in polymer fabrication by assembling several monomers. The characteristics of these Si atomic lines not only meet but in some cases exceed the requirements for nanotechnology as described in the National Nanotechnology Initiative White House Report [41]. These systems represent model cases in nanophysics. Acknowledgments The author is especially grateful to his PhD students in particular to Fabrice Semond and Vincent Derycke, to his collaborators Victor Aristov, Ludovic Douillard and Hanna Enriquez, and to his graduate students Pascal Fonteneau, Nga-phuong Pham and Pierrick Condette. He also wants to thank Andrew Mayne, Gerald Dujardin and the Laboratoire de Photophysique Mol6culaire in Orsay where part of the room temperature STM measurements have been performed. Very high quality SiC samples have been provided by Thierry Billon, Lea di Ciccio and their group at LETI (Grenoble), by Yves Monteil and his group at LMI-Universite Claude Bernard (Lyon) and by Andre Leycuras at CRHEA-CNRS (Sophia Antipolis). References 1. 2.

3. 4. 5. 6. 7. 8.

H. Moisan, Comptes Rendus de I'Academie des Sciences (Paris) 139, 773 (1904). Silicon Carbide, A Review of Fundamental Questions and Applications to Current Device Technology, edited by W.J. Choyke, H.M. Matsunami, G. Pensl, Akademie Verlag, Berlin, Vol. I & II (1998); and references therein. Silicon Carbide Electronic Devices and Materials, Materials Research Society Bulletin, Vol. 22, March (1997); and references therein. IEEE Transactions on Electron Devices, special issue on Silicon Carbide Electronic Devices 46 (1999); and references therein. R.W. Keyes, Proc. IEEE 60, 225 (1972). E.O. Johnson, RCA Rev. 26, 163 (1965). Diamond Detector Devices and Materials, Materials Research Society Bulletin, Vol. 23, (1998). P. Soukiassian, F. Semond, in Surfaces, Interfaces of Advanced Materials, J. Physique IV (France) 10 (1997); and references therein.

226

9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24.

25. 26. 27. 28. 29. 30.

P. Soukiassian, G. Dujardin, La Recherche 321, 38 (1999); and references therein. V.M. Aroutiounian, V.V. Bouniatian, P. Soukiassian, Sol. Stat. Electronics 43, 343, IEEE Transactions on Electron Devices, special (1999) issue on Silicon Carbide Electronic Devices 46, 585 (1999). U. Starke, J. Schardt, J. Berhardt, M. Franke and K. Heinz, Phys. Rev. Lett. 82,2107(1999). V.M. Bermudez, Phys. Stat. Sol. (b) 202, 447 (1997); and references therein. P. Soukiassian, Mat. Sci. Engineering B 61, 506 (1999) ; and references therein. F. Semond, P. Soukiassian, A. Mayne, G. Dujardin, L. Douillard and C. Jaussaud, Phys. Rev. Lett. 11, 2013 (1996). P. Soukiassian, F. Semond, L. Douillard, A. Mayne, G. Dujardin, L. Pizzagalli, C. Joachim, Phys. Rev. Lett. 78, 907 (1997). P. Soukiassian, F. Semond, A. Mayne, G. Dujardin, Phys. Rev. Lett. 19, 2498 (1997). J.M. Powers, A. Wander, P.J. Rous, M.A. Van Hove, G.A. Somorjai, Phys. Rev. B44, 11159(1991). J.P. Long, V.M. Bermudez, D.E. Ramaker, Phys. Rev. Lett. 76, 991 (1996). F. Semond, Ph.D. Thesis, (Universite de Paris-Sud/Orsay, 19 December 1996). V. Derycke, P. Soukiassian, A. Mayne, G. Dujardin, J. Gautier, Phys. Rev. Lett. 81, 5868 (1998). V. Derycke, P. Soukiassian, A. Mayne, G. Dujardin, Surf. Sci. Lett. 446, LI 01 (2000). P. Soukiassian, V.Yu. Aristov, L. Douillard, F. Semond, A. Mayne, G. Dujardin, L. Pizzagalli, C. Joachim, B. Delley, E. Wimmer, Phys. Rev. Lett. 82, 3721(1999). L. Douillard, V.Yu. Aristov, F. Semond, P. Soukiassian, Surf Sci. Lett. 401, L395 (1998). H.W. Yeom, M. Shimomura, J. Kitamura, S. Hara, K. Tono, I. Matsuda, B.S. Mun, W.A.R. Huff, S. Kono, T. Ohta, S. Yoshida, H. Okuski, K. Kajimura, C.S. Fadley, Phys. Rev. Lett. 83, 1640 (1999). V.Yu. Aristov, L. Douillard, O. Fauchoux and P. Soukiassian, Phys. Rev. Lett. 19, 3700(1997). H. Yan, A.P. Smith, H. Jonsson, Surf. Sci. 330, 265 (1995). M. Sabisch, P. Kriiger, A. Mazur, M. Rohlfing, J. Pollmann, Phys. Rev. B 53, 13121(1996). P. Kackell, J. FurthmUller, F. Bechtedt, G. Kresse, J. Hafner, Phys. Rev. B 54,10304(1996). A. Catellani, G. Galli, F. Gygi, Phys. Rev. Lett. 11, 5090 (1996). A. Catellani, G. Galli, F. Gygi, F. Pellacini, Phys. Rev. B 57, 12255 (1998).

227

31. L. Douillard, F. Semond, V.Yu Aristov, P. Soukiassian, B. Delley, A. Mayne, G. Dujardin, E. Wimmer, in Silicon Carbide, III-V Nitrides,elated Materials, Trans Tech Publications (Switzerland), Materials Science Forum 264, 379 (1998). 32. W. Lu, P. Kriiger, J. Pollmann, Phys. Rev. Lett. 81, 2292 (1998). 33. V. Derycke, Ph.D. Thesis, (University de Paris-Sud/Orsay, 6 November 2000). 34. H. Enriquez, V. Derycke, V.Yu. Aristov, P. Soukiassian, G. Le Lay, A. Cricenti, C. Croti, L. Ferrari, P. Perfetti, Appl. Surf. Sci. 162, 559(2000). 35. V. Derycke, P. Fonteneau, P. Soukiassian, Phys. Rev. B 62,12660 (2000). 36. V.Yu. Aristov, H. Enriquez, V. Derycke, P. Soukiassian, G. Le Lay, C. Grupp, A. Taleb-Ibrahimi, Phys. Rev. B 60, 16553 (1999). 37. V. Derycke, Nga Phuong Pham, P. Fonteneau, P. Soukiassian, P. AbouletNze, Y. Monteil, A.J. Mayne, G. Dujardin, J. Gautier, Appl. Surf. Sci.162. 413 (2000). 38. V.Yu. Aristov, L. Douillard, P. Soukiassian, Surf. Sci. Lett. 440, L 285 (1999). 39. L.J. Whitman, J.A. Stroscio, R.A. Dragoset, R.J. Celotta, Science 251, 1206(1991). 40. T.C. Shen, C. Wang, G.C. Abaln, J.R. Tacker, J.W. Lyding, Ph. Avouris, R.E. Walkup, Science 268, 1590 (1995). 41. Nanotechnology Research Directions: Vision for Nanotechnology R&D in the Next Decade, National Science, Technology Council, (The White House, September 1999). 42. R. Martel, T. Schmidt, H.R. Shea, T. Hertel et P. Avouris, Appl. Phys. Lett. 73, 2447 (1998); R. Martel, H.R. Shea et P. Avouris, Nature 398, 299 (1999).

GIANT MAGNETORESISTANCE IN ELECTRODESPOSITED NANOGRANULAR THIN FILMS S.C. KASHYAP Thin Film Laboratory, Department of Physics Indian Institute of Technology Delhi, New Delhi-110016, INDIA E-mail: skashyap62@hotmail. com In the present paper, an attempt has been made to briefly describe giant magnetoresistance (GMR) in electrodeposited nanogranular binary thin films with greater emphasis on the work carried out in the author's laboratory. High quality nanogranular thin films of Cu-Co system showing excellent metallic luster were galvanostatically electrodeposited in a single sulphate bath, under optimized processing parameters (i.e. deposition current density, bath temperature and pH). These parameters influence the composition and microstructure of the resulting films. Magnetoresistance measurements were carried out on both unannealed and annealed thin films deposited directly on n-Si or on conducting-glass/-AI203 (used as second electrode). The rnicrostructural and magnetic measurements have suggested that an optimum distribution of size and separation of the magnetic particles can result in maximum GMR in such systems.

1. Introduction Magnetoresistance (MR), the fractional change in resistance of a material in the presence of a magnetic field, is a well-known physical phenomenon. It can be defined as the relative change in resistance in presence of a magnetic field, MR = [(RH - Ro) / Ro ] x 100 = [(pH - p 0 ) / Po ] x 100 where PH (RH) is the resistivity (resistance) of the sample in presence of the magnetic field, p 0 (Ro) is the resistivity (resistance) of the sample at zero field. Recent measurements of very high values of MR in epitaxial multilayers, has redefined the effect as giant magnetoresistance (GMR). The GMR effect was simultaneously discovered in late 1980s by Peter Grunberg of Germany [1], and Albert Fert of France[2]. Grunberg's group reported a change (4%) in a Fe/Cr/Fe epitaxial sandwich structure and Fert's group reported a high value of 50% in (Fe/Cr/Fe)40 multilayered sandwiches at low temperatures. Since then a great deal of attention has been focused on the study of GMR in thin films. GMR has been observed in wide variety of transition metal magnetic multilayers [3]. Also, the GMR was found to be oscillatory with spacer layer thickness [4]. Although GMR was first discovered in antiferromagnetically coupled magnetic multilayers, subsequently it was discovered that the AF coupling and the ultrathin multiplayer structure are not essential. All that is necessary is that there 228

229

should not be ferromagnetic coupling to start with, and there should be some way to align the magnetic moments ferromagnetically [4]. It was in 1992, that for the first time GMR was observed [5,6] in granular films consisting of nanometric ferromagnetic (e.g. Co, Fe, NiFe) clusters embedded in a nonmagnetic (e.g. Ag, Cu, Au) metallic matrix. This was soon followed by a few other reports on granular systems [4,5]. Another variant of MR is colossal magnetoresistance (CMR). CMR was first observed [7] in 1994 in pervoskite manganites (Ri.xAx Mn0 3 ). The observed huge magnetoresistance (4 to 6 orders of magnitude higher) in the presence of very high magnetic field and at low temperatures was rightly called colossal magnetoresistance [8]. Besides high values of magnetoresistance, GMR is characteristically different from ordinary magnetoresistance. The GMR is always negative. It is anisotropic in multilayers and isotropic in granular systems. In some multilayers one magnetic layer moves in a small field, whilst the other does not, and is used as a reference magnetic moment. Such application specific multilayer structures are known as spin valves [9]. The research activities in GMR materials have picked up because of scientific and technical interests. The GMR effect is being applied to magnetoelectronics [10] for applications in information storage systems [2,5] and to magnetoresistive sensors [11]. Giant magnetoresistance randomaccess memory (GMRAM) is a nonvolatile memory that uses magnetic storage in a magnetic multilayer to store binary information and GMR effect to read stored data. The IBM has introduced new GMR read head products [12]. In thin film read heads for magnetic disk and tape recording, where response at low frequencies is not required, single MR elements are used and the output offset is removed by a high-pass filter [13]. The sensitivity of MR read heads could be increased by using nanometric magnetic multilayered (i.e. spin valve) structure [14]. With shrinking geometries forced on the (read head) designer by industry, the demands for higher density and demagnetizing effects in very narrow horizontal sensor stripes will become a major challenge. The advantages of magnetoresistive field sensors over others include high sensitivity, low source resistance, high operation temperature (up to 150°C), operation over wide frequency range, metal film technology, low sensitivity to mechanical stress and ease to miniaturize [11,15]. The first commercial GMR sensors, which were introduced in 1995, used multilayer GMR. The use of shields and flux concentrators allow the sensors to operate at fields of 10 - 100 Oe using GMR material with saturation fields of 200 - 300 Oe. Magnetic granular films have two immiscible phases with ultra fine magnetic particles dispersed in a nonmagnetic matrix, and constitute a special

230

class of artificially nanostructured materials. Freestanding ultra fine metallic particles are notoriously susceptible to environmental degradation (e.g. oxidation) and have a strong tendency to conglomerate into larger entities. But this difficulty is removed in granular films. These nanostructures can either exhibit GMR or tunneling magnetoresistance (TMR). It is known that the tunneling current between two metallic magnetic layers separated by a very thin insulating barrier (magnetic tunnel junction, MTJ) depends on the relative orientation of the magnetization in the adjacent magnetic layers. This magnetotunneling effect has been named Tunneling magnetoresistance (TMR) or Junction Magnetoresistance (JMR). Usually TMR is positive; this is called the normal TMR effect. For the inverse effect, which can occur in special cases when different magnetic materials are used on either side of the interlayer, TMR is negative. Besides GMR or TMR, the granular systems display very interesting novel magnetic and transport properties like spin glass behavior, superparamagnetism (SPM), extraordinary Hall resistivity, magnetoconductivity etc making them more attractive for fundamental scientific investigations. Although the origin of GMR is not yet fully clear, it was explained on the basis of spin dependent scattering of conducting electrons i.e. different scattering cross-section for spin up and spin down electrons [2]. Such spin dependent scattering is a well-known phenomenon in magnetic metals [6]. In multilayer films the magnetoresistance is explained on the basis of two currents model [16,17] in which the electrical current in the stack of layers is divided into two currents resulting from spin-up and spin-down electrons. These two types of electrons have different scattering probabilities at the interfaces and in the bulk of the layers. In general an electron will have higher scattering probability when its spin direction is opposite to the direction of local magnetization [18]. It is generally believed that spin-dependent scattering at the interface between the matrix and the magnetic particle is the most probable candidate [19,20] for explaining GMR in magnetic granular thin films. Since the discovery of GMR in granular films, several researchers have investigated a number of nonmagnetic magnetic metal combinations like Cu-Co[21-23], Ag-Co[24,25], Cr-Fe[26], CuFeNi[27] etc. Electrodeposition is the process of electrochemical precipitation through the reduction of metal ions at electrode/electrolyte interface under the influence of electric field, giving rise to the metal coating. Electrodeposition is one of the simpler, cheaper and older processes available for the fabrication of high quality thin metal and alloy films. Electrodeposition, also known as electroplating, came into existence in the early 19th century as a decorative and protective coating process. The process has recently been revived because of better controllability of the process parameters, improvement in film quality, comparable to those

231

prepared by UHV techniques [28], and large-scale industrial adoption as a high technology technique within the microelectronic industry for interconnections and packaging [29]. The electrochemical deposition also enables tailoring of electrical, mechanical, and magnetic properties, brightness, color, and resistance to corrosion of a material in thin film form. The advantages of the process not only include the low cost and the ambient conditions of pressure and temperature but also its ability to deposit thin films in complex geometries, where the conventional deposition processes would either fail or prove very difficult. Electrodeposition has successfully been used to prepare a variety of nanostructures including metal/metal superlattices with repeat distances down to 15 A [30] and nano-wires into pores of alumina and nuclear track-etched polycarbonate membranes [31-34] during last few years. It has made significant contribution towards the fabrication of microelectronic devices/components, magnetic devices and microelectromechanical systems (MEMS) [35]. In 1997, IBM reported [36] the fabrication of first working microprocessor using electroplated Cu interconnections. Since conducting cathodes are needed for the current flow through the cell, thin films are electrodeposited on the insulting substrates which are pre-coated with an extremely thin metallic layer. The use of doped semiconducting substrates with reasonably good conductivity, as a cathode, will eliminate the above-mentioned processing step (of pre-depositing a conducting layer). In order to make it easier to integrate electrodeposited magnetic nanostructures and conventional semiconductor electronics, some groups have studied the electrodeposition of magnetic metals, multilayers and spin valve structures on doped- Si and GaAs [37-42]. The binary system Cu-Co has attracted the maximum attention. In our laboratory also we have investigated GMR in granular Cu-Co films [43-45] and for the first time in electrodeposited Cu-FeNi films [46]. This is because largest GMR is observed in the multilayer of this combination and extensive studies have been carried out on this system in multilayers [47-50]. There are, however, few reports on electrodeposited granular Cu-Co [51-53] thin films prior to our work. In addition, the equilibrium phase diagram [54] shows that Cu and Co are essentially immiscible below 500°C. Thus, at equilibrium a mixture of Cu and Co phases is expected. Annealing of this metastable alloy at an elevated temperature can lead to the formation of a granular magnetic system consisting of single domain ferromagnetic Co-rich clusters in a nonmagnetic Cu-rich metallic medium. Recently Anton and coworkers have again reported the electrodeposition [55] and laser ablation [56] of Cu-Co system, and Champion and coworkers [57] the fabrication of Cu-Co nanomaterial bulk material by cryo-melting. There are couple of more reports in the literature on the

232

electrodeposition of Cu-Co thin films [58-59]. Typical values of GMR in some of the systems, as reported in literature are summarized in the following table: Table 1. Representative values of GMR in various samples in thin film form

Sample Cu 8 ] Co,9 CU80CO20

Cu 70 Fe 3 o Cu 8 oFeioNi 10 Co 30 Ag 7 o Cu 80 Co 2 o Ni-Co-Cu/Cu C o x Cllioo-x C1190C010

Cu 70 Co 3 o Co 18 Cu 72 Cu 8 4 Co ] 6 Co 2 oCu 80

Method of preparation Sputtering Sputtering Sputtering melt-spun mechanical alloying (bulk) Sputtering Electro-deposition Electrodeposition Laser Ablation Cryo-melting Electro-deposition Electrodeposition Ion beam sputtering

GMR (%) 22 17 9 17 7.7

Temperature (K) 10 5 5 4.2 5

H (Tesla) 2 5 5 7 1.5

Ref.

15 7 5 10 0.6 4 4.5 6.2 3

4.2 300 300 77 300 4.2 300 300 300

14 0.8 1.3 0.9 1.0 1.0 1.5 1.5 1.5

[261 [50] [53]

[51 [6] [61 [171 [24]

[561 T571 [591 [651 [66]

As already pointed out electrodeposition of heterogeneous granular alloy various workers have studied films consisting of ferromagnetic granules in a nonmagnetic metallic matrix. However, the process of electrodeposition is yet to be fully understood. In our case the deposition parameters were varied and optimized in order to obtain good quality, adhesive and compact thin films of Cu-Co. Several substrates namely, Cu-coated glass, ITO coated glass, Cu-coated alumina, doped semiconductor substrates like n-Si, metals like Cu and Ti etc. A growth mechanism has been used for depositing the films. The role of substrates and various process parameters in effecting the film properties have been investigated by employing several analytical techniques. The magneto resultant and macrostructure have also been corrected. 2. Methods Electrodeposition of Cu-Co films was carried out at a constant current density, using a programmable constant current source (Figure 1) from an aqueous electrolyte of sulphates of Cu and Co. Using different amounts of Co, and keeping Cu unchanged, the bath composition is varied. More details regarding film preparation are given in [43,60,61]. The films were vacuum annealed in Ar atmosphere at different temperatures (300-700°C) and for different durations (15 in - 1 hr). The thickness of the electrodeposited films is estimated from the

233

deposited mass according to Faraday's law (assuming 100% current efficiency) i.e., m = /e//96500 where m is the mass deposited in grams, I represents current in amperes, e and t are chemical equivalent weight and time in seconds, respectively, and then compared with the value obtained directly by employing a surface profiles i.e. Talystep (Taylor-Hobson, UK). An atomic absorption spectrometer was used to analyze the composition of the films. The technique of energy dispersive analysis of X-rays (EDAX) was also employed for the compositional analysis.

Keithley 224 programmable constant current source Voltmeter or Oscilloscope .

.

•

I 1 f

Reference electrode

Anode

Cathode

Figure 1. Schematic of electrodeposition process [Courtsey G.R. Pattanaik, Ph.D. Thesis (2002)]. The crystallographic structure of the films is investigated by glancing angle X-ray diffraction (GAXRD), using Cu Kct radiation obtained from a rotating anode (model RB-RU200, Rigaku, Japan) operating a 40 kV. The average grain size has been estimated using the classical Scherrer formula. Dhk, = ia/(p m 2 -p s 2 ) 1 / 2 cos2e where Dhki is diameter of the crystallite corresponding to the peak (hkl), K, X and (3m are shape factor (-0.89), wavelength of X-rays source (0.154 nm) and FWHM of the peak (hkl), respectively. ps is FWHM of the peak (near the same 20 value of the peak under consideration) of the standard sample to take care of the instrument broadening.

234

The surface topography of the films is studies using a scanning electron microscope (STM-RHK 635, RHK Technology Inc., USA) [60]. Room temperature magnetoresistance (MR) measurement in four-terminal van der Pauw geometry was carried out using a Keithley 224 programmable current source and a Keithley 181 nanovoltmeter at magnetic fields up to lOkOe. Both, the current and magnetic field were in the plane of the film and parallel to each other. A closed-cycle He cryostat (APD Crygoenics) was used for lowtemperature MR measurements in the same four-probe configuration at a fixed magnetic field of 3kOe. The room-temperature magnetization of the films was measured using a vibrating sample magnetometer (digital magnetic systemsDMS 880) with an applied magnetic field H in the range ±7 kOe[43]. 3. Results and Discussion 3.1. Electrodeposition of Nanogranular Cu-Co Thin Films on Conducting and Semiconducting Substrate In order for the two metals to co-deposit, their ions must coexist in an electrolyte in which the individual deposition potentials are the same or nearly the same. If the deposition potentials of the two individual metal components are far apart, the use of suitable complexing agents allows the codeposition of both the metals at an intermediate potential. This is because when a complexing agent is added to the solution, the complex ions are formed which take part in the deposition process. The standard electrode potentials of Cu and Co are - 0.277 and +0.34 V respectively. To codeposit the two we have used tri-sodium citrate as a complexing agent. The deposition parameters like bath composition, deposition current density bath temperature and pH of the electrolyte, were varied to obtain good quality, adherent, compact thin films of Cu-Co with metallic luster on n-Si substrate [60], Let us name these Cu-Co films on Si as 'S'. An increase in the pH from 4.7 to 6 resulted in an increase in the Co concentration in the film with improved physical properties like adhesion and luster. The SEM micrographs revealed a decrease in grain size with an increase in pH. At pH of 4.0, the observed coarse surface topography was quite remarkable, and it changed to a smooth surface at pH of 5.0. A fine-grained microstructure was observed at a pH value of 6.0 and hence, most of the depositions were carried out at this value. Concentration of Cu2+ and Co2+ in the bath is an important variable governing the composition of the electrodeposited films. The ratio of the two metals in the film is usually different from that in the electrolyte [62]. It is noted that Co concentration in the film is always less than that in the plating solution for all the concentrations of Co2+ employed by us. This indicates that the more

235

noble metal (Cu) is preferentially depositing in the film. The X-ray diffractograms revealed a single phasic face-centered cubic (fee) structure corresponding to an alloy of Cu and Co with the values of average lattice parameter lying between those for pure fee Cu and fee Co, which are 3.615 and 3.545 A, respectively. As the Co2+ concentration is increased in the bath, the peaks in the diffractograms shift toward the higher diffraction angles, and hence average lattice parameter shifts toward that of pure Co. This too implies that the concentration of Co is increased in the film with the increase in concentration of Co2+ in the bath. The linear variation of lattice parameter with Co concentration implies a solid-solution-like behavior of the films following Vegard's law. Since Cu and Co are practically immiscible, the as-prepared electrodeposited film is believed to be a supersaturated metastable solid solution of Cu and Co [43, 53, 63]. However, our magnetic and magnetoresistance measurements clearly indicate that the phase segregation exists even in the as-deposited films. Therefore, it is proposed that the Cu-rich matrix contains the ferromagnetic entitles (Co) in the form of very fine clusters such that these are not detected in the X-ray measurements, and this composite system exhibits a linear variation of lattice parameter with Co concentration in the film. The average crystallite size as deduced from the XRD line widths (full width at half maximum, FWHM) was 16, 15, 15 and 14 nm for the films deposited from electrolytes with [Co2+]=0.06, 0.072, 0.09 and 0.11 M, respectively. The Cu-Co films, are deposited from a bath with [Co2+]=0.072 M at current densities in the range 2-8 mA/cm2 at a pH of 6.0 and at 20CC. The Co content of the film increases from 12 to 45-atom % with an increase in current density from 2 to 8 mA/cm2. The rise in the deposition potential with increase in current density, as determined by chrono-potentiometric measurements, supports the reduction of more Co2+ at the cathode, thereby increasing the Co concentration in the film. The average grain size in the films deposited at 2,3, and 8 mA/cm2 is determined (from XRD linewidths) to be 17, 16 and 13 nm, respectively. SEM micrographs of the films also revealed a decrease in the grain size with the increase in the current density. This is because at higher current density the deposition rate is high and hence the ad atoms get largely immobilized and are incorporated in the film with little surface migration, thereby limiting the grain size [60]. Electrodeposition of Cu-Co alloy thin films is carried out at different bath temperatures ranging from 20 to 50°C at 3 mA/cm2 in a solution with [Co2+]=0.11 M and pH of 6.0. The XRD data revealed that there is more incorporation of Cu with increase in bath temperature. The grain size increases with the increase in bath temperature. Also the Cu concentration increases with the increase in bath temperature. The enhanced Cu concentration in the film can

236

be explained on the basis of the simple diffusion theory. The considerable variation of composition of Cu-Co alloy films due to increase in the bath temperature supports that the codeposition is a regular one. The nanogranular films of Cu8iCoi9 of two different thickness i.e. 60 and 120 nm, deposited in a bath with [Co2+]=0.072 M at a current density of 3mA/cm2 are investigated using scanning tunneling microscopy (STM). From the STM images it was inferred that the grain size at the surface increases with thickness. The surface roughness also increased from 10.2 to 27.3 nm as the film thickness increases from 60 to 120 nm. Also, the surface roughness is observed to increase from 27.3 to 47.2 nm as the current density is increased from 3 to 8 mA/cm2. It may be noted that the films for XRD studies were 300 nm thick, but the grain sizes calculated for such thicker films from XRD line widths are much lower in comparison to those seen by SEM or STM of 120 nm thick films. The grain sizes observed in transmission electron microscopy (TEM) studies are comparable to those estimated from XRD. This indicates that the granules seen by SEM or STM consist of a number of nanograins, as has been reported in another study [64]. Thus, the growth of the films seems to be via coalescence of smaller islands to form larger islands. It is well known that the roles of substrate temperature and deposition rate are quite important in determining grain size, surface roughness, and crystallinity of the thin films grown by physical vapour deposition (PVD) technique. In the case of electrodeposition of Cu-Co thin films, it is found that a small change (20-50°C) in bath temperature could yield changes in microstructure similar to those observed by large increase in substrate temperature. Similarly, the increase in the deposition current density decreases the grain size much like the effect of increase in supersaturation (deposition rate) in the PVD technique. It can, therefore, be inferred that bath temperature and deposition current density play the same role in electrodeposition as substrate temperature and supersaturation in physical vapour deposition techniques. Furthermore, the preferential deposition of more noble metals is similar to higher vapour pressure constituents in PVD techniques [60]. The CU75C025 films electrodeposited on glass or alumina coated with a vacuum evaporated copper film (to act as a cathode) (Films C) revealed a district single phasic FCC structure, implying that the as grown film is a metastable alloy of Cu and Co. With increase in annealing temperature to 700°C distinct peaks corresponding to Cu (fee) and Co (fee) are observed in the X-ray diffractograms. The analysis of the diffractograms is presented in Figure 2.

237 jo*: -

D

3.60D u

l_ 4> 4-1 4>

E CO

nU \

metastable phase

3.58-

Curich

\\ \ \ \1 \\ \

L.

CO Q. o>

ttic

Cu

n-a

*•>

s
20wt%) only broadened double film peak is observed.

(53

30,0

40,0

50,0

60,0

70,0

80,0

29(°) Figure 1. X-ray diffraction pattern of Ag-Ni alloys with different Ni concentrations.

The Ag-Ni phase diagram has a simple form without solid solutions and intermediate compounds, the maximum solubility being under 1%. If one assumes that the alloy is single-phase for a larger composition range as in the Ag-Cu system [12], this would explain the X-ray data but contradict the currently known Ag-Ni phase diagram. In order to establish the real grain structure, we employed electron microscopy techniques. We performed electron diffraction measurements on a 30wt% Ag-Ni specimen at its edges. The typical electron diffraction pattern is given in Fig.2. It shows unambiguously that the grain size is small (Ag+N_p+Agp,

p = \,2,...,N-2.

(l)

For doubly charged clusters, the decays can proceed via two different processes. The first one is the evaporation process

Ag2N+Âg2N+_p+Agp,

p = \,2,...,N-3

(2)

in which one of the products is neutral; and the second one is fission into two charged products

Ag2N+Âg+N_p+Ag+p,

p = 2,3,...,[N/2].

(3)

273

In evaporation processes, the negativity of the difference between total energies before and after fragmentation in a specific channel, Dz(N,p) = Ez(N-p) + E°(p)-Ez(N) (4) is sufficient for the occurrence of the decay in that channel. In the above equation, Ez (N) and E°(N) are the total energies of z-ply ionized and neutral N -atom clusters, respectively. However, in fission processes a negative value for the difference energy is not a sufficient condition for the decay of the parent cluster. This is because; the competition between the short-range surface tension and the long-range repulsive Coulomb force may give rise to a fission barrier. The situation in a fission process is shown in Fig. 1.

dO

Oo

N2,z2\D

„ O N1,z1

Reaction Coordinate

Figure 1. Fission path for the decay of a parent cluster.

In Fig. 1, the fission of a Z-ply charged TV-atom cluster into two clusters of respective sizes NlfN2 = N- JV, and respective charges zx,z2 =Z — z, is schematically shown. Qf is the energy release, Bc is the fusion barrier which is the maximum energy of the Coulomb interaction of two positively-charged conducting spheres, taking their polarizabilities into account. Bf is the fission barrier height which is defined as Bf=-Qf+Bc.

(5)

274

The coulomb interaction energy of two charged metal spheres has been numerically calculated using the classical method of image charges [13]. The calculations show that the maximum of the interaction energy is achieved for separations d0 > i?, + R2. The equality applies for equal cluster radii and charges. The situation for different sizes is shown in Fig. 2.

10

12

14

16 ' 18

20

22

24

26

28

30

d (a.u.)

Figure 2. Coulomb energy as a function of distance.

The most favored decay channel in evaporation processes is defined as the channel for which the dissociation energy attains its minimum value

Dz(N,p*) = mmD(N,p)

(6)

and the most favored decay channel in fission processes is defined as the channel for which the fission-barrier height attains its minimum value

Bf(N,p,)

=

mmBf(N,p).

(7)

The organization of this paper is as follows. In section II we explain the method of calculating the total energies of the clusters. In section III, we discuss the results, and finally, we conclude this work in section IV.

275

2. Calculational Schemes The total energy of a given cluster is obtained by solution of the self-consistent Kohn-Sham (KS) equations [14] in the density functional theory [15] (DFT) with local spin density approximation (LSDA) for the exchange-correlation energy functional. In the context of the SJM, the average energy per valence electron in the bulk with density parameter rs and polarization g is given by [16]

e{r, ,$,rc) = t, (r,, g) + sxc (rs ,g) + wR (rs ,rc) + eM (r,)

(8)

where, ts (rs, g) and sxc (rs, g) are the single-electron non-interacting kinetic and the exchange-correlation energies, respectively. For the correlation energy we use the Perdew-Wang parametrization [17]. In a z-valent metal the average Madelung energy, sM is defined as £u — -9z 110r0 , in which r0 = z rs is the radius of the WS sphere. All equations throughout this paper are expressed in atomic units (h = e = m = 1, the units of length and energy are bohr and hartree, respectively). In Eq. (8), the polarization is defined as g = (n^ — Wj,) /(«f + Hj,) in which n^ and n^ are the spin densities of the homogeneous system with total density n = (n^ + n^ ) . The quantity WR is the average value (over the WS cell) of the repulsive part of the Ashcroft empty core [18] pseudo-potential w(r) =

+ wR (9)

yvR(r) =

r

+-0{rc-r)

and is given by wR = 3rc / 1rs , where z is the valence of the atom, 0(x) is the ordinary step function which assumes the value of unity for positive arguments, and zero for negative values. The core radius is fixed to the bulk value, rc

by setting the pressure of the unpolarized bulk system equal to zero

at the observed equilibrium density n = 3 / 47r[rs ] :

276

dr.

= 0.

£(rsArc)

(10)

Here, rs = rs (g = 0) is the observed equilibrium density parameter for the unpolarized bulk system' and takes the value of 3.02 for Ag. The derivative is taken at fixed rc, and the solution of the above equation gives

B\in,

rf(r?) = Hr?y

2ts (r, ,0) - sx {rs ,0) + r, — ec (r, ,0) - eM (r,) or. (11)

The SJM energy for a cluster in the LSDA is given by {eM{rs) + wR(rs,rcB))\d'rn+{r) (<Sv)BS(rI,r;)(iV^-r)[«W-«t(r)]

+ (12)

where, E

JM ["t >"l'n* 1 = T.K•"J

+ E*c K . n J + - jd3r ([nf,nJ;r)[n(r)

- n+ (r)]

(13) and J

\r — r \

(14)

r-r

Here, n = n^ +n^ and n+ is the jellium density. 6{R — r) takes the value of unity inside the jellium background and zero, outside. Thefirstand second terms in the right hand side of Eq. (13) are the non-interacting kinetic energy and the exchange-correlation energy, and the last term is the Coulomb interaction energy of the system. The quantity ( A ) is the average of the difference potential over the Wigner-Seitz cell and the difference potential, „'«rx^ „ :

v^>rio; î r ^ K \^ - :-" '

--.«v* v* *v *"i- - " . * *'*^S& t^v; ,-v ' f -

30

^rf-M-/ |\ ;>

s "s > * "" ^ :V ^ -V*.- \ ^ ^ *v#^v ; ; r ; v

'* :

s / ^^K ^ h> -H? \^0-v^V .

* 1

•w^

+^v

^ A ;

j

• i

v

-s

v

s

!

;; w; 40

i

*Lk

l

«l*m

:;>!*'

15 fan

Figure 3. AFMtopographsof discontinuous films. Each row represents a pair of simultaneously deposited samples. The deposition time is specified for each row. The left column shows perytene films deposited on SAM-substrates, therightcolumn depicts films deposited on Au-substrates. The image size for each column is indicated.

291 .

87-

6

1 5 5 0 « m ) . The chosen materials for this filter were cryolite (Na3AlF6, n = 1.35) and cadmium telluride (CdTe, n = 2.916). The spectrum of the designed filter which consisted of 12 layers is plotted in Fig. 2,

313

and the spectrum of the fabricated filter is shown in Fig. 3. Since part of the light is absorbed and scattered in the glass substrate, the actual spectrum of the filter (Fig. 3a) has lower transmittance compared to the designed one. However, when the effect of the substrate is eliminated (Fig. 3b), the actual and designed spectra become quite close. This was done by placing a similar glass substrate (with no coatings) in the "reference" position in the double-beam photospectrometer, while recording the spectrum of the actual filter. 100

« 60

50 40 J20 10 h 1300

1350

WOO 1450 1500 1550 Wavelength (nm)

1600

1650

f. 00

Figure 2. Designed spectrum of the edge filter with cutoff wavelength at 1550.

1M

:

•

. • ••"

<S~

;

(b) fa)

, -'

60

1400

1500

1«B

ITflO

Wavelength (nm) Figure 3. Transmittance spectrum of the fabricated edge filter, (a) without eliminating the substrate effect, and (b) after removing the substrate effect (as explained in the text).

314 As a few examples, designed spectra of an antireflection filter, a highreflectance filter, a polarizing beam splitter, and a narrow-bandpass filter are given here. The antireflection filter (Fig. 4) consists of 14 layers of ZnS and cryolite on a glass substrate. This filter reduces the normal 4.25% reflectance of the glass substrate to less than 0.1% over the whole visible range. Figure 5 shows the spectrum of a high-reflectance filter (i.e. all-dielectric mirror), designed to cover the whole visible region. This consists of 41 layers of ZnS and cryolite, and reflects more than 99.9% of light over almost all of the visible wavelengths (as compared to ~ 93% for a freshly-evaporated aluminum frontsurface mirror). The polarizing beam-splitter filter (Fig. 6) was designed for a central wavelength of 632.8 nm and incidence angle of 45°. With 9 layers of ZnS and cryolite, it separates S and P components of light (at K=632.S nm) to around 99.3%. Figure 7 shows spectrum of a narrow-band-pass filter with central wavelength of 1550 nm. Much narrower band-pass filters can also be designed, but this filter has an almost rectangular shape, which is a good feature for its use in optical communication applications. It consists of 29 layers of CdTe and MgF2, and has transmission of more than 99% at its central wavelength. The actual spectra of the fabricated filters were generally quite close to the designed ones. 101

n

i

1

1

1

s

1

?

Wavdeqgh (nm) Figure 4. Spectrum of an antireflection filter designed for the visible region.

r

315 100

VT"

T

1

1

i—

—r • —

" i "V "

I

90

eo

i so ts.

:

:

:

:

i

:

:

:

:

I Li! \ L±.

H

; • • ' •

20

':

:

:

i

i

i

i

i

i

400

450

500

' : •

350

\

i

•

,

I ! |-

10 300

:

550

i 650

600

i . 750i f , , 800 700

Wavelength (nm) Figure 5. Spectrum of a high-reflectance filter designed for the visible region. 100

500

600

700

Wavelength (iim) Figure 6. Designed spectrum of a polarizing beam-splitter filter.

800

900

316 I

-

'1

1 "1-1

j

:

:

•

i

;

:

:

:

1 !

i ]

"1

1

1

:

i

:

i i

i i

i i

^-i 1S00

i 1700

i 1800

[j.. L

!

!

\

«00

1400

1500

Wavelength (ran) Figure 7. Designed spectrum of a narrow-bandpass filter with central wavelength at 1550 nm.

5. Conclusions Using the software package developed in Ferdowsi University, various multilayer thin-film optical filters were designed. The software based on genetic algorithm was found to be a strong tool for the design of optical filters with any desired optical characteristics. Usually, this method produced very good results. Its only drawback was the long computer time needed for achieving an optimized design. Despite shortcomings of the vacuum evaporation technique (as explained in section 3), it can be used satisfactorily for fabrication of multilayer optical filters. With careful monitoring of film thicknesses, the actual spectra of various fabricated optical filters were quite close to their respective design spectra. Acknowledgements The authors would like to thank the National Scientific Research Council of Iran and Iran Telecommunication Research Center for their financial supports. References 1. 2. 3.

H.A. Macleod, Thin Film Optical Filters, 3 rd edition, (Institute of Physics Publishing, London, 2001). J.D. Rancourt, Optical Thin Films User Handbook, (SPIE Optical Engineering Press, Bellingham, 1996). J.A. Dobrowolski, "Optical properties of films and coatings", in Handbook of Optics, M. Bass, ed., (McGraw-Hill, New York, 1995), pp. 42.1-42.130.

317 4.

5. 6. 7. 8.

9. 10. 11.

12. 13. 14. 15.

16.

17.

J.A. Dobrowolski, "Usual and unusual applications of optical thin films an introduction," in Thin Films for Optical Coatings, R.F. Hummel and K. H.Guenther, Eds. (McGraw-Hill, New York, 5,1995). J.A. Dobrowolski, Optics News 6, 24 (1997). R.J. Pegis, Journal of the Optical Society of America 51, 1255 (1961). E. Delano, Journal of the Optical Society of America 57, 1529 (1967). L. Sossi, "A method for the synthesis of multilayer dielectric interference coatings," Translated by J.A. Dobrowolski, (National Research Council of Canada, 1974). J.A. Dobrowolski, D. Lowe, Applied Optics 17, 3039 (1978). J.A. Dobrowolski, S.H.C. Piotrowski, Applied Optics 21, 1502 (1982). S.H. Kazemi-Riabi, M.M. Mirsalehi, S.H. Keshmiri, 8th Iranian Conference on Electrical Engineering, Isfahan, Iran, (Proc. of Electronics papers, 2000 in Persian), pp. 67-74. W.H. Southwell, Applied Optics 24, 457 (1985). D.E. Goldberg, Genetic Algorithms in Search, Optimization and Machine Learning (Addison-Wesley, MA, 1989). E. Michielssen, S. Ranjithan, Mitra, IEE Proceeding Journal 139, 413 (1992). M. Shokooh-Saremi, M. Noorian, M. Mirsalehi, S.H. Keshmiri, 6th Iranian Conference on Electrical Engineering, Tehran, Iran, Proc. of Communications papers, 2.31,1998) M. Shokooh-Saremi, S.H. Keshmiri, M.M. Mirsalehi, M.M. BagheriMohaghgheghi, 7th Iranian Conference on Electrical Engineering, (Tehran, Iran, Proc. of EM Fields and Applications papers, 139-145, 1999). B. T. Sullivan, et al., Vacuum 51-4, 647 (1998).

THIN FILMS FOR OPTICAL RECORDING A. KIKINESHI University of Debrecen, Hungary E-mail: kiki@tigris. kite, hu The main photophysical effects and the types of suitable thin film materials for amplitude-phase optical recording, surface pattern formation are reviewed. The photoinduced phenomena in amorphous chalcogenide layers are discussed in more details since these materials are excellent models for a number of fundamental processes of optical memory and are useful for a variety of applications in optoelectronics, data recording.

1. Introduction The development of analog-type optical recording, art photography during the last century was strongly connected with the photochemistry of silver halides and resulted in a high-quality color photography for the mass-media. But it is not suitable for high-density, high-speed digital data recording, holography and technical photography (photolytography) as well. The increasing interest to the nowel scientific and technical problems of optical recording undoubtedly is caused by the development of information technologies, first of all by the demands of data storage and processing. A number of non-silver photomaterials were developed on the basis of photophysical processes in organic and inorganic materials like polyvinilcarbasol, selenium, amorphous hydrogenated silicium, LiNb03 and others, which satisfied the requirements of archival or reversible optical relief formation mostly in a one-step recording-readout process, or in a repeated cycles of reproduction [1-5]. Electrophotography, laser ablation, lightinduced structural transformations, photorefraction, phase transitions should be mentioned amongst the well-known processes. Magnetooptical effects and materials like MnBi, Y3Fe20i2 were considered as promising for digital optical recording [1], but the simplicity of phase-transition based recording processes and the cheapness of appropriate materials and technologies made the last more acceptable for production. The general problem of optical recording, data storage at the present time can be determined as more, faster, reliable and cheaper. The physical limits of these requirements are known but really they do not restrict, most probably stimulate the development of above mentioned materials nowadays. These problems, physical processes and materials, connected with thin film technology and applications are the content of this lecture.

318

319 2. Optical Recording: Basic Physical Processes Basically the interaction of light (quantum energy hv * 1-3 eV) with a matter (with thin solid films in our case) results in the direct excitation of electrons and in the direct or, more probably, indirect stimulation of atomic displacements, which in turn cause the rearrangement of chemical bonds, change of the interatomic distances. Both are responsible for the stimulated changes of defect states (photocroms), electrical conductivity, polarization (photoconductors, ferroelectrics), mechanical parameters (viscosity, stress), for different phase transitions (amorphysation-crystallisation, evaporation) and other photophysical processes.. These processes determine the stimulated changes of optical parameters of the recording material i.e. the optical memory effects. The schematic diagram of some possible recording processes are presented in the Fig.l together with the most known types of appropriate materials. light

tttt material

election mobility

-> atomic mobility

electrophotography (a-Se, a-Si:H, ZnO, organics)

phase transitions, ablation (VOi, SmS, polymers)

photocroms (KC1, glasses )

structural transformations (amorphous-amorphous, crystalline- amorphous, density) (AsSe, Ge-Te-Sb, GeS2)

photorefractive effect (LiNb03:Fe'3,Sn2P2S6) magnetooptical effects (MnBi, CrTe, Y3Fe2Ol2)

diffusion, change of composition (Ag-As2S3, Se/As2S3 multilayers)

THE CHANGES OF OPTICAL PARAMETERS : Aa (absorption). AR (reflection). An (refraction). AP (polarization) amplitude or amplitude-phase recording Figure 1. Possible recording processes and some appropriate materials.

320

The next important question is the conformity of technical requirements, desirable limits of stimulated changes of optical (or electrical, in electrophotography) parameters with the realizable parameters in the given material or thin film structure. The next parameters should be provided and optimized for a given situation: a 0 =10' - 105 cm"1, Aa/a 0 = 10 - 100(photocroms, AsSe for example), n = 1,5-3,0, An=10- 4 -10" 1 (LiNb03, As 2 S 3 ), Ro = 5 - 20 %, AR/R = 0,1 - 0,5 (V0 2 , Ge-Te-Sb, MnBi), AP: (magnetooptical, electrooptical materials), a: 10"10-10"15Ohm"1cm"1 (electrophotography, a-Se, ZnO), Sensitivity S = 10 "8 - 10 5 J/cm2 in the spectral range AX: Xt < 0.3 - l.Oum £ X2, Resolution: 10 3 — 10 6 mm"1, Stability or reversibility: 10 +12 s — many years. 3. Optical Data Recording: Basic Configurations There are two basic possibilities for modern optical data recording: in the digital and holographic form. It means that one can divide the picture to points (bits) and write them continuously in a linear or matrix form (Fig.2), or create the interference of two coherent beams ( the object and the reference beams) and write this distributed intereference picture (the hologram) in the recording media (Fig.3).

Figure 2. Digital optical recording.

Both types of optical recording have a number of advantages and difficulties, which are connected with the parameters of recording materials, data areal density and transfer rates, technical conditions.

mirror Figure 3. Holographic recording.

object

hologram

321

4. Thin Films for Optical Recording: Technology and Parameters A wide number of thin film deposition methods are used for production of light-sensitive structures, which are rather well known in microelectronics. These are: thermal or electron-beam evaporation in vacuum (chalcogenide semiconductors, glasses), magnetron sputtering (chalcogenide glasses, amorphous silicon, metals, oxides), chemical vapour deposition (a-Si:H), spin coating (organics). The critical parameters usually are the initial structure, stresses, adhesion, surface roughness, optical homogenity. The influence of the deposition conditions, as well as of the structure and composition on the parameters of amplitude-phase optical recording may be demonstrated for the versatile type of optical recording media - chalcogenide glasses (see Fig.4).

OPTICAL RECORDING PROCESSES IN CHALCOGENIDE GLASSES ILLUMINATION

ELECTRON PROCESSES

->

DEFORMATION, CRYSTALLISATIONRELAXATION -> AMORPHYSATION

DIFFUSION

CliG Metal

?^£„ •

im lw Charge

innnnnn

-

«

Recoriliii|>

Recording

j^sassas

3

Recording

Reamiing

> Aj4"iw*rrj4wv4"'

Development Etching Kelief

Figure 4. Optical recording processes in chalcogenide glasses.

Two typical examples of the novel optical recording processes in lightsensitive chalcogenide glasses are presented in the next figures. The lightsensitive material is a Ge-Te-Sb-type chalcogenide glass, the focused light pulses are provided by the diode laser.

322

(a)

intensity record in erasing readout time

(b) track be tore erasing

1

after recording

1

(c) detected signal

o o o

o

1

o o

1

o

TTTT time

Figure 5. Schematic process of digital data recording on the rewritable optical dies.

The actual problems of the optical disc development are to increase the areal density (possibly up to 5>85 % with good figure of merit. 3.2. Structural Properties Figure 2 show the X-ray spectra of Al doped ZnO films deposited on glass at different Al doping level. An undoped x-ray spectrum is added for the comparison. From the figure it is seen that all films are polycrystalline and correspond to a Wurtzite structure with lattice parameter values a = 3.238 A and c = 5.205 A [1]. No metallic Al or its oxides are observed. From the X-ray spectra, it also appears that the preferred orientations of planes are sensitive to

347

1.2-i 1.0-

\ \

•£ 0 . 8

450

550

500

+ A t P H = 0.8x10 ":'itiibar • AtT

s

= 523 K

10 19

t

:

W i. 03

a -S

10 18 -

05

-

Vl

10 17

++

o

t+ + +

• •

+

a V

Q vacuum chamber

Mixed gas

Figure 1. The schematic of gas mixture system. Table 1. The correlation between deposition parameters and some physical properties of Ti02 magnetron sputtered thin films. Total Pressure (mTorr)

Thickness (nm)

4

305

8 16 32

280 280 240

Phase

Rutile + Anatase Anatase Anatase Amorphous

Grain Size (nm)

Eg on glass (eV)

Eg on FTO (eV)

28

3.12

3.14

28 30 -

3.17 3.18 3.29

3.22 3.22 3.30

2.2. Spray Pyrolysis Deposition A solution of alcoholic titanium (IV) oxide acetyl acetonate TiO[C5H702]2 (Ethanol, 100 cc; HC1, 5cc) of 0.08 M concentration was prepared. 25 cc of this solution was sprayed onto the heated glass substrates (300, 350, 400 and 450°C). A spraying period of 1 s was followed by an interruption time of 30 s to avoid excessive cooling of the substrate during the spray. The spraying process was performed with a glass nozzle using compressed air.

365

2.3. Characterization XRD analysis of Ti0 2 thin films was carried out on Bruker AXS D8 Advance X-ray diffractometer with Cu Kcc radiation. The film thickness was measured with a Sloan Dektak IIA profilometer. Optical transmission of the films was measured with an UV-visible Agilent 8453 spectrometer. The surface morphology of the films was studied by atomic force microscopy (AFM) in a Jeol JSPM-4210 microscope. We characterized photocatalytic properties of Ti0 2 thin films with decomposition of methylene blue (Ci6H18N3SCl-3H20). We used the same method reported by Zeman and Takabayashi [2]. Ti0 2 films were immersed in methylene blue solution 1 mM for 1 h and afterwards dried for 30 min in dark room. The surface of Ti0 2 covered with methylene blue was irradiated with UV light from a 20 W sterilamp for 30 min. From comparison of optical transmittance of 650 nm light before (T,) and after (7}) the UV irradiation, we obtained a quantitative evaluation of the degradation of the methylene blue (A ABS = In TilTf). We further characterized photocatalytic activity of these films with decomposition of an aqueous methylene blue solution (MBS). The films with surface area of 7.5 cm2 were dipped into a 10 ml of MBS 0.05 mmol/ml and irradiated with a sterilamp. The transmittance of the solution at 650 nm was measured with 2 h interval for a total irradiation time of 10 h. 3. Results and Discussion 3.1. Ti02 Film by Magnetron Sputtering In Figure 2, we present the evolution of Ti0 2 films from X ray spectra for different total gas pressure values; it can be observed that crystallinity of our films increases with decreasing pressure. While the anatase phase was observed in all films, a rutile phase was found only at 4 mTorr. In Figure 3, we observe the behavior of Ti0 2 films when were deposited on different substrates for a 16 mTorr gas pressure. The strongest peak corresponds to Sn0 2 (on FTO substrate). The spectrum in the lower part of the graph corresponds to Ti0 2 deposited on glass substrate. The substrate effect on the structure is due to the mobility of the ad-atoms on the substrate surface which is different. This fact changes the type of nucleation on the substrate. We can see that same thickness Ti0 2 thin films on tetragonal substrates (FTO) are more crystalline than films on amorphous substrates (glass).

366

Figure 2. XRD patterns of Ti0 2 thin films on glass prepared by D.C. reactive magnetron sputtering at different total pressures of an equimolar gas mixture Ar/0 2 .

on FTO substrate

AC101)

AI200) | AI211)

,

> on glass substrate A(101)

A(Z11)

40 50 60 2H Figure 3. XRD patterns of Ti0 2 thin films on different substrates prepared by D.C. reactive magnetron sputtering at 16 mTorr.

10

20

A(200)

30

367

The particle size was calculated from anatase (101) reflection and rutile (110) reflection, using the Scherrer equation [6]. The average particle size is around 30 nm for all films. For a 4 mTorr pressure, we found a mixture of anatase and rutile phases in T0 2 films; it is possible to calculate the weight percentage of the anatase phase, WA, using the following equation [5]: W A = l / [ 1+1.265 I R /I A ]

(1)

where IA denotes the intensity of strongest anatase reflection and IR is the intensity of strongest rutile reflection. The films prepared at 4 mTorr have a percentage of the anatase phase equal to 25% as calculated with relation (1). In Figure 4, we present the optical transmittance spectra for Ti0 2 films deposited by magnetron sputtering on glass and glass coated with FTO substrates; the highest values are obtained in both cases when a gas pressure of 32 mTorr was used. In the visible range the transmission of the films on glass is around 80%, whereas the Ti0 2 films on FTO present a decrease in the transmission due to the absorption by FTO. The transmission decreased as the wavelength decrease due to the fundamental absorption of the light [5]. The optical band gap of our Ti0 2 thin films was calculated according to the method described by Mardare et.al. [5]. An optical band gap, Eg, of about 3.2 eV has been observed by the films prepared at 8 and 16 mTorr due to the anatase phase, these values are agrees with the values reported in literature [7]. The 3.12 eV value for Eg of thin films prepared at 4 mTorr is situated between the values reported for anatase and rutile phases [7,9] because of its mixed structure. Ti0 2 films prepared at 32 mTorr present a highest value of Eg due to amorphous nature [10]. 3.2. Spray Pyrolysis Deposition Structure properties of titanium dioxide thin films at substrate temperatures varying from 300 to 450 °C were studied using the X-ray diffraction analysis. Anatase structure with the (101) predominant plane of crystallization has been identified for the films deposited at 400 and 450 °C as shown in Fig. 5 the films deposited at substrates temperatures below 350 °C shows an amorphous nature. The peak intensity, i.e. the degree of crystallinity, increases when temperature is rised.

368

(a

1100

1 -

'

0.6 -

0.4

0.2 "

w

fK0^My^ (b

',£' •li Ti JJ

-

4 mTorr 8 mTorr 16 mTorr 32 mTorr

jj 300

500

700 >.inm)

900

1100

Figure 4. UV-visible transmission spectra of Ti0 2 thin films prepared by D.C. reactive magnetron sputtering at different total pressures of an equimolar gas mixture Ar/0 2 . On different substrates (a) on glass and (b) on glass coated with FTO.

Fig. 6 shows the as measured transmission curves for sprayed Ti0 2 coatings at different substrates temperature. The transmission in the visible region increases with the substrate temperature. The increase in transmission value could be attributed to the well adherence and to the crystallized nature of the film, which is due to the evaporation of the undesired bi-products and improvement in the crystallinity [3]. In the visible range the transmission of the

369

films is around 70% and the spectra show waveforms that are characteristic of the interference light [4]. The transmission decreased as the wavelength decrease due to the fundamental absorption of the light [5]. According to the method described by Mardare et.al. [5], the optical band gap of Ti0 2 thin films was calculated. An optical band gap, Eg, of about 3.2 eV has been observed for the films prepared by spray pyrolysis method with different substrate temperatures. '

1

'

(a)

•"I

'

1

i

'

'

(101) T s =450 'C (200j

•:2lii

Ik\\MJ

WwWv /

4^4^^=350 "C

I

10

I

20

I

I

I

.

I

30

I

I

.

40

I

i

I

,

50

60

2H Figure 5. XRD patterns of Ti0 2 thin films on glass prepared by spray pyrolysis method at different substrate temperatures.

Table 2. The correlation of physical parameters of the Ti0 2 films deposited by spray pyrolysis for different substrate temperature is observed in this table.

TS°C

Phase

Thickness (nm)

350 400 450

Amorphous Anatase Anatase

260 270 285

Grain Size (nm) 32 30

E g (eV) 3.20 3.25 3.26

370 1

1

1

'

•

1

i

•

y

\

0.8

'

••-•-... /

v

/

\

- - ' ,

/"" 0.6

0.4

1

300°C •-—• 350 "C — 400 "C

yS

—

0.2

J

'300

, .

500

700 A(nm)

450"C

900

1100

Figure 6. Optical transmission spectra of Ti0 2 thin films prepared by spray pyrolysis method at different substrate temperatures.

3.3. Photocatalytic Activity Photocatalytic process is initiated by the absorption of a photon with energy equal to or greater than the band gap of Ti0 2 (-3.2 eV in anatase phase), producing an electron-hole pair. The resultant electron-hole pair has lifetime in the space charge region that enables its participation in chemical reactions. The postulated reactions are [8]: OH",., + h+

> *OH ad

(2)

0 2a d + e"

• 02"ad

(3)

Hydroxyl radicals ("OH) and super-oxide ions (02~) are highly reactive species that will oxide the organic compounds adsorbed on the semiconductor surface. Many kinds of organic pollutants can be oxide by Ti0 2 [1]. Fig. 7 shows the degradation of methylene blue film formed on Ti0 2 surface. The change of the absorbance (A ABS) characterizing the decomposition of methylene blue. In Fig. 8 we present the degradation of methylene blue when the film is dipped on a MBS and irradiated with UV light.

371

From Figures 7 and 8 we can see that the highest photocatalytic activity is achieved for the Ti0 2 thin films deposited by spray pyrolysis method at 400 °C and one deposited by D.C. reactive magnetron sputtering at 16 mTorr total pressure. By AFM both films present an open structure and surface porosity (see figure 9 (a) and (b)), anatase phase oriented along (101), (200) and (211) planes. The film prepared at 4 mTorr is characterized by very low photocatalytic activity due to the rutile structure with a high density and an absence of surface texture. When we used an amorphous film (prepared by spray pyrolysis method at 350 °C), we found a low photocatalytic activity. We found that the photocatalytic degradation of methylene blue was decreased in the same order found by Yumoto et al. [11] in the photocatalytic decomposition of N0 2 by Ti0 2 thin films prepared by arc ion plating technique. Then the photocatalytic efficiency was decreased in the order anatase, amorphous, rutile + anatase. Atomic force microscopy (AFM) was used to characterize the uniformity and particle size of sprayed and sputtered films. As shown in Fig. 9, the Ti0 2 crystals of the films has a particle size of 15-30 nm, only by spray pyrolysis, the Ti0 2 thin film present particle size of Ti0 2 crystals as large as about 30-50 nm. By AFM, we found that the films are suitable for photocatalytic applications since the porosity of the films resulted in improvement of the photocatalytic efficiency due to increase in their effective surface area.

Temperture fC) 350 400 450

«, d

50nm LSMO has major contribution to the negative MR and the MR ratio gets larger when the temperature is getting closer to its M-I transition temperature 360K

20

30

d

40

(nm)

LSMO

Figure 8. t dependence of MR ratio for YBCO(150nm)/LSMO(t) heterostructures at T= 50, 77 and 300K. As to the nature of low-temperature normal state for high-Tc materials, different observations of insulating and metallic behaviors for p(T) seem contradictory to each other. The observation that the electron state of La2. x Sr x Cu0 4 undergoes a gradual evolution from metallic to weakly localized and eventually variable-range hopping by increasing the magnetic field [34] was attributed either to a normal insulating ground state or to the magnetic-field-

399

induced localization of the metallic ground state. In the present YBCO/LSMO herterostructures, instead of the high magnetic field, the spin injection destroys the superconductivity of YBCO and the resulting insulating p(T) behavior could also be simply described by the equation of variable range hopping. Such localization behavior should be a bulk effect. It is unlikely generated by the disorder at the interface, because if the YBCO upper layer remains metallic, the transport properties of the present YBCO/LSMO structure should be dominated by YBCO such that neither the variable range hopping behavior of resistivity nor negative CMR effect can be observed. Therefore, the present result indicates that the ground state of the upper YBCO layer in the absence of superconductivity appears to be insulating. Our result is consistent with the data obtained from the previous pulsed high field experiment [33] on (Yo.6Pro.4)Ba2Cu30I and YBa2Cu3Ox, both exhibiting an insulating-like ground state at low temperature. 3.3. Critical Currents in YBa2Cu307/Lao.7Sro.3Mn03Heterostructures Thickness dependent V-I curves at 1.9 K for YBCO(150nm)/LSMO(t) herterostructures with t = 0, 10, 20, 30 and 40 nm are shown in Fig. 9. It displays a typical symmetrical V-I characteristic of YBCO with a plateau centered at V = 0, and the width of plateau decreases with increasing t. It is noted that the plateau width for t = 10 nm is smaller than those of t = 20 and 30 nm, which maybe due to an additional contribution of pairbreaking from the high interface strain. Figure 10 is the field dependent V-I curve for a single YBCO(150nm) layer, demonstrating a change in the width of plateau with increasing field. By comparing Fig. 9 and 10, we found an interesting phenomenon that the influence of 40nm LSMO on Ic is stronger than that of 5 Tesla field. The critical current (Ic) of each herterostructure sample is extracted as the current value corresponding V = luV. The field dependent Ic for YBCO (150nm)/LSMO(t) with t = 0, 10, 20 and 30 nm is plotted in Fig. 11. Based on Fig. 11, Ic is suppressed with increasing magnetic field but the suppression rate becomes smaller for higher t. For t = 30 nm, Ic is almost insensitive to magnetic field. Interestingly, the magnitude of Ic-suppressing by 30-nm LSMO is 20% higher than applying a 5 tesla magnetic filed. It is worthy to note that although the Icvalue was suppressed from 4 to 0.8 mA by inserting a 30-nm layer below YBCO film, the Tc of YBCO persisits at 74 K. The model of gapless superconductivity5 3 in which there is a finite superconducting order parameter but a vanishing energy gap can explain this phenomenon. In practice, this result suggests that proximity effect of YBCO/LSMO plays an important role in polarizing the spins and consequently serves as a spin injector.

400

0.9 40nm

0.6 lOhn,

03

20ml 30tn i

0.0 >

YBOO

-03 -0.6 -0.9 5 - 4 - 3 - 2 - 1 0

1

2

3

4

5

I(m\) Figure 9. Current-voltage characteristic for YBCO(150nm)/LSMO(t) with t = 0, 10, 20, 30 and 40 nm.

j,5T

0.04

A4*

Mi

0.02

w—

2 " 0.00 > -0.02 -0.04 .

-5

1

,

1

.

1

.

1

.

1

,

1

4 - 3 - 2 - 1 0

.

1

1

.

2

1

.

3

1

.

4

5

I(mA) Figure 10. Current-voltage characteristic for YBCO(150nm) under different field of 0, 1, 2, 3, 4 and 5 Tesla.

401

4.0 3.5 3.0 2.5 2.0

YBCO

1.5

t=20nm 1.0

't = 30nm

H(D Figure 11. Critical current vs. field for different YBCO/LSMO(t) bilayer with t = 0, 10, 20, and 30 nm.

4. Conclusions We have investigated simultaneously the physical properties of LSMO(t) films and YBCO(150nm)/LSMO(t) heterostructures with different t. It could be seen that a number of novel phenomenon were involved with a mixture/competition of localized/FM-metallic states. LSMO has an insulator to metallic transition while YBCO/LSMO has a superconductor to insulator transition when t = 50 nm. The field dependent and thickness dependent of V-I curves for YBCO(150nm)/LSMO(t) heterostructures further indicated an Ic-suppresion could be obtained by applying filed and by increasing the thickness of bottom LSMO layer. Since the phenomenon of Ic-suppresion is much abrupt by adding a 30 nm LSMO than by applying a 5 tesla field, the proximity effect at YBCO/LSMO may serve as an efficient spin-polarizer to break the superconducting pairs for the future device applications. Acknowledgements J.G. Lin wishes to thank her student Miss S.L. Cheng and her posdoc. Dr. A. Debnath for their contribution to this work. She also wishes to thank Profs. D.Y.

402

Xing and C.R. Chang for their theoretical inputs. The National Science Council and Ministration of Economics of R.O.C. support this project under the grant No. 91-EC-17-A-08-S1-0006 and NSC-91-2112-M-002-049. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25.

Y. Shimakawa, Y. Kubo, T. Manaka, Nature (London) 379, 53 (1996). P. Schiffer, A.P. Ramirez, W. Bao, S.W. Cheong, Phys. Rev. Lett. 75, 3336 (1995). C.W. Chang, A.K. Debnath , J.G. Lin, Phys. Rev. B 65, 024422 (2002). C.W. Chang, A.K. Debnath, J.G. Lin, J. Appl. Phys. 91, 2216 (2002). J.H. Park, E. Vescovo, H.J. Kin, C. Kwon, R. Ramesh, T. Venkatesan, Nature (London) 392, 794 (1998). H.L. Ju, C. Kwon, Q. Li, R.L. Greene, T. Venkatesan, Appl. Phys. Lett. 65, 2104(1994). W. Prellier, A. Biswas, M. Rajeswari, T. Venkatesan, R.L. Greene, Appl. Phys. Lett. 75,397 (1999). A. Asamitsu, Y. Tomioka, H. Kuwahara and Y. Tokura, Nature 388, 50 (1997). C.N.R. rao, A.R. Raju, V. Ponnambalam, Sachin Parashar, N. Kumar, Phys. Rev. B 61, 594 (2000). V. Markovich, E. Rozenberg, Y. Yuzhelevski, G. Jung, G. Gorodetsky, D.A. Shulyatev, Ta. M. mukovskii, Appl. Phys. Lett. 78, 3499 (2001). J.G. Lin, C.Y. Huang, S.Y. Lin, P.C. Kuo, Physica C 364, 668 (2001). X. Guo, S. Dai, Y. Zhou, G. Yang, Z. Chen, Appl. Phys. Lett. 75, 3378 (1999). P. Grunberg, R. Schreiber, Y. Pang, M.B. Brodsky, H. Sowers, Phys. Rev. Lett. 57,2442(1986). I.A. Garifullin, J. Mag. Mag. Mat. 240 571 (2002). A.M. Goldman, V. Vao'ko, P. Kraus, K. Nikolaev, A. Larkin, J. Mag. Mag. Mat. 200, 69 (1999). M.A. Clogston, Phys. Rev. Lett. 9, 266 (1962). B.S. Chandrasekhar, Appl. Phys. Lett. 1, 7 (1962). P. Fulde , A. Ferrel, Phys. Rev. 135, A550 (1964). A. Larkin, Y. Ovchinnikov, Sov. Phys. JETP 20, 762 (1965). T. Kontos, M. Aprili, J. Lesueur, X. Grison, Phys. Rev. Lett. 86, 304 (2001). M. Zareyan, W. Belzig, Yu. V. Nazarov, Phys. Rev. Lett. 86, 308 (2001). G.Y. Sun, D.Y. Xing, J.M. Dong, M. Liu, Phys. Rev. B 65, 174508 (2002). S. Jin, T.H. Tiefel, M. McCormack, R.A. Fastnacht, R. Ramesh, L.H. Chen, Science 264,413 (1994). M.K. Wu, J.R. Ashburn, C.J. Torng, P.H. Hor, R.L. Meng, L. Gao, Z.J. Huang, Y.Q. Wang, C.W. Chu, Phys. Rev. Lett. 58, 908 (1987). J.H. Park, E. Vescova, H.J. Kim, C. Kwon, R. Ramesh, T. Venkatesan, Phys. Rev. Lett. 81, 1953 (1998).

403

26. 27.

28. 29. 30. 31. 32.

33. 34.

35.

35. 36. 37. 38. 39. 40.

41. 42. 43. 44. 45.

V.A. Vas'ko, V.A. Larkin, P.A. Kraus, K.R. Nikolaev, D.E. Grupp, C.A. Norman, A.M. Goldman, Phy. Rev. lett. 78, 1134 (1997). Z.W. Dong, R. Ramesh, T. Venkatesan, M. Johnson, Z.Y. Chen, S.P. Pai, B. Talyansky, R.P. Sharma, R. Shreekla, C.J. Lobb, R.L. Greene, Appl. Phys. Lett. 71, 1718 (1997). Y. Gim, A.W. Kleinsasser, J.B. Barner, J. Appl. Phys. 90, 4063 (2001). A.M. Goldman, P.I. Kraus, K. Nikolaev, V. Vas'ko, A. Bhattacharya, W. Cooley, J. Supercond.: Incorporating Novel Magnetism 14, 283 (2001). P. Mikheenko, M.S. Colclough, C. Severac, R. Chakalov, F. Welhoffer, CM. Muirhead, Appl. Phys. Lett. 78, 356 (2002). Y. Ando, G.S. Boebinger, A. Passner, N.L. Wang, C. Geibel, F. Steglich, Phys. Rev. Lett. 11, 2065 (1996). Y. Ando, G. S. Boebinger, A. Passner, T. Kimura, K. Kishio, Phys. Rev. Lett. 75, 4662 (1995); G. S. Boebinger, Y. Ando, A. Passner, T. Kimura, M. Okuya, J. Shimoyama, K. Kishio, K. Tamasaku, N. Ichikawa, Uchida, Phys. Rev. Lett. 77, 5417 (1996). J. Vanacken, L. Trappeniers, P. Wagner, L. Weckhuysen, V.V. Moshchalkov, Y. Bruynseraede, Phys. Rev. B 64, 184425 (2001). K. Karpinska, A. Malinowski, M.Z. Cieplak, S. Gershman, G. Kotliar, T. Skoskiewicz, W. Plesiewicz, M. Berkowski, P. Lindenfeld, Phys. Rev. Lett. 11, 3033 (1996). A. Malinowski, M.Z. Cieplak, A.S. van Steenbergen, J.A.A.J. Perenboom. K. Karpinska, M. Berkowski, S. Guha, P. Lindenfeld, Phys. Rev. Lett. 79,495 (1997). M.N. Kunchur, B.I. Ivlev, D.K. Christen, J.M. Phillips, Phys. Rev. Lett. 84, 5204 (2000). M. N. Kunchur, Phys. Rev. Lett. 89, 137005 (2002). A. Debnath, J. G. Lin, Phys. Rev. B 67, 064412 (2003). LI. Balcells, J. Fontcuberta, B. Martinez, X. Obradors, Phys. Rev. B 58,R14697(1998). A. Biswas, M. Rajeswari, R.C. Srivastava, Y.H. Li, T. Venkatesan, R.L. Greene, Phys. Rev. B. 61, 9665 (2000). A.M. Haghiri-Gosnet, J. Wolfman, B. Mercey, Ch. Simin, P. Lecoeur, M. Korzenski, M. hervieu, R. desfeux, G. Baldinozzi, J. Appl. Phys. 88, 4257 (2000). R.A. Rao, D. Lavric, T.K. Nath, C.B. Eom, L. Wu, F. Tsui, Appl. Phys. Lett. 73, 3294(1998). J. Aarts, S. Freisem, R. Hendrikx, H.W. Zandbergen, Appl. Phys. Lett. 72,2975(1998). H.S. Wang, E. Wertz, Y.F. Hu, Qi Li, D.G. Schlom, J. Appl. Phys. 87, 7409 (2000). A.P. Ramirez, J. Phys.: Condens. Matter 9, 8171 (1997). S.Y. Yang, W.L. Kuang, CH. Ho, W.S. Tse, M.T. Lin, S.F. Lee, Y. Liou, Y.D.Yao, J. Magn. Magn. Mater, 226-230, 690 (2001).

404

46. N.A. Babushkina, L.M. Belova, D.I. Khomskii, K.I. Kugel, O. Yu, Gorbenko, A.R. Kaul, Phys. Rev. B. 59, 6994 (1999). 47. A. Avila, R. Asomoza, Solid State Electron 44, 17 (2000). 48. For review, see, e.g. J.M. Ziman, Models of Disorder (Cambridge University Press, Cambridge, U.K., 1971). 49. T.R. Kirkpatrick, D. Belitz, in Electron Correlation in the Solid State, edited by N.H. March (Imperial College Press, London, 297, 1999). 50. P.W. Anderson, Phys. Rev. 109, 1492 (1958). 51. For review, see Localization, Interaction and Transport Phenomena, (Solid-State Sciences 61, Springer Series, edited by B. Kramer, G. Bergmann, Y. Bruynseraede). 52. J.G. Lin, S.L. Cheng, P.C. Kuo, C.R. Chang, D.Y. Ding, submitted to Phys. Rev. B (2003) 53. Guoya Sun, D.Y. Xing, Jinming Dong, Phys. Rev. B 65, 174508 (2002).

DOMAIN STRUCTURE OF YBa2Cu3Ox FILMS ON NdGa0 3 SUBSTRATES I.K. BDIKIN Institute of Solid State Physics, Chernogolovka, Moscow distr., 142432, Russia E-mai: [email protected]

P.B. MOZHAEV, G.A. OVSYANNIKOV Institute of Radio Engineering and Electronics, Moscow, J 03907, Russia P.V. KOMISSINSKI Institute of Radio Engineering and Electronics, Moscow, 103907, Russia and Department of Microelectronics and Nanoscience, Chalmers University of Technology, Gothenburg, S-41296, Sweden I.M. KOTELYANSKII Institute of Radio Engineering and Electronics, Moscow, 103907, Russia Structure, orientational features and twinning of epitaxial superconductive YBa2Cu3Ox (YBCO) thin films and YBCO/Ce02 heterostructures on (110) NdGaC>3 (NGO) and tilted-axes NdGa0 3 substrates with an inclination of normal of the substrate from the [110] axis were investigated by X-ray diffraction methods. Orthorhombic structure of NdGa03 results in an increase of the angle between (110) and (110) twinning planes in YBCO films to 90.20° and in a difference in volume of two twin domain systems. The orientation of epitaxial YBCO thin films was shown to be influenced by the deposition rate and the presence of symmetrical-equivalent directions [110] and [HO] in the substrate and [100], [010], and [001] in the Ce0 2 layer. Domain structure of the YBCO thin film surface changes with an increase of the inclination angle. The YBCO thin films on the tilted-axes substrates are twinned in the same way as on (110) NGO substrates. However, formation of one or both twinning complexes is suppressed with an increase of the inclination angle.

1. Introduction Modern technology of deposition of the superconductive YBa2Cu307_x (YBCO) provides thin films of crystal quality close to that of single crystals. The major part of YBCO thin films structure investigations was performed using thefilmswith the c-axis normal to the surface of different substrates. The twin structure in the a-b plane is very important for interpretation of the YBCO thin film electrical transport properties. Similar to single crystals, twinning in c-oriented YBCOfilmsoccurs in accordance with the {110}/ 405

406

scheme with an angle of twinning about 1° [1,2]. The type of twinning structure of YBCO films correlates with the structure of the substrate. We present results of comparative studies of domain structure in YBCO films on standard and on tilted-axes substrates (TAS) of NdGa03 (NGO) both with and without a Ce0 2 seeding layer. 2. Experimental Ce0 2 films with typical thickness of 300-400 A, acting as seeding layer for YBCO thin films, were grown by RF reactive magnetron sputtering or by electron beam evaporation. Crystal structure of Ce0 2 seeding layer differs from that of NGO substrate, while keeping good lattice match between the film and the substrate (0.4% lattice mismatch is observed for YBCO/NGO heterostructures, increasing to 0.8% for YBCO/Ce02/NGO heterostructures). Such a combination provides a good possibility to check the effect of substrate lattice structure on the domain structure of the YBCO thin films. YBCO films with thickness about 1500 A were grown using DC sputtering at high oxygen pressure and laser ablation techniques. The typical YBCO deposition rates were 300A/min for laser ablation and lOA/min for DCsputtering at high oxygen pressure [3-5]. X-ray studies were performed using Siemens D500 and DRON-3M diffractometer systems. Both symmetric and asymmetric diffraction geometry were used. To study twinning in c-oriented YBCO films it is necessary to observe reflections from crystallographic planes, tilted to (001) planes. The chosen (103) and (113) reflections are of the most intensive reflections in the YBCO structure. Relative positions of diffraction peaks from corresponding planes provide the twinning angle value and the value of the angle between substrate and YBCO crystallographic planes. These angles were calculated using: 8 = acos(y) + p-sin(y), (1) where 8 is the measured misorientation angle between film grains (twin domains); a and p represent angles of misorientation between these grains in selected mutually perpendicular planes, y is angle between planes corresponding to 8 and a. a and p axes were chosen in substrate plane and normal to substrate surface. Grain misorientation spread is present in two perpendicular directions: normal to the substrate surface and in the substrate plane. These two values can be decomposed as follows: 2 2 2 2 2 A8 =Aa -cos (x) + AP -sin (x), (2)

407

where A5 is angle of grain misorientation spread from the 8 axis, Aa and Ap are grain misorientation spread angles from the mutually perpendicular axes; % is angle between axes 5 and a. A special technique of rocking curve measurement in a wide angular scanning range was implemented for precise measurement of the angles between the crystallographic planes of the layers in the heterostructures. This technique is illustrated on Fig.l with a sample rocking curve and the geometry of the measurement presented on the inset. No filtration of incident X-rays was implemented, resulting in a broadband X-rays beam. The 28-angle is set constant so that the characteristic line of the X-ray source provided a Bragg reflection peak from one of the film planes. Rotation of the sample around the [001] axis of NGO results in additional Bragg peaks on the rocking curve, corresponding to strong reflections from the substrate at different wavelengths of the X-ray irradiation, due to wide X-ray radiation spectra. 1600

X-RAY SOURCE

,(222)Ce02

010)NGO (22Z)Ce02

3. Low c parameter level supposes high oxygen contents in the films [6,7].

408

3.1. YBCO Films on (110) NGO Substrates The following epitaxial relations take place for all YBCO thin films grown on the (110)NdGaO3 substrates: (001)Y || (H0) N , and [100]Y || [001]N or [010]Y|| [001] N . Some films contain inclusions with epitaxial relations (100)Y || (110)N, Indexes "Y", "N", and "C" mark the crystallographic directions in the YBCO film, NGO substrate, and Ce0 2 interlayer, correspondingly. 0-scan diffraction pattern in vicinity of (113) YBCO film reflection is shown on Fig. 2 and presents a distinctive picture of YBCO twinning. The obtained pattern can be arithmetically decomposed into four diffraction curves. A and A' curves of this decomposition correspond to different twinning systems of (1J,0) plane, B curve results from other two twinning orientation on (IIP) plane, showing no splitting, and C curve corresponds to (020) NGO planes. Relative peak positions reflect misorientation of corresponding planes. Diffraction patterns in Fig.2 allow determination of mutual orientation of NGO and YBCO atomic planes as well as mutual orientations of twinning parts. Application of equation (1) in a similar way allows evaluation of the angle between twinning YBCO planes from mutual positions of A, A' and B peaks, being equal to 90.20°. Such discrepancy from 90° was also observed in YBCO film on NGO previously [2]. The increase of the angle between twinning planes can be explained by symmetry of NGO lattice. Lattice in (110) plane of NGO is close to tetragonal with angle between (111) and (001) planes of 44.94° (measured NGO lattice parameters were a = 5.428(2)A, b = 5.499(2)A, c = 7.711(3)A).

A 0 (°) Figure 2. X-ray diffraction 9-scans axis of (113) reflection of c-oriented YBCO films on (110) NdGaC>3. On the insert: scheme of the experiment.

409

3.2. YBCO Films on NGO TAS The YBCO films, deposited on the NGO TAS, follow the epitaxial relation (001) Y || N in the studied range of the inclination angle y = 5-26°. Thermodynamically the c-oriented growth of the YBCO thin films is preferential at chosen deposition conditions [2,8] and in further discussion we can neglect the growth of YBCO films with the a-axis normal to the surface. The presence of two symmetrical-equivalent planes in the lattice of the TAS NGO - (110)N and (1K))N - both satisfying conditions of c-oriented epitaxial growth, results in formation of two domain systems in the YBCO film. Following the traditional notation for the YBCO thin films on the (110) NGO substrates, we call "pseudo-c oriented" domains with the axis [001]Y close to the normal to the substrate plane; and "pseudo-a oriented" domains with the axis [001]Y close to substrate plane. The relative contents of pseudo-c oriented and pseudo-a oriented domains changes with y. The rocking curves in wide scanning range of the YBCO (3 0 10) planes for different y are shown in Fig. 3. Both for pseudo-a oriented and for pseudo-c oriented domains this peak is close to the (010)N reflection, allowing estimation of the volume ratio of these domains from the integral intensity of the peaks. At small y the formation of pseudo-a oriented domains is suppressed, but the contents of these domains increases rapidly with an increase of y. The experimentally determined parts of pseudo-a oriented domains at different y are given in the Table 1. 3 0 10,

(001)Y,=(110)N 26" &"

• • • • •

I 1

8 - "

11°

8

n = (110)N

(a)

AG(°) Figure 3. Rocking curves in a wide scanning range of the (3 0 10) peaks of the YBCO thin films on TAS NGO at different inclination angles. For small inclination angle (curves a, b) formation of pseudo-a oriented domains (Y2) is suppressed, while for large inclination angle the part of the pseudo-a oriented domains is almost equal to that of pseudo-c oriented domains (Yl).

410 Table 1. Contents of domains of different orientations in YBCO thin films on NGO TAS.

Substrate (110) NGO ~(571)NGO (120) NGO (130) NGO

Inclination angle (degree) 0 11 18 26

1(3 0 10)pseudo.a 1(3 0 10)Dseudo.c 0 0 0.8 1.6

Part of pseudo-a oriented domains (%) 0 0 45 60

Similar to YBCO films on the (110) NGO substrates, twinning in the pseudo-c oriented domains of the YBCO films on NGO TAS follows the scheme {110}/ and was observed in X-ray diffraction experiments as splitting of the corresponding reflections (Fig. 4a). Increase of inclination angle (y>15°), however, results in suppression of twinning (Fig. 4c). Rotation of the substrate surface around the [001]N axis results in the same inclination of both possible twin boundaries to the substrate surface, and both twin domains are suppressed equally. Rotation of the substrate surface around the [1U] N direction leaves one of the twin boundaries perpendicular to the substrate surface, inclining only the second one. On such a substrate (the studied substrate surface was close to the (571)N crystallographic plane with an inclination angle c about 10.6°) only the inclined twin complex is suppressed (Fig. 4b). Even in the YBCO thin films on TAS with large inclination angle, when the twinning is completely suppressed, the volume of pseudo-c oriented domains with different orientation of the [100]Y axis is almost equal. 3.3. YBCO Films on NGO TAS with a Ce02 Seeding Layer The orientation of YBCO films, deposited on a thin epitaxial Ce0 2 layer on the NGO TAS essentially depends on the deposition technique. The YBCO thin films grown by the pulsed laser deposition technique are c-oriented independently on the orientation of the substrate and Ce0 2 interlayer. The YBCO thin films, deposited by DC sputtering at high oxygen pressure, instead, are oriented along the axes of the Ce0 2 interlayer, independently on the substrate orientation: [001]Y || c. Cubic symmetry of Ce0 2 leads to growth of YBCO in three orientations; preferential orientations at chosen deposition conditions are those with the minimal angle between the [001]Y axis and the normal of the substrate. This effect is similar to the preferential c-oriented film growth compared with a-oriented on the (110) NGO substrates at high deposition temperatures [9, 10]. As a result of suppression of YBCO film orientation with an axis [001 ] Y close to the substrate surface, different number of

411 83.6

(a

138.0

(b

(c

Figure 4. Two-dimensional X-ray diffraction spectra of the (103) and (013) reflections of the YBCO thin film on NGO TAS in coordinates 0, 26. The 8-26 line is marked with the dotted line, (a) (110) NGO substrate: the twinning planes are (110) (A and A' spots) and (M0) (B and B' spots), (b) (571) NGO substrate: single twinning plane (110) (A and A' spots) can be seen, (c) (130) NGO substrate: no twinning observed.

412 YBCO domain systems (one, two or three) are formed, depending on the Ce0 2 orientation. The morphology of such films, investigated with the AFM, showed pyramidal, edge or needle structures, formation of which probably being caused by different number of domain systems in the film [11]. The strong dependence of YBCO thin film orientation on the deposition rate can be explained as insufficient oxygenation of the YBCO thin film on the seeding stage of growth. At low deposition rate (DC-sputtering) the deposited material is oxidized completely and the oxygen sublattice determines epitaxial relations of the growing film. Continuation of the oxygen sublattice from the Ce0 2 layer into the YBCO thin film results in alignment of the axes of both materials ([001]Y || c). At high deposition rates (pulsed laser deposition) the significant differences of crystal structure result in film growth with rotation of the planes of the minimal energy (the (001 )Y planes) parallel to the substrate plane. Independently on the mutual orientation of the YBCO film and Ce0 2 seeding layer, the YBCO thin film contains a significant (up to 60%) nontwinned part, probably due to small size of the seeding layer grains. Twinning complexes are following the same scheme, as on bare NGO substrates. The effect of twinning suppression with increase of inclination angle was not observed. 4. Conclusions Domain structure and twinning of the YBCO films on NGO substrates and on Ce0 2 /NGO heterostructures were investigated. The crystallographic parameters of the YBCO thin films on these different substrates were close, but the films domain orientation and twinning showed particularities resulting from the nature of substrates. Twinning orientation features of the YBCO film on (110) NGO correlated with the symmetry of the substrate. The orientation of the YBCO epitaxial films on the tilted-axes substrates is determined by an existence of the symmetrically-equivalent directions in the substrate and in the Ce0 2 seeding layer. The presence of a thin Ce0 2 epitaxial layer essentially changes orientation of the YBCO thin film growing on the NGO TAS. At high deposition rate the superconductor film grows in orientation (001). At small deposition rate the YBCO thin film grows with directions [001 ] Y along symmetrical - equivalent directions c of the Ce0 2 layer.

413 Acknowledgements Authors would like to thank I.V. Borisenko, O.G. Rybchenko for help in experiments as well as Yu.Boikov, T.Claeson, Z.Ivanov and E.Stepantsov for fruitful discussions. This study was supported by INTAS (grant no. 2001-0249). References 1. 2.

3.

4. 5. 6. 7. 8. 9. 10. 11.

T. Scheme, P. Marienhoff, R. Herwig, M. Neuhaus, W. Jutzi, Physica C 197,79(1992). T. Steinborn, G. Miehe, J. Wiesner, E. Brecht, H. Fuess, G. Wirth, B. Schulte, M. Speckmann, H. Adrian, M. Maul, K. Petersen, W. Blau, M. McConnel, Physica C 220, 219 (1994). P. B. Mozhaev, G. A. Ovsyannikov, S. N. Polyakov, E.K. Kov'ev, N.P. Kukhta, Superconductivity, Physics, Chemistry and Technology (in Russian) 9, 304 (1997). I. M. Kotelyanskii, V. A. Luzanov, Yu. M. Dikaev, Superconductivity, Physics, Chemistry and Technology (in Russian) 7, 1306 (1994) A. D. Mashtakov, K.Y. Constantinian, G. A. Ovsyannkov, E. A. Stepantsov, Techical Physics Letters 25, 249 (1999). M. S. Osofsky, J. L. Cohn, E. F. Skelton, M. M. Miller, R. J. Soulen, Jr., S. A. Wolf, T. O.Vanderah, Phys Rev. B 45, 4916 (1992). J. D. Jorgensen, B. W. Veal, A. P. Paulikas, L. J. Nowicki, G.W. Crabtree, H. Claus, W. K. Kwok, Phys. Rev. B 41, 1863 (1990). I. K. Bdikin, A. D. Mashtakov, P. B. Mozhaev, G. A. Ovsyannikov, Physica C 334, 168(2000). F. Vassenden, G. Linker, J. Geerk, Physica C 175, 566 (1992). F. Miletto Granozio, M. Saluzzo, U. Scotti di Uccio, I. Maggio-Aprile, Oe.Fischer, Phys. Rev. B 61, 756 (2000). I. K. Bdikin, P. B. Mozhaev, G. A. Ovsyannikov, P. V. Komissinskii, I. M. Kotelyanskii, Physica C, 377, 26 (2002).

RAMAN ACTIVE APICAL OXYGEN MODES IN Cui.xTlxBa2Ca3Cu4Oi2-8 SUPERCONDUCTOR THIN FILMS N.A. KHAN Materials Science Laboratory, Department of Physics, Quaid-i-Azam University Islamabad, Pakistan E-mail: [email protected] H. IHARA Electrotechnical Laboratory 1-1-4 Umezono, Tsukuba, Ibaraki 305-8568, Japan Raman spectroscopy of extremely pure Cui.sTUE^CajCiuOu-a superconductor thin films is studied in the Ein // to c-axis and Ei„ 1 to c-axis configurations. The samples were prepared by amorphous phase epitaxy method and were characterized by resistivity, susceptibility, electron microscopy and EDX measurements. The resistivity and susceptibility measurements have shown the lc of the material to be 113K and the X-ray diffraction confirmed the films to be c-axis oriented and predominantly single phase. In the Raman spectroscopy in Ein // to c-axis configuration, we have observed the Raman active modes at 598, 527, 304, 232 & 151 cm"1 while in the Ein ± to c-axis configuration at 520, 232 & 150 cm"1. We have assigned 598 and 527 cm"1 modes to the apical oxygen of types CU(1)-OA-CU(2) and Tl -0A-Cu(2). The 304 cm"1 mode is Ag type and is assigned to the motion along the c-axis of Ca atom. The 232 cm"1 mode is Eg type and is due to motion in the ab-plane of Ca atoms. The 151 cm"1 mode is due to planer motion of the Cu(2) atoms.

1. Introduction The superconductivity in CuBa2Ca3Cu40i2.6 (Cu-1234) [1] system is a best choice in the cuprates family due to its low superconductor anisotropy (y~1.6) and long coherence length along the c-axis. The low superconductor anisotropy and long coherence length give this compound ability to carry very high current density (5xl07 A/cm2). The normal pressure synthesis of this compound, have not yet become possible, which makes it unsuitable material for commercial device fabrication. However, very close derivatives of this compound in the form of thin films of Cui.xTlxBa2Ca3Cu40i2.6 (CU|.XT1X-1234) have been prepared at normal pressure [2]. This is achieved by amorphous phase epitaxy method (APE-method), which is thallium treatment of the amorphous phase at the thermal stability temperature of Cui.xTlx-1234. The presence of thallium in the charge reservoir layer of Cu,.xTlx-1234 gives this material a relatively higher superconductor anisotropy y~4 [3] than Cu-1234 material which has y=1.6. However, the superconductor anisotropy of the former compound can be

414

415

decreased by removing thallium from the charge reservoir layer. The detailed preparation and growth kinetics of this compound have been studied and reported [2,4], The Raman active phonon modes are reported in this paper. The phonon modes assignment is done by comparing our results with the results based on lattice dynamic calculations [5] and observed data on T1-1223& Bg«° 1223 superconductors [6,14]. Our results have shown that this compound has two apical oxygen modes, at 530 and 598 cm 4 , due to the coupling of apical oxygen (0 A ) with the Cu(l) and Tl atoms of the Cui.xTlxBa204„$ charge reservoir layer. 2* Experimental The samples of Cui„xTlx«1234 superconductor thin films were prepared by amorphous phase epitaxy method (APE). In this method, the amorphous phase was deposited on SrTi0 3 substrate and treated with precursor thallium pellet of composition Cuo.sTlojBaaCaaCu^y, as reported elsewhere [2]. The XRD spectra of Gii„xTlx-1234 films showed a predominantly single phase c-axis oriented material The pole figure measurements of (103) reflections showed the crystals to be oriented along the a-axis, the films are bi-axially oriented. The surface of the films was analyzed by scanning electron microscopy, which showed the surface roughness to be less than 0.2|im. The composition of the films was measured by energy dispersive x-ray spectroscopy (EDX). For the Raman spectroscopy, we have used 2x2 mm wide and 1.2|un thick sample. The spectrum was taken choosing a Ifim spot with the incident laser power of 2 mW. A home made sample holder was used to hold the sample in the Ef„ // to c»axis and E»„ 1 to c-axis configurations, the geometry of measurements is shown in Fig. 1. SrTiQ 3

Thin Film Figure 1. The geometry of the sample in different incident electric field of the laser.

416

3. Results and Discussions The XRD of the Cui.xTIx-1234 films is shown in Fig. 2, the material is predominantly single phase and the films are oriented along the c-axis. The resistivity measurements showed the Tc to be 113K, as shown in the inset of Fig.2. The Raman spectra of the films are shown in Fig. 3 for both Ein // to caxis and Ejn 1 to c-axis configurations. In the Ein // to c-axis configuration, the phonon modes are observed at 598, 527, 304, 232 & 151 cm"1 while in the Ein 1 to c-axis configuration at 520, 232 & 150 cm'1. In the previous studies performed on Tl&Hg-based superconductors [6-10, 14], the phonon modes at the frequencies above (0>300cm"' are suggested to be due to the vibrations of lighter atoms (such as oxygen) and the phonon modes at the frequencies below Kx300cm"' to be due to the vibrations of heavier atoms (such as calcium, barium and copper). We have assigned the Raman active modes at 598 and 527 cm"1 to the apical oxygen (Ag type). In the previous studies on Cu-1234, a single apical oxygen mode (Ag mode) is observed at 490 cm"'[ll]. The splitting of this Raman mode into two apical oxygen modes appearing at 598 and 527 cm"1 in our Cui.xTlx-1234 is possibly due to changed charge reservoir layer. The Cu1234 has CuBa203.8 while Cu,.xTlx-1234 has Cu,.xTlxBa203.6 type of charge reservoir layer. There are possibly two types of apical oxygen atoms expected in our Cu|.xTlx-1234 material, the one connected with the chained thallium atom, such as Tl-0A-Cu(2) and the other with chained copper atom, such as Cu(l)-0 A Cu(2). Such types of apical oxygen modes have also been observed in Hg|.xTlx1223 superconductors [14]. The Raman mode at 520-527 cm"1 is assigned to the Tl-0A-Cu(2) apical oxygen. This mode is present in Ein //c-axis and Ein J. c-axis configurations, however, the intensity of this mode in the Ejn ± c-axis configuration is 50% reduced than in the E;n //c-axis configuration. The other apical oxygen mode observed at 595 cm"1, appearing only in the Ein //c-axis configuration, is possibly due to Cu (1)-0A-Cu (2) type of apical oxygen. To justify the appearance of this mode at 595cm"1 we have compared the bond lengths of apical oxygen bound with Tl and Cu atoms in Tl-1234 [12] and Cu1234 [1] superconductors. The bond length of Tl-0A-Cu (2) in Tl-1234 superconductor is 4.75A corresponding to the c-axis length 19.1 A, while Cu (1)0A-Cu(2) has bond length 4.20 A in Cu-1234 corresponding to its c-axis length 17.9A. The apical oxygen mode of type Cu (1)-0A-Cu (2), therefore, will oscillate at relatively higher frequency (co~595 cm"1) and Tl-0A-Cu (2) will oscillate at lower frequency ((0-520 cm"1). In single Thallium layer superconductors (such as Tl-1212, Tl-1223, Tl-1234 etc) a single apical oxygen mode had been observed previously. The reason for this discrepancy in our

417

deposited Cui.xTlx-1234 is possibly due to many different local environments of apical oxygen (0 A ) generated by the presence of the interstitial oxygen Og [14]. 10000

eooo

6000

4000

2000

_L-^_W 20

^SA-LJ 40

28 (degree)

Figure 2. Typical XRD of Cui.sTlxBa2Ca3Cu4Oi2^ superconductor thin films prepared at 900°C. The inset shows the resistivity measurement as a function of temperature for these samples.

45x10 527cm-1 598cm-1 (a)-EinWc-axis

(b)-Eiruc-axis 15

200

_L 300

_L 400

500

_L 600

700

Wave Number (cm-1) Figure 3. The Raman spectra of the Cu,.xTlxBa2Ca3Cu.,C>i2^ superconductor thin films taken under the laser power of 2 mW for different incident electric fields.

418

The other two Raman modes appearing at 304 and 232 cm"1 are due to the vibrations of Ca atoms. The former mode appears only in the Ein //c-axis configuration, while the later in the both Ein //c-axis and Ein 1 c-axis configurations. The mode at 304 cm"1 is Ag-type and is due to vibration along the c-axis of Ca atom. This mode is theoretically predicted in Tl-1223 at 293 cm"'and observed experimentally at 260 cm'1 in Tl-1223 [12] and at 300 cm'1 in Hg-1223 system 285cm"'[6]. The mode at 232cm"1 is Eg-type and is assigned to the planner motion of the Ca atoms. This mode is predicted theoretically at 251 cm"1 [13] and observed in Hg-1223 superconductor at 285 cm"1 [6]. The Raman mode at 151cm"1 is assigned to the planner motion of the Cu(2) atoms. This mode is theoretically predicted at 131cm'1 [13] and is observed in Tl-1223 at 150 cm'1 [13] and in Hg-1223 at 151 cm"1 [6]. 4. Conclusions In conclusion, we have successfully assigned the possible Raman active modes in pure Cui.xTlx-1234 superconductor thin films. The apical oxygen mode observed at 520 cm"1 is due to Tl- 0A-Cu (2) and 595 cm"' due to Cu(l)-0 A Cu(2) apical oxygen atoms. Of these two former is active in both, the Ein//c-axis and Ein 1 c-axis configurations, while the later is active only in the Ein //c-axis configuration. The two modes due to the out of and in plane vibrations of the Ca atoms are at 304 cm"1 and 232 cm"1, respectively. Of these two the former is of Ag-type while the later has Eg symmetry. The vibrations of the planner Cu(2) atoms are observed at 151 cm"1. References 1. 2. 3.

4. 5. 6.

H. Ihara, K. Tokiwa, H. Ozawa, M. Hirbayashi, A. Negishi, H. Matuhata, Y.S. Song, Jpn. J. Appl. Phys. 33, L503, (1994). N.A. Khan, Y. Sekita, N. Terada, H. Ihara, Supercond. Sci. Technol. 14, 603 (2001). H. Ihara, K. Tokiwa, K. Tanaka, T. Tsukamoto, T. Watanabe, H. Yamamoto, A. Iyo, M. Tokumoto, M. Umeda, Physica C 282-287, 957 (1997). N.A. Khan, Y. Sekita, H. Ihara, A. Maqsood, Physica C 377, 43 (2002). A.D. Kulkarni, F. W. deWette, J. Prade, U. Schroder, W. Kress, Phys. Rev. B 41, 6049 (1990). A. Sacuto, A. Lobon, D. Calsan, A. Bertinohi, J.F. Marucco, V. Iallet, Physica C 259, 209(1996).

419 7.

C. Thamson, M. Cardona, Physical properties of high temperature superconductors, ed. D.M. Ginsberg (World Scientific, Singapore, 1989), p. 409. 8. K.F. McCarty, J.Z. Liv, R.N. Shelotn, H.B. Radousky, Phys. Rev. B 41, 8792(1990). 9. M.C. Krontz, C. Thomsen, Hj. Mathauch, M. Cardona, Phys. Rev. B 50, 1165(1994). 10. A. Sacuto, C. Julien, V.A. Shehukin, M. Makhtari, C. Perrin, Phys. Rev. B 52,7619(1995). 11. K. Tokiwa, N. Terada, A. Iyo, Y. Tsubaki, K. Tanaka, J. Akimoto, Y. Oosawa, M. Hirabayashi, M. Tokumoto, S.K. Agarwal, T. Tsukumoto, H. Ihara, Physica C 298, 209 (1998). 12. H. Ihara, R. Sugise, K. Hayashi, N. Terada, M. Jo, M. Hirabayashi, A. Negishi, N. Atoda, H. Oyanagi, T. Shimomura, S. Ohashi, Phys. Rev. B 38, 11952(1988). 13. K.F. McCarty, B. Morosin, D.S. Ginley, D.R. Boehme, Physica C 157, 135(1989). 14. I.S. Yang, H.G. Lee, N.H. Hur, J. Yu, Phys. Rev. B 52, 15078 (1995).

GENERATION AND AMPLIFICATION OF ELECTROMAGNETIC RADIATION BY SUPERCONDUCTING FILMS - A SUPERCONDUCTOR MASER A.N. LYKOV P.N. Lebedev Physical Institute, Leninsky pr., 53, 119991 Moscow, Russia E-mail: [email protected]

A new method for the designing of active superconducting elements is developed. The mixed state in the superconducting films is influenced by an alternative magnetic field directed perpendicular to the film surface. The transition of the vortex system into ground state synchronized via electromagnetic interaction with external resonant circuit causes the generation of the electromagnetic radiation. Coherent microwave radiation has been directly detected from superconducting thin films excited by the method described in literature [1-3]. Harmonic mixing and rf amplification are also detected using this approach.

1. Introduction An interesting problem in modern superconductivity is the creation of microwave oscillator. This was predicted by Josephson [1]. But the short current and voltage range of the observed Josephson ac effect precludes the practical application of the junctions as microwave oscillators. The power emitted by a single junction or coherent arrays of the junctions into a broadband system is very small. Moreover, phase-lock of a large number of Josephson junctions appears to be very serious problem. A novel approach was proposed [2] for designing a superconducting microwave oscillator. A superconducting film was placed in a low-frequency oscillating magnetic field directed perpendicular to the film surface. The field sets up a vortex structure in the film. The interaction of the vortices with planar pinning centres can lead to a metastable mixed state. The probability of the vortices or its bunches to jump from one pinning centre to another adjacent center is proportional to the relation: P ~ Clexp(-U/kBT), where T is the temperature; kB is the Boltzmann constant; Q is a depinning attempt frequency with which vortices try to escape from the pinning well, and U is the activation energy for flux jumps. Such an incoherent vortex motion gives rise to an electromagnetic noise generation by the superconducting film. The key feature of the approach is magnetic coupling between the film and a resonant circuit, mounted in such a way that the inductive coil is located in the vicinity of the film and creates an additional high-frequency oscillating magnetic field directed transverse to the film surface. The transition of a metastable vortex 420

421

lattice into its ground state causes vortex jumps. Under a change in the applied magnetic field, the high-frequency field periodically helps the vortices to overcome energy barriers and to escape from the pinning centre. Thus this field can synchronize the jumps of the vortices in the superconductor. In turn, every vortex jump induces an electromagnetic pulse in the coil of the circuit, and can increase the energy of the electromagnetic oscillation in the circuit. Thus a positive feed back coupling arises. A change in the magnetic field during each period of high-frequency oscillation gives a new and critical condition for jumping of the pinned vortices, and, thus, a new fraction of the trapped vortices are included in the process. Coherent microwave radiation has been directly detected from Nb [2] and GdBa2Cu307.x films [3] in the frequency range up to 600 MHz at 4.2 K and 10 MHz at 77.4 K, respectively. It is evident that the vortex jumps in superconducting films can be stimulated by external high-frequency electromagnetic field because the probability of the vortices or its bunches to jump from a pinning centres to an adjacent centre depends on the applied electromagnetic field. In this case the hysteresis formed in the magnetisation curve of the superconducting films leads to an increase in the electromagnetic field energy and to the amplification of the signal [4]. There is a close similarity between the amplification and the generation of the electromagnetic field. High-frequency operation of the amplifiers experimentally verified up to MHz range. For amplification of the electromagnetic radiation, the superconducting film is magnetically coupled to the two coils. The first drive-current coil connected to an external oscillator creates high-frequency magnetic field directed transverse to the film surface. The second pickup voltage coil is used for detection of the field in the range of 0.210 MHz. The amplification proves the realization of the positive feedback coupling when the superconducting films are used to generate coherent electromagnetic radiation. The purpose of the present work is to investigate the influence of the feedback coupling on the properties of the oscillator. 2. Method The experimental technique used for investigating the electromagnetic radiation emitted by the superconducting films is described in [2]. The sketch of the experimental set-up is shown in Fig. 1. The measurements were carried out using Nb films at 4.2 K. The films were prepared by electron-beam evaporation in a high-vacuum system. The low-frequency oscillating magnetic field, directed perpendicular to the film surface, varies in time (0 according to the law

HL=H10sm(2nfej)

(l)

422

where the amplitude of the oscillation H±0 reaches 1 kOe, and the frequency/^ is within 17 Hz to 1 kHz. The field sets up a vortex structure in the film. The interaction of the vortices with planar pinning centers, such as grain boundaries in the film, can lead to the metastable mixed state, for example, to the existence of the vortices even in zero magnetic field H±=0. The film is magnetically coupled to a resonant circuit, so that the inductive coil is located in the vicinity of the film and creates additional high-frequency oscillating magnetic field directed transverse to the film surface. The radiation power spectra were measured by selective microvoltmeters (SMV) in the range of 0.15 MHz - 0.6 GHz.

0

SMV

0

superconducting film

r L

' A

Figure 1. Equivalent circuit diagram (SMV is a selective microvoltmeter).

3. Results An example of the coherent radiation spectra emitted by superconducting films is shown in Fig 2. Coherent microwave radiation was detected in the frequency range of 0.15 MHz - 0.6 GHz from superconducting thin films. The radiation was excited by an oscillating magnetic field directed perpendicular to the film plane. In [2], the radiation frequency was found to be equal to the resonant frequency of the LC-circuit:y,= l/27i(ZC)05, where L and C is the inductance and capacity of the circuit, consequently. It will be noted that the radiation frequency is a half million times more the frequency of the excitation. By varying the parameters of the circuit it is possible to vary the frequency of the generated electromagnetic radiation. Thus our results show that the superconducting films are highly nonlinear elements, which make it possible to efficiently multiply the frequency of the external radiation. The radiation emitted by the films increases

423

with the increase of the film surface, and decreases with the increase of the radiation frequency. The energy source of the field is irreversibility of the magnetization curves of the superconducting films. It is observed that the power of emitted radiation increases proportionally with the increase in the amplitude and frequency of the exciting magnetic field. In order to explain the features, we should take into account that the energy output of the radiation is proportional to the total quantity of the vortex jumps in a unit time. Evidently, this quantity is proportional to both the H10 and fexc in agreement with the experiment. So the emitted power can reach a large value. The radiation disappears in the vicinity of the critical temperature of the superconducting film. It results from decreasing pinning forces, which set up the metastable vortex state in the film.

440

442

444

446

448

450

f, M H z Figure 2. An example of the frequency dependence of the amplitude of the A/voltage at high/, for //±o=35.3 Oe, and/„c=910 Hz.

In this study, the results of the detailed analysis of the radiation spectra are presented in the case of their high power. In a broader frequency range not one resonant peak is observed but also higher harmonics are presented. A spectrum of the coherent radiation emitted by superconducting films is shown in Fig.3. As it is seen, there appear a number of new maxima in addition to the main peak at the f=fr, in the figure. The positions of the maxima are f=nFr, where n is an integer. It was found that the amplitudes of the odd harmonics are higher than those of the even ones. The new maxima appear only if the inductance coil of the

424

circuit is located near the superconducting film. Thus their appearance results from back effect of the electromagnetic radiation on the superconducting film. These hysteresis loops are symmetrical with respect to the origin, thus generating only odd harmonics of the applied frequency. Moreover, it is found that the radiation depresses the noise radiation, so that frequency noise transforms into a generated signal at f=fr. In this case we measured the signal-noise ratio as a function of the frequency for two position of the inductance coil: near and far away the superconducting film. In the first case the back effect of the electromagnetic field generated in the resonance coil is more prominent than in the last case.

v.. nv

Figure 3. An example of the frequency dependence of the rf signal at low/, for //io=l 11.6 Oe, and /,«=183Hz.

To demonstrate the role of the back effect we measured the ratio: (2)

V/(f)/V/(fr) where Vf(f) and V((j) are the frequency dependences of the signal when the resonance coil is located near and far away the superconducting film, respectively. Fig.4 shows an example of the r(f) dependence. The figure shows that the r(f)2a3 cos 3cot, (5) 32 32 where a and § 6000 H o 3000 -

(003) (007)

L

(002)

(001)

JL i

10

T

20

JL.

(004) IN

I 30

Thin4

J

1L

I

40

50

29 Figure 9. The XRD patterns for Thin3 and Thin4 samples.

Figure 10. Temperature dependence of resistivity for Thin3 and Thin4 samples.

.AThin3 60

437

4. Conclusions The structural and physical properties of YBCO and BSCCO films made by PLD and sputtering techniques were investigated. It was found that a lower post annealing temperature (800 °C, 3h) improves Tc (zero) and the metallic behavior of normal state resistivity of YBCO films deposited by sputtering. For the BSCCO films, the proportions of 2223 phase will increase as the substrate temperature increases. The results show that for the fabrication of high quality YBCO thin films, a post-deposition heat treatment in oxygen atmosphere is necessary. Acknowledgment The authors would like to thank Isfahan University of Technology for supporting this project. We also would like to thank Dr. Soleimani, Dept. of Electrical Engineering, Tehran University, and Dr. Doroudian, Materials and Energy Research Lab., Karaj, for their help and assistance in the experiments. Part of this project was financially supported by the Ministry of Science, Research and Technology under the grant project 503495. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13.

14.

J. Bednors, K. Muller, Z Phys. B 64, 189 (1986). R. Wordenweber, Suprcond. Sci. Technol. 12, R86 (1999). D.J. Rogers, et al., Supercond. Sci. Technol. 12, R75 (1999). H. Ota, et al., Physica C 311, 42 (1999). S.J. Golden, F.F. Lange, D.R. Clarke, L.D. Chang, C.T. Necker, Appl. Phys. Lett., 61,351(1992). T. Sugimoto, et al., Appl. Phys. Lett. 63, 2697 (1993). Z. Mori, E. Minamizono, S. Koba, T. Doi, S. Higo, Y. Hakuraku, Physica C339, 161(2000). S. I. Karimoto, S. Kubo, K. Tsuru, M. Suzuki, Jpn. J. Appl. Phys. 36, 84 (1997). G. Balestrino, et al., J. Appl. Phys. 70, 6939 (1991). L. Ranno, et al., Phys. Rev. B 48, 13945 (1993). H. Salamati, P. Kameli, Physica B 321337 (2002). R. Gilbert, et al., J. Vac. Sci. Technol. 17, 389 (1980). S. Rossnagel, J. Cuomo, Thin Film Processing and Characterization of High-Temperature Superconductors, (Am. Inst. Phys. Conf. Proc. 165 New York, 106, 1988). A. Z. Moshfegh, O. Akhavan, H. Salamati, P. Kameli, M. Akhavan, in: Magnetic and Superconducting Materials (MSM-99), Eds. M. Akavan, J. Jensen, K. Kitazawa, (World Scientific, Singapore, vol. A, 585,2000).


VIII. MAGNETIC THIN FILMS


SOME ASPECTS IN THIN FILM MAGNETISM M. FARLE Universitat Duisburg-Essen, Experimentalphysik-AG Farle, Lotharstr. I/ME, 47048 Duisburg, Germany

The study of magnetic anisotropy energy in magnetic monolayers has provided an understanding of its microscopic origin. Some examples of the thickness and temperature dependence of the magnetic anisotropy energy are discussed in this contribution, mainly pointing the reader to more extensive literature. It is emphasized that while the microscopic anisotropy of the orbital moment can persist even above the Curie temperature the macroscopically measured anisotropy vanishes. Another consequence of magnetic anisotropy, i.e. the shape and size of magnetic domains, is discussed using the example of Ni monolayers on Cu(OOl).

1. Introduction Research on magnetic mono- und multilayers has provided a better understanding of the microscopic mechanisms which determine macroscopically observable quantitities like the magnetization, different types of magnetic order (ferro-, ferri- and antiferromagnetism), magnetic anisotropy, ordering temperatures (Curie, N6el temperature) and exchange coupling (see for example [1-9]). Aside from these basic research orientated investigations the technological exploitation of thin film magnetism has lead to huge increases in the hard disk's magnetic data storage capacities and new types of magnetoresistive angle and position sensitive sensors in the automotive industry , for example. The technological important aspects of exchange biased [10], spinvalve [11], or exchange-coupled multiplayer [12] structures will not be discussed here, however. In this contribution I will discuss some aspects of thin film magnetism dealing mainly with the variation of magnetic anisotropy energy (MAE) as a function of temperature and film thickness and its importance for the understanding of so-called spin reorientation phase transitions, in which the magnetization changes its easy axis either within the film plane or from in- to out-of-plane. 2. Magnetic Anisotropy Energy It is experimentally observed that a ferromagnet can be magnetized more easily along certain crystallographic directions than in others. One finds so called easy, intermediate and hard axes of magnetization, e.g. for bcc Fe these are , 441

442

, and for fee Ni , , at room temperature. The energy difference associated with different directions of M, that is the magnetic anisotropy energy MAE, is a small contributionon the order of a few ueV/atom to the total energy (several eV/atom) of a bulk crystal. To estimate the magnitude of the MAE one can use as a rule of thumb, that the lower the symmetry of the crystal or of the local electrostatic potential (crystal or ligand field) around a magnetic moment, the larger the MAE is. This becomes evident, if one remembers that in a crystal field of cubic symmetry the orbital magnetic moment is completely quenched in first approximation. Only by calculating in higher order (2nd) or by allowing a slight distortion of the cubic crystal a small orbital magnetic moment, i.e. a non-vanishing expectation value of the orbital momentum's z component is recovered. Without the presence of the orbital momentum which couples the spin degrees of freedom to the spatial degrees of freedom the MAE would be zero, since the exchange interaction is isotropic. One should also note, that the easy axis can deviate from crystallographic directions as for example in the case of Gd whose easy axis is temperature dependent and lies between the c-axis and the basal plane at T = 0 K. 2.1. Microscopic Origin Experimentally, it is often overlooked that information on the intrinsic origin of the macroscopically measured MAE can be detected by straightforward SQUID magnetometry measurements along different crystallographic axes. The saturation magnetizations along the easy and hard axes are different! The effect is very small - on the order of 10"4, but measurable, and well documented (Table 1). This means in other words that the magnetization vector changes its length, it is "longer" in the easy direction. Since the spin moment is usually assumed isotropic, the MAE arises from the anisotropy of the orbital moment. The anisotropy Au/u,ot (with An = ueasy-uhard) of the magnetic moment (it0, is related to the magnitude of MAE. Both are larger for the lower symmetry of bulk hep Co and smaller for cubic fee Co, fee Ni and bee Fe. The same observation holds for Rare Earth elements. Even for the S-state ion Gd which crystallizes in the hep structure an anisotropic moment (uc-ua)/uc = 10"3 (u ca :moment parallel to c, a-axis) has been measured (Table 1). For Tb with its large orbital moment (L=3) a much larger difference is found. There are only interactions which cause an magnetic anisotropy energy: a) the dipole - dipole interaction: H = D, and

withD,, = *

1

1

4 * ^

443

b) the spin - orbit coupling: HLS = —XLj • St. Both interactions couple the spin vector S to the lattice vector R. Different orientations of the spin with respect to the lattice vector yield different energies. The exchange interaction does not contribute to MAE, since the scalar product of the spin vectors is independent of the angles with respect to the crystal axes. The long range dipolar interaction is the source of the so-called shape anisotropy, which senses the outer shape of the sample. For homogeneously magnetized samples the dipolar anisotropy is given by Fd = 1/2 u0 (Nx Mx2 + Ny My2 + Nz Mz2) with the components of the demagnetization tensor Nx + Ny + Nz = 1. One should note that the dipolar interaction in the near field (only nearest neighbors) is sometimes called the pseudo-dipolar or anisotropic exchange energy. This name is misleading, since exchange interaction is isotropic. This contribution vanishes for a cubic crystal. Allowing a spontaneous magnetostrictive deformation of the lattice yields only a slight dipolar anisotropy which is 1/1000 of the MAE measured in Ni or Fe. Also, in bulk hep Co this contribution is negligible, since the c/a ratio deviates only by 0.67 % from the ideal ratio. Table 1. Anisotropic orbital moments, direction of easy axis of the magnetization, and magnetic anisotropy energy at T= 0 K for the four elemental ferromagnets as taken from standard references like Landolt-Bornstein.

Afi/ntot

Easy axis

MAE (0 K) (jieV/atom)

Fe

1.7xl0'4

[100] bec

+1.4

Co

4.5X10"4

[0001] hep

+65

[111] fee

1.8

Co Ni

1.8xl0'4

[111] fee

-2.7

Gd

~io-3

[tilted] hep

+50

The more important interaction is the spin-orbit coupling, which couples the spin to the charge (orbital) density distribution in the crystal. Thus, the spontaneous magnetization "gets the feel" of the crystal via the orbital motion of the magnetic electrons. Two kinds of microscopic energies may be produced as a result of this mechanism: a) spin-orbit coupling which depends on the spin

444

states of two or more ionic carriers of magnetic moment (pair model of magnetic anisotropy). b) coupling which depends on the effective spin state of individual ions {single-ion model of magnetic anisotropy). The magnitude of MAE is related to an increase of the difference in the orbital moment between the easy direction [001] and the hard direction. From this the following conclusions on the correlation of MAE and orbital moment are manifested: a) The orbital magnetic moment is anisotropic. b) The larger orbital moment is parallel to the easy axis of magnetization. c) The magnitude of MAE is related to the difference of the orbital moments parallel and perpendicular to the easy axis. d) The anisotropy of the orbital moment increases for lower symmetries, e.g. fee ->fct. A more quantitative discussion of the intrinsic origin of magnetic anisotropy can be found for example in [6,13]. 2.2. Phenomenology of Magnetic Anisotropy Energy In the phenomenological the crystallographic easy axis of the magnetization is determined by the minimum of the free energy density F, which in an external magnetic field can be expressed as the sum: F = Fex + F^ + Fei + Fmag.ei. + F„ + Fa + Fzee with Fex, the energy of exchange interaction; Fan, the energy of crystallographic magnetic anisotropy; Fei, the internal elastic energy of the crystal; Fmage|., the energy of magnetoelastic interaction; F„, the energy of external stresses associated with magneto-striction; Fd the energy of the demagnetizing field of the sample; and Fzee the energy of the magnetized sample in the external magnetic field. For the moment we will regard the free energy density for the case of zero applied magetic field, zero externally applied stress and for a spherical sample shape. Furtherore, we will consider only the single domain state. The exchange energy and the demagetizing energy for a sphere are isotropic. The easy axis of the spontaneous magetization is deterined by the minimum of the sum F^ + Fei + Fmagei.- One should note that a gain in magneoelastic energy depends linearly on the elastic deformations while Fei is a quadratic function of the deformations. Hence in general, the crystal will spontaneously deform, if a magnetization develops at the para- to ferromagnetic phase transition. The equilibrium lattice constant is determined by the minimum of Fan + Fe| + F mage i. For a cubic crystal Fan is written in the form:Ean = K0 + K4' (ax20y2 + Oy2az2 + az2ax2) + K6' ax2ay2az2+ where U\ are the direction cosine with respect to the crystallographic axes. This represents the anisotropic energy at constant volume. If one allows the crystal to deform, the anisotropy parameters Kj' are replaced

445

by coefficients K; given by Kj = K;' + /i(Cy, ay) where cy are the elastic moduli, and ay the magnetoelastic coupling constants. One has to distinguish between anisotropy parameters at constant volume (Kj') and the ones at constant stresses (Kj). The latter case always occurs in practice in bulk crystals. In bulk systems the differences between Kj and Kj' are very small and can be neglected in most cases. 2.3. Surface and Interface Anisotropy From the discussion above it has become clear that any change in the local symmetry of a magnetic moment causes a change of the magnetic anisotropy energy. At surfaces the change of symmetry is especially spectacular and causes much larger magnetic anisotropy energies per atom then in the respective bulk environment. This aspect has been demonstrated in studies of magnetic monolayers for which the surface contributions become dominant. In such a case the anisotropic part of the free energy is usually written as Kj = KjV + (Kjsurf+ Kjint) / d where the index i refers to the second and fourth order anisotropy coefficients and d is the film thickness. The superscripts "surf and "int" refer to the "surface" contributions of the vacuum/film and film / substrate interface. Since most measurements average over both contributions, Ki is replaced by the average over both interfaces 2 KjS. In most experimental studies the Kj have been analyzed in second order only, that is K2V and K2S is determined from a plot of K2 as a function of 1/d or by plotting K2-d versus d. Both contain contributions of dipolar (shape) and spin-orbit (intrinsic) origin. Changes in symmetry (lattice structure) and lattice distortions can lead to a tremendous enhancement of K2V. This effect is due to spin-orbit interaction and can be orders of magnitude larger than the shape anisotropy. Any epitaxially grown film on a non-magnetic single crystal substrate is strained. Depending on the lattice mismatch between substrate and film different thickness regimes must be distinguished: the thickness range of coherency strain and for strain relaxation by incorporation of misfit dislocations. Both regions are connected at the critical thickness dcs which is given by the energetic minimum of the sum of elastic energy, which increases proportional to the strain volume, and the energy to form a dislocation. For a lattice mismatch r\ < 3% a critical thickness dcs 1020 ML is calculated. This thickness dependence has important consequences for the MAE. In the coherent growth regime the lattice constants are assumed thickness independent and different from the bulk structure, i.e. the film assumes the lateral lattice constant of the substrate and relaxes vertically according to the Poisson ratio.

446

Usually the anisotropy is given as energy per volume, e.g. J/m3. The surface and Stepp anisotropy have the dimensions of energy per area (J/m2) and energy per line (J/m) which makes numbers hard to compare. A better way is to give the anisotropy in energy per atom, which means that the atomic volume in a sample consisting of N atoms must be estimated. The different faces ((111) versus (100)) contain a different number of atoms per unit area. This gives a different surface anisotropy in eV/atom, when Ks in J/m2 is the same for both. A dimension of energy/atom allows a convenient comparison of volume-, surface-, and step-type anisotropy. Also the correlation to calculated values is much more straightforward. The dipolar interaction which senses the shape of the sample has for magnetic moments located on a two-dimensional sheet the lowest energy when all the moments are aligned in the film (x,y) plane. The magnetization of a thin film lies in the plane along a crystallographic direction determined by in-plane anisotropics. To produce a perpendicular magnetization the shape anisotropy Fd = Vz 11 o(Nj_-N||)M2 must be compensated. F 2) for bcc(001) layers, A= 0.2338 (n>2) for fcc(001) layers and 0.15 (n>3) for hcp(0001). The deviation is largest for the most open structure, a "bcc" film. On the other hand, susceptibility measurements indicate that the discrete summation of point dipole fields may give questionable demagnetization factors. The classical continuous thin disk approach seems to agree better with the experiment, but a satisfactory conclusion has not been reached yet. The shape anisotropy contributes in second order only. Some groups include the shape anisotropy in K2eff, that is they use K2cff = K2 - Vi ii oM2 + 2K2s/d. In our work the shape anisotropy is always subtracted, before the intrinsic (spin-orbit) anisotropy (K2 or K4) is discussed.

447

3. Temperature Dependence of MAE 3.1. Phenomenological Description The magnetic anisotropy vanishes above the ordering temperature T c of the ferromagnet. Despite the close relation with the anisotropy of the orbital magnetic moment one must not conclude that the magnetic moment per atom or the orbital magnetic moment vanishes. Also, the difference of the orbital magnetic moment along the easy and the hard magnetic axis persists above T c . What is the experimental evidence for this? First and foremost the temperature dependence of the magnetic susceptibility follows a Curie-Weiss law. The Curie constant is NOT zero, proving the existence of a magnetic moment also in the intinerant ferromagnet Ni above T c . As the magnetic moment above T c is the same (except for some additional polarization of the conduction electrons) as the one measured below T c (for T=0 K), it is reasonable to conclude that the orbital magnetic moment is unchanged in the paramagnetic state. Consequently, the ratio of orbital-to-spin magnetic moment is temperature independent. A direct proof of the existence of the orbital magnetic moment and its anisotropy in the paramagnetic state is given by the deviations of the spectroscopic splitting factor, the g-factor, from the pure spin value g=2.0023. This has been measured by paramagnetic resonance and theoretically described for example in crystal field theory. If it is experimentally proven that the anisotropy of the orbital moment exists above T c , but the magnetic anisotropy energy vanishes above T c , there seems to be a conceptual problem in the microscopic origin of the MAE. To explain that one has to consider the temperature dependence of the magnetization. The magnetization also decreases as the temperature increases from T= OK due to the excitations of spin waves or in other words fluctuating moments. This does not mean that the magnetic moment vector vanishes, but in fluctuates so quickly and uncorrected to other moments above T c that spatially and timely averaged moment vanishes. Hence, also the macroscopically measurable MAE vanishes, it averages out above T c . This hand-waving argument has been quantified [14] by expanding the MAE in a series of Legendre polynomials with temperature dependent coefficients k((T) according to MAE = ki(T)Y"(P) + ki(T)Y*(0) + magnetization by

• The coefficients are related to the

M Z 2 x M D ! ! ^ l , i . e . k 2 o c M ( T ) 3 , k4 0cM(T)'° *,(0) M(0) Assuming a typical temperature dependence of the magnetization one can plot the anisotropy coefficients as shown in Fig. 1. One sees that the k; decrease monotonically with increasing temperature and vanish at T c . If one confuses these temperature dependent kj with the usual magnetic anisotropy parameters

448

Kj one would draw the conclusion that a temperature change of the easy axis of magnetization is not possible. However, as discussed in [6] for example, one finds that if one rearranges the cos and sin terms in the legendre polynomials in terms of increasing powers that the new parameters K, (actually the ones used in the experimental analysis) can vary in sign and their temperature dependent change of sign in Co or Gd can be quantitatively understood. An illustrative example is plotted in Fig. 1 showing that K2 = 1.47 k2 - 3.3 lct(T) changes its sign.

Kf/ arb.units

—o—K

K

41

' ^

K2(T) = 1.47 k 3 - 3.3 k„(T) K, l (T) = 3.85k4(T)

100

200 300 Temperature (K)

400

Figure 1. Theoretical temperature dependence of the anisotropy coefficients. Note, that a change of sign of K2 anisotropy parameters measured in the experiments can be constructed.

3.2. Surface Anisotropy as a Function of Temperature As mentioned before, the magnetic anisotropy at an interface or surface differs dramatically from the respective interior. Hence, also the temperature dependence of bulk and surface/interface anisotropy is different. Before one can discuss the different temperature dependencies one needs to realize that the Curie temperature changes as a function of film thickness. What is the effect on the MAE? In Fig. 2 a possible scenario is sketched. The reduction of T c causes a "compression" of the MAE along the temperature axis. Along the vertical axis one may expect an increase of the MAE due to the reduction of the effective coordination number in thin films. For the experimental analysis one has the following evident problem. To determine surface anisotropics one usually measures at constant absolute temperature for different thicknesses and plots the result as a function of reciprocal film thickness. If the Curie temperature,

449

however, approaches or even decreases below the measuring temperature, this analysis becomes wrong. The correct way is to use the thermodynamic relevant temperature T/T c , as has been confirmed in several studies. One may note that this correct way of determining the Temperature dependence of Ks and K requires knowledge of the Curie temperature for each film thickness d. Examples have been presented in [6].

ferromagnet ^K-2

....

.•

•,

paramagnet

Figure 2. Schematics of temperature dependence of MAE when the Curie temperature changes as a function of film thickness from a 3D to a 2D value.

4. Magnetic Domain Structure So far the discussion has been restricted to single domain behavior. Due to the competition of magnetostatic energy which tries to minimize the stray field energy and the exchange energy which tries to keep the magnetic moments parallelly aligned one finds characteristic domains in ferromagnetic layers. As an illustrative example, the domain formation of an in-plane magnetized film is presented. The easy axis of Ni layers on Cu(OOl) is known to be in-plane up to a thickness of approximately 8 monolayers. One finds [15] large (several micrometer) magnetic domains separated by a Neel wall as shown in Figure 3 for a 4.8 ML Ni/Cu(001) film prepared at 300 K and measured in-situ at 100 K. Our microscopic observation of in-plane magnetization at that thickness is in good agreement with reference [16]. We observed large domains with sizes of several 10 urn, substantially larger than our 10 um maximum field of view. To confirm that our images of these domain walls are representative of typical configurations, we traced domain walls over extended distances. In Figure 3 the imaged area (circles) was moved in several steps to trace the domain wall. The correct alignment of the magnetic images was unambiguously verified by

450

comparing surface-step patterns in the corresponding low energy electron microscopy (LEEM) images, which are not shown here. Close inspection of the "spin-up' and % spin-down' LEEM topographic images and comparison with the corresponding magnetic images reveals no correlation between the topography of the Ni film and its magnetic domain structure. For example, we did not find evidence for domain wall pinning at atomic step bands. The increased wall roughness of the film measured at 100 K compared to measurements at 300 K can be attributed to two effects: a) when cooling from 300 K to 100 K in 10"8 Pa trace amounts of residual gases (CO, C0 2 ) may adsorb, which reduce the total magnetic anisotropy in the 5 monolayer regime considerably, b) the formation of a zig-zag domain wall which is not completely resolved. Such domain walls are known [17] (to originate from "head-on" 180° domain walls). One should note, however, that we find no statistically significant differences between the domain structures at 100 K and 300 K for up to 8 monolayers thickness when comparing many images recorded in different areas of films grown at 100 K and 300 K.

Figure 3. Large magnetic domains of Ni/Cu (001) imaged by spin-polarized low energy electron microscopy. Three images with 10 urn fteld-of-view (circles) are added to trace a domain wall of a 4.8 ML thick Ni film prepared at 300 K and measured at 100 K. The domain sizes are several 10 pm. No preferred direction of the domain wall with respect to crystallographtc axes was found.

The structure of the domain wall in an 8 ML Ni/Cu (001) film was analyzed in greater detail in Figure 4. The spin polarized low energy electron microscopy allows an unambiguous determination of all components of the magnetization vector M by rotating the spin polarization P of the reflected low energy electron beam. First, while keeping the azimuthal angle fixed at ^=0°? the polar angle 0 of P was varied in 10° steps from 90° (spin-polarization in-plane) to 8=0° (spin-

451

polarization along the surface normal). Diminishing contrast in this series confirms the absence of out-of-plane magnetization components, i.e., the local magnetization vector lies in the surface plane in both domains. After that, the polar alignment of the illumination beam polarization was fixed at 8=90° (spin» polarization in-plane) and the azimuthal polarization orientation was swept through an angle of -135°. The magnetic contrast between the domains can be seen to decrease in this series, it finally vanishes when the beam polarization is perpendicular to the magnetization vectors of the two domains. The absence of magnetic contrast between the two domains for |=-90° confirms that the two domains are anti-parallel and separated by a 180° domain wall. In all in-plane orientations, except at the 4 =0°* additional contrast can be discerned in the region of the domain wall Most clearly at #=»90°? a large section of the domain wall appears brighter, and a shorter segment near the top of the image appears dark. Our interpretation of this contrast is that the domain wall has a N6elstracture, in which the spin-reorientation between the two anti-aligned domains takes place within the film plane. The fact that different sections of the wall show opposite contrast is consistent with the expectation that N6el-walls must occur in two degenerate chiralities, as indicated schematically in the figure.

Figure 4. SPLEEM images of an 8 ML Ni/Cu (001) film at 300 K as a function of the polar angle 0 (top, #=0°) and azimuthal angle t n e exchange constant A had to be assumed one order of magnitude smaller than the bulk value and Ken- turned out to be more than two orders of magnitude larger than typical Fe film values. 5. Conclusions Some aspects of the physics of magnetic monolayers where discussed in this contribution mainly focussing on the issues associated with magnetic anisotropy energy. It has been emphasized that to obtain reliable data it is important to consider the temperature and thickness dependence of the magnetic quantities, which are investigated.

453

Acknowledgement The fruitful collaboration and many discussions with so many colleagues working in the field of thin film magnetism as well as the financial support by the Deutsche Forschungsgemeinschaft is thankfully acknowledged. References 1. 2. 3. 4. 5.

6. 7.

8.

9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.

J.A.C. Bland, B. Heinrich (eds.), Ultrathin magnetic structures I & II (Springer Verlag, Berlin Heidelberg, 1994). H.J. Elmers, Ferromagnetic Monolayers, Intern. Journal of Modern Physics B 9, 3115 (1995). U. Gradmann, Magnetism in ultrathin transition metal films, in: Handbook of Magnetic Materials, Vol.1 (Elsevier Science Publishers B. V., 96,1993). D. Sander, The correlation between mechanical stress and magnetic anisotropy in ultrathin films, Rep. Prog. Phys. 62, 50 (1999). CM. Schneider, J. Kirschner, Magnetism at Surfaces and in Ultrathin Films, in Handbook of Surface Science, Vol. 2 Elsevier Science B.V., 668 (2000). M. Farle, Ferromagnetic Resonance of ultrathin metallic layers, Rep. Prog. Phys. 61, 755 (1998). B. Hillebrands, Brillouin Light Scattering from Layered Magnetic Structures, in Topics in Applied Physics, Vol. 75 (Springer-Verlag Berlin Heidelberg 1752000). K. De'Bell, A.B. Maclsaac, J.P. Whitehead, Dipolar Effects in magnetic thin films and quasi-two-dimensional systems, Reviews of Modern Physics 72, 225 (2000). P. Poulopoulos, K. Baberschke, Magnetism in thin films J. Phys.: Condens. Matter 11,9495(1999). J. Nogues, Ivan K. Schuller, Exchange bias, Journal of Magnetism and Magnetic Materials 192, (2) 203 (1999). V.S. Speriosu, D.A. Herman, Jr., I.L. Sanders, T. Yogi, Magnetic thin films in recording technology, IBM J. RES. DEVELOP. 44, 186 (2000). I.K. Schuller, S. Kim, C. Leighton, Magnetic superlattices and multilayers, Journal of Magnetism and Magnetic Materials 200, 571 (1999). M. Farle, Magnetic Thin Films, in NANOSCALE MATERIALS eds. L.M. Liz-Marzan, P. V. Kamat, (Kluwer Academic, 395, 2003) H.B. Callen et al. , Phys. Rev. 139 (1965) A455 ; W.J. Car, Jr.; Phys. Rev. 109, 1971 (1958). R. Ramchal, A.K. Schmid, M. Farle, H. Poppa, Phys. Rev. B 68, 054418 (2003). K. Baberschke, M. Farle, J. Appl. Phys. 81, 5038 (1997). A. Hubert,R. Schafer, Magnetic Domains, Springer (1998). A. Berger, H. P. Oepen, Phys. Rev. B 45, 12596 (1992). M. Pratzer et al., Phys. Rev. Lett. 87, 127201 (2001).

MICROSCOPIC MECHANISMS OF MAGNETOOPTICAL ACTIVITY IN EPITAXIAL GARNET FILMS A.I. POPOV Moscow Univ. of Electronics - MIET, Russia E-mail: [email protected]

The microscopic mechanisms of magnetooptical activity in dielectric compounds with fand d-ions are described. The results of experimental and theoretical study of the Faraday effect and the magnetic linear birefringence in the various materials are presented and analyzed.

1. Introduction At the present time, the magnetooptical effects play an important role in the study of magnetic thin films. Magnetooptical characterizations of surface magnetism began a little over a decade ago. One of the first surface magnetooptical studies, concerned with the magnetic properties of ultrathin Fe films grown epitaxially on Au [1]. Hysteresis loops of the Fe film with atomic layer sensitivity were successfully obtained. Since then magnetooptical effects have emerged as a premier surface magnetism technique. Magnetooptical effects have been applied to various topics in lowdimensional magnetism ranging from the detection of magnetic order to the characterization of critical behavior, magnetic surface anisotropics and the oscillatory antiferromagnetic coupling exhibited by giantmagnetoresistance heterostructures. Additional interest to the magnetooptical effects is generated by the recent commercialization of high-density magnetooptical information storage media [2]. 2. The Permittivity Tensor Magnetooptical effects are manifested themselves in changing of the light phase, its intensity and polarization, which arise as the result of interaction between light and magnetic matter. The reason of magnetooptical effects comes about from the appearance of the magnetic linear and circular birefringence and dichroism. The magnetooptical phenomena are governed by dielectric permittivity tensor (in the visible and UV region of the spectrum) and by magnetic permittivity tensor (in the infra-red region). We consider the dielectric materials 454

455

with f and d - ions. In this case, the wave functions of the f and d electrons are localized. The components of the polarizability tensor a^ of the ion in electronic dipole approximation are determined by Kramers-Heisenberg formula

^=hpg\ heg

d egdJge s

'

.r +

[G)eg-(o-ireg

digdge

.r 1 .

(i)

+iregj

where the sum is taken over all of the ground (g) and excited (e) states of the ion , pg is the Boltzmann factor giving the population of the energy level E„, d'eg is the matrix element of the /-th component of the dipole moment connecting the g - and e -states, fi(Oeg = Ee-Eg

and Teg is

the half-width of the spectral band of the g - e transition. The contribution of magnetic ions of concentration N to the materials permittivity is SSjj = 4nN

n +2

(Zjj ,

(2)

where n is the average refractive index. In the most cases we actually deal with the "long wavelength wings" of the allowed transition between the groups of levels (multiplets, terms, etc.). In this case Ee-Eg

= Tio)ov + (AEe - AEg ) ,

where C0OV is the mean frequency of the transition between ground and exited groups of levels, AZse, AEg are the splitting of these groups. In the frequency range, where AEe - A £ J « (1) as a power series [AEe - AEg ) «

h\d) - Q)ov\ we can develop

ti{p) - coov)

456

*"i£*

"g&g

4e4 •g(reg 1+

OH-C0OV

c*-%i

^-+..

(3)

ti(a)-%v)

The effects of absorption are neglected. Let us consider the magnetic ions which orbital angular momentum is not equal zero. Neglecting the linear (diamagnetic) and higher order terms we can find [3-5]

+a

(4)

2 < QV(L) > ,

where < LK > and < Qtj > are the components of the orbital momentum L and quadrupled moment Q. The coefficients a0, ax, a2 are defined in terms of the matrix elements of the optical transitions considered [3,6]

an=-2h~1Y.A(L)(co2

-aLr^LyPra.,

L here

A(L) = (2L0+l)(L0\\dI\\L)2, P0L=1/3,P1L C1(L0+1)

=-(2o)Lvr1o>C1(L),P2L

= (L0+1)-},C1(L0)

=

=C2(L), [L0(L0+1)]-1,

Ci(L0 -1) = L0 , C2(L0 +1) = -[(L0 + 1)(2L0 +1)]'1, C2(L0) = Cj(Lo), C2(L0-1)

=

-[L0(2L0-J)]-1.

The first term in (4) is the isotropic contribution to the polarizability of the ion, the second term is the gyrotropic contribution, and the third term represents the even magnetooptical effects. It should be noted that the gyrotropic contribution to the polarizability tensor of an ion is determined by the average value of orbital momentum of the ion, and that the last term in ( 4), which represents the even magnetooptical effects, is proportional to the average quadruple moment < Q^ > of the ion. In other words the even magnetooptical effects like

457

magnetic linear birefringence are governed by the response of the quadruple momentum of the magnetic ion on the magnetic field. This leads to the conclusion that the odd effects like Faraday effect and magnetooptical Kerr effects measure the average orbital momentum of an ion (if it is non-zero and allowing a correction for small diamagnetic and magnetic dipole contribution), whereas using the even effects one can measure the quadruple moments of ions [3-6]. This result is applicable not only to the f - ions, but also to the d - ions, to clusters of MeOn type, where Me is a transition metal, etc. It is very important that in the case of d-ions the orbital momentum is usually frozen, i.e., in the first approximation < L > = 0. It is partially unfrozen under the influence of spinorbital interaction, but in this case the diamagnetic contribution can be of the same order as or exceed the "unfrozen" contribution from . In a number of cases one should take into account the magnetic dipole transitions. It can be important for Faraday effect in the infra-red and near visible spectral region (gyromagnetic contribution to the Faraday rotation [7]. As has been shown by Krinchik and Chetkin, gyromagnetic Faraday effect does not depend on the radiation frequency

aM=CM^M, mc

(5)

where M is magnetization, g is the g-factor of the magnetic ions. 3. The Odd Effects in Rare-Earth Materials The rotation angle of polarization plane in the visible and ultraviolet spectral ranges is predominantly determined by the gyroelectric contribution due to the electric susceptibility of the medium. For rare-earth ions, which occur in dielectric media, the gyroelectric contribution is formed, for the most part, by the 4fN - 4fN~l5d electric dipole transitions, except the narrow spectral regions close to the resonant frequencies of the forbidden//transitions. It was demonstrated [3-5] that the contribution to the rotation angle of polarization plane from the 4f - 4fN~' 5d transitions for the magnetic ions with the nonzero orbital angular momentum is equal to

aF=a{Lz)

+ VDH

(6)

The first term is the combination of the paramagnetic contribution and mixing contribution, i.e., the contribution due to the multiplet interaction (J-J mixing), and the second term is the diamagnetic contribution. Let us consider the Faraday rotation in the organic glass containing the Eu3+ ions. The Faraday effect in the organic glass containing the Eu3+ ions can be analyzed with expression (6), in which the second term is the combination of

458

diamagnetic contributions of europium ions and the matrix. It is important that the deviation of the Faraday rotation from the linear dependence with an increase in the external magnetic field H is determined solely by the nonlinear behavior of the average orbital angular momentum of the Eu3+ ion as a function of H. In [8], the average orbital angular momentum (Lz) of the Eu ion in ultrastrong magnetic fields at T = 30 K was calculated, and the experimental dependence obtained at X = 0.85 um was compared with the theoretical curve described by formula (6). It is seen from Fig. la that the theoretical results, which were obtained at VD =3.14 x 104 deg/(cm Oe) and a = 75.67 deg/cm, are in good agreement with the experimental data. In order to determine the contribution from the orbital angular momentum of europium ions to the Faraday effect, we subtracted the diamagnetic term (VD), which linearly depends on the magnetic field, from ocF and, thus, determined the dependence of the mixing contribution (J-J mixing) on the magnetic field. Then, this dependence was compared with the field dependence of the average orbital angular momentum of europium ions, which was also calculated (Fig. lb). It is seen that the theoretical results, as a whole, are in reasonable agreement with the experimental data. Note that such a treatment of the results requires the high accuracy of the measurements of H and a F at a low concentration of the Eu3+ ions. For 8 = Ah/ h = A(Xp /(Xp = 2 %, the < L^ > values obtained from the treatment of the experimental results (Fig. lb) and the calculated data coincide within the limits of error. Now, we describe the Faraday effect in ultrastrong magnetic fields in the vicinity of the resonant frequencies of the forbidden//transitions. This situation is apparently realized under laser radiation at the wavelength X = 0.63 um. To accomplish this, it is necessary to add the contribution of the adjacent forbidden absorption line ccR to expression (6). Then,

aF=a(Lz)

+ VDH + aR(H)

(7)

The resonant frequency G)Q of the actual forbidden optical transition depends on the magnetic field strength H and, in a certain field, can achieve the frequency of the laser radiation used. In other words, we believe that the optical resonance induced by the magnetic field takes place in this case. For the qualitative description of the contribution of the field-induced optical resonance to the Faraday rotation, let us use the linear approximation of the dependence of the resonant frequency on the magnetic field co(h) = 0)Q + yH, where y is the rate of frequency change. In this case, it follows from formula (1) that

aR=f(H)-f(-H), where

(8)

459

C(co20(H)-coj-r2)co2

/YW) J(")

=

y

y

y

2

y J

(coo(H)-co, -r ) + 4a)2r2

Here, T is the width of the forbidden line (for rare-earth ions, the spectral lines are very narrow r ~ 1013 s"1), CO] is the frequency of the laser radiation (for the radiation at the wavelength X= 0.63 um, cox = 3 x 1 0 s'1), and C is the coefficient proportional to the oscillator strength for this transition. Parameters y and Aco = COQ — CO] can be expressed through the magnitudes of the fields Hi and H2 [where H, is the field at which f(H) = 0 (according to [9], H, = 10 MOe), and H2 is the field at which f(H) reaches the maximum value (according to [9], H2 = 9 MOe)], as follows Aco = rH]/(H]-H2) y = r/(H]-H2). Then, f(H) takes the form .,„. . 2(H]-H )(H]-H) 2 2n ] f(H) = A ' '— (9) (H,-H)2+(H]-H2)2 where A = CCOQ / 4T. Expressions (7)-(9) permit us to describe the experimental field dependence of the contribution of the Eu3+ ions to the Faraday rotation. Figure 2 shows the comparison between the experimental data obtained in [9] and the results of calculations at //, = 8.8 MOe, H2=10.2 MOe, a = 55, A = 74 deg/cm, and VD=\3 deg/(cm Moe). Thus, the anomalies of ap(H), which were found in [9] with the use of the laser radiation at the wavelength X= 0.63 urn, are most likely due to the magnetooptical resonance induced by the ultrastrong field. Now let us consider the cases of the S-ions. Thefirstexample is the Faraday effect in KMnF3 [10]. There are three contributions of the Faraday rotation of KMnF3

aF where

a^

=ad+ap+amat,

is the diamagnetic

contribution

of the allowed

S(3d j - > p(3d 4p) electron transition, having the resonance frequency tico0 « 3 x l 0

cm'1, Op is the paramagnetic contribution of the forbidden

460

S-> Tig(G)

electron

transition,

having the

hcox « ( 1 3 - 1 7 ) x l 0 3 cm1 and the line width

resonance

fcr«(4-5)xl03

frequency cm"1,

Figure 1. (a) Dependence of the Faraday rotation angle of the organic glass containing europium ions (5 wt %) on the magnetic field at T = 30 K. Crosses correspond to the experimental data obtained under laser radiation at the wave-length X = 0.85 urn, and the solid line is the linear approximation of the dependence of the Faraday rotation on the magnetic field, (b) Dependence of the mixing contribution (J-J mixing) to the Faraday effect of europium ions on the magnetic field. Crosses correspond to the processing of the experimental data, and the solid line is the calculated field dependence of the average orbital angular momentum for europium ions at T= 30 K.

461 0CF(EU3+), deg

120-

Figure 2. Dependence of the Faraday rotation of europium ions on the magnetic field at room temperature. Points correspond to the experimental data obtained in [9] under laser radiation at the wavelength X = 0.63 urn, and the solid line is the result of calculations by formulas (7)-(9) at H} = 8.8 MOe, H2=10.2 MOe, a = 55, A = 74 deg/cm, and F 0 =13 deg/(cm Moe). a

mat ' s t n e diamagnetic contribution of the KMnF3 matrix. The energy of the

laser radiation in experiment is hco, « 1 6 x 1 0 6

S^>4Tlg(G)

cm"1, i.e. near transition

• The Faraday rotation ap can be expressed in the simple form

aF =a-H + b-m(H,T), (10) where a and b are constants, and m(H,T) is the relative magnetization. We have calculated the dependence aF(H) and have determinate constants a and b . The experimental and theoretical results of the investigations of Faraday rotation in KMnF3 are presented in the Fig. 3. It should be noted that the Faraday rotation does not get saturated with increasing field up to 400 T. The relatively small value of Faraday rotation in this matter can be explained by the almost complete compensation of the negative diamagnetic contribution (due to allowed s-p - transition) and positive contribution (due to forbidden 6S-**T1 (G) transition). In this case, the role of the Faraday rotation matrix is very important. The second example is the Faraday effect in GdGG. There are three contributions of the Faraday rotation of GdGG F = CgM(H,

T) - CpG)2M(H,

T) + a2yH,

(11)

462 60 •

Figure 3. Magnetic field dependence of the Faraday rotation in the 3.6 mm thickness KMnF3 sample; points - experiment, lines theory; The dotted lines show the field dependence of the second term bm(H,T) of the Eq.(lO).

where the first term is the gyromagnetic Faraday effect (that does not depend on frequency), the second term is the paramagnetic contribution (that approximately depends on frequency as 0) ), and the last term is the combination of diamagnetic contribution of Gd3+ ions and the matrix, M(H,T)is magnetization C „ , Cp, y are constants. The results of the experimental and theoretical study of thefielddependence of the Faraday rotation in GdGG at T= 4.2 K, A = 0.47 um and the magnetization are presented in Fig. 4 [5,11]. It should be noted that there is a strong field dependence of the Faraday effect. First, Faraday rotation increases with the field and then in then in the field region where the magnetization is practically saturated it starts to decrease. Second, the Faraday effect changes its sign at higher temperature. In [11], it has been predicted that at T = 70 K the Faraday rotation is independent of frequency because of the compensation of the paramagnetic contribution (the second term in (11)) and the diamagnetic contribution (the last term in (11)) and at X = 0.5 um the temperature dependence of Faraday effect is vanished because of the compensation of the gyromagnetic contribution (the first term in (11)) and the paramagnetic contribution (the second term in (11)). These peculiarities are naturally described by (11) and its cause is the competition of the above considered contributions to the Faraday rotation of GdGG.

463

H IkOe)

Figure 4. The field dependence of the Faraday rotation (curve 1) and the magnetization (curve 2) for Gd3Ga5Oi2 (A=0.47 urn, r = 4.2 K).

4. The Even Magnetooptical Effects The understanding of the microscopical mechanisms of magnetooptical activity in epitaxial rare-earth films is based on the knowledge of the electronic structure of magnetic ions. Especially it is essential for the even magnetooptical effects like magnetic linear birefringence. As a rule, growth induced magnetic anisotropy of thin epitaxial magnetic garnet films is uniaxial. It is manifested in the fact that the ground state of the rare-earth ions in such films is a doublet. In this case, at low temperature the magnetic linear birefringence can be presented in form [3]

AnozH-M(H,T), where M(H,T) is the contribution of the ground doublet to the magnetization. It is interesting to state that the magnetic linear birefringence has a linear field dependence in the region of magnetic saturation for the ions with doublet ground state. On the other hand the magnetic linear birefringence in systems of S-ions (GdGG) can be described by classical Akulov-Callen-Callen theory which yields the quadratic dependence on magnetization. In [12], the results of the experimental and theoretical study of the magnetic linear birefringence in heavy rare-earth garnets R 3 M 5 0i 2 (R= T b, Dy, Ho, Er, Tm, Yb; M = Al, Ga) in magnetic fields up to 5T and in temperature range from 4.2 K to 50 K are presented. A complete quantative agreement with the whole complex of experimental data have been obtained (Fig. 5-8).

464

HIT)

\ . 600

V*

\.

\ 0

\ 200

\

TbGG

TbAO 20

*0T(K|

TbGG

TbAG fc) 20

40 TIKI

Figure 5. The MLB [X=0.63 urn, kl(110)], and magnetic moment in Tb - based garnets; experiment: -m- , H||[lll]; A A A , H|| [110]; • • • ; H||[001]; theory: solid line (a) Field dependence of relative magnitudes of MLB coefficient 5n in TbAG at 4.2 K; inset show the magnetization curves; (b) the same dependence as in (a) for TbGG ;(c) temperature dependences of the absolute MLB coefficients An in TbAG and TbGG in H=4 T; inset shows the temperature dependences of magnetization.

465

HIT!

.i 6 i 15

2i /

f

S~

0

/ 2

/

HIT)

/

DyGG

/

0.5^

Ufl

'H'TI

300

\

OvAG, DVGG

.-200

100

0

20

\

DyGG

«T|K|

(c)

20

40 T(K)

Figure 6. (a) - (c) The same as in Fig 5 (a) - 5(c) for Dy - based garnets.

466

ooo •k

SOO -

\

H11 [HI]

"\ BOO -

4 \ 400-

+S. o

"**•.

4 ^

200 -

••

200 -

*^"' i

4 0 0 -1

T. 1

100

1

1

150

1

[-

ZOO

- i

I

I

250

i

300

Figure 7. The temperature dependence of MLB for SmIG , H || [111], H=l7 kOe, A=l.l5 urn; experiment: + + +, theory: solid line, *** contribution of the ground multiplet of Sm3+, - - contribution of J-J~ mixing.

4 0 0

^ 2 0 0

-

- ZOO

-

HIIC1003

St,

-••OO

- BOO

-

H

>

-

,.

Figure 8. The temperature dependence of MLB for SmIG , H || [100], tf=17 kOe, A=1.15 (am; experiment: + + +, theory: solid line, *** contribution of the ground multiplet of Sm3+, - - contribution of J-J- mixing.

467

5. The Magnetooptical Kerr Effect Optical anisotropy of magnetized medium manifests itself also in the reflection of light from its surface. Phenomena arising here are generally referred to as the magnetooptical Kerr effect. It is significant that the magnetooptical Kerr effect is dependant on the magnetic film thickness [13]. In Fig. 9, the results of the measurements and calculations of the Kerr rotation and ellipticity as functions of the iron layer thickness are presented [13]. •••

r

i

••

I

•

•

1

-

i

1

& 0.3 a

Ellipticity

/ • oo ° o ... 20

40

60

80

Iron thickness (nm)

-i

1

p

a

1

-

••

Ellipticity

Rotation

0 0 1 *-

20

40

60

80

Iron thickness (nm) Figure 9. The measured saturated values of the Kerr rotation (.) and the ellipcity(o)as function of fte iron layer thickness for s polarization (a) and p polarization (b). The curves show the calculated Kerr rotation (full lines), and Kerr ellipticity (dotted lines).

468

References 1. E.R. Moog, S.D.Bader, Superlattices and microstructures 1, 543 (1985). 2. S. Klahn, P. Hansen, F.J.A.M Greydaus, Vacuum 41,1160 (1990). 3. A.K. Zvezdin, A.l. Popov, H.I. Turkmenov, Sov. Phys. Solid State 28, 974 (1986). 4. A.K. Zvezdin, A.S. Ovchinnikov, V.I. Plis, A.I. Popov, JETP 82, 939 (1996). 5. 6. 7. 8. 9. 10. 11. 12. 13.

A.K. Zvezdin, V.A. Kotov, Modern Magnetooptics and Magnetooptical Materials (IOP Publishing, UK, 1997). N.P. Kolmakova, A.I. Popov, Physica B, 179, 19 (1992). G.S. Krinchik, M.V. Chetkin, Sov. Phys. JETP, 13, 509 (1960). M.I. Dolotenko, A.K. Zvezdin, G.G. Musaev et al., Phys. Solid State, 42, 726 (2000). A.I. Pavlovskil, V.V. Druzhinin, O.M. Tatsenko, et al., JETP Lett. 31, 622 (1980). A.A. Mukhin, V.V. Platonov, V.I. Plis, A.I. Popov, O.M. Tatsenko, A.K. Zvezdin, Physica B, 246 -247, 195 (1998). A.K. Zvezdin, S.V. Koptsik, G.S. Krinchik et al., Pis. Zh. Eksp.Teor. Fiz. 37,331(1983). N.P. Kolmakova, R.Z. Levitin, A.I. Popov, N.P. Vedernikov, A.K. Zvezdin, V. Nekvasil, Phys. Rev. B 40, 6170 (1990). K. Postava, J.F. Bobo, M.D. Ortega, et al., J. Mag. Mag. Mater. 163, 8 (1996).

DESIGN, FABRICATION AND APPLICATIONS OF MULTILAYER THIN-FILM SQUID SENSORS D. RASSI, Y.E. ZHURAVLEV School of Health Science, University of Wales Swansea, Singleton Park Swansea SA2 8PP, UK

Over the last thirty years, the Superconducting Quantum Interference Device (SQUID) has undergone considerable development in terms of basic design and fabrication methods as well as its applications, which now include such diverse fields as brain science and oil exploration. The SQUID is by far the most sensitive detector of magnetic fields (with an inherent field sensitivity down to femtotesla levels), can operate from near DC to tens of kilohertz and has a virtually unlimited dynamic range. In order to obtain the full benefit of such extreme sensitivity in the presence of much larger ambient magnetic noise (both from natural and man-made sources), sensor design and fabrication must be optimized. Also required are advanced digital signal detection and conditioning techniques, as well as a deep understanding of the nature of the signals to be detected and the environment in which the measurements are to be made. Over the last eighteen years our group in Swansea has been using and developing SQUID magnetometers for biomagnetic and geophysical applications with particular emphasis on systems and techniques for unshielded operation. In this paper, we review the optimization of SQUID parameters for practical applications, which are illustrated with specific examples.

1. Introduction There are two main types of SQUID sensors: the RF (radio frequency) SQUID with a single Josephson junction in the superconductive loop, and DC (direct current) SQUID with two Josephson junctions incorporated into the superconductive loop. The general rule is that the noise of the RF SQUID is proportional to its electronics operational frequency which can be hundreds of megahertz. This causes significant design problems for the SQUID electronics which, as with most electronic devices operating at such high frequencies, is subject to external noise interference at a similar frequency range. The DC SQUID by its nature operates at the intrinsic Josephson quantum effect generated GHz frequency and does not require high frequency electronics. For the DC SQUID, a modest 100 kHz modulation to allow reduction of 1/f noise works well. Therefore, the DC SQUID is the preferred choice as a practical lownoise magnetic sensor. Fabrication of SQUID sensors from high-temperature (mixed metal oxide) superconductors was probably one of the first practical achievements in this exciting new field. And indeed in the last ten years there have been significant improvements in the sensitivity and quality of these sensors. However, our 469

470

experience of operating one of the best commercially available high-temperature (HTc) SQUID magnetometers with a white noise of 40 fT/VHz highlighted several shortcomings. It was shown, for example, that the sensor can become 10 times more noisy and unstable when operated in an open unshielded environment because it is affected by the Earth's magnetic field. Hence, in spite of the complexity and cost associated with the use of liquid helium needed for conventional low-temperature (LTc) superconductors, at present the LTc DC-SQUID offers the best performance as a magnetic field detector in terms of stability, sensitivity and low noise. 2. Practical SQUID Design Years of trial and error by different research groups around the world have resulted in the conclusion that the best practical SQUID sensor is a thin film structure incorporating Josephson junctions made of traditional Nb-AlOx-Nb materials. Such sensors possess ultra-low noise and a high level of stability of parameters. DC SQUID with very low inductance of the order of 1-10 pH can have extremely high intrinsic energy sensitivity down to just a hundred Planck constant. But it is not an easy task to efficiently couple such a device to the outside world and for this reason they are not really suitable for practical use. Several multi-loop designs were suggested to overcome this problem but this creates the problem of parasitic resonances due to the additional cross-layer capacitors. As a result, the SQUID sensor together with the coupling coil considered as one system is prone to microwave resonances [1] that increase the low-frequency noise of the sensor with standard electronics. Because of these resonances, the high resolution of a multi-loop sensor can only be realized for a small interval on the voltage-flux SQUID characteristic. In order to improve the performance of multi-loop SQUID sensors with integral input coil, a common design practice is to add a dumping circuit of resistor and capacitor connected to the SQUID superconductive loop [2]. A practical SQUID design, first suggested by Ketchen [3], consists of a SQUID base electrode in the shape of a square washer of a few mm side and a spiral integral or hybrid input coil on top of this washer. The problem with integral input coil is that the thickness of the insulating layer cannot be made bigger than hundreds of nanometers which gives rise, just as in the multi-loop SQUID, to parasitic capacitance and microwave resonances necessitating damping circuits. A hybrid input coil, on the other hand, is fabricated on a separate substrate which makes it possible to increase the insulation thickness between the SQUID washer and the spiral coil. The decrease of coupling

471

coefficient associated with this approach is compensated by the decrease of parasitic resonant noise and in some cases good coupled energy sensitivities over a substantial part of the SQUID dynamic range can be obtained. 3. Optimization of SQUID Parameters For practical SQUID sensors, the most important parameter is coupled energy resolution Ec Ec=o*)=j^jMutytf

CD

where W represents surface roughness, h is height of i-th column on the surface at time t and h is average surface height. CE surface roughness has measured by utilizing AFM observation and the above expression resulting a roughness of about 0.56 nm. This value indicates an almost smooth interlayer (spacer) for the Co/Cu/Co structure that can be used in GMR active region. To evaluate surface electrical resistance of the deposited layers, we have employed Rs technique. Based on our measurements, the average sheet resistance for the deposited Cu/Co/NiO/Si(100) structure was obtained 131 Q/D.

Figure 3. 3D-AFM micrograph of the Co surface in the Cu/Co/NiO/Si(100) structure.

Figure 4 illustrates AFM micrograph of the Co surface topography deposited on the Cu/Co/NiO/Si(100) structure in 3x3 |im scale . The electrical property of deposited Co/Cu/Co/NiO/Si(100) structure was measured as compared with the

504

other systems. The average sheet resistance for the Co/Cu/Co/NiO/Si(100) structure was obtained 124 Q/D. Also, we have studied the degree of cobalt surface roughness by AFM method and measured its value of about 1.52 em. According to our AFM analysis, the surface roughness of the second Co layer in the present structure is approximately the same as the roughness of the first Co layer deposited on the NiO surface. Recently, a modified Co surface layer was fabricated in the Cu/Co/NiO/Si(100) magnetic multilayer structure by applying an optimum negative bias voltage (Vb=-60) during its sputtering growth employing AFM5 Rs and SEM techniques [11].

Figure 4. 3D- AFM micrograph of the Co surface in the Co/Cu/Co/NiO/Si(100) structure.

In order to study corrosion resistance of the deposited Co/Cu/Co/NiO/ Si(100) system , we have measured corrosion rate of the top Co surface layer. Potential dynamic method was used to measure that in the H20 electrolyte. Figure 5 depicts the corrosion rate of the Co/Cu/Co/NiO system. A corrosion rate of about 0.22 mpy was obtained indicating stability of Co layer against humid media. To compare properties of the deposited subsystems, our major experimental results are summarized in Table.2.

505

1.5

1

> LU

> LU * *~w\*w.r.

0

-0.5, -3

,

, , r— -7 -5 Log I/Area (A/cm2)

,

, -3

Figure 5. The variation of voltage with surface current density for the Co/Cu/Co/NiO/Si(100). Table 2. Results for the Co/Cu/Co/NiO/Si(100) structure.

System Co/NiO/Si(100) Cu/Co/NiO/Si(100) Co/Cu/Co/NiO/Si(100)

(Q/D)

Surface Roughness (nm)

433 131 124

1.57 0.56 1.52

Sheet Resistance

4. Conclusions We have fabricated the Co/Cu/Co/NiO/Si(100) magnetic multilayers system using combined sputtering-evaporation methods. Different techniques including AFM, Rj and surface roughness measurement were utilized to determine surface properties of the deposited structures. Based on our AFM, Rs and roughness analysis, the Cu layer formed on the bottom of the second Co layer under our experimental conditions exhibits an appropriate spacer in the Co/Cu/Co structure with good property that can be used as an active region in a GMR structure. Acknowledgments The authors would like to thank the Research Council of Sharif University of Technology for financial support of the project. Useful discussions with Dr.

506

Rahimitabar and Mr. Jafari for AFM analysis as well as the assistance of Mr. Gholami for Rs measurements and Mr. Azimirad are greatly acknowledged. References 1. 2. 3. 4.

B. Dai, J.N. Coi, W. Lai, J. Magn. Magn. Mat. 27, 19 (2003). H.W. Jiang, M.H. Li, G. H. Yu, J. Magn. Magn. Mat. 242-245, 341 (2002). B. Dieny, J. Magn. Magn. Mat. 136, 335 (1994) H. Yu. C.L. Chai, H.C. Zhao, F.W. Zhu, J.M. Xiao, J. Magn. Magn. Mat. 224,61 (2001). 5. W. Guda, K. Shiiki, J. Magn. Magn. Mat. 205, 136 (1999). 6. J. Pelegri, J.B. Eje, D. Ramirez, P.P. Freitas, Sens & Act. A 25, 132 (2003). 7. D. G. Hwang, CM. Park, S.S. Lee, J. Magn. Magn. Mat. 166, 265 (1998). 8. H. Chingtong, F. Liu, K. Stoeu, Y. Chen, X. Shi, C. Qian, J. Magn. Magn. Mat. 239, 106 (2002). 9. S.S. Parkin, R. Bhadra, K.P. Roche, Phys. Rev. Lett. 66, 2152 (1991). 10. Porqueras, E. Brtran, Thin Solid Films 370, 10 (2001). 11. A.Z. Moshfegh, P. Sangpour, Physica Status Solidi(c), (2004) (accepted).

OPTIMISATION OF THIN FILM MULTI-LAYERS BY MICROMAGNETIC SIMULATIONS FOR MR APPLICATIONS P. GORNERT, D.V. BERKOV, N.L. GORN Innovente.V., Prussingstr. 27B, D-07745 Jena, Germany E-mail: [email protected]

The understanding of the equilibrium magnetisation structure and quasistatic remagnetization processes in single- or polycrystalline layers is extremely important when developing technical applications such as magnetoresistive (MR) sensors and Magnetic Random-Access Memories (MRAM's). To optimise the design of singleand/or multilayer elements, we have developed a programme package, which allows to calculate quasistatic remagnetisation processes (e.g., the hysteresis loops) and corresponding magnetic domain structures for definite geometric arrangements, edge roughness, and material parameters of magnetic layers on the base of finite-difference approximations of the energy contributions. The equilibrium magnetisation state for the given external conditions is found by minimising the total magnetic free energy of the system, which includes four standard contributions: external field, anisotropy, exchange and demagnetising energies and - for multilayers - the interlayer exchange energy (ferroor antiferromagnetic interlayer coupling). Various examples of simulations will be presented to illustrate some important topics of the MRAM design and the abilities of the program package.

1. Introduction Magnetoresistive (MR) sensors (AMR - anisotropic magnetoresistive, GMR giant magnetoresistive, TMR - tunnel magnetoresistive) and Magnetic RandomAccess Memories (MRAM's) play an increasing role in science and technology. These devices consist of single- (AMR) or multi-layer (GMR, TMR, MRAM) magnetic thin film elements. The functioning of all these elements is based on the magnetism in thin films and at interface effects where both charge and spin of electrons play an important role. In this article, we will focus our attention on the quasistatic switching behaviour of MRAM elements, which are considered as the memory elements of this century with a world market of about 40 billion US$. The first commercial products are expected in 2005 and the maximum storage density should be 400 Gbit/inch2. To satisfy the general requirements and to optimise the details of the construction of such MRAM elements we model their hysteresis loops and the corresponding magnetic domain structure by means of micromagnetic simulations under realistic experimental conditions. In fact, it is well known that computer simulations help to find out optimum system parameters and 507

508

processess if the simulations (1) reflect the reality well enough and (ii) they can be done quick enough - preferably with a normal PC. That means in our case, micromagnetic simulations [1-3] have to be carried out on a PC in a short time [4-10] for thermally activated processes [11,12]. In tie following we present some general aspects of MRAM's and some examples of simulations of switching behaviour of single and double-layers of soft and hard magnetic thin films in the inhomogeneous magnetic field of crossed conductors as used in MRAM's. For the calculations we consider polycrystalline films with cubic or uniaxial anisotropy of the grains and ferromagnetic or antiferromagnetic interlayer coupling. The films may possess arbitrary shape and defined edge roughness. Different layers in a multi-layer system can have different magnetic parameters. All calculations with the MicroMagus programme package [13] were carried out with a normal PC on the base of finite-difference approximations of the energy contributions. The equilibrium magnetisation state for the given external conditions is found by minimising the total magnetic free energy of the system (c.f. section 3). Various examples of simulations will be presented in section 4 to illustrate some important topics of the MRAM design, and the abilities of the programme package.

MRAM cells (cf. Fig. 2) apply the TMR arrangement and consist therefore of a hard (CoPe, CoPt) and a soft (NiFe) magnetic thin film separated by a nonconducting tunnel barrier (A1203, A1N) as shown in Fig. 1. i

1

-r

50-

•

j

30

i

!

20-

1

10-

i

40-

1

,

!

soft ferromagnet

...,

0-

,

__ ., -0

-20

: r-—

—T

,

•

1

20

hard ferromagnet

field (Om)

Figure 1. Typical TMR arrangement with an amplitude of 50% [14].

Figure 2 illustrates the MRAM cell with TMR elements as shown in Fig. 1. Moreover, the conductors above and below the elements can be seen. The conductors produce the magnetic switching field during writing and measure the resistance during reading. Parallel magnetisations with resistance Rft means "0" and antiparallel ones with Rfj, means "1", where Rti > Rft- The TMR amplitude

509

is defined as TMR = (Rti - Rtt) / Rtt» where the TMR amplitude in Fig. 1 is 50%.

Figure 2. Schematic arrangement of a MRAM cell.

3. Theoretical Background Our micromagnetic calculations are based on the minimisation of the total energy Et0, consisting of four energy contributions fc'tot = êxt ""• t-an "*" texch "*" E

Physics and Technology of Thin Films IWTF 2003: Proceedings of the International Workshop, Tehran, Iran 22 February - 6 March 2003

Finite Groups 2003: Proceedings of the Gainesville Conference on Finite Groups, March 6-12, 2003

php|architect (February 2003)

Scientific american (March 2003)

Computational linguistics and intelligent text processing: 4th international conference, CICLing 2003, Mexico City, Mexico, February 16-22, 2003: proceedings

Finite Groups 2003: Proceedings Of The Gainesville Conference On Finite Groups, March 6 - 12, 2003 (De Gruyter Proceedings in Mathematics)

Computational Linguistics and Intelligent Text Processing: 4th International Conference, CICLing 2003, Mexico City, Mexico, February 16-22, 2003. Proceedings

Tinear(2003-07-22)

php|architect (March 2003)

Experimental and Efficient Algorithms: Second International Workshop, WEA 2003, Ascona, Switzerland, May 26-28, 2003, Proceedings

Digital Watermarking: Second International Workshop, IWDW 2003, Seoul, Korea, October 20-22, 2003, Revised Papers

Preparation of Thin Films

Computational Methods in Sciences and Engineering 2003: Proceedings of the International Conference (Iccmse 2003)

Algorithms in Bioinformatics: Third International Workshop, WABI 2003, Budapest, Hungary, September 15-20, 2003, Proceedings

Cooperative Information Agents VII: 7th International Workshop, CIA 2003, Helsinki, Finland, August 27-29, 2003, Proceedings

Waveguide Spectroscopy of Thin Films, Volume 33 (Thin Films and Nanostructures) (Thin Films and Nanostructures)

Asian and Pacific Coasts 2003: Proceedings of the 2nd International Conference, Makuhari, Japan 29 February - 4 March 2004

High Performance Computing: 5th International Symposium, ISHPC 2003, Tokyo-Odaiba, Japan, October 20-22, 2003, Proceedings

Advanced Parallel Processing Technologies: 5th International Workshop, APPT 2003, Xiamen, China, September 17-19, 2003, Proceedings

Functional Imaging and Modeling of the Heart: Second International Workshop, FIMH 2003, Lyon, France, June 2003, Proceedings

Information Processing in Sensor Networks: Second International Workshop, IPSN 2003, Palo Alto, CA, USA, April 22-23, 2003, Proceedings

Новый Мир ( № 6 2003)

Databases in Networked Information Systems: Third International Workshop, DNIS 2003, Aizu, Japan, September 22-24, 2003, Proceedings

Intelligent Virtual Agents: 4th International Workshop, IVA 2003, Kloster Irsee, Germany, September 15-17, 2003, Proceedings

Multiple Classifier Systems: 4th International Workshop, MCS 2003, Guilford, UK, June 11-13, 2003, Proceedings

Formal Modeling and Analysis of Timed Systems: First International Workshop, FORMATS 2003, Marseille, France, September 6-7, 2003, Revised Papers

Noncommutative geometry and physics : proceedings of the COE International Workshop, Yokohama, Japan, 26-28 February, 1-3 March, 2004

Principles and Practice of Semantic Web Reasoning: International Workshop, PPSWR 2003, Mumbai, India, December 8, 2003, Proceedings

Computational Processing of the Portuguese Language: 6th International Workshop, PROPOR 2003, Faro, Portugal, June 26-27, 2003. Proceedings

Thin films of soft matter

International Investment Perspectives 2003

Physics and Technology of Thin Films IWTF 2003: Proceedings of the International Workshop, Tehran, Iran 22 February - 6 March 2003

Finite Groups 2003: Proceedings of the Gainesville Conference on Finite Groups, March 6-12, 2003

php|architect (February 2003)

Scientific american (March 2003)

Computational linguistics and intelligent text processing: 4th international conference, CICLing 2003, Mexico City, Mexico, February 16-22, 2003: proceedings

Finite Groups 2003: Proceedings Of The Gainesville Conference On Finite Groups, March 6 - 12, 2003 (De Gruyter Proceedings in Mathematics)

Computational Linguistics and Intelligent Text Processing: 4th International Conference, CICLing 2003, Mexico City, Mexico, February 16-22, 2003. Proceedings

Tinear(2003-07-22)

php|architect (March 2003)

Experimental and Efficient Algorithms: Second International Workshop, WEA 2003, Ascona, Switzerland, May 26-28, 2003, Proceedings

Digital Watermarking: Second International Workshop, IWDW 2003, Seoul, Korea, October 20-22, 2003, Revised Papers

Preparation of Thin Films

Computational Methods in Sciences and Engineering 2003: Proceedings of the International Conference (Iccmse 2003)

Algorithms in Bioinformatics: Third International Workshop, WABI 2003, Budapest, Hungary, September 15-20, 2003, Proceedings

Cooperative Information Agents VII: 7th International Workshop, CIA 2003, Helsinki, Finland, August 27-29, 2003, Proceedings

Waveguide Spectroscopy of Thin Films, Volume 33 (Thin Films and Nanostructures) (Thin Films and Nanostructures)

Asian and Pacific Coasts 2003: Proceedings of the 2nd International Conference, Makuhari, Japan 29 February - 4 March 2004

High Performance Computing: 5th International Symposium, ISHPC 2003, Tokyo-Odaiba, Japan, October 20-22, 2003, Proceedings

Advanced Parallel Processing Technologies: 5th International Workshop, APPT 2003, Xiamen, China, September 17-19, 2003, Proceedings

Functional Imaging and Modeling of the Heart: Second International Workshop, FIMH 2003, Lyon, France, June 2003, Proceedings

Information Processing in Sensor Networks: Second International Workshop, IPSN 2003, Palo Alto, CA, USA, April 22-23, 2003, Proceedings

Новый Мир ( № 6 2003)

Databases in Networked Information Systems: Third International Workshop, DNIS 2003, Aizu, Japan, September 22-24, 2003, Proceedings

Intelligent Virtual Agents: 4th International Workshop, IVA 2003, Kloster Irsee, Germany, September 15-17, 2003, Proceedings

Multiple Classifier Systems: 4th International Workshop, MCS 2003, Guilford, UK, June 11-13, 2003, Proceedings

Formal Modeling and Analysis of Timed Systems: First International Workshop, FORMATS 2003, Marseille, France, September 6-7, 2003, Revised Papers

Noncommutative geometry and physics : proceedings of the COE International Workshop, Yokohama, Japan, 26-28 February, 1-3 March, 2004

Principles and Practice of Semantic Web Reasoning: International Workshop, PPSWR 2003, Mumbai, India, December 8, 2003, Proceedings

Computational Processing of the Portuguese Language: 6th International Workshop, PROPOR 2003, Faro, Portugal, June 26-27, 2003. Proceedings

Thin films of soft matter

International Investment Perspectives 2003

Recommend Documents