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

Physics and Technology of Thin Films

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Proceedings of the International Workshop on

Physics and Technology of Thin Films I W T F 2 0 0 3

Tehran, Iran 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

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THE PHYSICS AND TECHNOLOGY OF THIN FILMS

Copyright © 2004 by World Scientific Publishing Co. Pte. Ltd.

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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

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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 F. Falahi M. Alempour M. Kianpisheh A. Azarm S. Mohajer R. Azimirad F. Nasiripour E. Bagheri K. Ourami M. Bashlideh P. Rajai H. Bayat P. Sangpour L. Chahoshizadeh S. Sanjabi M. Dashti F. Shahbandi M. Ebadi

.2 e B

en

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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 nano-structural 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,

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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 3 Inaugural Address by: S. Sohrabpour 5 Closing Address by: A.Z. Moshfegh 6 Panel Discussion 7

II. DEPOSITION PROCESSES

Vacuum Technology: Principles and Applications (Invited) 11 A. Z. Moshfegh

PVD Growth Method: Physics and Technology (Invited) 28 A. Z. Moshfegh

Introduction to Semiconductor Epitaxy (Invited) 54 H. von Kdnel

Semiconductor Superlattices (Invited) 70 H. von Kdnel

Oxide Thin Film Growth on Silicon Carbide Surfaces (Invited) 85 P. G. Soukiassian

Al-W Amorphous Thin Films (Invited) 101 N. Radic, T. Car, A.Tonejc, J. Ivkov, M. Stubiear and M. Metikos-Hukovic

Heat and Mass Transfer During ZnSe CVD Deposition Process 119 V. G. Minkina

III. CHARACTERIZATION TECHNIQUES

Thin Films Analysis Using Photoelectron Spectroscopy (Invited) 129 S. K. Kulkarni

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Passivation Investigations of GaAs (100) Surface 149 R. Purandare, B. A. Kuruvilla, S. M. Chaudhari, D. M. Phase and S. K. Kulkarni

Correlation Between Microscopic and Macroscopic Properties of 158 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

IV. SURFACE PROCESSES

Reliability and Failure of Electronic Materials and Devices (Invited) 171 M. Ohring

Diffusion in Multilayers 180 S. Luby, E. Majkova and A. Luches

Dynamics of Interacting Adatoms on Complex Surfaces 187 Z Chvoj

Copper Surface Segregation During V205 Thin Film Deposition 198 M. M. Ahadian, A. Iraji-zad, M. Ghoranneviss and M. Hantizadeh

The Preparation and Surface Studies of Fe/Pt Thin Films 205 G. Varghese

V. NANOMATERIALS

lD-Nanostructures on Silicon Carbide Thin Films (Invited) 213 P. G. Soukiassian

Giant Magnetoresistance in Electrodesposited Nanogranular Thin Films 228 (Invited)

S. C. Kashyap

Self-Assembled Quantum Dots: Structural and Optical Properties, 244 and Device Applications

M. Henini

XIII

Preparation and Characterization of Ultrathin Films and Film 256 Coatings for Microelectronics

Y. A. Pogoryelov

Nanocrystalline Films in the Ag-Ni System 265 /. K. Bdikin, G. K. Strukova, D. V. Matveev, S. A. Zver'kov, V. V. Kedrov and G. V. Strukov

Fragmentation of Positively Charged Metal Clusters in Stabilized 271 Jellium Model With Self-Compression

M. Payami

VI. OPTICAL MATERIALS

Organic Films for Optoelectronic Applications (Invited) 285 X. Liu, T. Michely and M. Wuttig

Development of Highly Reactive Photo-Catalytic Ti02 Films (Invited) 297 S. H. Mohamed, R. Drese, M. M. Wakkad and M. Wuttig

Multilayer Thin-Film Optical Filters: Design, Fabrication, 306 and Applications (Invited)

S. H. Keshmiri and M. M. Mirsalehi

Thin Films for Optical Recording 318 A. Kikineshi

Diffusion of Atomic Hydrogen and Passivation of Structural Defects 324 in Silicon and in Transparent-Conducting Thin Films (Invited)

S. H. Keshmiri

The Effect of Particle Size on Optical Properties of CdS Films 337 Formed by Photochemical Technique

S. M. Mahdavi, A. Iraji-zad, F. Razi and M. Rezaesmaeili

Enhancement in Physical Properties of ZnO Transparent Conducting 344 Coating by Al Incorporation

B. N. Pawar, S. R. Jadkar, K. C. Mohite andM. G. Takwale

XIV

Optical Energy Gap of Magnetically Confined Arc Discharge 355 D.C. Sputtered Hydrogenated Amorphous Silicon

M. C. Abdulrida, H. A. Hamed and B. A. Hassan

Photocatalytic Study of Ti02 Thin Films Deposited by DC Reactive 363 Magnetron Sputtering and Spray Pyrolysis Methods

A. I. Martinez, D. Acosta and A. Lopez

Novel Transparent and Highly Conductive ZnO-Based Coatings 374 B. M. Ataev, A. M. Bagamadova, I. K. Kamilov, V. V. Mamedov, A. K. Omaev and S. Sh. Makhmudov

Low-Temperature CVD Growth of ZnO Films Stimulated 380 by RF-Discharge Plasma

B. M. Ataev, A. M. Bagamadova, I. K. Kamilov, V. V. Mamedov, A. K. Omaev and S. Sh. Makhmudov

VII. SUPERCONDUCTIVITY

Physics and Applications of YBa2Cu307/Lao.7Sro.3Mn03 387 Heterostructures (Invited)

/. G. Lin

Domain Structure of YBa2Cu3Ox Films on NdGa03 Substrates 405 /. K. Bdikin, P. B. Mozhaev, G. A. Ovsyannikov, P. V. Komissinski and I. M. Kotelyanskii

Raman Active Apical Oxygen Modes in Cu1.xTlxBa2Ca3Cu4012_5 414 Superconductor Thin Films

N. A. Khan and H. Ihara

Generation and Amplification of Electromagnetic Radiation by 420 Superconducting Films — a Superconductor Maser

A. N. Lykov

Fabrication of YBCO and BSCCO Thin Films 428 H. Salamati, P. Kameli and M. Akhavan

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VIII. MAGNETIC THIN FILMS

Some Aspects in Thin Film Magnetism (Invited) 441 M. Fade

Microscopic Mechanisms of Magnetooptical Activity in Epitaxial 454 Garnet Films (Invited)

A. I. Popov

Design, Fabrication and Applications of Multilayer Thin-Film 469 SQUID Sensors (Invited)

D. Rassi and Y. E. Zhuravlev

Deposition and Characterization of Fe/Si Multilayers 479 S. Kharrazi, S. Ashtaputre, S. Kulkarni, R. Choudhary, S. Shinde andS. Ogale

Surface Nonlinear Magneto-Optical Effect in Antiferromagnetics 489 A. K. Zvezdin, A. R Pyatakov, V. I. Belotelov and V. A. Kotov

Fabrication and Characterization of the Co/Cu/Co/NiO/Si( 100) 499 Magnetic Multilayer

A. Z Moshfegh, P. Sangpour, O. Akhavan, G. Kavei and A. Iraji-zad

Optimisation of Thin Film Multi-Layers by Micromagnetic Simulations 507 for MR Applications

P. Gornert, D. V. Berkov and N. L. Gorn

Conference Photos 515

Author Index 529

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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.

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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.

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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

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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 thin films in Prague, Czech Republic in 2006.

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II. DEPOSITION PROCESSES

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VACUUM TECHNOLOGY: PRINCIPLES AND APPLICATIONS

A.Z. MOSHFEGH

Department of Physics, Sharif University of Technology,

P.O. Box 11365-9161, Tehran, Iran

E-mail: moshfegh@sharif.edu

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,

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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

Low Vacuum

Medium Vacuum

High Vacuum

Very High Vacuum

Ultra High Vacuum (UHV)

Extreme High Vacuum (XHV)

Pressure Range (torr)

~10 3 -1

1-10"3

10"3 - 10"6

10"6-10"9

10"9-10"12

<10"12

2. Properties of Vacuum

Vacuum environment possesses different properties. Some of the important properties are including: lowers materials melting point, reduces the gas impingement rate, increases particles mean free path and as a result increases time of monolayer (ML) formation.

Concerning the role of vacuum in the film deposition, it reduces the vapor pressure thus lowering the evaporation temperature of materials and provides the ultimate clean environment resulting improvement in purity of the deposited film. There exist many reasons for the deposition of materials in a vacuum environment. Vacuum conditions increase the mean free path for atoms; eliminate the presence of gases that could react with the deposited materials. For example, Al-based intermetallic alloys are good candidate for high temperature

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application. But, these classes of alloys are brittle in air at room temperature. However, according to mechanical testing studies, they are quite ductile under high vacuum conditions. It is believed that the chemical interaction between the moisture in air and the alloy surface is the main cause for the brittleness of this intermetallic compound.

3. Pumping Vacuum System

It is generally known that there are two broad vacuum pumping mechanisms: gas-transfer and entrapment. Gas transfer pumps remove gas molecules from the pumped volume and transfer them to the ambient via one or several stages of compression. This type of pumps may be subdivided into positive displacement and kinetic vacuum pumps. Rotary mechanical and turbo molecular pumps are examples of former and latter division, respectively. Positive displacement pumps operate based on this principle that repeated volume of gas are transferred from inlet to the outlet usually with some compression. Entrapment pumps are those, which retain molecules by sorption or condensation on internal surfaces. Sputter ion and cryogenic pumps are two typical examples of this type of pump.

Selection of a pump depends on required background pressure and working conditions for a process. Rapid increases in pumping speed and decreases in the base pressure of commercial vacuum systems resulted in increasing reproducibility of processes.

Gases flow in response to pressure differences. The quantity that measures ease of flow is called conductance defined by

C= Q/(P,-P2) (2) where (Pi-P2) is the pressure difference between two ends of a tube and Q is a

throughput which is a quantity related to particle flow rate (dN/dt). Based on the correlation between system dimension or diameter of a tube (D) and mean free path of particles (X), there are three different flow regimes: a) viscous (A.«D), b) intermediate (k~D) and c) molecular (X»D).

At low pressures, flow is characterized by the molecular region. This region can be defined in terms of the dimensionless ratio called the Knudsen number Kn=A/D. At high pressures, the flow is determined by the Reynold's number defined by a dimensionless quantity as given in the following expression:

1 in this equation p is density of fluid, v represents flow velocity (volumetric flow rate/ cross-sectional area) and n is viscosity of fluid. The viscous flow

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becomes turbulent when R,.>2100 and it is entirely laminar when Re<1100 [1] and it is mixed when Re takes a value between these two limits. Figure 1 shows how the conduction of a pipe changes from viscous to molecular flow through a transition (intermediate) region as pressure is reduced.

The Reynold's number can be related indirectly to the throughput (Q) defined as the quantity of gas flowing a pipe expressed in PV (pressurexvolume) units per unit time (see below). Therefore, we have Q=PV(JTD2/4). It is clear that for the larger tube diameter the higher throughput is achieved.

For the viscous laminar flow, the total flow (J) of a gas in particles per second can be obtained by the well-known Poisseulle formula given below.

na AP J = ( )pav —

SrjkT L (4)

where a is the tube radius, L is tube length, AP is pressure difference between its two ends (assuming AP is very small), k is Boltzman constant, T represents temperature and Pav is average pressure in the tube .

o o c re o 3 •a c o o

Viscous,, f low.

molecular

Transition region

Pressure

Figure 1. Dependence of conductance on pressure in various flow regimes.

If we focus on large mean free path (molecular flow), the pump down in this region is limited by equilibrium between the gas entering the system and pumping speed itself. The gas load Q0 of the pump consists of two parts. First, the leakage into the pump and then the back streaming of the pumping fluid. If the S be theoretical pumping speed of the pump and P represents pressure at the inlet to the pump, then the throughput is defined by Q = SP-Q0. Thus, the lowest pressure of the pump can be obtained if Q=0 resulted in Q0= SP. In addition to

15

the back-streaming contribution to the pressure increase in a vacuum chamber, there are several mechanisms that they are responsible for this pressure rise. Figure 2 shows all these contributions. Therefore, total load (throughput) for any vacuum system can be defined by the following equation:

Q=Qp+Qv+Qo+Qi (5) where QP,QV,Q0 and Q, represent load due to dimension of pipe, materials vapor, outgassing from chamber wall, and leak throught the system. It is normally observed that the quality of vacuum required dictates the connections used. The gases in a vacuum system result from leaks in the chamber evaporation of moisture and other materials from the walls, and outgassing from the materials that makes up the vessel. The rate of outgassing from a surface can be defined as:

Qr = qA (6) in here, A is surface area and q represents the specific outlasting rate for the material which is unique for a given materials [see for example 2]. A well-written review article on calibration and the use of leaks with some recommendations of standardization of connections and safety is given in reference [5].

L , 1 r * 1

Vat V " ior

^eak

J i

Process

Outgassin &

Pipe

Pump

Figure 2. Various sources of pressure rise in a vessel and the pumping process.

To analyze a large and complex high vacuum system with outgassing consideration, a numerical modeling of vacuum system was applied based on an electronic circuit. For example, vacuum variable and electrical variable are volume (V) and capacitance (C); pressure (P) and voltage (v); throughput (Q) and current (i); NkT and charge (q), respectively. The model was suggested to obtain the response of a system to transient and time-varying loads, and to analyze vacuum systems such as accelerators, storage rings and process lines [6].

16

In achieving high vacuum environment, one can examine the conservation of mass equation for a vacuum system of volume V pumped by a pumping speed S with a leak rate Q|. The relationship between these parameters is described by the following expression:

dP S ,„ _ Q, = -(P-PU) + ^L

dt V V (7)

where P is instantaneous pressure and Pu is ultimate pressure of the vacuum system that is attainable by the pump . Practically, the lowest possible pressure is about P ~ 10'12 Torr (lTorr =133.3 Pa). This pressure corresponds to particle density of about 105 particles/cm3. The pump down time may be calculated using this equation (assuming no leak) which reduces to P=P0exp[-(S/V)t] where P0 is initial system pressure at t=0. The pressure reduction by this equation (~e ) is attributed to chamber volume. Other evacuation processes during pump down time is illustrated in Figure 3. It is clear that the rate-limiting processes, which determine the ultimate vacuum attainable is different for each, time scale. According to this figure, once the volume removal of gas is completed, the vacuum is controlled by gas desorbing mechanism from the chamber surface, which shows a ~ t"1 dependency.

Volume

101 103 105 107 109 1011 1013 1015 10'7

Time (Sec)

Figure 3. Log-Log plot of pressure versus time for a typical unbaked vacuum system (from [7]).

17

Since pumping speed is determined at the intake cross section of the pump, the pump intake should be as close to the main vacuum chamber as possible. Therefore, to optimize pumping, interconnecting tubes and valves should be as large as possible resulting a maximum conductance.

An evacuated chamber can be used for a process involve evaporation of particles from a heated source in a special environment for a certain experiment. For example, a beam of Ag particles is sent into nitrogen or air. The amount of Ag particle deposited by the beam in a time t on a surface located at a distant d from the source can be computed. If I0be the intensity of the beam at the source, then the intensity at the collector is:

/ = / . exp(-d/L) (8) L is the mean free path of Ag atoms in the gas at the pressure in the collecting chamber. Therefore, the intensity at the collector is reduced as d-value increased.

To summarize this section, knowledge on gas flow, leaks, design and construction of vacuum systems as well as types of pumps and their selection is important in optimizing quality of vacuum system for a desired process. A complete description on vacuum properties, types of pumps, vacuum productions and measurements as well as vacuum applications is discussed in some important reviews [1-2,7-11].

4. Gas-Surface Interaction

One of the important concepts in vacuum technology is the energetic interaction between gas molecules and solid surfaces. Considering gas-surface heat exchange, when an atom impinges onto a surface from the adjacent gas phase an exchange of energy may occur. For instance, if the atom comes from a gas with temperature T; and if the surface be at temperature Ts where Ts>Ti then, the atom usually gains some kinetic energy corresponding to its new equivalent temperature Tr such that Tj<Tr<Ts. The correlation between these temperatures is linked by the accommodation coefficient defined by:

a=(Tr-Ti)/(Ts-Ti) ( 9 )

On average, the magnitude of this coefficient is in a range 0<a<l and mainly depends on Ts and Tj.

When the gas molecules interact with a solid surface, several phenomena can occur. Among them adsorption, back-scattering, desorption, diffusion, molecular dissociation, surface reactions and product formation. Figure 4 shows the schematic representation of some of these phenomena. A further description of these processes has been explained elsewhere [12,13].

18

Chemical nature, crystal structure and physical characteristic of the surface are important parameters for the extent of these processes. Kinetically, the adsorption process can be described by the following general equation:

S = PlP2P3v (10)

where S is the rate of adsorption, Pi is the probability that the particle collides with the surface finds a free site, P2 is the probability that particle has enough energy for the adsorption to take place and P3 is the probability that when the two above conditions are fulfilled, the particle is actually adsorbed and V is the collision frequency of the particle on the surface, the product quantity S=P1P2P3

is called sticking probability or sticking coefficient. The magnitude of S depends on the surface coverage, the size of the adsorbate, the dissociative (nondissociative) character of the adsorption and so on.

C~*)-(~} REAGENT ^^^ ^^^ MOLECULE

o-o BACK-SCATTERING REMOVAL OF

NONREACTIVE > RESIDUE

SURFACE DIFFUSION

NONDISSOCIATIVE CHEMBORPTION

1

PHYSICAL DISSOCIATIVE ADSORPTION CHEMBORPTION

2 3 Figure 4. Schematic representation of major surface processes.

DESORPTION 5

(K)-CK)-0-0 O-FORMATION OF

PRODUCT MOLECULE 4

One of the major applications of UHV is studying of adsorption/desorption phenomena. These surface processes at different conditions are shown in Figure 5. In Fig. 5a), the surface is covered with single, isolated gas particles. This state corresponds to an extremely ultra high vacuum (UHV) that occurs at pressure of less than 10"9 Torr. In Fig 5b), the surface is covered with a 1ML (one ML is described previously in section 1). This state corresponds to a pressure of 10"6

Torr and finally, in Fig 5c), the surface is covered with multilayer adsorption. This last state is given at all pressures higher than 10"6 Torr. At normal high vacuum (~10"5 Torr) on the surface, there exist 20-30 layers and at atmospheric pressure, there are up to 150 layers.

19

a) b) c)

Figure 5. Three states of adsorption a) adsorbed single particles, b) adsorbed monolayer and c) adsorbed multilayer.

Considering the adsorption state where the gas particles adsorbed on solid surface, they will be desorbed if the energy is high enough to overcome the barrier. However, they will desorb after a certain period of time. This description will be explained on the basis of residence time later.

There are two types of adsorptions, physical adsorption (Physisorption) and chemical adsorption (chemisorptions). Physical adsorption involves only attractive forces that are the van der Waals or dispersion type in nature. It occurs at or below the boiling temperature of the adsorbate. Their adsorption energy values are in the range 0.1-5 kcal/mole (leV=23.1 kcal/mole). In the other words, in physisorption, there is weak interaction energy between the adsorbed gas and a solid surface (e. g. Helium adsorption on many transition metals). Chemisorptions involves systems in which hydrogen bonding, covalent chemical bonding or metallic bonding. The interaction energy in this type of adsorption is very strong and is in a range of 5-150 kcal/mole. Another distinction between physisorption and chemisorption states, the equilibrium distance between the surface and an adsorbed atom or molecule are shorter in chemisorption (-1-3 A). However, for physisorption is about ~4 A. In addition, chemisorption is highly specific to the nature of the adsorbent whereas for the case of physisorption is not. Figure 6 shows potential energy diagram model for both types of adsorption as a function of distance between the adsorbate and a surface. In this figure, EP and Ec are heats of adsorption for physisorption and chemisorption states, respectively, and Ea is an effective energy barrier that

20

controls the rate of the transition from the physisorbed to chemisorbed state [14]. From kinetic view point, adsorption process can be described by several models including Henry, Langmuire and BET. A complete discussion on these models is given in [15]. In addition, for a review on microscopic approach to physisorption from theoretical and experimental views, one can use ref. [16].

> a ec in z UJ - I < z UJ H o a

CHEMISORBED MOLECULE

IC

DISTANCE FROM SURFACE (r)

Figure 6. Physisorption (curve P) and chemisorption (curve C) processes.

Concerning chemisorption process, one of the best understanding systems is CO chemisorptions on metals of group VIII and lb. It is generally accepted that the CO molecule is attached to the metal surface via the carbon atom. Observations of frequency shifts in C-0 stretching vibration of the adsorbed molecule and their analysis lead to a belief in the existence of two types of adsorbed complexes, " Linear " and " bridged" by analogy with the formation of linear and bridged metal carbonyls. Chemisorption and reactions of several diatomic molecules (CO, NO, H2 and 02) on group VIII metals surface have been investigated previously [17]. Further discussion on chemisorptions and reactivity of metals is reported in [18]. For the case of CO chemisorptions, which is an important step in many heterogeneous catalytic reactions, Yates [19] has investigated the nature of chemisorbed CO the Pt (111) surface using scanning tunneling microscopy (STM) observation.

On the concept of desorption, this process can be occurred by several methods including thermally, known as thermal desorption spectroscopy (TDS), electron stimulated desorption (ESD), photon stimulated desorption (PSD) or ion impact desorption (IID). The residence time (T) of a species stays on a surface before it desorbs mainly is dependent on substrate temperature. This quantity defined by:

T = T„ exp (AH/RT) (11)

21

where ro is inverse of vibrational frequency of surface atoms (~10"13 sec), AH is heat of desorption (desorption energy), R is universal gas constant and T is temperature in Kelvin. It is obvious that surface specie has a shorter residence time at higher temperatures.

In addition to adsorption phenomenon on transition metals as discussed above, investigation on adsorption/desorption properties of transition metal oxides is also important due to their unique surface properties that can be used in various technological applications specially for heterogeneous catalytic reactions. For example, H20 desorption from V205 surface is measured at temperature of TM=390 K for a surface coverage of 9H2o=0.20 ML, using temperature programmed desorption (TPD) spectroscopy [20]. This type of knowledge on surface chemical bonding was used to identify the rate determining step (RDS) for the catalytic decomposition reaction of isopropyl alcohol over V205 surface. More recently, adsorption and photocatalytic property of M0O3 (010) surface during direct conversion of methane to methanol is also investigated using computer simulation approach [21].

One of the important quantities that controls kinetics of the surface processes is desorption rate. The maximum desorption rate from unit surface area at temperature T can be written as:

N(t) = -— = k„a" exp(-£/RT) (12) dt

where kn is the rate constant, ex is the surface coverage (number of molecules/cm2) and E is the activation energy of desorption (kcal/mole). Further knowledge on adsorption/desorption [12,15,16] and kinetic modeling of surface rate processes [22,23] are well described in given literatures.

From atomic and molecular point of view, surface reactions can be categorized into two different types: a) Langmuir-Hinshelwood mechanism and b) Eley-Rideal mechanism. In the first type, reactions occur between adsorbed species or between an adsorbed species and a vacant site forming adsorbed product(s) on the surface. For the second type, the direct interaction of a gas-phase species with an adsorbed species takes place resulted in formation of a product which may either remain adsorbed on the surface or desorb into the gas phase. Several well-written and comprehensive literatures on chemisorption, diffusion and surface reactions are reported elsewhere [15,18,24-26].

5. Vacuum Applications

The degree or quality of vacuum required depends on the application. The properties and function of vacuum are routinely applied in various technological and industrial fields. Some of the important applications of vacuum are

22

including: thin film deposition, surface science, microfabrication, space simulation, particle accelerators, vacuum impregnation, lifting process, food packaging. Table 2 lists some other important applications of vacuum technology in different industries. Due to importance and the key role of surface science and microfabrication process in many advanced technologies, only these important fields as well as vacuum applications in particle accelerators and analytical techniques will be described briefly below.

5.1. Surface Science Studies

The field of ultra high vacuum is mainly used for investigating surface properties of solid materials. For a solid surface, one must determine chemical identity of the atoms present, the geometrical arrangement of these atoms and the distribution of electrons surrounding these atoms both in energy and space.

Table 2. Various Applications of Vacuum Technology.

Industry

Thin Films

Surface Science Studies

Electrical/Electronics, microelectronic and solar

Technology Aeronautics and space

exploration

Particle Accelerators

Metallurgical

Biological/Medical

Chemical, Pharmaceutical and food

Packaging

Application examples

Electrical and optical coating

UHV surface analytical techniques including, AES, XPS, UPS, SIMS. Production of Vacuum electronic

components, semiconductor components, lamps and solar cells

Development of nano structure

Storage rings

Melting, casting, annealing electron beam welding, milling and tool

manufacture STM and AFM observation of

DNA, RNA Mixing of materials, evaporation of

materials Food products, carton erectors,

drum labeling

In the other words, surface science deals with the relationship between the chemical composition and structure of surfaces as well as transition regions between phases and their properties (electronic, chemical, mechanical...).

The development of electronics and computer industries as well as rise of aerospace technology were connected to miniaturization of devices. Thus, a

23

surface with on atomic layer thick can play a key function in some surface processes. As technology evolves toward the use of systems with a large surface (A) to volume (V) ratio (A/V), knowledge and detail understanding of surface structure and its chemical composition become more essential in maximizing the yield for a surface process. Some of the important examples of evolutionary systems with high A/V ratio are including integrated circuits (IC) and catalysts [27]. Further information on the role of ever-increasing A/V ratio can be found elsewhere [22,28].

An important topic in fundamental surface research, is investigating the nature of the bond between an adsorbate and a surface. Surface science focuses on first few atomic layer (-10 A) of solid. A clean surface is defined when the impurity concentration on the surface is below the detection limit of current chemical analysis techniques (0.1-1%ML). This corresponds to 1012-1013

atoms/cm2. The above characterization of the static surface is the first step in

understanding the dynamics of the interaction of surface atoms with external systems such as atoms, molecules, photons or other surfaces. Thus, knowledge and determination of surface properties of materials will lead to development of new applications.

Since the presence of impurities or preadsorbed gases on a surface will notably change subsequent adsorption properties and behavior thus, the first step in a study of the adsorption of gases or vapors on a solid is the preparation of a clean surface. There are several methods, which commonly use to obtain clean surfaces including: inert gas (i.e. Ar) ion sputtering, chemical reactions, high temperature treatment, thin film deposition and cleavage in UHV environment. These generated surfaces can be applied in different fields including microelectronic devices, catalytic reactions and corrosion study. Recently, Goodman [29] carried out UHV spectroscopy of surface species under realistic reaction conditions (including pressure and temperature) by using well-defined model catalysts (single crystals and thin films) with known surface structure and composition. In addition, Sinfelt [30] have used surface science techniques for understanding kinetics and mechanism of heterogeneous catalytic reactions on solid surfaces. History, growth, trends and applications of UHV surface spectroscopic analytical techniques are well described recently in [31-34].

5.2. Microfabrication Process

The integration of micromechanics and microelectronics permits the production of more sophisticated micro devices such as micro spectrometers. In this context, vacuum plays a significant role in some steps of micro fabrication

24

technologies. These are complex and multidisciplinary such as the LIGA (German acronym for lithography, electroforming and molding). The LIGA can be also used for mass scale production. In addition, the deep x-ray lithography with synchrotron radiation can be applied in production of microelectromechanical systems or known as MEMS technology [35,36].

Concerning application of physical vapor deposition to semiconductor microfabrication processes, the vacuum requirement is one of the key issues to consider for a process with free contamination. Hashim et al. [37] have found the major sources of impurities in sputtered Al-alloy films for interconnect prepared by PVD method. They have observed that the origin of these impurities is from residual gases (H20, N2, 02, CH4) present in the vacuum system and from target materials.

5.3. Particle Accelerators

Generally, vacuum technology is used for particle accelerators to keep the particles on the desired trajectories. UHV conditions are mandatory only for storage rings for which beam lifetime in excess of 10 h is usually required. In addition, low pressures XHV are routinely obtained inside vacuum system to limit both beam neutralization and the beam induced pressure rise. The most recent development on this issue is achieved by CERN scientists and engineers where they utilized XHV technology for particle accelerators [38]. In addition, a complete description and evolution of particle accelerator as well as storage rings based on UHV development is reviewed very recently by Dylla [39].

5.4. Analytical Techniques

In vacuum environment, (as described in sections 2 and 3) mean free path of particles ( X) increases with decreasing pressure (P) of a system (A-l /P) . This property can be used in many analytical techniques to avoid interaction of probing electrons or photons with residual gases in the vacuum system. Therefore, vacuum technology must be used for surface characterization methods including AES1', XPS2), LEED3), RBS4) as well as some microscopic techniques including AFM5), TEM6), and SEM7). The trends and development of AES and XPS is described recently by Powell [40].

Auger Electron Spectroscopy ' Atomic Force Microscopy X-ray Photoelectron Spectroscopy Transmission Electron Microscopy Low Energy Electron Diffraction 7) Scanning Electron Microscopy Rutherford Back-scattering Spectrometry

25

Vacuum environment can be also applied for detail understanding of a special investigation. Recently, Nakamoto et al. [41] have observed freeze fractured red blood cell with high vacuum-low temperature using atomic force microscopy (AFM) integrated vacuum system to study the mechanism of the fracture. Table 3 lists some of the surface analytical techniques in the monolayer range.

Table 3. Procedures for surface analysis in the monolayer range.

^\Stimulation

Emission ^ ^ .

Electrons

Photons

Ions

Neutrals

Electrons

AES

EID, ESDIAD

Photons

XPS, UPS

Ions

SIMS

Heat

TDS

EID: Electron Induced Desorption ESDIAD: Electron Stimulated Desorption Ion Angular Distribution TDS: Thermal Desorption Spectroscopy SIMS: Secondary Ion Mass Spectrometry UPS: Ultraviolet Photoelectron Spectroscopy

Concerning the evolution and future aspects of vacuum science and technology, the routine production and measurement of vacuum at much lower pressures range (10"13-10"12 Torr) can be achieved in near future. It is expected that thin walled UHV stainless steel vacuum system will become widespread that UHV/XHV systems with nonevaporable getters coated on the interior wall will be used for special purposes [3].

6. Conclusions

In this article, some important and basic concepts of vacuum are introduced. Principles, divisions and properties of vacuum are reviewed. Different pumping mechanisms of a vacuum system are briefly discussed. Properties of three gas flow regimes namely viscous, transition and molecular type flow based on Reynold's number and Knudsen's number are described. A general equation for pump down of a vacuum system is presented based on pressure-time relationship. The fundamental of gas-surface interactions in particular two

26

adsorption states, physisorption and chemisorption is reviewed. Other surface processes including dissociation, surface chemical reactions, product formation and desorption phenomena are discussed. The product formation mechanisms are interpreted based on Langmuir-Hinshelwood and Eley-Rideal type surface reactions. At the end, different applications of vacuum technology in various industries with some examples are introduced.

Acknowledgements

The author is especially grateful to all his graduate students. He also wants to thank the Research Council of Sharif University of Technology and High Technology Organization (Ministry of Industries and Mines) for financial support. The assistance of Mr. Azimirad, for preparing some of the figures and Ms. M. Bashlideh and Ms. L. Chahooshi for typing the manuscript is greatly acknowledged.

References

1. A. Roth, Vacuum Technology (Elsevier, New York, 1990). 2. G.M. Lafferty, Foundations of Vacuum Science and Technology (John

Wiley & Sons, New York, 1998). 3. J.P. Hobson, J. Vac. Sci. Technol. A 21, S7 (2003). 4. P. A. Redhead, J. Vac. Sci. Technol. A 21, S12 (2003). 5. CD. Ehrlich, J. A. Basford, J. Vac. Sci. Technol. A 10, 1 (1992). 6. S.R. Wilson, J. Vac. Sci. Technol. A 5, 2479 (1987). 7. T.A. Delchar, Vacuum Physics and Techniques, (Chapman & Hall, London,

1993) p 168. 8. J. M. Lafferty, "Vacuum From Art to Exact Science", Physics Today 34,

211(1981). 9. H.G. Tompkins, An Introduction to the Fundamentals of Vacuum

Technology (American Vacuum Society publishing, 1984). 10. J.P. Hobson, J. Vac. Sci. Technol. A 2, 144 (1984). 11. D.M. Hoffman, B. Singh, J.H. Thomas, Handbook of Vacuum Technology

(Academic Press, Boston, 1997). 12. A. Zangwill, Physics at Surfaces (Cambridge University Press, Cambridge,

1988). 13. K. Reichelt, Vacuum 38, 1083 (1988). 14. M. Ohring, Material Science of Thin Films (Academic Press, New York,

second edition, 2002) p 374. 15. J.B. Hudson, Surface Science: An Introduction, (Butterworth-Heninemann,

Boston, 1992).

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16. U. Landman, G. G. Kleiman, "Microscopic Approaches to Physisorption: Theoretical and Experimental Aspects", Surface and Defect Properties of Solids, Vol. 6, (The chemical Society, London 1977) p 1.

17. B.E. Nieuwenhuys, Surface Science 126, 307 (1983). 18. B. I. Lundqvist, "Chemisorptions and Reactivity of Metals", Many-Body

Phenomena at Surfaces, eds: Langreth and Suhl, (Orlando Academic Press, 1984) p 93.

19. J.T. Yates, Jr, Surface Science 299/300, 731 (1994). 20. A.Z. Moshfegh, A. Ignatiev, Surface Science 275, L650 (1992). 21. A.Z. Moshfegh, M. Dashti, Surface Reviews & Letters 11(1) (2004). 22. H.C. Kang, W.H. Weinberg, "Kinetic Modeling of Surface Rate

Processes", Surface Science 299/300, 755 (1994). 23. H.C. Kang, W.H. Weinberg, Chem. Rev. 95, 667 (1995). 24. R.P.H. Gasser, An Introduction to Chemisorption and Catalysis (Clarendon

Press, Oxford, 1985). 25. G. A. Somorjai, Surface Science 299/300. 849 (1994). 26. F. Zaera, Surface Science 500, 947 (2002). 27. A.Z. Moshfegh, "Surface Science and its Application to Industry", The

Proceedings of the 5-th World Seminar on Heat Treatment and Surface Engineering (Isfahan, Iran, 1995) p 42.

28. G.A. Somorjai, Introduction to Surface Chemistry and Catalysis (John Wiley & Sons, New York, 1994).

29. D.W. Goodman, Journal of Catalysis 216, 213 (2003). 30. J.H. Sinfelt, Surf. Sci. 500, 923 (2002). 31. J.A. Venables, Introduction to Surface and Thin Film Processes

(Cambridge University Press, Combridge, UK, 2000) p 128. 32. P.J. Feibelman, J. Vac. Sci. Tech. A 21, 564 (2002). 33. S. Ferrer,Y. Petroff Surf. Sci. 500, 605 (2002). 34. C.B. Duke, J. Vac. Sci. Technol. A 21, S36 (2003). 35. E.W. Schmidt, J. Vac. Sci. Technol. B 16, 3526 (1998). 36. F. Mazzdini, Vacuum 65, 239 (2002). 37. I. Hashim, I.J. Raaijmakers, S.E. Park, K.B. Kim, J. Vac. Sci. Technol. A

15, 1305 (1997). 38. C. Benvenuti, IEEE Proceedings of the Particle Accelerator Conference

(Chicago, USA, 2001) p 260. 39. H.F. Dylla, J. Vac. Sci. Technol. A 21, S25 (2003). 40. C.J. Powell, J. Vac. Sci. Technol. A 21, S42 (2003). 41. K. Nakamoto, C.B. Mooney, S. Kitamura, Jeol News 37E, 62 (2002).

PVD GROWTH METHOD: PHYSICS AND TECHNOLOGY

A.Z. MOSHFEGH

Department of Physics, Sharif University of Technology P. O. Box 11365-9161, Tehran, Iran

E-mail: moshfegh@sharif.edu

In this review, the foundation of thin film technology namely fabrication, characterization and application is described. Classification of physical vapor deposition (PVD) is presented based on evaporation and sputtering methods. The physics and technology of three main branches of PVD deposition techniques including sputtering, pulse laser deposition (PLD) and molecular beam epitaxy (MBE) along with their characteristic differences are compared. The application of bias sputtering in producing thin films with modified properties is presented. A correlation between deposition variables and parameters of nucleation and growth is discussed. The initial stages of PVD growth modes such as layer by-layer, island, and mixed layer-island growth mechanisms are reviewed. At the end, the applications of PVD in microelectronics with several recent examples especially in the metallization process are presented.

1. Introduction

The progress in thin films science and technology accelerated following the development of vacuum technology for the vapor phase synthesis of layers with controlled and reproducible properties. Thin film applications rely strongly on the electrical, optical and physical properties of the materials chosen. These properties depend on the deposition methods and processes to fabricate the films. Thus, there is a relationship between film growth dynamics and specific film property, which in turn depends on deposition method.

There are several methods to deposit thin film via physical or chemical processes namely two main classes of deposition, physical vapor deposition (PVD) and chemical vapor depositions (CVD). It is well established that PVD differs from chemical vapor deposition. In CVD process, a solid material is deposited onto a heated substrate surface as a result of chemical reactions in the gas phase in a proper reactor. The deposition reaction can be of several types including pyrolysis, reduction, oxidation, compound formation, disproportionation and etc.

The notion of PVD includes all methods in which the vapor particles are ejected from a source into the vacuum by a physical process. PVD is a simple and powerful technique to grow wide variety of materials over large substrates. It is used for deposition of various materials in different technological applications. For example, in semiconductor manufacturing and device fabrication, thin films are utilized for interconnections, multilevel metallization

28

29

(MLM), diffusion barriers, adhesion or seed layers and so on. Recently, Rossnagel [1] has investigated and reviewed directional PVD sputter deposition for semiconductor manufacturing process, in particular, the use of metal-rich plasmas fed by sputtering technique known as ionized PVD (I-PVD) for filling deep features is introduced in that report. The application of PVD in semiconductor manufacturing and microelectronic industry will be discussed at the end of this report.

Generally, thin film technology is based on three foundations namely fabrication, characterization and applications. Figure 1 illustrates correlation between these three processes. It is obvious that the use of thin films is implemented for a desired application, if they are high quality and possessing special properties.

Characterization ^ w 4 ^

Applications

Figure 1. Basic operations and their relations in thin films technology.

PVD deposition technique is divided into two main classes a) evaporation and b) sputtering. Figure 2 illustrates all branches for these two techniques. In all PVD branches, phase transformation from vapor to solid is resulted in film nucleation and subsequent growth. Unlike epitaxial deposition, both evaporation and sputtering methods yield polycrystalline films with typical submicron grain size. To understand and describe PVD processes, some fundamental concepts of this technique is reviewed below.

2. Fundamentals of PVD

In PVD processes, atoms or small clusters of atoms are removed from a source (solid or liquid). These atoms are traveling in a vacuum chamber reaching a substrate to form a thin film. The rate of deposition depends mainly on system pressure, source-substrate distance and substrate temperature. Generally, PVD deposition processes contain at least four steps namely 1) a source of film material is first provided, 2) then the material is transported to the substrate with an uniform arrival rate under proper vacuum environment, 3) the deposition takes place and as time goes on, a thin layer is formed on the substrate and 4) sometimes the deposited film is subsequently annealed in a proper temperature and suitable environment.

One of the important physical parameters that controls the quality of thin films is pressure. The lower the pressure (P) is associated with the larger mean free path (k). These two parameters have a fundamental relation with each other

30

(p ~l/X). Therefore, in PVD deposition technique, the background pressure must be reduced in order to have a high purity thin film with good quality.

Physical Vapor Deposition

Evaporation

Thermal

Electron Beam

Ion Plating

Laser Ablation

Sputtering

Direct Current

Radio Frequency

Magnetron

Reactive

oias

Figure 2. Classification of PVD method.

The quality of the vacuum is very decisive and a practical consideration in determining the deposition rate, since gaseous impurities in the deposition chamber can impinge on the growing film resulted in a non-stiochiomctirc film. Therefore, complete evacuation of a deposition chamber with optimizing growth parameters can lead to high quality of thin films. In addition to the minimization of the base pressure, types of substrate are also important. To grow an epitaxial film, a single-crystalline substrate is required to provide a template. Substrates can be divided into three types: metal, semiconducting and insulating materials which each of them has its own advantages. The selection of a substrate type depends on depositing materials and applications.

The motion of particles in PVD deposition methods (see Figure 2) is different. During evaporation, molecular motion is more or less non-randomized and there is line-of-sight deposition. However, during sputtering, there is a

31

considerable randomization of travel direction (in the absence of bias sputtering) leading to better uniformity of deposition on stepped surfaces. Although evaporation and sputtering processes are physically very different, certain behavior of the gas phase species are governed by the same principles. Atomic or molecular scattering is one of the important phenomena that takes place during the process. This type of scattering occurs due to collisions with atoms or molecules in a vacuum chamber.

Due to simplicity of thermal and electron beam evaporation as well as ion plating techniques, they will not be discussed here. For further knowledge about thermal and e-beam evaporation techniques, readers are referred to excellent reviews [2-7]. In additions, various aspects of reactive ion plating specially for deposition of optical films is described in [8].

Therefore, in this manuscript, physics and technology of sputtering, pulse laser deposition (PLD) and molecular beam epitaxy (MBE) will be discussed, respectively.

3. PVD Techniques

3.1. Sputtering

Before we begin to describe the physics of sputtering process, a deep understanding of the ion-surface interaction is necessary. In this interaction, several phenomena occur depends on incident energy, mass of ion and the nature of surface atoms. Figure 3 shows some of these phenomena with their corresponding energies. It is obvious that the sputtering process takes place usually in the ion energy range of about 102-104 eV depending on the gas pressure.

Sputtering is a PVD process that involves the removal of material from a solid cathode. It is a surface phenomenon that is accomplished by bombarding the cathode with positive ions originated from a rare gas discharge. The positive inert gas ions (usually Ar+) are generated according to the following ionization process:

Ar + e" • Ar+ + e"+e" (1) The sputtering process can be performed either by gas discharge or ion beam (single or dual as a collimated source). In gas discharge sputtering process, the generated ions accelerated within a dark space (sheath) toward the target surface held at cathode side (negative potential), and the substrate for the film deposition is placed at anode side which is in opposite side of the target that is usually grounded. The accelerated ions transferring their momentum to surface target atoms leading to atomic collision cascade within the target material. This

32

Ion implantation at surface

Ion accumulation Deep ion implantation

Thermal effect

1 1 i M i 1—' r r

io-2 IO-1 io0 101 io2 io3 io4 io5 io6 io7 E(eV)

Figure 3. Ion - surface interaction and created phenomena with their energy range.

process resulted in ejection of atoms from the target surface. Figure 4 illustrates these processes schematically.

Taiget

Figure 4. Schematic representation of sputtering process consists of 1) acceleration of ion across the cathode sheath, 2) an atomic collision cascade within the target material and 3) ejection of a target atom, n is the angle of incidence and 8 is the emission angle (with permission from [9]).

33

The ejection of atoms from surface requires a threshold energy (E0). The threshold energy is the minimum energy required to remove one atom from the target surface. For sputtering to occur, the magnitude of E0 must be higher than the surface binding energy of target atoms (Eb). The threshold energy is about four times of the binding energy of atoms to the surface. Theoretically, Eb can be assumed approximately equal to the heat of sublimation of target atom, that is, between 4 - 8 eV. Therefore, the magnitude of the threshold energy for most of target atoms is in a range of 15 < E0 < 40 eV. To understand these sequential steps, computer simulation approach can be used. Motohiro [10] and Heberlein [11] have developed a Monte Carlo simulation to analyze a sputter-deposition process.

The potential of several 100 volts applied between these plates leads to the ignition of a plasma discharge for typical pressure of 10"'-10'3 Torr. The discharge is maintained as the accelerated electrons continuously produce new ions by collisions with the sputtering gas atom (i.e. Ar). To improve the ionization rate, magnetic fields can be used which force the electrons move to helical paths close to the cathode resulting a much higher ionization probability and subsequently yields a higher sputtering rate. In magnetron sputtering, the magnetic field is located parallel to the cathode surface. Secondary electrons, which are emitted from the cathode due to ion bombardment, are constrained by this magnetic field to move in a direction perpendicular to both the electric field (normal to the surface) and magnetic field.

Magnetron sputtering systems operate with an unusual voltage-current relation given below:

I=kVn (2) where I is the discharge current, V the voltage and k and n are system material and gas-dependent constants, respectively. For typical magnetron systems, the magnitude of n is between 5-10. For dc diode (non-magnetron) the value n is two or less.

The most important parameter that characterizes the sputtering process is sputtering yield (Y) defined as the number of atoms (molecules) ejected from a target surface per incident ion. The quantity Y depends on mass and energy of incident ion, type of discharge gas as well as mass and binding energy of target atom. Figure 5 shows the variation of sputtering yield for six different metallic targets under He+, Ne+, Kr+ and Xe+ normal ion bombardment (n=0) as a function of their incident energy using computer simulation method [12]. For all cases, Cu exhibits a much higher yield that is consistent with reported experimental results [13]. The accepted theory to describe the sputtering yield based on collision cascade is described by Sigmund [14,15].

34

0.5—i

1 1 1 l l l i H

V .

^ ^ ^ S - - ^

i 1 1 i i n i |

Cu

X Co \ .

^ ^ T a N*>

1 1 l l l l l j 1 1 1 Hlli

Mo

^ s

(a

Energy (keV)

(b

0 I i 1111 ii|—i i i mill—i i 111 MI | M I inn—i i i mill—i i ii ni] 18"* 10" 1 10 10* 10' 10"

Energy (keV)

C 1 0 -

I I l l l f | I I T I ' l l l l l I I I I Uli I I I I I i 11 j 1 I | | i l l l j 1 10 1 8 ' 1 0 a 10-

Energy (keV)

"1 I I I llll| I I I I llll| 1 I I I llll I I I I Illlj I I I I 1111| 1 10 10a 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 3xl05 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. cos^ n where 8<n<12 [7]. Figure 6 shows the schematic arrangement of a typical PLD deposition system. It is obvious that a high concentrated beam is focused on a target material resulting the plume formation under special conditions and geometry.

Aperture Focussing

Target Port

Laser Beam

Vacuum Chamber

Beam Focal Plane (Target)

Substrate Port Substrate Plane

Figure 6. Depiction of relative position of basic components in PLD technique.

PLD is a versatile thin film growing method. During the deposition, laser ablation uses a high-power pulse to evaporate a small area of target material. The output pulse energy of typically about one J/pulse leads to the immediate formation of plasma due to the high energy density of 3-5 J/cm2 at the target surface. UV excimer laser delivering 0.1 to 1 joule pulses of 15 to 45 ns duration time are commonly used with 1 to 100 Hz repetition rates to produce a high deposition rate of 0.1 to 100 nm/s.

The plasma contains energetic neutral atoms, ions and molecules and they reach the substrate surface with a broad energy distribution in a range 0.1-15 eV. During the deposition, the target and substrate are rotating for better film uniformity. The basic PLD process usually requires the use of a background gas during deposition. The addition of a background gas assists the interaction of the

38

plume with a reactive gas that plays an important role in producing the atomic and molecular precursors required for the growth of the compound phase [6]. During the film growth, gases such as 0 2 (for oxides) or N2 (for nitrides) are often introduced in the deposition chamber to enhance surface reactions or obtain better film stoichiometry.

It is generally known that the most nonmetallic materials that are evaporated show a strong absorption of UV radiation (200-400 nm). Absorption coefficients (a) tend to increase at the shorter wavelengths meaning reduced penetration depth. Thus, a deeper penetration depth is associated with a larger a"'. Consequently it allows one to concentrate the effects of irradiation in the process i. e. melting the target atoms very near the surface. In the other words, the usage of short wavelengths is necessary in pulse laser deposition of thin films. Therefore, UV lasers are commonly used to evaporate atoms on a target surface and as a result the stiochiometry of a deposited film is similar to the target composition during the process. The usual wavelength for laser deposition that can be used is in the range of 200-400 nm. Some of the important ultraviolet (UV) excimer lasers with their corresponding operating wavelength are including: ArF(193nm), KrF (248nm) and XeCl (308nm). A further knowledge on operation and performance limitations of the excimer lasers can be found elsewhere [28, 29].

The optical absorption length is the inverse of the optical absorption coefficient of the target sample (L0=l/a), in here; L0 is the penetration depth of the laser light. For metals and strongly absorbing semiconductor, this value is about L0 ~ 100 A for UV laser irradiation. On the other hand, the thermal diffusion length is described by:

Lt=(25tX/cnraol)1/2 (6)

where 5t is the pulse duration length, x is the thermal conductivity, c is the molar heat capacity and nmoi is the molar density of the target under investigation. It is to note that x/cnmoi is the thermal diffusivity. The depth of this heated volume is fixed by the pulse duration of the laser but not the power of the laser. The typical value of Lt is ~ 2um [9]. For most metals and semiconductors this, value is much higher than L0.

In the PLD deposition process, a tremendous amount of energy is transported from the laser to the target surface. During the laser heating process, the total loss from the target is given by:

Depth I Pulse = — (2mnk)~X'2 J f exp(-A//v / kT)T~l'2 dt (7)

39

where AHV is heat of vaporization of target atoms, nc is the number density of condensed phase, T is temperature and P° is a constant. The above equation gives the asymptotic ( '- ') result as following [30]:

Depth/Pulse ~(Patmfl/2r/Ml/2AHv)x\.53xl06 nm/Pulse (8)

in here Tis the maximum surface temperature (ideally determined experimentally) M is the molecular weight of vaporizing species, r is laser pulse length. This equation is one of the fundamental relations in the PLD process.

Some of the major physical parameters that affects the properties of PLD deposited thin films are including: 1) target-substrate distance, 2) growth rate, 3) background gas pressure and 4) substrate temperature. Zheng and kwok [31] have investigated the effect of target-substrate distance (d) on the growth rate of Indium-Tin Oxide (ITO) films. They have found that the deposition rate varied with 1/d2, which is a characteristic of a point-source evaporation in vacuum. Moreover, they have shown that the electrical property of ITO depends on the magnitude of the distance d. \

PLD can be used to grow a wide class of materials including metals, HTSC, graphite, tribological coatings, giant magnetoresistance (GMR), Ferroelectric, ferromagnetic, piezoelectric, polymers, biocompatible and other compound materials. Among these materials, the fabrication of HTSC thin films and ultra thin films materials have been investigated intensively using PLD deposition technique [32,33]. The primary advantage of PLD is its ability to produce the film composition with similar target composition. However, a major problem in PLD is droplet formation. The expulsion of droplets from the laser - induced melt on the target surface is very critical for the ablation of some materials. Mechanism and applications of laser ablation deposition is well documented in literature (see for example [30]). A detailed description on PLD equipments, processes and growth mechanism as well as applications can be found in [34,35].

3.3. Molecular Beam Epitaxy

Molecular beam epitaxy (MBE) has evolved from simple thermal evaporation method by using UHV techniques. It is a layer-by-layer atomic growth of special materials that is characterized by continuation of crystal structure from the substrate to the film. The deposited film is essentially single crystalline. The requirements of MBE growth is at least four folds a) clean room laboratory, b) UHV deposition- surface analysis chambers with attaining a vacuum level in the 10"'° Torr range. The UHV environment minimizes contamination of the growing surface. In the UHV chambers, the beam of atoms

40

and molecules travel in nearly collision-free path until arriving a substrate, c) the use of single crystal substrate and d) materials should have low vapor pressure and low sticking coefficients. Therefore, the above conditions allow the deposition of high quality films.

One of the important characteristic that must consider for a substrate is that it should possesses a minimum lattice mismatch with growing film on the substrate at the initial stage. This quantity is defined by

ctc—a TJ f (9)

af where as and af represent unstrained lattice parameters of substrate and film, respectively. For epitaxial growth this quantity tends to n—*0.

Modern MBE system is modular with separate chambers to perform individual functions such as substrate preparation, metal depositions, dielectric deposition, film characterization and so on. Figure 7 shows the schematic arrangement of a typical MBE system. This deposition method is able for independent control of beam sources and film growth, cleanliness and real-time structural and chemical characterization capability. A review article on classification and applications of metal vapor sources can be found in reference [36]. In MBE process, there are several sources with varying flux rates that can be used to grow complex compounds under control and precise conditions. It is usually used for the semiconductor heterostructure thin film materials, i.e. III-V compounds with a typical MBE growth rate in the range of 1 um/h.

Plasma Etching

MBE Compound System ]

With Mask

Multi-Sample Load Lock

( l0" 1 I Implantation I

MBE Compound System 2

Figure 7. A typical multichamber of MBE system.

41

MBE can be used to construct layer-by-layer growth of a hetrostructure material. This type of structure consists of thin stacked layers of semiconductors including GaAs, AlxGai.xAs, InP, InGaAsP, SiGe with different band gap energy. By using these types of heterustructure, it is possible to tailor materials under special experimental conditions. In the other words, one can engineer the electrical and optical properties of material by choosing their proper layer composition, thickness and doping. For example excellent quality of AlxGai.xAs was grown in a temperature range of 550-680° C as reported in [7]. Other examples for MBE growth thin films are including InSbxAsi.x and InxGa!.xAs. Therefore, by changing the composition of x, one can obtain different properties of these types of materials. This process is called " band gap engineering ". It is reported that band gap (Eg) of the AlxGa!.xAs is a function of x [37]. Moreover, the index of refraction (n) for light guiding properties in lasers also varies with x as measured earlier [38]. Thus, it is possible to fabricate ternary alloys with Eg larger than GaAs and n smaller than GaAs while maintaining an acceptable lattice matched high quality heterojunctions.

In MBE process, we consider a substrate positioned at a distance L(cm) from a source aperture of area A (cm2) with (j) =0, then the number of evaporate species striking the substrate can be obtained by the following expression:

R = 3.51 x lO 2 2 PA lnl}(MT)112 (molecules/cm2-sec) (10) In here, P is pressure in Torr and M represents molecular weight and T is source temperature (°C). The basic processes involve adsorption and surface diffusion resulting in nucleation, growth and coalescence of islands or motion of terraces. A detailed description about nucleation kinetics of metal-on-metal can be found elsewhere [39].

The advantage and a distinguishing feature of MBE is the possibility of implementing all UHV surface analysis techniques for controlling growth process and characterizing the deposited films in situ. One of the important technique for monitoring the structure and composition of epitaxial thin films during in situ growth is reflection high-energy electron diffraction (RHEED) that provides useful information about the growth mechanism of epitaxial semiconductor thin films. In the RHEED technique, a high-energy electron beam (5-30 keV) directed at a glancing angle (~2°) with respect to a single crystal held in the center of an UHV chamber. The electron energy is related

to the wavelength by the de Broglie equation x(A) = hi p = 12.25 l{V 2 ) where V is in volt. After the beam strikes the growing film, electron reflects off the surface atoms. If the top surface is periodic with a period greater than the wavelength of the beam, diffraction occurs, the diffracted beam intensity is relatively immune to thermal attenuation arising from lattice vibrations that

42

resulted in observation of RHEED oscillation during MBE growth at higher temperatures. In this technique, one period of oscillation corresponds to the growth of one monolayer on a suitable substrate.

RHEED technique is the essential surface science tool for routine MBE deposition, it is usually used to study thin film growth mechanism in situ (at the earliest stage of the growth) of the III-V of II-VI semiconductor compounds. Recently, Shimoyama et al. [40] have implemented RHEED observation to investigate stoichiometric Ba/Ti=l epitaxial growth of thin film BaTi03 on SrTi03 substrate as a function of oxygen pressure. Figure 8 shows RHEED oscillation of BaTi03 thin film system at different oxygen partial pressures during the growth of BaTi03 epitaxial thin film.

i I f s o 0) D.

500 1000

Time (sec)

Figure 8. RHEED oscillation during epitaxial growth of BaTiOj on SrTiOj substrate at various oxygen partial pressures of: a) 10"4 Pa, b) 10"6 Pa and c) 10"7 Pa (Reprinted with permission from [40]).

In analysis of RHEED pattern, in -plane lattice constant a is described by:

a = 2(h2 +k2)V2AL/A (11)

43

where A is the wavelength of incident electron, h and k are miller indices, L is the distance between the substrate and fluorescent screen and A is the distance between diffraction rods in RHEED pattern. MBE has now become an extremely useful technique to construct a well-ordered thin film and interfaces. It is able to produce wide variety of devices including transistors, optoelectronic devices (such as laser diodes) photometers, solar cells, integrated circuits and other advanced devices. Comprehensive studies on principles, physics, technology and application of MBE technique with several examples is described in recent review articles [41-45].

To summarize about PVD deposition methods, each technique contains some advantages and disadvantages and can be used for the special applications. Table 1 lists characteristic and differences exist between sputtering, PLD and MBE deposition methods. An extensive description on materials and equipments of the MBE deposition method has been reported recently [46].

Table 1. Some characteristics of vapor deposition processes [7].

Parameter

Mechanism of production of

depositing species Process Pressure

(torr) Atom removal rate

from source (atoms/cm2-sec) Average velocity

of species (cm/sec) Energy of

deposited species (eV)

Evaporation /MBE

Thermal energy

io-5-io-10

~1017

~105

Low (0.1 -0.2)

Gas incorporation None

Sputtering

Ion bombardment and momentum transfer

lO-'-lO"3

~1016

~5xl04

High (<10) higher with substrate bias

Some

PLD

Thermal energy

~10-3

~1018

~106

0.1-100

Some

4. PVD Growth Modes

The nucleation and the growth of thin films generally occur via condensation of individual atoms or polyatomic species striking the substrate surface. Nucleation steps and the growth mechanisms of PVD films determine the microstructure and morphology of the deposited materials, which ultimately alter physical properties of PVD films such as microhardness, residual stresses, surface

44

roughness, mass density, reflectivity and etc. Therefore, it is important to understand the growth mechanisms of films.

Generally, a proper choice of deposition variables (vacuum pressure, deposition rate, substrate temperature) is crucial and decisive to the film to be deposited. The influence of these variables on film nucleation and growth can be understood by the same effect of the parameters on the film nucleation and growth. For example, reducing deposition pressure it corresponds to an increase of film purity. Therefore, there is a correlation between the deposition variables and the parameters of nucleation and growth on one-to-one basis. Table 2 shows this link during film deposition.

Table 2. A linkage between deposition variables and nucleation and growth parameters.

Deposition Variables

Vacuum pressure

Deposition rate

Substrate temperature

Substrate structure

Parameters of nucleation and growth

Impurity content

Supersaturation

Undercooling

Interfacial energy

The selection of a substrate allows and controls the quality of thin films. Thus, according to the above discussions, the combination of poor vacuum, high deposition rate and low substrate temperature will result in the formation of finegrained thin film (or amorphous film) on a glass substrate. It is often grain size varies with film thickness. According to microstructural evolution study, for PVD Al film, the mean grain size (d) increased with film thickness (t) as relation d ~ t09 by using transmission electron microscopy (TEM) observation [47].

Thin film formation involves basic processes of nucleation and growth. Concerning thin film nucleation on substrate surface, there are three primary growth modes that they occur at the earliest stages of film formation. They can be distinguished based on bonding strength between depositing species and the substrate surface atoms. According to many observations, these 2D growth mechanisms are including 1) two-dimensional layer-by-layer (Frank-Van der Merwe) mode. This growth mode valids when the deposited atoms have a greater bonding strength and is applied for the homoepitaxial growth i.e. metal/metal or semiconductor/semiconductor type structure. 2) island growth mechanism (Volmer-Weber) mode is applied when the smallest stable clusters nucleate on the substrate and grow in three dimensions to form islands on the

45

surface. This type of growth mode happens when atoms or molecules in the depositing films are more strongly bound to each other than to the substrate atoms. The islands formed by this process are merging by a coalescence phenomenon, which resulted in reduction of island density. The initial growth of many metals and semiconductors on oxide substrates follow such mechanism and 3) Mixed layer-island (Stranski- Krastanov) mode is an intermediate mixture of the preceding two modes. In this mechanism, first one or more layers are formed followed by the growth of 3D islands. Figure 9 shows the schematic illustration of these three thin film growth modes in terms of substrate surface coverage (6) defined as a number of adsorbed atoms (na) per total surface sites (n0) available on the substrate (9=no/na) in unit of monolayer (ML).

The growth modes can be also described by interfacial tension y according to the following inequalities:

l)For layer by-layer growth: ysv > y^+y^

2) For island growth: ysv <yfi + y fi

3) For Mixed layer-island growth: ysv > y fs + y^

where the subscripts v, s and f represent vapor, substrate and films, respectively. A pair of subscripts refers to the interface between the indicated phases. For more information about the film formation mechanism, readers are referred to [3,7,48]. An excellent review on nucleation theory and the early stages of thin film growth along with rate equations and kinetic Monte Carlo simulation is reported recently [49].

Figure 9. Simplified schematic view of three growth modes of thin film in different 8 (from [6]).

46

After the nucleation, coalescence and growth, grains are formed on the substrate surface. The film morphology (shape and size) can be different depends on substrate temperature, its orientation, type of depositing material and etc. For example, columnar grains have been observed in high melting point materials (Cr, Be, Si) or high binding energy compounds (TiC, TiN) where the deposited atoms have limited mobility on the substrate surface. For this type of structure, a so-called Tangent Rule is expressed the geometry of columnar grains as following:

tana=2tan/? (12) where a is defined the angle between the source direction and substrate normal (0<a<90°) and /? represents the angle between the column axis and the substrate normal ( / ? < « ) . For detail of deposition geometry, readers are referred to [7]. A comprehensive review on various aspects of nucleation, growth and structure of thin films with several examples with observations can be found elsewhere [50]. Moreover, for the metallic films, a phenomenological description of grain growth mechanisms based on morphology and substrate temperature is proposed by Grovenor et al. [51]. Recently, microstructural evolution during polycrystalline film growth is described in ref. [52].

5. PVD in Microelectronics

It is well established that the low resistance interconnection materials such as Al, Cu, and Ag are used mostly for metallization in integrated circuits. One of the important techniques to fabricate these interconnection in microelectronic technology is PVD. Concerning thin film application in microelectronic, PVD has historically been used to deposit conductors for contacts, diffusion barriers, liners, and wiring in multilevel metallization.

For metallizations process, sputtering is preferred in most semiconductor depositions because stoichiometry of alloys is maintained and step coverage is usually superior [53]. Since 1990, copper (Cu) has been suggested as a new metallization material for ultra large-scale integration (ULSI) and giga scale integration (GSI) technology [54-56] owing to its lower bulk resistivity value ( p c = 1.67/iQ-cm) and higher electromigration resistance as compared to Al metal (PAI = 2.70//Q - cm). Despite such excellent properties of Cu, there are several problems with Cu including, i) fast diffusion of Cu into Si substrate that results in silicide formation, ii) poor adhesion to dielectric material, iii) fast reaction of Cu with environment even at low temperature [54]. To improve some of these problems particularly inter-diffusion between Cu and Si substrate, the use of high melting point material (Ta,W,Mo,Ti,C) is suggested as a diffusion barrier between Cu-Si system. Recently, DC sputtering technique was

47

used during Cu metallization process in the Cu/Ta/Si(lll) structure to study thermal stability of Ta diffusion barrier layer. Figure 10 shows the variation of sheet resistance of both Cu/Ta (unbiased)/Si (111) and Cu/Ta (biased)/Si (111) thin film structures in different annealing temperatures [57,58]. A sudden sharp increase in sheet resistance at 650 °C can be related to inter diffusion and/or reaction between Cu and Ta layers. Copper metallization for ohmic contacts, gate metal and interconnections is well studied and reviewed by Murarka et al. [54]. In addition, a well studied report on various aspects of silver metallization and encapsulation for future device fabrication is discussed recently [59]. Furthermore, an up dated review on surface and interface characterization of microelectronic materials and their progress is reported in [60].

350

250

G 150-Ui O z

tn CO \u tc i -iii UJ x CO

200 400 TEMPERATURE (° C)

600 800

Figure 10. The dependence of sheet resistance on different annealing temperatures for multilayer structures a) Cu/Ta (unbiased)/ Si(l 11) and b) Cu (Ta (biased)/Si(l 11).

In the last four decades, semiconductor technology has been grown at exponential rate. It is argued that this advancement in both productivity and performance do not continue endlessly. Thus, for giga scale and perhaps trascale integration (TSI) in early 21 st. century, there are a hierarchy of limit whose five levels are codified as 1) fundamental, 2) material, 3) device, 4) circuit and 5) system which govern future GSI technology [61]. Concerning the fundamental limits on interconnect, the equation that correlates the material properties with

48

miniaturization process is R/L= p /F2 where p is the effective resistively of a conductor, F is a minimum feature size, and R/L is resistance per unit length of interconnects. It is obvious that as F decreases R/L is drastically increased. As feature size (F) gets smaller and smaller, the Al interconnects cross section and thus the resistance of the line increases. On the other hand, the increase in resistance resulted in increasing RC time delays. Therefore, it is necessary to use the lower resistance interconnects (i.e. Cu, Ag) and dielectric materials with lower dielectric constant (i.e. Si02) for future GSI/TSI device fabrication [62].

One of the important applications of PVD is in metallization process. For example, contacts to doped semiconductor electrodes such as source are necessary in the device fabrication. Al or PtSi are two examples of these contacts. On the other hand, interconnects such as Cu or Cu-Si can be also made by PVD technique. Figure 11 shows various thin film materials produced by PVD technique for wide use in microelectronic industry [63].

1000

800

0 O) g 600 =} E i l Q 400 > Q.

200

0 Al Ti TiN MoSix TiW WSix other

Figure 11. Histogram of PVD thin film applications in microelectronic industry.

Although Cu is one of the best candidates for metallization in submicron scale (< 0.18//m) in ULSI technology, passivation treatment of its surface is also required since Cu oxides in air even at low temperatures. In addition, Cu demonstrates poor adhesion to most dielectric materials. Thus, inert coatings such as SiC can protect Cu from degradation resulting from high temperature and corrosive environments.

II

49

There are many reasons why PVD has been successful for microelectronic fabrications. Some of these are including: 1) sputtering can be used to deposit all conducting films that currently used in interconnect metallization schemes [63]. 2) PVD deposition rate is also well matched to the throughput needs of wafer fabrication (-40 wafers per hour). 3) PVD is able to produce global film uniformity. In this context, advanced PVD sources can deposit films with 3 a (three-sigma) nouniformity of 3-5% over 200 mm diameter (8-inch) Si wafers in a production environment [63]. In addition, PVD utilizes nontoxic targets and low pressures of inert gas 4) PVD has shown an acceptably low cost-of-ownership (CoO) consistent with the economic demands of production-line wafer fabrication [63].

There are only two thickness ranges of interest for PVD films used in microelectronics a) they are 5-50 nm for contacts, barriers, liners and anti-reflection coating (ARC) layers and b) 0.5-1 ju m for contact plugs, via plugs and interconnect lines. It is important to note that the ratio of height-to-width (i.e. the aspect rations, AR) to features to be coated or filled with metal can varied from zero for a planar surface to AR=10:1 for a contact hole in advanced devices [63].

Concerning film thickness, a control of conductive film thickness is essential, because a film thinner than desired one can lead to excess current density (J) and resulted in device failure during the operation. On the other hand, excessive thickness can result in difficulties in etching the film. For example in considering the first case, in electromigration of thin wire conductor film, a thin film yields a passage of higher J that resulted in lower mean time to failure (MTTF) which causes device failure. For further and comprehensive knowledge on the device failure, the readers are referred to a recent and well-written text [53].

Metal silicides have also received much attention in the last decades due to their extensive application in microelectronic devices such as contacts, gate electrodes and interconnect materials. This is mainly due to their low electrical resistively, high thermal stability and compatibility with current processing techniques [64].

Among refractory metal silicides, CoSi2 is regarded as a good choice for application in microelectronic devices as compared to conventional TiSi2

compound used in very layer scale integration (VLSI) technology. Recently, a high quality single crystalline CoSi2 thin film has been prepared in the Co/W/Si(100) multilayer structure using a combined evaporation-sputtering techniques [65,66]. A nearly single crystalline structure of CoSi2 with a dominant (200) texture and low sheet resistance (Rs) value of about 1 Q/n was formed under the controlled experimental conditions [67].

50

6. Conclusions

The principles of thin film technology namely fabrication, characterization and applications and their interrelation is reviewed. Physics and technology of three important physical vapor deposition (PVD) techniques a) sputtering, b) pulse laser deposition (PLD) and c) molecular beam epitaxy (MBE) are described. In sputtering techniques the physical processes, sputtering yield at different energies as well as magnetron and bias methods are explained. A linkage between deposition variables and parameters of nucleation and growth of thin films is emphasized. The initial stages of nucleation and growth modes: layer by- layer, Island, and mixed layer-Island with illustrations are presented. Finally, several examples of PVD thin films in metallization process applied in microelectronic industry. Are reviewed.

Acknowledgement

The author is especially grateful to all his graduate students. He also wants to thank the Research Council of Sharif University of Technology and High Technology Organization (Ministry of Industries and Mines, Iran) for financial supports. The assistance of Mr. Azimirad and Mr. A. Azarm for providing useful and most recent information on PVD subjects as well as Ms. M. Bashlideh and Ms. L. Chahooshi for preparing the manuscript is greatly acknowledged.

References

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New York, 1970). 3. L. Eckertova, Physics of Thin Films (Second Edition, Plenum Press, New

York, 1984). 4. D.L. Smith, Thin-Film Depositions: Principles and Practice (McGraw-

Hill, New York, 1995). 5. A. Elshabini-Riad, F.D. Barlow III., Thin Film Technology Handbook, ed.

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11. T. Heberlein, G. Krautheim, W. Wuttke, Vacuum 42, 47 (1991). 12. A.Z. Moshfegh, S. Behnamfar (unpublished). 13. R. Behrisch, Sputtering by Particle Bombardment II, Topics in Applied

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19. G.N. Jackson, Thin Solid Films 5, 209 (1970). 20. J. S. Logan, Thin Solid Films 188, 307 (1990). 21. W. D. Westwood, Physics of Thin Films (Academic Press, Vol. 14, 1989),

pi. 22. S.M. Rossnagel, J. Vac. Sci. Technol. A 21, S74 (2003). 23. M.V.R. Murty, Surf. Sci. 500, 523 (2002). 24. A.Z. Moshfegh. Y.Q. Wang, Y.Y. Sun, A. Mesarwi, P.H. Hor, A. Ignatiev,

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63, 1963, Sandia Corp. 26. A.Z. Moshfegh, P. Sangpour, Phys. Stat. Solidi (c) 2004 (accepted). 27. A.Z. Moshfegh, O. Akhavan, H. Salamati, P. Kameli, M. Akhavan, The

Proceedings of the First Regional Conference on Magnetic and Supercomputing Materials, eds. M. Akhavan, J. Jensen, K. Kitazawa, (World Scientific Publisher, Vol. A, 2000) p. 585.

28. T. A. Znotins, Solid State Technology 99 (September-1986). 29. J.T. Cheung, "History and Fundamentals of Pulsed Laser Deposition",

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30. R. Kelly, A. Miotello, "Mechanism of Plused Laser Sputtering", in Pulsed Laser Deposition of Thin Films, eds. D.B. Chrisey, G.K. Hubler, John (John Wiley & Sons, New York, 1994) p. 55.

31. J.P. Zheng, H.S. Kwok, Thin Solid Films 232, 99 (1993). 32. Q. Li, T. Venkateson, X.X. Xi, High Tc Superconductor Thin Films, ed. L.

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(Cambridge University Press, Cambridge UK, 2000). 49. C. Ratsch, J.A. Venables, J. Vac. Sci. Teehnol. A 21, S96 (2003). 50. R. W. Vook, Inter. Metals Reviews 27,209 (1982). 51. C.R. M. Grovenor, H.T.G. Hentzell, D.A. Smith, Acta Metall. 32, 773

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Kohl, IBM. J. Res. & Dev. 46, 245 (2002). 63. R.A. Powell, S.M. Rossnagel, PVD for Microelectronics: Sputter

Deposition Applied to Semiconductor Manufacturing (Academic Press, Vol. 26, 1999).

64. S.P. Murarka, Silcides for VLSI Application (Academic Press, New York, 1983).

65. A.Z. Moshfegh, S.J. Hashmifar, O. Akhavan, R. Azimirad, Thin Solid Films 433, 298 (2003).

66. A.Z. Moshfegh, S.J. Hashmifar, O. Akhavan, Solid State Communcation 128, 239 (2003).

67. O. Akhavan, A.Z. Moshfegh, S.J. Hashemifar, R. Azimirad, Appl. Surf. Sci. 2004 (accepted).

INTRODUCTION TO SEMICONDUCTOR EPITAXY

H.VON KANEL INFM and L-NESS, Dipartimento di Fisica del Politecnico di Milano, Polo Regionale di

Como, Via Anzani 52,1-22100 Como, Italy E-mail: hvkaenel@como.polimi. it

The principles of semiconductor epitaxy are reviewed with emphasis on hetero-epitaxy between dissimilar materials. The role of strain in lattice-mismatched systems is outlined, and the ways in which strain can be relaxed. Illustrative examples stem from SiGe/Si heteroepitaxy. Depending on misfit, plastic relaxation by means of misfit dislocations or elastic relaxation through nanostructure formation can be observed.

1. Introduction

The term epitaxy is used for the deposition of a crystalline material on a likewise crystalline substrate, where the crystal axes of the overlayer bear some definite relation to those of the substrate. Often a somewhat more precise terminology is used when it is to be emphasized that two materials grown on each other are the same or different from each other, i.e., homoepitaxy or heteroepitaxy, respectively.

There exist a huge number of material systems, which can be grown epitaxially on one another, by a great variety of deposition techniques. It is therefore necessary to limit ourselves to an outline of the basic principles, together with a few specific examples, without being able to do justice to the wealth of information available for specific combinations of materials and techniques.

A schematic representation of the atomic processes occurring during epitaxial growth can be seen in Figure 1. This figure actually refers to the conceptually simplest process, namely molecular beam epitaxy (MBE). In MBE, the heated substrate onto which the epitaxial film is grown is contained in an ultra-high vacuum (UHV) chamber. The source material is usually evaporated from so-called effusion or Knudsen cells. For materials with high boiling points, electron beam evaporation is often used instead. Since the process takes place in UHV, the evaporated molecules (or atoms) travel along straight paths without being scattered, before they impinge onto the heated substrate, where they may dissociate and eventually be incorporated into the growing film. The technique of MBE has experienced a remarkable period of expansion since the early seventies, when resonant tunneling [1] and quasi-two dimensional excitons [2] in semiconductor quantum heterostructures were first demonstrated in GaAs

54

55

based heterostractures. Experimental and theoretical work on MBE extends much farther back, however. Ref. [3] gives a comprehensive review of the early studies of MBE-growth.

Figure 1. Fundamental processes occurring during epitaxial growth. Atoms arriving on the substrate can be adsorbed (a), diffuse on the surface (b), encounter each other to form islands (c), attach to pre-existing islands or steps (d) and (g). Atoms can also detach from islands (e) or even desorb again CO.

Although MBE does have some industrial applications, such as the production of GaAs lasers for compact disc players, it has the disadvantage of source depletion: the effusion cells or e-beam evaporators need to be refilled regularly, which requires breaking of vacuum.

This disadvantage does not exist in gas phase processes in which the deposition material is supplied in a continuous fashion. The most important gas phase process is Chemical Vapour Deposition (CVD). It is schematically outlined in Figure 2 for the example of Si and SiGe epitaxy. In CVD the reactive precursors, such as silane (SiH4) or germane (GeH4), are fed into the reactor by means of gas lines equipped with mass flow controllers. These molecules stick to the heated substrate with some probability, and diffuse around until they encounter some reactive site, where a more reactive radical is formed by stripping off one or more hydrogen atoms. After losing all H-atoms, Si and Gc are finally incorporated into the crystal lattice.

There exist a large number of different CVD techniques, distinguishing themselves basically by the pressure regime under which the reactor is operated. This may range from atmospheric pressure in APCVD [4] to approximately 10'5

mbar in gas source MBE [5]. There are many CVD techniques with different names, operating at pressures in between these extremes, such as reduced pressure CVD, low pressure CVD (LPCVD), etc.

56

Tie reader interested in an overview of different deposition techniques should consult, e.g. Ref. [6].

chem. vapor deposition (CVD) molecular beam epitaxy (MBE) Figure 2. Schematic representation of chemical vapour deposition (left) and molecular beam epitaxy (right) for the case of SiGe heteroepitaxy.

According to Ernst Bauer, one of the pioneers in this field, epitaxial growth of material B on material A may occur in either of three modes [7]. The Frank-van der Merwe (FM), or layer-by-layer mode, leads to films with the smoothest surfaces, since each new monolayer (ML) virtually starts growing only when the previous one has been completed. Very often, whenever films with flat surfaces and Interfaces are required, this Is the desired mode of growth. As we shall see, however, the FM mode forms the exception rather than the rule. Much more common are the Volmer-Weber (VW) and the Stranski-Krastanow (SK) growth modes. In the former, the epitaxial growth occurs In the form of three-dimensional islands nucleating all over the substrate right from the very beginning. By contrast, in the SK mode the substrate is covered first by one or more MLs of a two dimensional film, the so-called wetting layer, before three-dimensional clusters start to form.

The three growth modes are schematically displayed in Figure 3. They can be qualitatively understood by considering the various surface free energies <JA, 0B and on?. Here aA and aB stand for the solid/vapour interface of material A and B, respectively, while <J!F describes the interfaclal energy between the two materials. VW growth is expected when aA < CJB + aw. Qualitatively, It Is clear that the VW mode will be favoured when the film/substrate interaction is weaker than the film/film interaction.

57

Figure 3* The three fundamental growth modes according to Bauer 7]. The Frank-van der Merwe mode leads to a flat, two-dimensional film of material B on substrate A, whereas Volmer-Weber and Stranski-Krastanow modes result in islanding.

When oA > 0B+^IF* the often desirable layer-by»layer growth becomes possible In principle. The distinction between the FM and SK modes is slightly more delicate, however. Before we go on with our discussion, let us define an important quantity, namely the geometric misfit f:

Here, a and b are the lattice parameters of substrate and overlayer, respectively. When f = 0 the overlayer Is said to be lattice-matched to the substrate. In homoepitaxy this is of course always the case. In heteroepitaxy, the FM growth mode is more likely to occur when this condition is fulfilled.

As an example of a closely lattice-matched system., we might mention GaAs/Gai.xAlxAs, with x < 0.4, which is the prototypical system for epitaxial semiconductor heterostractures (see Refs. [1] and [2]).

At finite misfit a sufficiently thin film is characterized by a uniform elastic strain, such that its lattice parameter parallel to the interface plane exactly matches that of the substrate [8-9]. In such a case the interface is said to be coherent or the film to be pseudomorphic. Such a case Is illustrated on the left hand side of Figure 3 for a simple cubic lattice. There, the lattice parameter of the film has been assumed to be larger than that of the substrate. The film is hence under compressive strain, as a result of which the lattice expands in the direction perpendicular to the film. One example of practical significance is epitaxial growth of Ge on top of Si, the Ge lattice parameter being larger by 4.2 %. The presence of misfit Is essential for the occurrence of the SK growth mode.

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In fact, molecular dynamics simulations predict [10] that in the absence of long-range forces a uniform film with a thickness exceeding a few MLs cannot be the equilibrium state for any finite value of the misfit f. The term "equilibrium" is crucial here, since, as we shall see below, uniform films can very well be grown in a metastable form for growth conditions far from equilibrium. A very qualitative understanding of SK growth may be reached in the following way. Assume that the film/substrate interaction is very strong compared with the film/film interaction. It is evident then that the first monolayer deposited will be likely to cover the substrate completely. The atoms of the next ML will experience a potential with the same periodicity, but with a smaller amplitude, because of the weaker interaction with the substrate. For the next ML the geometrical situation is similar, but with the atoms "feeling" even less the presence of the substrate. Then, the strain energy accumulated with each successive monolayer provides the driving force for the nucleation of islands [10].

Let us now consider a case in which a film with a lattice parameter different from that of the substrate starts to grow in a layer-by-layer mode. It is evident that, as the film grows thicker, more and more strain energy will be accumulated. A pseudomorphic film must therefore become increasingly unstable as growth proceeds. There are essentially two ways in which the elastic energy can be lowered:

• by elastic relaxation, i.e., by island formation or, more generally, by roughening of the film

• by plastic relaxation, i.e., by the introduction of misfit dislocations In an island under compressive strain, elastic strain relaxation is possible

because the lattice planes can bend outwards, as schematically shown in Figure 4 (a). By contrast, a misfit dislocation is obtained by removing one atomic plane from the film, such that its average lateral lattice parameter is allowed to expand as in Figure 4 (b).

Both the nucleation of islands and the formation of misfit dislocations are thermally activated processes. In practice, they are therefore often governed by kinetics rather than by equilibrium thermodynamics, and thus depend on the details of the growth procedure.

A more detailed view of a compressively strained and relaxed film can be seen in Figure 5. We use this example of two simple cubic lattices with finite misfit to illustrate the important definition of the Burgers vector. The additional substrate lattice planes terminating at the interface define the dislocation lines. For the simple geometry considered here each dislocation line runs perpendicular to the image plane. A similar set of dislocation lines will be oriented perpendicular to those just discussed, if we take into account the biaxial

59

nature of the problem. By counting the number of nearest neighbour distances traversed as one moves once around a closed loop encircling the dislocation line, it is easy to see that there is a missing piece, which defines the Burgers vector. In the example of Figure 5 the Burgers vector lies in the interface plane and is perpendicular to the dislocation line. This is the simplest case of a so-called pure edge dislocation. More generally, dislocations may have a Burgers vector component parallel to the dislocation line (screw component), and the Burgers vector need not lie in the interface plane. In such a case, the misfit dislocation is less effective in relieving strain, as only the edge component in the interface plane may contribute.

Figure 4. Schematic representation of strain relaxation by island formation (a) and by the introduction of misfit dislocations (b).

a a

Figure 5. Coherent (left) and incoherent (right) interface for two simple cubic lattice with finite misfit f, such that b > a (compressively strained film). The dislocation lines in the relaxed film run perpendicular to the image plane. A Burgers circuit and the Burgers vector are indicated on the right.

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After having outlined the general principles of epitaxial growth, we shall consider as an example one of the material systems of practical importance: Ge and Sii.xGex alloys grown epitaxially on Si(001).

2. Epitaxial Growth on Silicon

2.1. The Clean Si(001) Surface

The Si surface is normally covered with a thin natural oxide, which has to be removed before any epitaxial growth is possible. The methods of oxide removal depend on the growth technique. Most common are

• chemical etching and passivation by a dilute HF dip • thermal oxide removal in UHV, typically above temperatures of 800° C.

The clean surface of Si(001) is characterized by the presence of dimer lines (Figure 6 (b), (c)), dimer formation being a way to reduce the number of dangling bonds on the surface, and hence reduce its energy. On a bulk terminated surface, each atom would have two dangling bonds. Dimerization halves the number of dangling bonds at the expense of a weakening of the backbonds due to the induced distortion. The ground state actually consists of buckled dimers which cannot be seen at room temperature because of rapid flipping between the two possible buckled configurations. When observed by scanning tunneling microscopy (STM) at room temperature, the dimers therefore look symmetric, unless the buckling is stabilized by some surface defects (Figure 6 (c)).

Since in practice a surface is never completely flat, exact knowledge of the initial step structure and its evolution are indispensible in order to understand epitaxial growth.

The surface steps on singular (i.e., nominally exactly oriented) or slightly misoriented clean Si(001) are of monolayer (ML) height. The characteristic (2x1) reconstruction results in dimer rows oriented perpendicular to one another on adjacent terraces. The terraces are bounded by two distinct types of steps: SA

steps to which the upper-terrace dimer rows are oriented parallel, and SB steps to which they are perpendicular.

These feature can be seen in the STM images of Figure 6, obtained on a homoepitaxial Si(001) surface grown by magnetron sputter epitaxy (MSE), yet another epitaxial growth technique based on an industrial process adapted to UHV conditions [11], We shall not go into any details here. Suffice it to say that by appropriately choosing the experimental conditions (working gas pressure, target-substrate distance, etc.) it is possible to thermalize sputtered particles to such an extent that excellent control of interfaces and surfaces can be exercized also by MSE.

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Figure 6. STM image of a homoepltaxial Si(OOI) film grown by magnetron sputter epitaxy (MSE). The images are oriented along the [001] direction. Dimer rows ran along <110>. Monoatomic steps are labelled as SA and SB in (b). Dimer rows are indicated by DR5 dimer vacancy lines by DVL m

(c).

22. Epitaxial Growth Under RHEED Control

MBB in particular offers excellent control over the thickness of epitaxial layers, since the flux of Impinging molecules or atoms can be almost instantaneously stopped by fast shutters interrupting the beams, at typical growth rates of 1 A/s. By using Reflection High Energy Electron Diffraction (RHEED) even more precise thickness control can be realized. In this technique, a focused electron beam with mer& of typically a few tens of keV Is directed onto the sample at a shallow angle of the order of one degree.

Under such conditions, we have to bear in mind the following aspects:

• Despite of the large energy of the beam, the electrons will penetrate only a small distance Into the solid, say of the order of 1 ran.

• The radius of the Ewald sphere will be huge in comparison to that common in X-ray scattering, because the electron wavelength is much smaller than typical lattice spacings.

In view of the small penetration depth we can consider the electron beam to be diffracted from the surface atoms alone in a first approximation. That means, however, that we are considering scattering from a periodic 2D object. If we take the z»axls to be pe^endicular to the surface for convenience, this means that the reciprocal lattice vector Qz will be equal to zero. Instead of a three dimensional lattice of reciprocal lattice points (as Is the case In ordinary X-ray scattering), we therefore obtain an array of reciprocal lattice rods perpendicular to the surface.

The Ewald sphere will cut these reciprocal lattice rods in a series of concentric circles as indicated in Figure 7.

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The main merits of RHEED are that the technique (as opposed to low energy electron diffraction or LEED) can be used in situ, that is during the epitaxial growth.

Let us now consider a case where the FM or layer-by-layer growth mode is present. We assume that the substrate surface has been well prepared to be clean and flat before the shutter is opened, and that it is well oriented in the sense that the distance of the ubiquitous surface steps is large. This corresponds to a maximum of the specular beam intensity in RHEED. The moment the shutter is opened, nucleation of 2D islands will begin, the surface roughens and the specular RHEED intensity drops. The islands continue to grow until a full monolayer of material is grown. At this point, the surface is perfectly flat again, and, upon further growth, the process is going to repeat itself. In this simple scenario, a minimum of the RHEED intensity is expected at half monolayer coverage, because at that point the surface is maximally rough (see, e.g., Ref. [12]).

Figure 7. Reciprocal lattice for a two dimensional lattice (a), and corresponding Ewald construction (b). Only one Laue zone (one circle) is shown. In general, there will be several concentric circles.

2.3. SiGe Nanostructures on Silicon

As we have pointed out in the Introduction, one way to relieve misfit strain is by means of island formation (Figure 4). SiGe grown on top of Si is a typical SK system. Si and Ge are two elements which are miscible in any proportion due to their chemical similarity. To a good approximation, the lattice parameter of a Sii_xGex alloy can be interpolated linearly between that of Si, 5.43 A, and that of Ge, 5.65 A. Neglecting the deviations from linearity, the misfit of a Si^Ge* alloy on Si with a Ge concentration of x is thus given by

/ = 0.042 • x (2)

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Whereas for low misfit, I.e., low x, relaxation takes place mainly by dislocation formation, the driving force for island formation can be shown to dominate for high Ge concentrations [13]. For pure Ge in particular, 2D growth persists to a film, thickness of approximately 3 ML only, above which 3D islands start to form [14].

Most work to date has been carried out on Si(001), since this is the technologically most important Si surface. The Ge/Si(001) system actually turns out to be surprisingly complex, as can be seen in the scanning tunnelling microscopy image (STM) of Figure 8. Islands do not just form on top of the wetting layer and then simply grow in size upon further deposition. Instead, they exhibit a number of shape transformations as the growth proceeds. Some of these shape changes are apparent from Figure 8. The first faceted islands, appearing after the wetting layer thickness is exceeded, are so-called pyramids and hut clusters [14] with very typical {105} facets. Recent calculations have actually shown that it is not just the elastic strain relief (Figure 4), which causes islands on top of a wetting layer to be more favourable than a flat film. It is rather the particular surface reconstruction on the {105} facets which makes islands more stable [15].

[1001 4£nm

Figure 8, STM image of a sample obtained by depositing 7 ML of Ge on Si(001) at a substrate temperature of 500° C. Four kinds of islands can be seen; small {105} faceted pyramids, elongated {105} faceted hut clusters, larger multifaceted islands (domes), and a large dislocated island, a so-called super dome.

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Strain relaxation Is more efficient for the larger dome-shaped Islands to Figure S because of their steeper facets. This Is the reason for which these islands become more numerous with larger coverage, at the expense of pyramids and hut clusters.

Figure 9 shows a series of STM images, obtained on different islands, from which the probable pathway followed In the transition from pyramid to dome can be guessed. According to these images, steps first appear near the pyramid apex. This can be understood by assuming that the less strained apex (Figure 4) attracts Ge atoms from the base of the pyramid. A schematic representation of the process leading from pyramid to dome can be seen In Figure 10.

Figure 9. 3D views of STM images illustrating the probable from a pyramid (a) to a dome (f). Islands with kinetically determined intermediate shapes are shown in (b)-(e). The indicated dimensions are in nanometers.

Nanostractures of the kind considered here have become Increasingly popular in recent years. The hope is that the use of small self-organized islands as quantum dots or quantum boxes in micro- and opto»electronic devices might be a way In which expensive lithography steps can be avoided. There are, however, a number of obstacles to be overcome to make such applications possible:

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Figure 10. Schematic picture showing the transition from pyramid (a) to dome (d) through intermediate transition domes (b), (c). The transition begins with the formation of incomplete {105} facets (b). Atoms, instead of sticking to the edge of incomplete facets, hop onto the existing {105} terraces, causing step bunching (c), finally resulting in the reconstructed facets composing the dome surface.

• Shape and size uniformity are usually far from satisfactory • The process of embedding the islands within a matrix is far from trivial,

since it can strongly modify their shape and composition. The shape changes can be quite dramatic, as is evident from Figure 11,

obtained after depositing Si at a substrate temperature of 450° C onto a sample originally composed of domes. The domes have been grown by depositing 7 ML of Ge at a substrate temperature of 550° C, followed by rapid quenching, such that additional shape changes during cool-down due to ripening can be excluded. Figure 11 shows that the dome shape has changed appreciably after as little as 1 ML of Si deposition. It is hence clear that once the process of embedding the islands is completed, the latter do not have much in common with those present before the capping process.

That does not mean that shape changes of islands during capping are inevitable. The thermally activated nature of processes like surface diffusion and segregation does in fact pave the way for a solution of this problem by lowering the substrate temperature during the capping process [16].

2.4. Virtual SiGe Substrates

Up to now, we have only been considering pseudomorphic structures, i.e., structures which either adapt their lateral lattice parameter to that of the substrate (2D films), or structures which relax elastically (islands).

Just as important are structures with low or negligible residual strain that can be maintained in the form of 2D films. One very important example are relaxed SiGe alloys on top of Si. They can be considered as virtual substrates, since they can assume the role of a substrate during any following epitaxy steps.

66

Figure 11* STM Images of representative islands during Si capping, (a) Uncapped dome, (b)-(f) Si coverage of 1,2,4,8*16 ML, respectively.

The plastic relaxation we are describing here is fundamentally different from the elastic relaxation treated in section 23 (see also Figure 4). It is accomplished by means of misfit dislocations in the interface plane, each of which can relieve some amount of strain. The easiest way to visualize a misfit dislocation is by introducing or removing an additional lattice plane or by removing one, depending on whether the epitaxial layer has a smaller or larger lattice parameter than the substrate.

A fully relaxed layer is characterized by a network of misfit dislocations lying in the interface plane. For a square network of dislocation lines, such as is expected in the case of epitaxy on Si(001), the spacing L of the dislocation lines is given by

67

Here, b Is the component of the Burgers vector lying in the interface plane and perpendicular to the dislocation line, and f is the misfit defined In Eq. (1).

In reality, the situation is complicated by the geometry of the dislocations in Si and SiGe (Figure 12). For the most part, these are so-called 60° dislocations, where the Burgers vector lies in a {111} plane and forms a 60° angle with the misfit dislocation line [17].

Figure 12. Geometry of 60° dislocations TD is the threading segment lying in the (111) glide plane, and MD the misfit segment in the (001) interface plane (after [17]).

At this point it is appropriate to remind ourselves that we are considering a case in which an epitaxial film starts to grow in a layer-by-layer fashion (Figure 3). In other words, the film starts to grow fully strained until at some critical thickness the introduction of strain relieving misfit dislocations becomes energetically favourable. How can an entire lattice plane be removed from the film once the critical thickness is reached? In principle, this could take place by diffusion, which, however, requires far too high an activation energy. Alternatively, a dislocation half loop may be formed at the surface, which expands on a {11.1} plane until the dislocation line reaches the interface, forming a misfit segment there. Evidently, the expansion of a dislocation half loop results not only in a misfit dislocation in the interface plane, but also to two dislocation lines penetrating the surface of the film. They are called threading arms or threading dislocations (IDs). In Figure 12 just the threading dislocations on one side of the misfit segment are represented.

68

<iio>

Figure 13. Schematic representation of a dislocation half loop nucleating at the surface and expanding until it reaches the interface at which point a misfit segment starts to form.

Ideally, the threading arms would continue to move in opposite directions under the influence of the misfit stress, until they finally reach the edges of the sample. Were this to happen? a square network of misfit dislocations would ultimately form, and the epitaxial film would be fully relaxed and defect-free* Unfortunately, this does not happen because dislocations tend to interact and block each other in practice. This leads to shorter misfit segments and the presence of a high density of TDs, which for SiGe can be as high as 1011 cm**2, when a film of constant composition is grown on Si.

Here, the breakthrough came with the discovery that grading the Ge composition from zero to some final value results in much lower TD densities 18[18]. The reason is that dislocation interaction is strongly reduced in this approach, because the misfit segments are no longer confined to one single plane. This can be seen in the transmission electron microscopy image of Figure 14 in which the dislocations are confined to the graded part of the material.

Figure 14. Cross-section transmission electron microscopy image of a Yirtual SiGe substrate linearly graded to a final Ge composition of 40 % at a grading rate of 10 % per pm. The upper part of the material is defect-free on this scale.

69

References

1. L.L. Chang, L. Esaki, R. Tsu, Appl. Phys. Lett. 24, 593 (1974).

2. R. Dingle, in Festkorperprobleme XV, ed. H.J. Queisser (F. Vieweg & Sohn, Braunschweig, 1975).

3. Epitaxial Growth, ed. J.W. Matthews (Academic Press, New York, 1975).

4. P.D. Agnello, T.O. Sedwick, M.S. Goorsky, J. Cotte, Appl. Phys. Lett. 60, 454 (1992).

5. J.M. Hartmann, B. Gallas, R. Ferguson, J. Fernandez, J. Zhang, J.J. Harris, Semicond. Sci. Technol. 15, 362 (2000).

6. Handbook of thin-fdm deposition processes and techniques, ed. K.K. Schuegraf (Noyes Publications, Park Ridge, New Jersey, 1988).

7. E. Bauer, Z. Kristallogr. 110, 372 (1958).

8. F.C. Frank, J.H. van der Merwe, Proc. Royal Soc. A 198, 205 (1949).

9. J.H. van der Merwe, J. Appl. Phys. 34, 117 (1963).

10. G.H. Gilmer, M.H. Grabow, J. Metals 39, 19 (1987).

11. P. Sutter, D. Groten, E. Muller, M. Lenz, H. von Kanel, Appl. Phys. Lett. 67,3954(1995).

12. T. Kawamura, T. Natori, T. Sakamoto, P.A. Maksym, Surf. Sci. 181, L171 (1987).

13. J. Tersoff, F.K. LeGoues, Phys. Rev. Lett. 72, 3570 (1994).

14. Y.-W. Mo. D.E. Savage, B.S. Swartzentruber, M.G. Lagally, Phys. Rev. Lett. 65, 1020(1990).

15. P. Raitieri, D.B. Migas, L. Miglio, A. Rastelli, H. von Kanel, Phys. Rev. Lett. 88, 256103 (2002).

16. A. Rastelli, E. MUller, H. von Kanel, Appl. Phys. Lett. 80, 1438 (2002)

17. A.E. Blakeslee, Mat. Res. Soc. Symp. Proc. 148, 217 (1989).

18. F.A. Fitzgerald, Y.-H. Xie, M. L. Green, D. Brasen, A. R. Kortan, J. Michel, Y.-J. Mii, B.E. Weir, Appl. Phys. Lett. 58, 811 (1991).

SEMICONDUCTOR SUPERLATTICES

H.VON KANEL

INFM and L-NESS, Dipartimento di Fisica del Politecnico di Milano, Polo Regionale di Como, Via Anzani 52, 1-22100 Como, Italy

E-mail: hvkaenel@como.polimi. it

The formation and electronic structure of semiconductor superlattices is treated at an elementary level. Structural characterization by X-ray diffraction and Raman scattering is discussed on the example of strained-layer Si/Ge superlattices grown onto Si(OOl). Possible opto-electronic applications are outlined.

1. Introduction

1.1. Quantum Wells and Superlattices for Conduction Band States

We have seen that it is possible to control epitaxial film thickness to within fractions of a ML, e.g., by applying the method of RHEED oscillations to control the shutter actions in MBE. Precise thickness control is particularly important when artificial quantum heterostructures are to be grown, such as quantum wells (QWs) and superlattices (SLs).

A QW is simply a structure in which a classical particle is confined to stay between two parallel planar boundaries, while it is free to move in any direction along the bounding planes defining its width. Quantum mechanically, the same applies to the wave function describing the particle, although to a lesser extent, because of the tunnel effect which allows the wavefunction to penetrate slightly into the barrier material surrounding the well.

The effect does not depend on the physical nature of the particle or the associated wave, although the usual case considered is that of electrons or holes in semiconductor quantum wells. In order to act as a QW, a structure must be sufficiently thin, such that phase coherence can be maintained between a few bounces back and forth between the boundaries and interference becomes possible. For electrons this means that their de Broglie wavelength must be on the order of the thickness of the well.

Other examples are photons in so-called photonic crystals or phonons in laminar heterostructures.

Let us consider as an example a one-dimensional square quantum well with finite barriers of height V0 (Figure 1). The solution of this problem can be found in standard text books on quantum mechanics. Most importantly, it can be shown that such a QW always contains at least one bound state, no matter how

70

71

small the potential barrier V0. In the example we have assumed the QW width d and barrier height V0 to be such that there are just two bound states, a symmetric ground state n = 1 and an asymmetric excited state n = 2.

The problem at hand differs in one important aspect from the simpler problem of a square potential well with infinite barrier height: In contrast to the latter, the wavefunctions penetrate somewhat into the barrier regions, and they penetrate further the higher the energy of the bound state. This can easily be explained in terms of tunneling. There are exponentially decaying solutions of the Schrodinger equation in the barrier regions which have to be matched to the sinusoidal solutions inside the well. Since tunneling becomes more pronounced with decreasing tunneling barrier height, the excited states penetrate further into the barriers (see Figure 1).

E'

n=2 >

CJSX

i

i

'

i

v.

'-+> _ l 1 1 1__^

-ill ill Z Figure 1. Example of a quantum well of width d and height V0. The QW contains two bound states, the wave functions and energies of which are schematically indicated.

The significance of this can best be seen by considering two QWs coupled by a tunneling barrier in between (Figure 2). This time we present the solution of the quantum mechanical problem using realistic parameters typical for GaAs QWs and Gai.xAlxAs barriers. Each of the two QWs is 100 A wide, and the barrier is assumed to be 0.2 eV high. Instead of the free electron mass m0 we have chosen effective masses equal to the conduction band effective mass of GaAs. In other words, the example applies to coupled QWs for electrons in GaAs. The parameter for the barrier height, V0, is hence equal to the conduction band offset between GaAs and Gai.xAlxAs.

In the example we see that the uncoupled wells indeed contain two bound states E! and E2. In other words, the system of two uncoupled QWs contains two doubly degenerate energy levels. As the distance between the two gets smaller, the two sets of degenerate states start to interact because of the increasing overlap of the wavefunction tails penetrating into the barrier region: The

72

probability of tunneling between the two wells gets larger and larger with decreasing barrier thickness db.

As a result of this interaction, each degenerate state is split into a lower lying bonding state and an antibonding state at higher energy. This is the exact analog to the covalent interaction responsible for bonding in molecules, the simplest of which is the hydrogen molecular ion.

0.16

0.14

^ . 0-12

<a o.io g) 0.08 <E c 0.06 LU

.04

0.02

0

E4(2df

i i A 4 * * nin»o«E2(d)

E3(2d)

E

^

(2d) * 8 l i , « J 1 _ • • • • • •

' ^ (d )

20 60 0 20 40 60 80 100 db

Thickness of tunneling barrier [A] Figure 2. Bound states of two coupled square QWs as a function of the thickness db of the tunnel barrier between the two (see insert). The parameters are d = 100 A, mi* =m2* = 0.067 m0, V0 = 0.2 eV, and k± = 0.

From Figure 2 it is evident that the bonding-antibonding splitting is larger for the excited state. This is of course the result of the stronger wavefunction overlap expected from Figure 1. The limiting case of vanishing tunnel barrier width, db = 0, corresponds to a single QW with width 2d. It evidently contains four bound states in our example.

Clearly, our example can be generalized to the case of n coupled QWs separated by n - 1 tunneling barriers of equal width. In this case, each bound state of an isolated QW will be split into a so-called miniband containing n states.

Figure 3 shows a schematic real space representation of the resulting superlattice (SL). The minibands are indicated by shaded regions penetrating the whole SL. Note that the bandwidth of the minibands increases with their energy. This is again a consequence of the stronger wavefunction coupling as the effective height of the tunnel barrier decreases.

73

When the number of coupled wells becomes very large, it becomes appropriate to neglect the finite nature of the SL and to treat it as truly periodic. This facilitates the computation of the minibands, as now periodic boundary conditions and the Bloch theorem can be applied.

" : • ! • • .

-W..

! - * • ' •

3Er.:

; h i K 4 ^

• / * • * '

•?•

>

* .

r

• " ? • •

llfjS.'j.i'.V

E

• • • < . , . • :>

•••"[;;. ^

i!i.|i.'i--.-=i'S^B'

_̂

k

VI I

_̂_

• r r . t ; ; . - i . . . | t -

: , - ; •

rvr'

y ^

-f.rtr*-

• • • • . .

V " •"! *

M — ^ H " •

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

c

1.0 n

0.8

0.6

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 two-dimensional. 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.

-

\

1 '

V V * x

1

/ E 2

VE - = ! - •

1 0 L z 2 u 2

J L

0 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) - Ec(z)

"Ur E„(z)

(a)

: r - i j EcW

(b)

(c) EvW

Ev(z)

Ec(z)

(d) E„(z)

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

2n/d

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.

- • 3 - • 2 • 1 # 0 • 1-• 2-• 3-•

reciprocal space

79

(400)

10'

103

10*

10'

f 10°

10"'

10"2

103

10J

lo-'h

10"

jWX -0.06 -0.04 -0.02 0.00

e-eB (rad)

0.02

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

inelastic shift [cni ] 2 0 0 O lOO

inelastic shift [cni 2 0 0

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 u 1 - 1 — 1

UUL O 1 0 0 2 0 0 400

inelastic shift [cm1] 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 <100> 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 ^ J U J .

A \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 high-resolution 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 L. Esaki, R. Tsu9 IBM J. Res. Dev. 14, 61 (1970). 2. G. Bastard, in Wave mechanics applied to semiconductor heterostructures,

(les Editions de physique, Paris, 1992). 3. E. Carlino, C. Giannini, L. Tapfer, K.A. Mlder, H. von Klnel, Microsc.

Microanal Microstruct. 6, 473 (1995). 4. W. Bacsa, H. von Klnel, K.A. MSder, M. Ospelt, P. Wachter,

Superlattices and Microstructures 4, 717 (1988). 5. M. Rytov, Akust. Zk 2, 71 (1956). 6. F. Cerdeira, A. Pinczuk, J.C. Bean, B. Batlogg, B.A. Wilson, Appl Phys.

Lett 45, 1138(1984). 7. A. Fasolino, E. Molinari, Journal de Physique C5, 569 (1987). 8. P. Sutter, C. Schwarz, E. Miller, V. Zelezny, S. Goncalves-Conto, H. von

Klnel, Appl. Phys. Lett 65, 2220 (1994). 9. 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, France E-mail: psoukiassian@cea.fr

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, high frequency and 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 6H-SiC(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],

104 103 102 101 100 99 98

Binding Energy (eV)

C l s

Clean

./ \

d)

H < 1 ' 1 1 1 1 1 1 1 1 1 -

0.25 L O 2 ; \

j i

0.5 L 0 9 .-i 2 i\ j i

J \

e)

H 1 1 1 1 ' 1 1 1 1 1 1 1-

f)

286 285 284 283 282 281 280

Binding Energy (eV)

Figure 2. Si 2p core level for the 6H-SiC(0001) 3x3 a) clean surface and oxygen exposed at b) 0.25 L and c) 0.5 L in the surface sensitive mode (hv = 150 eV and emission angle and Ge = 30°). d), e), f): C Is core level for the same surfaces as in a), b) and c) (hv = 330 eV and emission angle and 6e=60°).

90

4. Initial SiO2/6H-SiC(0001) 3x3 Interface Formation

4.1. SiOySiC(0001) 3x3 Interface Formation by Direct Surface Oxidation at 500°C

We now look at the SiC oxidation process in a totally different regime, focusing on the initial insulator/SiC interface formation. CLPS is an especially appropriate tool to study silicon oxide growth on the 6H-SiC(0001) surface. We study the effect of a 1000 L of O2 exposure at 500°C on the 6H-SiC(0001) 3x3 surface reconstruction. Figure 3a shows the Si 2p recorded in a surface sensitive mode at a photon energy of hv = 150 eV and grazing emission angle. The three peaks A, B and C correspond to the clean Si 2p surface (SSI and SS2) and bulk shifted components respectively [20]. As can be seen from Figure 3a, the silicon dioxide (Si^+) is the dominant spectral feature [29]. However, we can also notice the presence of important amounts of sub-stoichiometric oxide components clearly indicating a non-abrupt SiO2/6H-SiC(0001) 3x3 interface [29]. These sub-stoichiometric oxide components have binding energies corresponding to the Si^+, Si^+ and Si+ oxidation states [29]. Looking in Figure 3b at the Si 2p in the bulk sensitive mode (hv = 300 eV and normal emission angle), one can see that the peak corresponding to S'fi+ is now the dominant spectral feature. However, such a large intensity does not result from S'fi+ only but from oxygen atoms bonded to both C and Si species, forming a carbonate mixed oxide product [29], as previously observed for the cubic 3C-SiC(100) 3x2 surface oxidation [28,33]. Further support for this assignment to the presence of mixed oxides can be found by looking in Figure 3c at the C Is core level for the same oxygen exposure with a C Is broadening by 26 %. Therefore, this indicates the strong involvement of carbon atoms in the oxidation process and, together with the Si 2p core level results, suggests the formation of carbonate oxidation products with Si02 presence at the top surface of the passivation layer [29]. Such a non-abrupt Si02/SiC interface would probably result in a high density of interface states.

4.2. SiOySiC Interface Formation by Oxidation of a Pre-deposited Si Thin Film on The 6H-SiC(0001) Surface

In order to improve the interfacial quality, we now proceed to the oxidation of a pre-deposited Si thin film on 6H-SiC(0001), using the same conditions as in IV-b). As can be seen from Figure 4a, the Si 2p core level recorded in the surface sensitive mode, silicon dioxide (Si^+) is the dominant oxide product, but the amount of sub-stoichiometric (Si^+, Si^+ and Si+) and mixed (Si-O-C) oxides has decreased dramatically [29]. Unlike the case of direct oxidation (Fig. 3b),

91

1 000 L 0 2 /6H-SiC (0001) 3x3 at 500°C

Relative Binding Energy (eV)

Figure 3. 1000 L of O2/6H-SiC(0001) 3x3 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 surface and bulk shifted components A, B, and C (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 surface and bulk shifted components A, B, and C (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 eV and grazing emission angle (surface sensitive mode).

92

1 000 L 0 2 /Si/6H-SiC (0001) at 500°C

Inte

nsity

£>

nsi

1

Surface ^

Si2p A / Si4+

Bulk

Si2p A /si4+ \

^ s i 2 A y . ;

Si-O-C.

\SJT A .-Clean!

a)

b)

A Clean ~

0.68 e V v ^ ^ _ ^

284 283 282 Binding Energy (eV)

8 7 6 5 4 3 2 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)

4H-SiC

7 6 5 4 3 2 1 0 - 1

Relative Binding Energy (eV)

Si 2p Bulk hv = 300eV,9 =60°

b)

4H-SiC

6H-SiC

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 6H-SiC(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 Si4 + , Si3 + , 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 02 / 6H- 4H-SiC(0001)-(3x3) at 500°C oxidation

Si 2p Surface hv=150eV,6 =30°

a)

>\ 4H-SiC

Si2+ i 6H-SiC

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

Si hv

2p Bulk = 300 eV, 9 e

/ s i 4 + )

= 60°

A ' 1

) J J )

/ S i " , s

b)

"*•

\4H-SiC

\ 6H-SiC Mean V V.

1 i • i

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

Figure 6. Comparison between the Si 2p core level spectra for the 4H-SiC(0001) 3x3 and 6H-SiC(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 Si4 + , Si3 + , Si2 + 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

O2/SiC(0001)-(3x3)

•I

12

10

8

6

4

2

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

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)

4H

6H

b)

Si 2p Bulk hv = 300eV,6e = 60°

J/KBISSA

4H / y\ i / 7B* A \sspk

A

i n / 1

6H \ \ j J Wf\

c)

d)

104 102 100 98 96 Binding Energy (eV)

104 102 100 98 96 Binding Energy (eV)

Figure 8. Si 2p core levels for clean 4H-SiC(0001) 3x3 and 6H-SiC(0001) 3x3 surfaces: a) 4H-SiC(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

O2/Si/6H-SiC(0001)-(3x3): post-oxidation

Si 2p Surface hv=150eV ~ 9 =30°

3000 L 0 2

at 650°C

2000 L 0 2

at 500°C

Si4+

4 1000 LO, \Si3+ Si2+

at500°C

•9A/ \\4Si-°-C

106 104 102 100 98 Rinrlirirr Ttr\(*rcT\r (t

Si 2p Bulk. hv = 300eV 9 =60°

3000 L 0 2

at 650°C

. 2000 L 0 2

/> at500°C

v..v~--..

1000 L 0 2

at 500°C

106 104 102 100 98 R i n H i n r r T?n<=»rrr\7 (p

06 104 102 100 98 106 104 102 100 91 Binding Energy (eV) Binding Energy (eV)

Figure 9. Thermal oxidation of a pre-deposited Si overlayer onto the 6H-SiC(0001) 3x3 surface followed by a post-oxidation sequence: Si 2p core level spectra for the Si/6H-SiC(0001) surface exposed to 1000 L O2 at 500°C and then re-exposed to a total of 2000 L O2 at 500°C and 3000 L O2 at 650°C recorded in a) surface sensitive mode (hv = 150 eV and emission angle 0e = 30°) and b) bulk sensitive mode (hv = 300 eV and emission angle Ge = 60°). The total oxides thickness (st) is 9 A for the Si/6H-SiC(0001) surface exposed to 1000L 0 2 at 500°C and 12 A for the 3000 L 0 2 at 650°C. The Si 2p core level peak decomposition for the O2/Si/6H-SiC(0001) surface show the 4

oxidation products Si4+, SP+, Si2+ 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).

98

6. Conclusions The atomic scale hexagonal SiC oxidation and oxide/SiC interface formation have been investigated by advanced experimental probes including atom resolved scanning tunneling microscopy, core level photoemission spectroscopy using synchrotron radiation and infrared absorption spectroscopy. For the Si-rich 6H-SiC(0001) 3x3 surface reconstruction, the results indicate a very high reactivity to oxygen (about 3 orders of magnitude larger than for Si surfaces). Oxygen atom insertion takes place below the surface. A carbon free Si02 ultrathin film (« 10A) could be grown onto a 6H-SiC(0001) 3x3 surface by low temperature oxidation at 500°C of a predeposited Si overlayer leading to abrupt Si02/6H-SiC interface formation. Most interestingly, this Si overlayer, which plays a central role in abrupt insulator/SiC interface formation, is found to grow in a novel and unexpected cubic 4x3 on the hexagonal 6H-SiC(0001) 3x3 surface reconstruction. The 6H polytype is found to be much more reactive to oxygen when compared to the 4H, likely due to the presence of two bulk domains for the latter mimiting oxygen insertion into the SiC lattice. All these findings show that "gentle" oxidation is a promising approach to SiC surface passivation.

Acknowledgments

This work was supported in part by THALES, by DGA (Paris, France), by the Institut Francais de Taipei (IFT) and by the Taiwanese National Science Council (NSC, Taipei). The author is grateful to his PhD student Fabrice Amy and collaborators Yeu-kuang Hwu and Christian Brylinski. The synchrotron radiation measurements have been conducted on the 3rd generation storage ring at the National Synchrotron Radiation Research Center (NSRRC), Hsinchu, Taiwan. The NSRRC staff is especially acknowledged for expert and outstanding technical assistance.

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17. P. Soukiassian, F. Semond, L. Douillard, A. Mayne, G. Dujardin, L. Pizzagalli, C. Joachim, Phys. Rev. Lett. 78, 907 (1997).

18. V. Derycke, P. Soukiassian, A. Mayne, G. Dujardin, J. Gautier, Phys. Rev. Ie«. 81,5868(1998).

19. V. Derycke, P. Soukiassian, A. Mayne, G. Dujardin, Surf. Sci. Lett. 446, LI 01 (2000).

20. 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).

21. F. Amy, P. Soukiassian, Y.K. Hwu, C. Brylinski, Surf. Sci. Lett. 464, L691 (2000).

22. J. Schardt, J. Bernhardt, U. Starke, K. Heinz, Phys. Rev. B 62, 10335 (2000).

23. L.I. Johansson, F. Owman, P. Martensson, Phys. Rev. B 53, 13793 (1996).

24. I. Forbeaux, J.M. Themlin, J.M. Debever, Phys. Rev. B 58, 16396 (1998).

25. M. Naitoh, J. Takami, S. Nishigaki and T. Toyama, Appl. Phys. Lett. 75, 650(1999).

26. V.V. Afanas'ev, A. Stesmans, M. Bassler, G. Pensl, M.J. Schulz, C.I. Harris, Appl. Phys.Lett. 68, 2141 (1996).

27. V.R. Vathulya, D.N. Wang, M.H. White, Appl. Phys. Lett. 73, 2161 (1998).

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28. V.M. Bermudez, Appl. Surf. Sci. 84, 45 (1995). 29. M.R. Chudoba, P. Soukiassian, C. Jaussaud, S. Dupont, Phys. Rev. B 51,

14300(1995). 30. F. Semond, L. Douillard, P. Soukiassian, D. Dunham, F. Amy, S. Rivillon,

Appl. Phys. Lett. 68, 2144 (1996). 31. F. Amy, P. Soukiassian, Y.K. Hwu, C. Brylinski, Appl. Phys. Lett. 75,

3360(1999). 32. L. Simon, L. Kubler, A. Ermolieff, T. Billon, Phys. Rev. B 60, 5673

(1999). 33. C. Virojanadara, L.I. Johansson, Surf. Sci. Lett. All, L145 (2001). 34. G.Y. Chung, C.C. Tin. J.R. Williams, K. McDonald, M. Di Ventra, S.T.

Pantelides, L.C. Feldman, R. A. Weller, Appl. Phys. Lett. 76, 1713 (2000). 35. P. Soukiassian, in Reference 2, 257 (1998). 36. G. Lucovsky, H. Niimi, in Reference 2, 447 (1998). 37. C. Radtke, I.J.R. Baumvol, J. Morais, F.C. Stedile, Appl. Phys. Lett. 78,

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Dujardin, Appl. Phys. Lett. 79, 767 (2001).

Al-W AMORPHOUS THIN FILMS

N. RADIC, T. CAR

Rudjer Boskovic Institute, Dept. Materials Physics, Bijenitka 54, Zagreb, Croatia E-mail: radic@rudjer.irb.hr

A. TONEJC,

Department of Physics, Faculty of Science, Bijenitka 32, Zagreb, Croatia

J. IVKOV

Institute of Physics, Bijenitka 46, Zagreb, Croatia

M. STUBICAR

Department of Physics, Faculty of Science, Bijenicka 32, Zagreb, Croatia

M. METIKOS-HUKOVIC

Faculty of Chemical Engineering and Technology, Savska 16, Zagreb, Croatia

Completely amorphous Al-W thin films were prepared by magnetron codeposition in the AUjWig - AU3W37 composition range. Thermal stability of the amorphous films was investigated by the continuous in situ electric resistance measurements up to 720°C, and a prolonged annealing in vacuum at high subcrystallization temperatures. The change in film structure caused by isochronal and isothermal heat treatments was determined by the XRD method. The crystallization temperatures were determined within a whole compositional range of amorphicity. A pronounced structural relaxation of the Al-W amorphous films was attributed to specific electronic structure of aluminium. Mechanical properties of the Al-W thin films were improved, while corrosive properties in a very aggressive environments and in the presence of aggressive CI" anions became only slightly degraded upon crystallization of the amorphous Al-W thin films.

1. Introduction

Amorphous materials are interesting due to their increasing applications in

modern technology [1]. Their possible applications are large, based on

characteristic properties such as chemical reactivity or inertia, electronic-

excitation phenomena, soft magnetic properties, superconductivity etc. New

nanocrystalline and quasicrystalline materials can be obtained by thermal

treatment of amorphous alloys. Binary amorphous alloys of refractory transition

metals and highly electrically conductive metals seem to be promising materials

according to the specific properties of their components [2], The aluminium-

101

102

tungsten system is particularly interesting due to its current and possible applications: In ULSI and VLSI technology of integrated microelectronics the aluminium is a material of choice for interconnect lines, while its diffusion into silicon and electromigration is suppressed by tungsten antidiffusion barriers [3-8]. The aluminium-tungsten amorphous films exhibit excellent corrosion resistance [9-14]. Finally, the amorphous Al-W alloys might be precursors for the production of Al-W quasicrystals, which bear promises of more interesting properties [15,16].

The thermal stability of amorphous alloys is essential for any application at sustained elevated temperature. Here we report a summarized results of the investigation of Al-W amorphous alloys thermal stability, and variation of mechanical, electrical and corrosion properties up to, and beyond, crystallization temperature.

2. Experimental

The Al-W thin films were prepared by simultaneous d.c. sputtering of both pure Al and pure W in a two source sputtering apparatus described in more details earlier [17,18]. Aluminium and tungsten atoms, sputtered from the respective cylindrical magnetron cathodes, are allowed to leave the respective discharge compartments only through the rectangular windows on the cylindrical anodes. In order to assure homogeneity of the deposited films, the substrate holder rotates during the deposition at 5 rpm. The working gas was argon at a pressure of 0.7 Pa in a constant flow mode. The area-averaged cathode current density was 3-7 mA/cm2 for tungsten, and 7-13 mA/cm2 for aluminium. In this way, the composition of deposited films could be varied over a fairly wide range. A typical rate of the film growth was about 0.1-0.3 nm/s, and the final film thickness was in the 1-3 urn range. Samples were deposited onto circular (1 cm dia) substrates, held at temperature in the range from LN2 to 400°C. Glass, monocrystalline silicon, and fused quartz were used for low and room temperature (RT) measurements, while samples for thermal stability and temperature variation of mechanical and corrosion properties investigations were deposited onto alumina ceramic or sapphire substrates.

The composition of the alloys was determined from the measurements of respective deposition rates of pure components, and checked by Proton Induced X-ray Emission (PIXE) on several samples deposited onto monocrystalline silicon. The estimated uncertainty of the film composition is about 10% relative to the minor component.

The structure of the prepared and thermally transformed films was investigated by the XRD method using a Philips PW 1820 vertical goniometer with monochromatized CuKa-radiation.

103

The electron microscopy was performed by JEOL JSM 5800 instrument, without gold coating of the samples.

The electric resistivity of the Al-W films was measured by the standard a.c. technique. The samples deposited on glass, with silver flakes partially buried under the film as electric contacts, were used for measurements below room temperature. A high-temperature electrical measurements were performed by using tungsten rod electrodes spring-loaded onto the film surface, in order to avoid spurious effects due to possible reactions between Al-W alloys and contact material. Isochronal measurements of thermal stability were performed in vacuum of 10"^ Pa, while isothermal annealings were performed in tubular furnace filled with a protective argon gas (99.999%) in a continuous flow. Microhardness of the films deposited onto sapphire substrates was measured by the Leitz (Wetzlar, Germany) Miniload II apparatus. A loading pressure of the Vickers diamond pyramide was 0.25 N, and a minimum of twenty indentation traces and their measurements on each sample were made. The sequential isothermal annealing of the Al-W samples was performed in vacuum better than

10"3 Pa, at selected temperatures from the 150-1000°C range, in order to follow microhardness and structural changes.

Corrosion resistance of amorphous Al-W alloys was examined by the electrochemical methods in 1 M deaerated hydrochloride acid solution and artificial saliva solution (ASS). A standard two-compartment electrochemical cell with platinum gauze (for chloride solution) and graphite rods (for artificial saliva) as the counter-electrodes, and a saturated calomel as the reference electrode was used. The current-potential characteristics were obtained, using a EG&G PAR Model 273 potentiostat, controlled by a PC.

3. Results

3.1. Composition and Structure of The As-deposited Al-W Thin Films

The Al-W films with composition in the range from A188W12 to Al5oW5o were prepared. Completely amorphous films were obtained in the A182W,8 - AU3W37 composition range, on various substrates held at room temperature during deposition. No diffraction lines of pure Al, W or stable intermetallic compounds (A112W, AI5W, AI4W) were observed. Further on aluminium-rich side, two-phase films were obtained with pure aluminium mixed with amorphous Al-W phase. Beyond the amorphicity range towards a tungsten-rich side, a mixture of predominantly crystalline (a-W, (3-W) and remnant amorphous phase first appears, while further increase of tungsten content results in a crystalline solid solution of aluminium in tungsten (Fig. 1).

Scanning electron microscopy reveals a rather featureless appearence of Al-W amorphous film deposited onto monocrystalline silicon (Fig. 2a), while a

104

nodular surface morphology is observed at the alumina grain sides covered by Al-W film deposited under oblique angle (Fig. 2b). However, no significant difference in film composition was detected by the electron microprobe across the area containing several alumina grains, and the XRD of the respective films confirmed their XRD-amorphous structure.

Figure 1. The XRD patterns of Al-W films deposited onto alumina ceramic substrates, a) AlsoW2o, b) AI75W25, c) AI67W33, and d) AI50W50. Diffraction lines from substrate and stainless steel holder are marked (s) and (h), respectively.

105

(a

v . *

s H H H H B I

(b

Figure 2. a) A SEM micrograph of the amorphous AI75W25 film at the fracture, b) A SEM micrograph of the amorphous Al§oW2§ film deposited onto alumina ceramic.

3.2, Electrical Properties

Electric resistivity of the amorphous Al-W thin 11ms is an important parameter for their eventual applications in integrated microelectronics. Electric resistivity variation with temperature allows to test the amorphicity of the prepared samples, and to follow a phase transformations during their isochronal and isothermal annealing.

The electric resisitivity of the Al-W amorphous films is about 200 fiOcm at both ends of the amorphicity compositional range, and exhibits a maximum of about 500 fxOcm for Al8oW2o amorphous films. With these resistivity values, the Al-W thin films are within 1000 fiOcm limit set as an acceptable antidiffiision barrier resistivity by considering a Joule heat dissipation in a model calculation [3].

106

A temperature coefficient of the electric resistivity (TCR) at room temperature allows easy inspection of the one-phase metallic samples amorphicity: negative value of TCR is always a sign of amorphous phase. However, due to a possibility of various spatial distributions, in case of mixed phases this criterion should be corroborated by another method of structure determination (XRD). In Fig. 3, the values of thermal coefficient of the electric resistivity at room temperature for the Al-W films and pure aluminium and tungsten films prepared by the same method are shown. As seen, the TCR is negative in a whole composition range of the prepared Al-W films, although the XRD method proved that beyond the limits of complete XRD-amorphicity the Al-W films consist of phase mixture. This indicates a crystalline phase imbedded into an amorphous matrix as a structure of the Al-W samples consisting of two or more phases.

15

-T 101 *:

T o

H 5 p Q.

T3 Q. ^ 0

-5

-

k^

^ \

*-* \

\ \ *̂

o

I 1 I I I

'%

1 ;' / /

/ / / /

/ i i i

t i i i i

•

0 10 20 30 40 50 60 70 80 90 100

x (at%)

Figure 3. Variation of the temperature coefficient of electric resistivity at room temperature with the AlxWi.x film composition. Open symbols belong to the samples deposited onto alumina ceramic, filled symbols to the samples deposited onto glass.

3.3. Thermal Stability of Amorphous Al-W Films

Variation of the electric resistance of prepared films with temperature is an excellent indicator of the occurence of phase transformations in the film upon cooling and heating [19, 20]. An example is shown in Fig. 4, where the results for Al7gW22 on glass below RT, and A178W22 on alumina above RT, are normalized to the resistance value at room temperature and shown together.

107

1,5

1,0

Q.

"3.

0,5

0,0 0 200 400 600 800 1000

T(K)

Figure 4. Variation of the electric resistivity (normalized to the room temperature value) of the AI78W22 films in the 1.5-1000 K temperature intervals.

The main phase transition below room temperature is transformation to the superconducting state at about 1.8 K, preceded by the still unresolved resistance oscillations. Up to about 150°C, the thermal coefficient of electric resistivity retains a negative value. Above that temperature, the electric resistance exhibits variation which depends upon the heating rate: for slow heating rates (1-4 K/min), the resistance (steadily) increases up to some temperature in the 500-600°C range, at which a steep increase of electric resistance indicates a phase transformation. At higher heating rates, after an upward turn above 150°C, the resistance again turns downward at some temperature, and exhibit the same sudden change in the 500-600°C interval. To ensure that no gross phase transformation in amorphous films occurs in the 100-500°C interval, the up-down heating cycles up to the selected temperatures were performed, with film structure XRD-checked after each cycle. It was found that the thermal coefficient of the electric resistivity is strongly negative up to about 500°C, and the corresponding XRD patterns (Fig. 5) reveal that the film is still amorphous at that temperature.

Al W

dT/dt = 2 K/min

V

0 5 J ( K ) 10 15

108

+-»

X>

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 tungsten-rich samples, a pure tungsten separates alongside A14W. The temperatures of crystallization are shown in Fig. 6, for the whole Al-W amorphicity range.

700

650

600

-"

:

<B

lin

fl - t» : & . o

:

- , i

l •

! 00° * •

1 • ***

1*

lAmorphous

,1 I i 1 i

Al W M,100-X™X

Crystalline

1 , 1 , 1 , 1 , 1 , 1 ,

o 2- 550

500

450

400 Figure 6. Crystallization temperatures of the amorphous Al-W alloys. Filled symbols correspond to the results obtained with 2 K/min heating rate, while open symbols mark the results obtained with 10 K/min heating rate.

3.3.2. Isothermal Annealing at High Subcrystallization Temperatures

The Al-W amorphous film stability at high subcrystallization temperatures is important for applications at sustained elevated temperatures. The amorphous AI78W22 thin films, which crystallize at about 570°C, were investigated as a representative case for other amorphous A-W films with similar composition. The resistivity of the samples was monitored during isothermal annealing at

110

530°C, 540°C, and 550°C, respectively, for several tens of hours. At selected periods, the XRD examination of annealed film structure was performed. The temperature coefficient of electrical resistance as a function of annealing time for the A178W22 thin film held at 530°C for accumulated time of 66 hrs is given in Fig. 7. As seen, the formation and growth of crystalline phase resulted in a decrease of negative magnitude of TCR, which finally acquired a positive value.

1

T

(10

H XJ dpd/

y—

1

0

-1

-2

-3

A

-

_ -•

-D -•

D

•

J 1—1—

•

•

i i ,

20°C

• D

• •

, i I i i

D

•

. i 1

•

•

• •

D

•

a

•

530°C

• i i i i i

a

•

• • 1 • t i i

0 10 20 30 40 50 60 70

t(h) Figure 7. Variation of the temperature coefficient of the electric resistivity with the accumulated time of annealing at 530°C in vacuum for AI78W22 film. The TCR values at 530°C is marked by filled symbols.

The XRD examination confirms that a partial phase transformation occured, with AI4W as a dominant crystallization product. However, the asymptotic nature of the electric resistivity variation with a prolonged isothermal heating at subcrystallization temperatures indicates that the partial crystallization reaches a stationary state within several tens of hours, and no complete crystallization is achievable by prolonged annealing at that temperature.

3.3.3. Structural Relaxation

Several factors were found to be important in the magnitude of electrical resistance variation upon first heating of the Al-W amorphous films: heating rate (discussed above), alloy composition, substrate quality and substrate temperature [21]. The relaxation was more pronounced in aluminium-rich alloys, which might be due to the specific electronic structure of aluminium. In order to check

111

that assumption, a temperature dependence of the electric resistivity was measured on amorphous Al-Ti and Cu-Ti films prepared by the same method, and the results are shown in Fig. 8.

1.4

1.2

1.0

*•• 0 . 8 a

0.6

0.4

0.2

0.0

-

-

-

o oimm •aaxja a o I

AITi

a D a D D

I i

200 400

T(°C)

600 800

Figure 8. A temperature dependence of the resistivity for the amorphous Al-Ti and Cu-Ti thin films, prepared by the same method as the Al-W thin films. Cathode currents ratio for both samples was 1:3, while the heating rate was 4 K/min.

As expected, the relaxation is minimized by the increase of substrate temperature. However, the corresponding X-ray diffraction patterns of the A178W22 films revealed the formation of new phase in samples deposited at 400°C, which was tentatively identified as whiskers atop the amorphous films.

3.3.4. Whisker Phase

The AI75W25 the films deposited onto sapphire substrates held at 250-400°C (about 3 urn thick) exhibited the XRD pattern that corresponds to a two-phase structure (Fig. 9).

112

3518

f. counts 1

38BB

^®Hm¥^

deposition terapeipature : 400 C

leposition tine! 158 nin 3

^*^m^ w

— r — 28 4B 6 8 1&2Q1 6 8

Figure 9, The XRD pattern of the Al7sW22 film deposited onto sapphire substrate held at 400°C. Lines marked 1-3 correspond to the same ieterplanar distance (0383 nm) in three orders of diffraction.

The SEM examination of the samples with most pronounced two»phase XRD pattern, revealed a dense away/population of whiskers at the film surface (Fig. 10).

Figure l i . A SEM micuwph of use Al,>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 whisker-related 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 • 1 1 1 • 1 1 1 • 1 1 1 200 400 600 800 1000 1200 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].

1000 E 1

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 Al-W 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. Rapidly Quenched and Metastable Materials, Part I, ed. by T. Masumoto, K.Hashimoto (Elsevier Science, Tokyo 1994).

2. N. Radie, D. Gracin, Fizika A4, 233 (1995). 3. 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).

4. M. Tsukada, S.I. Ohfuji, J. Vac. Sci., Technol. A 13, 2525 (1995). 5. D.B. Bergstrom, I. Petrov, L.H. Allen, J.F. Greene, J. Appl. Phys. 82,

201(1997). 6. D.B. Bergstrom, I. Petrov, J.F. Greene, J. Appl. Phys. 82, 2312 (1997). 7. M.B. Takeyama, A. Noya, Jpn. J. Appl. Phys. Part 1, 38, 2116 (1999).

118

8. R. Pantel, J. Torres, P. Paniez, G. Auvert, Microelectron. Eng. 50, 284 (2000).

9. K. Hashimoto, N. Kumagai, H. Yoshioka, H. Habazaki, A. Kawashima, K. Asami, B.-P. Zhang, Mater. Sci., Eng. A133, 22 (1991).

10. H. Habazaki, K. Tahakira, S. Yamaguchi, K. Hashimoto, J. Dabek, S. Mrowec, M. Danielewski, Mater. Sci and Eng. A181/182, 1099 (1994)

11. 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).

12. A.Wolowik, M. Janik-Czachor, Mater. Sci. Eng. A267, 301 (1999). 13. M. MetikoS-Hukovic, N. Radie, Z. GrubaC, A. Tonejc, Electrochimica

Acta 47, 2387 (2002). 14. L. Iglesias-Rubianes, P. Skeldon, T.E. Thompson, H. Habazaki, K.

Shimizu, Corrosion Sci. 44, 751 (2002). 15. P.A. Bancel, P.A. Heinay, Phys. Rev. B 33, 7917 (1986). 16. H.S. Chen, J.C. Phillips, P. Villars, A.R. Kortan, A. Inoue, Phys. Rev. B

35,9326(1987). 17. N. Radie, B. Grzeta, D. Gracin, T. Car, Thin Solid Films 228, 225 (1993). 18. T. Car, N. Radie, Thin Solid Films 293, 78 (1997). 19. J. Ivkov, N. Radie, Solid State Commun. 106,273 (1998). 20. T. Car, N. Radic, J. Ivkov, E. Babic\ A. Tonejc, Appl. Phys. A 68, 69

(1999). 21. J. Ivkov, N. Radie, A. Tonejc, T. Car, J. Non-Cryst. Solids 319, 240

(2003). 22. M. Stubicar, A. Tonejc, N. Radic, Vacuum 61, 309 (2001). 23. N. Kimura, M. Nakano, K. Sato, Jpn. J. Appl. Phys. Part 1 37,

1020(1998). 24. B. X. Liu, Mater. Lett. 5, 322 (1987). 25. Y.F. Ouyang, X.P. Zhong, W.M. Wu, Sci China Ser. A 43, 184 (2000). 26. J.H. Mooij, Phys. Status Solidi a 17, 521 (1973). 27. A.R. Miedema, Z. Metalkd. 70,345 (1979). 28. K.J.H. Buschow, N.M. Beekmans, Solid State Commun. 35, 233(1980). 29. A.W. Weeber, J. Phys. F: Met. Phys. 17, 809 (1987). 30. J. Ivkov, N. Radie, A. Tonejc, 12th International Conference on Thin Films,

(Bratislava, Slovakia, September 15-20, Book of Abstracts, 164). 31. L. Succo, J. Esposito, M. Cleeves, S. Whitney, R.E. Lionetti, C.E.

Wickersham, jr., J. Vac. Sci. Technol. A 7, 814 (1989). 32. K. Hinode, Y. Homma, Y. Sasaki, J. Vac. Sci. Technol.A 14,2570 (1996). 33. H. Takatsuija, K. Tsujimoto, K. Kuroda, H. Saka, Thin Solid Films 343,

461 (1999). 34. 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: exper2@itmo.by

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 zinc-selenide 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 fed from the periphery.

4

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

< y* •,

t ^

t,

1

^ 2 l6*fefrM t,

r

Figure 1. Diagram of a reactor: 1 - reactor wall; 2,3 - inlet holes; 4 - hole for gas outlet.

The deposition of zinc selenide layer in the reactor with heated walls is a result of the reaction

Zn (C2H5)2 + H2Se = ZnSe (solid) +C2H6 (1)

The gaseous deposition of ZnSe is described by the Navier-Stokes conservation equations:

d dy/

dx\ dy

d( dy/\ CO

dy dx ,

d_

dx

( dy/)

dy

dx

d'

\_dy_

p dx

d2([i0) d2(juco) dp = Q - ( 2 )

dx2 dy2 fy

+ •

dy H

\

dy^\

dx

dy

d_

dx

p dy J + cv = 0;

f \ X dy/

C dx \ P

d_

dy

X dy/

~lT"~dy \ P

(3)

:0;(4)

J

121

d_

dx

( C

V

dy/) dy dy

C dy/

i dx ) dx 5 <PD &

i dx dy PD

dy/ > = 0, (5)

where v|/ - the stream function, co - the vortex strength, x and y - the longitudinal and transverse coordinates, Q - the weight concentration of the i-th components, g - the acceleration due to gravity, \i - viscosity, p - density, Cp - the heat capacity, X - the thermal conductivity, H - enthalpy, and D( - the diffusivity of the i-th component.

This system of differential equations was solved by the method of finite differences according to [1], with allowance made for the dependence of the physical properties of the gas mixture (such as the diffusivity, heat capacity, thermal conductivity, viscosity, and density) on its temperature, pressure, and chemical composition [2].

Let us consider the boundary conditions that should be satisfied by equations (2)-(5) describing two-dimensional flows. Then, the stream function \\i and the vortex strength a> at the inlet of the reactor in the central nozzle (from the axis to yi) are

y\ y/ = \pudy,

0

dv CO =

dx

du

~dy (6)

in the peripheral nozzle (from y2 to y3), they are

y/= \ pudy + y/[y\ dv (7) _ du

1 dx dy ' where u and v are the longitudinal and transverse components of the gas-mixture velocity, y, is the halfwidth of the central nozzle, (y3-y2) is the width of the peripheral nozzle, and h is the halfwidth of the reactor. At the outlet and in the symmetry plane of the reactor, 6F / dx = 0 and 5F / dy = 0 , respectively (here F are the desired functions). The stream function on the lower wall from yi to y2 is v|/=i|/(yi), from y3 to h it is v|/=p0u0h, and on the side and upper walls it is v|/=p0u0h (the subscript 0 refers to the reactor inlet).

Calculations were performed for the range of the technological parameters (temperatures, pressures, and concentrations) that are usual for experiments on the deposition of zinc selenide layers [3-5].

The boundary conditions on concentrations are based on the assumption that reaction (1) proceeds in the diffusion regime. The absence of kinetic restrictions is supported by the weak temperature dependence of the growth rate [3] and its high values obtained under low pressures [5].

The boundary conditions for the concentrations on the channel walls are

122

DLdC1 = £LE2. (8) Mj dy M2 dy

C,=0; (9)

DLdC1=DLdC1^ (1Q)

where Mi is the molecular weight of the i-th component; 1 stands for Zn(C2H5)2, 2 for H2Se, and 3 for C2H6.

The limiting component is ZDE, because hydrogen selenide is fed in threefold excess. The equality of the fluxes of ZDE and hydrogen selenide in (8) points to the condition of stoichiometry of the produced zinc selenide.

The boundary conditions on the inlet and outlet concentrations and temperatures are represented as dependencies (11)-(12).

At the inlet of the reactor (x=0): in the central nozzle (from the axis to yO Ci=C0i, C2=0, T=T0; in the peripheral nozzle (from yi to y3) Ci=0, C2=C02, T=T0; on the front wall (from y! to y2) T=T0; on the front wall (from y3 to h)

T = T3+(TS-T3)^-. (11)

At the outlet of the reactor (x = L): 8T n dC>

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<L)

T=TS. (13) By solving the system of equations (2)-(5) with boundary conditions (6)-

(13), we obtained the distribution of the rate of deposition of zinc selenide layers on the reactor walls.

123

3. Results and Discussions

The dependence of the rate of layer growth on the channel V (x) length has a maximum in all considered cases. This shape of the distribution V (x) is typical for reactors with separate feeding of the reagents [6]. The ascending branch of the curve V (x) describes the mixing of separately fed reagents, whereas the descending branch describes the depletion of the gas mixture.

Zinc diethyl and hydrogen selenide are fed in local streams through the central nozzle and from the periphery, respectively. This leads first to the increase of the deposition rate (from the reactor inlet) due to the mixing of the streams and then to its decrease, which is due to the depletion of the gas mixture caused by the deposition of zinc selenide. This circumstance ensures a maximum in the distribution of the rate of deposition of zinc selenide along the reactor (Fig. 2 and Fig. 3). The location of the maximum is determined by technological parameters. The rate of growth of a zinc-selenide layer was calculated by the flux of zinc diethyl mass on the surface of deposition.

Figure 2 shows the influence of the gas-mixture velocity at the reactor inlet, Uo, on the rate of deposition. When u0 increases with the corresponding shift of the ascending branch and increases on the descending branch. For the measure of uniformity of the layer growth, we took the quantity

1̂ V » max" » min/'v * max""" * min^

where Vmax and Vmin are the maximal and minimal values of the rate of deposition in the part of the reactor under consideration. The parameter r\ is equal to the maximal deviation from the mean deposition rate.

s c E 3 .

S a •w

s. CI •o

<*• e « « as

: 1 . 6 T

---1.2-

; 0.8-f

' : -0.4-|

0

x. cm

Figure 2. Rate of ZnSe deposition along the reactor as a function of the velocity of the gas mixture for Ts=623 K, T=523 K, h=5 cm: (1) u„=0.37; (2) 0.74; (3) 1.48 cm/s.

124

0 10 20 30 40 x, cm

Figure 3. Rate of ZnSe deposition along the reactor as the function of h and u0 for Ts=623 K, T=523 K: uo=0.37 cm/s: (1) h=5 cm, (2) 2 cm, (3) 1 cm; u0=1.48 cm/s: (4) h=5 cm, (5) 2 cm, (6) 1 cm.

In the region of the reactor lying at a distance of 15-45 cm from the inlet and under the conditions specified for Fig. 2, the minimal nonuniformity is rj = 0.127 for u0 = 0.74 cm/s (r\ =0.387 for u0 = 0.37 cm/s, and r| =0.443 for u0 =1.48 cm/s).

Comparison of the curves in Fig. 3 shows that when h decreases, the rate of deposition of zinc selenide increases, the maximum is shifted toward the reactor inlet, and the uniformity of distribution of growth rate decreases (curves 1-3 in Fig. 3: r|i = 0.387, r\2 = 0.843, and r)3 = 0.874). By varying the velocity of the gas mixture u0, we can (for small h) obtain layers of zinc selenide with fair uniformity and a high rate of deposition (curves 4-6 in Fig. 3: r|4 = 0.443, r|5 = 0.136, and % = 0.156).

The process of deposition is conducted under low pressure in order to avoid the reaction between zinc diethyl and hydrogen selenide in the gas phase. It follows from calculations that the variation of pressure from 20 to 200 Pa with a constant mass flux and composition of the gas mixture does not affect the rate of deposition much.

The variation of the inlet temperature of the gas mixture from 373 K to 523 K and of the temperature at the surface of deposition from 523 K to 623 K has a slight effect on the rate of deposition.

The narrowing of the outlet hole from 0.34h to 0.06h leaves the deposition rate practically unchanged, whereas the shift of the peripheral inlet hole from y2

= 0.8 to y2 = 0.4 increases the rate of zinc selenide deposition.

125

4. Conclusions

The distribution of the deposition rate along the reactor has a maximum, whose location is determined by technological parameters. By varying the feeding rate of the gas mixture and the temperature in the deposition zone, we can ensure a fair uniformity of the deposited layer of zinc selenide.

The rate of deposition of zinc selenide increases considerably when the width of the reactor decreases, and remains practically unchanged when the outlet hole is narrowed. The shift of the peripheral inlet hole increases the rate of zinc selenide deposition.

The variation of the inlet temperature of the gas mixture and of the temperature at the surface of deposition have a slight effect on the rate of deposition.

References

1. A.D. Gosman, W.M. Pun, Heat and Mass Transfer in Recirculating Flows, (Academic Press, London and New York, 1972), p. 93.

2. R.C. Reid, G. Prausnitz, T.K. Sherwood, Properties of Gases and Liquids, (New York, McGraw-Hill, 1971).

3. P. Blanconnier, M. Cerclet, Thin Solid Films 55,375 (1978). 4. W. Stutins, Appl. Phys. Lett. 38, 656 (1978).

5. B.V. Zhuk, V.K. Khamylov, et al., Poverkhnost 7, 112 (1982).

6. V.G. Minkina, Izv. Akad. Nauk SSSR, Seriya Khimicheskaya 7, 1625 (1996).

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III. CHARACTERIZATION TECHNIQUES

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THIN FILMS ANALYSIS USING PHOTOELECTRON SPECTROSCOPY

S.K. KULKARNI

Department of Physics, University ofPune, Pune - 411 007, India E-mail: skk@physics. unipune. ernet. in

Photoelectron spectroscopy has substantially contributed to the understanding of electronic structure of atoms, molecules and solids. Photoelectron spectroscopy using conventional X-ray or UV sources has proved to be advantageous by revealing the Electron Distribution Curves (EDC) of Solids. However in the last twenty years or so, the use of synchrotron radiation has become a more powerful source for carrying out photoelectron spectroscopy, with additional advantages due to its tunability to desired energy and polarization. This becomes very useful to understand electronic structure in thin films. In this article an overview of various analysis aspects of photoelectron spectroscopy using conventional and synchrotron radiation source are given.

1. Introduction

Surface in general differs from bulk material due to loss of periodicity in the direction normal to the surface. Surface atoms relax or move in such a way that surface structure changes dramatically. Additionally adsorption of the atoms from the ambient or diffusion or segregation from bulk to surface may take place changing the stoichiometry at the surface. In thin films surface to bulk atom ratio is quite large. Therefore the change in the stoichiometry or structure may result into changes in the electronic structure. Thin film properties are controlled by surfaces and need to be investigated using surface sensitive techniques. Fig.l shows a list of few techniques often used currently in structure, composition and electronic structure analysis of thin films. Amongst these techniques photoelectron spectroscopy is a technique which can be most suitable for thin films composition and electronic structure analysis. It is an atom specific technique. By suitably collecting the data it is useful even to determine the local structure (bonding distance and coordination) at particular atom. Photoelectron spectroscopy can be performed using X-rays and ultra violet rays or synchrotron radiation. There are many examples of thin films analysis using photoelectron Spectroscopy. Thin films growth [1-4], core and valence band studies of variety of materials [5-7], quantum dots and quantum well investigations [8-11], self assembled quantum dots [12], surface passivation [13-15], semiconductor heterojunctions[16-18], analysis of magnetic thin films [19-21] etc. form some of the topics in which photoelectron spectroscopy has

129

130 Thin Films / Surfaces

Structure Electronic Composition Structure

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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

Vacuum Level

EF=0

hv

Analyzer

1'

Ev measured

"sample

T analyzer

E,

hc = E e + EK+^

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 100-200 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.

<1.

a, V 4> U fa s

3 V

20

10

5

T 1

2r

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*y s A

< o

I 10 100

Electron Energy (eV)

1000

Figure 3. The dependence of mean free path on electron energy.

134

2.3. Experimental Set Up

In Fig. 4a simple schematic diagram is given to illustrate basic requirements of photoelectron spectroscopy set-up. Usually Al Ka or Mg Ka twin anodes with photon energies 1486.6 eV and 1253.6 eV respectively are used as source of X-rays. For valence band studies He I and He II having photon energies 21.2 and 40.8 eV are available. He I and He II are obtained under certain pressure using a UV discharge lamp with helium gas flowing in it. The lamp is differentially pumped because surface studies require Ultra High Vacuum (108 - 10"10 mbar) environment. Photoelectrons are analyzed using Concentric Hemispherical Analyzer (CHA) or double pass Cylindrical Mirror Analyzer (CMA). Analyzer is controlled using spectrometer control unit (SCU) so that electrons of certain energies pass through them, detected and amplified using channeltron or channel plate. Amplified signal after suitable noise filtering is used as an input for an X-Y recorder or computer. An X-Y plot or a spectrum of intensity versus electron kinetic/binding energy can be obtained. One can make use of readily available data books [23-25] to identify the elements. Fig. 5 shows a typical X-ray photoelectron spectrum of Bi2Sr2CaCu208 sample obtained using Al Ka source. Besides the prominent core levels due to composing elements, as marked on the figure, there also occur other peaks large in number. They content lot of information regarding the electronic structure of the material.

2.4. Information Content of X-ray Photoelectron Spectra

Photoelectron spectra may have following ingredients useful for understanding properties of surfaces and materials. i) Chemical shift ii) Spin - orbit splitting iii) Multiplet splitting iv) Plasmon loss v) Satellites vi) Valence band vii) Auger peaks Here a brief discussion on occurrence of above features is given. i) Chemical Shift - As mentioned earlier, core electrons of atoms in solids are very sensitive to their environment. Whenever there is a charge transfer between valence electrons of different atoms, core level electrons also adjust to these changes. Thus, there are changes in the binding energies of electrons in atoms. These changes can be observed by analyzing the kinetic energies of photoelectrons. Many examples can be found even in the standard books [23-25],

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Bi2Sr2C aCu208/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.

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-©--©-

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)

Where rop is plasmon frequency, n is number of electrons per cm3 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 cross-sections 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, ELn and ELm 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 K̂ , 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,,,

"T iK^

" A I I W

" M g K ^

11NaKai^

40ZrM5

9YM^

2He II

2 H e l

1 0Nel

1 8Arl

Tunability of photon

Energy (eV)

8048.00

4511.00

1487.00

1254.00

1041.00

151.40

132.30

40.80

21.22

16.85

11.83

Natural width (eV)

energy is achievable using synchrotron

2.5

1.4

0.9

0.8

0.7

0.77

0.47

0.1

0.1

0.1

0.1

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

21

\ hv=1486.6eV

20 19 43 42

As3d

41 Bintlfng Energy (eV) -

Ga3d As3d AS CSOj^.^) ,0-TOeV

B

21 20 19 43 * 2 41

Binding 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 - • 3p63d4 (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

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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 Si02 [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^C^I^O 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 of - 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 H2S04:H202:H20 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. Ga203 and As203 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

23.5

20-

16

2- 12

0-1

t ~* 'I

10 JO 30

T i m ( • •<)

etchant

. H2Sq4:H202:H20-.

BOO 830 860 Wave length (ran)-

890

Figure 1. PL spectra of p-GaAs (100) surface after etching ( a ) 'as received', ( b ) after etching for 5 sec, ( c ) 10 sec, ( d ) 25 sec. and ( e ) 35 sec Inset: ratio of PL intensity of etched surface (IP) with respect to the "as received' sample (Ic).

B. E.(eV) B.E. (eV)

Figure 2. Ga 3d and As 3d core level photoemission spectra for the GaAs (100) after Chemical cleaning by a solution of H2S04 : H202: H20 :: 4:1:1 before passivation (A) 'as received', (B) after 30 sec cleaning, (C) after 60 sec, (D) after 90 sec. and (E) after 120 sec.

153

structureless indicating the presence of suboxides. However due to limited resolution of our instrument (0.5eV) the suboxide could not resolved. Besides, it should be noted that inherent width of 3d levels also is large. After chemical etching as in Fig.l Ga 3d peak progressively shifts to smaller binding energy until 90 seconds, and then remains at the same position. However even at 90 second or 120 second etching an oxide component is present. Although there are no remarkable peak shifts in case of As 3d, the non-oxidic As and oxide peaks change in the intensity ratio quite dramatically. On the as received sample, arsenic oxide component is larger than non-oxide arsenic component is but on the etched samples reverse is the case with non-oxidic As component becoming quite large. Small energy shifts are however present even for non-oxide As component.

Earlier PC15, PC13, and P2S5 have been reported to have been used for GaAs passivation [24, 25]. However there are some difference in the chemical treatments compared to what we report have [26]. It is not possible to directly correlate PL results with XPS but only to some limited extent as transfer of samples from solutions takes longer period in case of XPS analyzed samples (typically 2-3 hours before vacuum is achieved). PL experiments are carried out immediately (few minutes) after chemical treatments and drying. Nevertheless one can understand to some extent the effects of various chemical treatments. In. Fig. 3, PL spectra after PCI5 treatment are showed. The inset shows the variation in intensity with respect to 'as received' GaAs (100) samples for different times. It can be seen that after a long treatment the sample becomes worst than the original sample indicating that defect state density has increased, trapping the charge carriers that should be available for luminescence. Interestingly as showed in Fig.4, Ga 3d and As 3d spectra suggest that oxides of Ga and As are minimal at the surface and even GaP layer should be present at the surface. Obviously PC15 must be incorporating some chlorine in the surface which may attack GaAs Surface and create defect centers responsible for reduced PL intensity. Energy Dispersive Analysis of X-ray (EDAX) analysis indeed showed that ~ 17 % chlorine was present as the sample surface.

However use of P2S5 is better than that of PC15. As showed in Fig.5, photoluminescence intensity increased by about five times in first two minutes and then decreased. However it is better than 'as received sample'. The reduction in PL intensity has been attributed to thick phosphorous film formed on the surface as well as to the oxidation to some extent. In this case we studied core levels of GaAs using Al Ka (1486.6eV). In Fig.6 one can see the changes due to P2S5 treatment. Surdu-Bob et al [27] have discussed some of the advantages, like sharp line shape of 2P3/2 peak in the context of GaAs passivation. Core level spectra in Fig. 6 show the effective removal of oxides

154

820 840 860

Wavelength Inm)

900

Figure 3. PL spectra of n-GaAs (100) surface after PCI5 treatment ( a ) 'as received', ( b ) after treatment for 1 min. and ( c ) 2 min. Inset: ratio of PL intensity of PCI5 treated surface (IP) with respect to the "as received' sample (Ic).

aE(eV)

Figure 4. Ga 3d and As 3d core level spectra of GaAs(lOO) surface after PC15 treatment for various times (A) 15 sec, (B) 30 sec, (C) 45 sec. and (D) 1 min.

155

£ SO

800 820 840 860 880 900

• W»venumber (nm) FigureS. PL spectra of n-GaAs (100) surface after P2S5 treatment ( a ) 'as received', ( b ) after treatment for 1 min., ( c ) 2 min. and ( d ) 2 min 30 sec. Inset: ratio of PL intensity of P2S5 treated surface (Ip) with respect to the "as received' sample (Ic).

na 1116 1119 1122 uw 1321 ms 1329

— — Binding Energy («V) *•

Figure 6. XPS spectra of P2S5 treated sample ( a ) 'as received' and ( b ) P2S5 treatment.

156

from the surface. As against the PC15 treatment there are no oxides associated with As. Ga 2P3/2peak was observed at 1117.3eV and As 2p3/2 at 1323.5 eV. P2p peak (not showed here) occurred at 133.2 eV. The broad shoulder on high binding energy side can be associated to As-S species on the surface. Thus, it can be seen that phosphorous can be effective passivating agent. The heat of GaP formation is rather low and fewer defects may be introduced [28].

GaAs + P - • GaP + As [-7.4 kcal/mol] GaP would provide a chemical continuity due to III-V compound (GaP) formation. Covalent radii of As (1.18) and P (1.10) are also very close to each other. Therefore, atomic arrangement of GaAs and GaP should be quite similar, and provide structural continuity.

4. Conclusion

We have carried out investigations of GaAs(lOO) surface after chemical cleaning and passivation using PC15 and P2S5. It should indeed be possible to use phosphorous as an effective passivant, though PC15 may attack adversely through chlorine. However, in the case of P2S5 as both P and S are passivants, improvement can occur in surface properties if thick film is avoided.

Acknowledgments

RP thanks IUC-DAEF for providing necessary facilities during the stay of the work at Indore. SKK thank UGCJndia . A special thanks are due to the help rendered by Mr. A.D.Wadikar in operation of the beamline.

References

1. S.M. Sze, Physics of Semiconductor Devices (Wiley Eastern Ltd, 1981). 2. L.T. Canham, Appl. Phys. Lett. 57, 1045 (1990). 3. D.L. Lile, Solid State Elect. 21, 1199 (1978). 4. T. Ito, Y. Sakai, Solid. Stat. Elect. 17, 751 (1974). 5. D.L. Thurmond, G.P. Schwartz, G.W. Kammlot, B. Schwartz, J.

Electrochem. Soc. 127, 1366 (1980). 6. C. J. Spindt, W. E. Spicer, App. Phys. Lett. 55, 1653 (1989). 7. G.P. Schwartz, W.A. Sunder, J.E. Griffiths, J. Electrochem. Soc. 120, 1361

(1982). 8. H. Hasegawa, H. Ohno, J. Vac. Sci. Tech. B4, 1130 (1986). 9. S.D. Offsey, J.M. Woodall, A.C. Warren, P.D. Kircgner, T.I. Chappell,

G.D. Petit, Appl. Phys. Lett. 48, 475 (1986). 10. T. Sawada, H. Hasegawa, H. Ohno, Jpn. J. Appl. Phys. 26, L 1871 (1987). 11. S. M. Beck, J.E. Wessel, Appl. Phys. Lett. 50, 149 (1987).

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12. G.F. Fountain, S.V. Hattangady, R.A. Rudder, R.J. Markunas, G. Lucovsky, S.S. Kim, D.V. Tus, J. Vac. Sci. Tech. A7, 576 (1989).

13. J.L. Freeouf, D.A. Buchanan, S.L. Wright, T.N. Jackson, B. Robinson, Appl. Phys. Lett. 57, 1919 (1990).

14. M.T. Cuberes, J.L. Sacedon, Appl. Phys. Lett. 57, 2794 (1990). 15. J. Reed, Z. Fan, G.B. Gao, A. Botchkarev, H. Morkoc, Appl. Phys. Lett. 64,

2706 (1994). 16. K.W. Vogt, P.A. Kohl, J. Appl. Phys. 74, 6448 (1993). 17. A.M. Green, W.E. Spicer, C. Kim, R. Cao, P. Pianetta, J. Vac. Sci. Tech.

A12, 1158(1994). 18. M. Sugiyama, S. Maeyama, M. Oshima, H. Oigawa, Y. Nannichi, H.

Hashizume, Appl. Phys. Lett. 60, 3247 (1992). 19. C. Kirchner, M. George, B. stein, W.J. Parak, H.E. Gauband, M. Seitz, Adv.

Fund. Mater. 12, 266 (2002). 20. P.S. Xu, F.P. Zhang, E.D. Lu, F.Q. Xu, H.B. Pan, X.Y. Zhang, Nucl. Instr.

Methods in Phys. Res. A 467-468, 1202 (2001). 21. N. Kuroda, H. Ikoma, Jpn. J. Appl. Phys. 40, 6248 (2001). 22. N. Esser, D.R.T. Zahn, C. Muller, W. Richter, C. Stephens, R. Whittle, I.T.

McGovern, S.K. Kulkarni, W.Braun, Appl. Surf. Sci. 56-58, 169-177 (1992).

23. S.M. Chaudhari, D.M. Phase, A.D. Wadikar, G.S. Ramesh, M.S. Hegde, B.A. Dasannacharya, Current Science 82 , 305 (2002).

24. K.C. Hwang, S.S. Li, J. Appl. Phys. 67, 2162 (1990). 25. Y. Wang, Y. Darici, P.H. Holloway, J. Appl. Phys. 71, 2746 (1992). 26. J.A. Dagata, W. Tseng, J. Bennet, J. Shnier, H.H. Harary, Appl. Phys. Lett.

59,3288(1991). 27. C.C. Surdu-Bob, S.O. Saied, J.L. Sullivan, Appl. Surf. Sci. 183, 126 (2001). 28. S.K. Krawezyk, G. Hollinger, Appl. Phys. Lett. 45, 870 (1984).

CORRELATION BETWEEN MICROSCOPIC AND MACROSCOPIC PROPERTIES OF YTTRIA-STABILIZED

ZIRCONIA THIN FILMS

M. HARTMANOVA

Institute of Physics, Slovak Academy of Sciences, 8451J Bratislava 45, Slovakia E-mail: fyzihart@savba. sk

M. JERGEL

Dept. de Fisica, CINVESTAV- IPN, Apdo. Postal 14-740, Mexico 07300, D.F., Mexico

V. NAVRATIL

Dept. of Physics, Faculty of Education, Masaryk University, 60300 Brno,Czech Republic

K. NAVRATIL

Inst, of Condensed Matter, Faculty of Set, Masaryk Univ, 61137 Brno, Czech Republic

K. GMUCOVA

Institute of Physics, Slovak Academy of Sciences, 84511 Bratislava 45, Slovakia

F.C. GANDARILLA

Departamento de Ciencia de Materiales, Instituto Politecnico Nacional, U.P. ALM Lindavista, Mexico 07738, D.F., Mexico

J. ZEMEK

Inst, of Physics, Academy ofSci. of the Czech Republic, 18121 Prague 8,Czech Republic

S. CHROMIK

Inst, of Electrical Engineering, Slovak Academy of Sciences, 84239 Bratislava, Slovakia

F. KUNDRACIK

Dept. of Radiophysics, Faculty of Mathematics and Physics, Comenius Univ., 84215 Bratislava, Slovakia

The investigated films were deposited by e-beam evaporation of yttria-stabiIized zirconia (YSZ) crystalline samples with the well-defined phase composition, on the n-doped Si (111) substrate at 750°C. The XRD of films exhibited their polycrystalline structure with a mixture of different phases, mainly the cubic phase. The electrical conductivity and the activation energy as the functions of yttria content indicated the present influence of isolated oxygen ion vacancies as well as the associated point defects. The relative effective permittivities (sr = 17-26) confirmed YSZ as a high-i gate dielectric also in the form of thin film. The measured microhardness data, evaluated with reference to the used substrate by Jonsson-Hogmark composite hardness model, (H = 5.9 - 10.8 GPa), and the high refractive index (n = 2.20 - 1.96) indicated YSZ to be a promising material for the protective coatings.

158

159

1. Introduction

Zirconia is one of the best examples in materials science where structure, microstructure, defects and phase composition/transformations are intimately connected with macroscopic properties such as, for instance, electrical conductivity and microhardness. The variety of industrial applications of zirconia requiring both high temperature and chemical stability stems from its basic property of polymorphism (monoclinic, tetragonal and cubic phases).

The purpose of the present study was to carry out a more detailed investigation of zirconia polymorphism based on the interest in a better understanding of the correlation between microscopic and macroscopic properties not only for the bulk form of this material as it was done in our recent study [1] but as far as it will be possible, also for zirconia thin films.

2. Experimental

Preparation of samples. Yttria stabilized zirconia (YSZ) films were deposited on n-doped Si (111) substrates by electron beam evaporation of Zr02 + x. Y203

(1.25; 3.24; 4.31 and 9.21 mol%) crystalline samples used as the targets. The temperature of deposition was 750°C. The details of the deposition can be found elsewhere [2].

The amount of yttrium ions in the deposited films was estimated by X-ray photoelectron spectroscopy (XPS). The measurements have been carried out in the ADES-400 spectrometer equipped with the hemispherical angle resolved energy analyzer and X-ray tube. The atomic concentration ratios Zr/Y obtained together with the mol % Y203, calculated from these ratios by stoichiometry can be seen in Table 1. The thicknesses of films estimated by the surface profilometry using Talystep are shown in the same Table 1.

Table 1. Concentrations (x) of Y203 in Zr02 crystals [1] and thin films, atomic concentration ratio (Zr/Y) and thickness t of thin films.

Xcrvst (mol%Y203) 1.25 3.24 4.31 9.21

(Zr/Y)fllm

26.5 13.1 8.1 4.3

xfilm (mol%Y203) 1.22 2.39 3.73 6.49

tfiun (nm) 304 204 103 116

X-ray diffraction (XRD) measurements were performed in the surface sensitive grazing incidence geometry on a 12 kW Rigaku rotating anode generator at the angles of incidence fixed at 0.5°, 1.0° and 1.5° using CuKa

radiation. A dedicated laboratory-made diffractometer with a graphite monochromator in the primary beam was used.

160

A.C. conductivity measurements were made using a Solartron SI 1260 impedance/gain phase analyzer. The impedance measurements were made in the frequency range of 10 Hz - 1 MHz at temperatures RT - 480°C in air. All electrical measurements were performed in the metal-insulator-semiconductor (MIS) configuration with aluminium as the top and the gold as the bottom electrodes.

Vickers microhardness measurements were made in air at room temperature by means of a Hanneman (Vickers) microhardness tester and a Zeiss-Neophot microscope. Specimen indentations were performed with the load of 1.1 N. Reflectance measurements were carried out by the reflectance spectrophotometry using a double-beam spectrophotometer Varian CARY5E in the spectral region of 200 - 800 nm in air at room temperature. The reflectance spectra were measured using the polished plate from crystalline silicon as a standard of reflectance.

3. Results and Discussion

Phase composition and structural parameters. The structure of all investigated oxide film configurations was found to be polycrystalline, the crystallite sizes obtained are 20.3 nm for sample with 6.49 mol% Y203 and 18 nm for sample with 3.73 mol% Y203. A general Rietveld analysis of the diffraction patterns discovered the presence of a face centered cubic (fee) phase (space group Fm3m) in the samples with 3.73 and 6.49 mol% Y203. No satisfactory phase indentation could be done in the samples with 1.22 and 2.39 mol% Y203. An example of the diffraction pattern is shown in Fig. 1. There is one unidentified peak in the diffraction pattern, which is the case also for other samples, however, the unidentified peaks occur at different angles in different samples. From the relative intensities of the fee peaks, a tendency to one axial texture with (100) planes oriented preferentially parallel to the substrate could be concluded which was the most obvious for the sample with 3.73 mol% Y203. Therefore in the second step, the Rietveld refinement was performed taking into account this texture, however, the fit quality was poor and prevented us from a reliable determination of the cell parameters and their variation with the film composition. The perfection of the fee phase in terms of Rietveld fitting decreases with decrease in the Y203 content, i.e. from the sample with 6.49 mol% Y203 down to one with 3.73 mol% Y203. Therefore, we can conclude only that the films contain at least two structurally different phases: an fee phase with increasing degree of perfection as Y203 content increases, and the other one which may be presumably a highly distorted form of some regular phase.

161

| 7 5 0

| 5 0 0

250

o = 0.5

iw^wfaffi* V # * N^umi' imldktt/ -^rY T*IU|M»I«IM j

30 40 60 70 50

2 0 ( d e g r e e )

Figure 1. Grazing incidence diffraction pattern of the film with 6. 49 mol% Y203 measured at the angle of incidence 0.5°. The indices refer to the face-centered cubic phase.

These results obtained for thin films are different from the results of Rietveld analysis obtained for powdered crystals [1], which were used as the targets for e-beam deposition. For instance, the crystalline target for the preparation of film with 6.49 mol% Y203 was identified as single-phase cubic, while that one for the preparation of film with 3.73 mol% Y203 as a mixture of cubic and tetragonal phases. On the other hand, the films deposited from them possess unambiguously only the cubic phase. However, the peak in the diffraction pattern not belonging to this cubic phase suggests still the presence of another phase. At lower Y203 content, even, no regular phase could be identified in the films deposited from the crystalline targets. These facts indicate, besides the others, a difference between elemental compositions of targets and the films deposited from them, probably due to the volatility of yttrium. This difference was confirmed also by XPS.

Electrical and dielectric properties: The structure-conductivity relation for the fluorite-related structures is strongly dependent on the phase composition and crystallographic structure also in the case of thin film form of material. The impedance diagrams of investigated films at higher temperature consist of two semicircles as it is shown in Fig. 2. The first semicircle, a high-frequency one, can be attributed to the bulk of sample. The relative permittivities calculated from this semicircle in the next are in a good agreement with the permittivity values obtained by the other independent methods. The origins of the second semicircle could be the crystallite boundaries, the presence of thin (~ 1 nm) layer of Si02 or the small isolated islands of Si02 on the surface of Si substrate and/or also the imperfect electrical contact due to these islands. The simultaneous influence of all these factors is also possible. The impedance diagram at the highest frequencies (the inset in the Fig. 2) is slightly influenced by the presence of a parasitic inductance. The effect of yttria content and phase

162

composition on the bulk electrical conductivity o and activation energy Ea can be seen in the form of Arrhenius plots (Fig. 2), the isothermal bulk conductivity vs yttria content (Fig. 3a) as well as in the form of a change of activation energy in dependence on the mol% Y203 (Fig. 3b).

E o

ioN

10"

ior7-

ior,

1.0 — I —

2.5 — I —

3.0 —I

4.0

1000/1 [K1]

Figure 2. Arrhenius plots of ac bulk conductivity a of zirconia thin films doped with various amounts of Y2Oj (with the inset of impedance diagram for Zr02 + 2.39 mol% Y2O3 at 261 °C): (•) 1.22 mol% Y20,, (A) 2. 39 mol% Y20,, ( • ) 3.73 mol% Y203 and (•) 6.49 mol% Y203, measured in air.

The bulk conductivity increases from the lowest value obtained for the samples having concentration of 1.22 mol% Y203 to its maximum value for samples containing 2.39 mol% Y203. The further increase in yttria content from 3.73 mol% Y203 up to 6.49 mol% Y203 in Zr02 results in the decrease of conductivity. The observed increase in conductivity is probably due to an increase of the fraction of more conductive phase and in such a way also the increase of amount of defects such as the isolated oxygen ion vacancies V 0

which control the bulk ionic conductivity of the lightly doped samples and are generated by the presence of substitutional Y3+ ions. The decrease of conductivity at the higher used amounts of yttria is probably attributable to the decrease of isolated oxygen vacancy concentration in the anion sublattice YxZri.x02.x/2 solid solution due to the increased interactions between the dopant cations and vacancies at low temperatures. The associated point defects (YZ rVoJ control the transport properties of samples with the higher concentrations of yttria at low temperatures.The calculated activation energy Ea

decreases with increasing small amounts of yttria. The minimum of Ea was observed at x = 2.39 mol% Y203.

163

> .2.0.3

\

xlmoCStYjOJ x[mol%Y203]

Figure 3. (a) Isothermal dependences of the bulk conductivity o of YSZ on the amount of Y203: (A) 201°C, (o) 261°C, (n) 323°C; (b) Influence of amount of Y20, on the activation energy Ea of YSZ; both in the low-temperature region.

Recently, Zr02 attracted much attention as a possible high-/: gate dielectric in microelectronic metal-oxide-semiconductor devices, promising to overcome the tunneling limits of Si02 and thus allowing further reduction of device dimensions [3]. The relative effective permittivity er of the YSZ thin films investigated in the present study was obtained by three independent ways using usual formula er = C.t / e0. S, directly from the corresponding:

Bulk parallel capacitance Cp measured at room temperature and frequency of 1 MHz Capacitance Cimp of bulk semicircle maximum in the impedance diagram (modelled by parallel combination of resistor and capacitor and using LSQ technique) Accumulation capacitance Cacc determined from C-V curves

where t is the thickness of film, S is a gate area and e0 is the permittivity of the free space (8.845 x 10"14 F . cm"1). The capacitances and the corresponding relative permittivities obtained can be found in Tab. 2.

Table 2. Capacitances Cp, Cimp, Cacc and relative permitivities cr of ZK32 + x. Y203 mol% thin films deposited by e-evaporation at 750°C on Si(l 11) substrate.

X

(mol%Y203) 1.22

2.39

3.73

6.49

Cp(pF)

599 985 1479

1113

Sr(Cp)

23 26 19 17

Cimp(pF)

635 943 1370

1208

ErV^imp/

26 25 21 18

Cacc(pF)

288 383 700

450

^rl^acc/

29 26 24

17

The er values obtained in the present study are in a good agreement with the literature data [4, 5] and confirm YSZ as a high-dielectric constant material also in the form of thin film.

164

3.1. Microltardness Tests

The hardness of films is usually an important mechanical characteristic for the surface coatings. However, due to the small thickness, the measured hardness is usually influenced by the substrate. In order to separate the properties of the film from those of the substrate (composite ones), several composite hardness models have been proposed. We have used in the present study one of them. The room temperature Vickers microhardness value of thin ductile film Hf was estimated using Jonsson-Hogmark model [6] based on the geometrical approach to combine the microhardness of film and substrate according to an area mixture model. This method is based on the calculation of the relative contributions of substrate and film to the effective microhardness. The effective microhardness is expressed by the area mixture model as the weight mean of the substrate hardness and the film hardness according to the relation

H f = H s + H c ~ H s , (1)

2 C - - C D

t

where Hf is the microhardness of ductile film, Hs is the microhardness of used substrate (9.81 GPa), Hc is the microhardness of all cell (i.e. the thin ductile film together with the used substrate), t is the thickness of thin ductile film and D is a depth of indentation (equalled to the 1/7 of measured indentation diagonal length). The constant C is determined by the geometry of indentation (the tip angle of diamond pyramid, 136°), C = sin 22° and the mechanical properties of our system (the harder, ductile and at the indentation also flexible film on the softer substrate). A comparison of the data obtained in the present case and the one obtained under the same conditions of measurements i. e. temperature, indentor and load for the corresponding crystalline samples [1] is presented in Table 3.

Table 3. Data from the microhardness test for Zr02 + x. Y203 films deposited on Si(l 11) substrate measured at the load 1.1N and compared with the data obtained under the same conditions for the corresponding crystaline samples of unknown orientation [1].

Xcwst (mol% Y203) 1.25 3.24 4.31 9.21

I W G P a ) 10.0 13.6 16.5 12.0

xfilm (mol% Y203) 1.22 2.39 3.73 6.49

Hf„m(GPa) 5.9 8.4 9.8 10.8

In general, the microhardness of thin films is varied widely depending on the substrate and the processing conditions used. It is a complex function of

165

many variable factors, besides the measured technique, such as a surface finish, microstructure, indentor load, configuration, etc.

The value of microhardness Hf in Table 3 increases with increasing amount of yttria in the sample. This is in an agreement with the increasing crystallographic ordering of structure. No load -dependence (in the investigated region), similarly as no cracking around the indent (impression) were observed. The microhardness of investigated YSZ films was found to be lower than that for the corresponding crystalline form of material in similar concentration range. However, the measured values indicate that YSZ thin films could be used as the protective coatings.

Refractive index: Zirconia is preferred for the optical devices because of its desirable properties, such as, for instance, a high refractive index. It is possible to obtain the information about it from the reflectance measurements. The reflectance spectra of the investigated YSZ films having different compositions as a function of wavelength X are very similar with the exception for the film with 1.22 mol% Y203, the data of which were not evaluated due to the unclear behaviour of this film (absorption, extremely low reflectance). As an example, the reflectance spectra of zirconia doped with 2.39 mol% Y203 and the substrate Si (111) with the inset of fitting curve are presented in Fig. 4. The measured reflectance spectra have shown that the films are not absorbing in the investigated spectral region (400 - 800 nm), however, they were found to be inhomogeneous along their thickness.

0.8-,

0 7 -

0.6-

0.5-

0.4-

Di

0.3-

0.2

0.1

o.o - i — i — i — i — . — | — i — i — i — i — 100 200 300 400 500 600 700 800 900

X[nrrj

Figure 4. Reflectance spectra of zirconia film doped with 2.39 mol% Y203 and the substrate Si (111) with the inset of fitting curve.

When the linear change of the refractive index along the thickness of film is assumed, we can obtain the values of normal refractive index n using the well-

166

known Cauchy equations for the refractive indexes at the air-film and the film-silicon interfaces. Such values of normal refractive index n obtained at the wavelength X = 550 nm are presented in Table 4.

Table 4. Refractive indices of Zr02 + x. Y203 thin films deposited by e-beam evaporation at temperature Ts = 750°C on Si(l 11) substrate and measured at the wavelenght X = 550 nm.

x (mol% Y203) 0.00 1.22 2.39 3.73 6.49

n -2.22

-2.20 2.14 1.96

Table 4 contains refractive index of undoped Zr02 film. This was obtained by other authors under approximately same conditions (Ts, X), however deposited by a different (sputtering) method [7]. According to literature [7], the presence of Y203 dopant in the stabilized zirconia reduces the effect of substrate temperature Ts on the film properties with regard to the undoped Zr02. As it was mentioned above, the films investigated in the present study are composed from the very small crystallites (grains) what implies a large fraction of surface regions and thereby an important role-played by the surface energy. The refractive index of YSZ films is relatively high and its value slightly decreases with increasing amount of Y203 dopant in zirconia from 2.20 for 2.39 mol% Y203 to 1.96 for 6.49 mol% Y203. This behaviour of refractive index is in a good agreement with the literature data [8, 9] and could be attributed to the change in crystallographic structure.

4. Summary

The purpose of the present study was to investigate the correlation between the crystallographic structure and some physical properties of yttria stabilized zirconia (YSZ) thin films deposited by e-beam evaporation of YSZ crystalline targets with a well defined phase composition on the n-doped Si (111) substrate at 750°C. Our aim was fulfilled only partially because in spite of the careful XRD measurements and the effort to utilize Rietveld analysis, we have been able to obtain only a rough insight on the phase composition and the crystallite sizes. The investigation has provided the following data:

-The structure of all investigated oxide film configurations was found to be polycrystalline, the crystallite sizes were estimated to be approximately 20.3 nm for sample with 6.49 mol% Y203 and 18 nm for sample with 3.73 mol% Y203. Rietveld analysis of diffraction patterns discovered at least two structurally

167

different phases, one cubic (fee) phase with the increasing degree of perfection as Y203 content increases and the other one which may be presumably a highly distorted form of some regular phase in the samples with 3.73 and 6.49 mol% Y203. No satisfactory phase indentation could be done at the samples with 1.22 and 2.39 mol% Y203.

-The total electrical conductivity of films consists of the contributions from the bulk and one of the following factors such as the crystallite boundaries, the presence of the thin (~ 1 nm) layer of Si02 or the small isolated islands of Si02

on the surface of Si substrate and also the imperfection electrical contact due to these isolated islands in dependence on the measured conditions (temperature and frequency). The simultaneous influence of all these factors cannot be excluded.

-The bulk conductivity and activation energy as the functions of yttria content indicate the present influence of isolated vacancies as well as the associated point defects, the conductivity maximum and the minimum of activation energy at ~ 2.39 mol% Y203.

-The relative permittivities er = 1 7 - 2 6 confirmed YSZ as a high-dielectric constant material also in the form of thin film.

-The measured microhardness data, H = 5.9 - 10.8 GPa, evaluated with respect to the present Si substrate using Jonsson - Hogmark composite hardness model increased with the increasing perfection of the fee phase.

-The relatively high refractive index, n = 2.20 - 2.96, the values of which could be attributed to the change of crystallographic structure and important role of surface energy due to a large fraction of surface regions of very small crystallites.

-The comparison of properties between the investigated YSZ films and the corresponding crystalline targets (with the well-defined phase composition), from which the films were prepared, has shown a good qualitative agreement of their behavior in dependence on the degree of perfection of their crystallographic structure.

Acknowledgements

The work was supported by the research grants No. 2 / 1119/ 21, 2 / 7103 / 20, 2 / 1013 / 21, 2 / 2068 / 22 of the Slovak Grant Agency and FRVS MSMT 996 / 2000.

References

1. Hartmanova M., Schneider J., Navratil V., Kundracik F., Schulz H., Lomonova E. E., Solid State Ionics 136 - 137, 107 (2000).

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2. Chromik S., Wuyts B., Vavra I., Rosova A., Hanic F., BenaCka S., Bruynseraede Y., Physica C 226, 153 (1994).

3. Sharma R. N., Rastogi A. C , J. Appl. Phys. 76, 4215 (1994). 4. Hartmanova M., Gmucova K., Thurzo I., Solid State Ionics 130, 105

(2000). 5. Croset M., Schnell J. P., Velasco G., Siejka A. J., J. Appl. Phys.48, 775

(1977). 6. Jonsson B., Hogmark S., Thin Solid Films 114, 257 (1984). 7. Uyama H., Oka N., Ono I., Matsumoto O., Denki Kagaku 58, 564 (1990). 8. Liu W. - Ch., Wu D., Li A. -D., Ling H. -Q., Tang Y. -F., Ming N. -B.,

Appl. Surface Sci.191, 181(2002). 9. Boulouz M., Boulouz A., Giani A., Boyer A., Thin Solid Films 323, 85

(1998).

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RELIABILITY AND FAILURE OF ELECTRONIC MATERIALS AND DEVICES

M. OHRING

Stevens Institute of Technology, Department of Materials Science and Engineering, Hoboken, New Jersey, USA, 07030

E-mail: mohring@stevens.edu

Advances in microelectronics in the last third of the 20th century have changed the world in unprecedented ways that will continue unabated at least early into the 21st century. Progress promises to be driven by an impressive array of computer and telecommunications equipment as well as other electronic products of all kinds that will become ever more powerful in function, while simultaneously smaller in size. At the root of these trends has been the exponentially increasing transistor count defined by Moore's Law (Fig. 1), namely a doubling of the number every 18 months or so. Because the individual chip area has also grown (Fig. 2), but at a smaller exponential rate, the transistor density has continued to rise. This has been accomplished by shrinkage of device features as exemplified by aluminum interconnects. It is to note that the dimensions have shrunk over the years from several microns to 0.13 um presently. By 2010, conductor line widths of 0.07 um are projected [1].

Adherence to Moore's Law has been achieved by a massive parallel effort to ensure that the ever-smaller device components and packages have the requisite reliability as measured by failure rates in service. Traditionally, the bathtub curve has provided a basis for understanding the time dependence of reliability in products [1]. Once weak devices are weeded out by testing (or early in service) failure rates are random, but then at longer times are governed by wear out effects. There are important economic implications associated with increasing reliability in each of these time periods. It is primarily the range of constant and increasing failure rates that we are concerned with here.

On the chip-level, two of the important reliability problems arise from the behavior of metallizations and dielectrics. A cross-section of the chip shown in Fig.3 reveals where metals and dielectrics are employed. The active complementary metal oxide semiconductor (CMOS) transistors are fabricated on top of the wafer surface. Because real estate on this surface is expensive, the metal conductors are located above these devices in a layered superstructure network consisting of interconnections embedded within Si02 on all sides. Each layer is stacked vertically with parallel levels connected to each other by means of metal vias. Properties required of the metallizations and interlevel dielectrics

171

172

are listed in Table 1. A major reliability problem associated with metallizations is electromigration.

Table 1. Metallization and dielectric issues.

Metallization

Aluminum (Al) still metal of choice High conductivity

Compatible with IC processing Reflowable at low temperature

Aluminum has low melting point Electromigration

Incompatible with processes above 500°C

Copper (Cu) Higher conductivity than Al

Higher resistance to electromigration than Al

For same R, smaller line to line coupling than Al

Other Conductors High melting point

Lower conductivity than aluminum Used mainly for local interconnects

Tungsten Refractory metal silicides Refractory metal nitrides

Dielectrics

Low dielectric constant for ILD

High dielectric constant for cell, gate

Low leakage

High breakdown strength

Low mobile ion concentration

Low water content

"Good" mechanical properties

Complete gap fill

Low contamination

Under DC powering at high current densities, mass is transported toward the anode as shown in Fig 4. This leads to mass depletion and voids in the cathode area, a region where interconnect failures are favored. Similar metal accumulations and depletions occur in vias (Fig 5) that represent sites where the metallization (Al-Cu) adjoins and contacts plugs (W) and via linings (TiN). In the case of dielectric, the critical device component is the gate oxide, which is subject to dielectric breakdown (see Fig. 6) and hot-electron effects. The latter occurs because high energy electrons scatter vertically into the gate oxide during transit from the source to drain. Dielectric breakdown and hot-electron damage

173

to Si02 as well as electromigration-induced degradation damage to conductors are generally worn out effects.

Irrespective of the kind of damage produced, the analysis of more than thirty years of electromigration reliability testing of interconnections has distilled the following general relationship between mean time to failure (MTTF) and operating current density and temperature, namely:

MTTF~X = BJ" exp[-£g IkT] (1)

where B is a constant, Ee is the activation energy for electromigration (usually between 0.25-1.0 eV) is attributed to Black [2] who found that the n = 2 value. MTTF dependence on the current density suggests that MTTF"1 is proportional to the product of the number of voids in a thin film strip and a rate at which voids grow supporting the current density power (n = 2). The electromigration lifetime depends on conductor length and width. Figure 7 shows the dependence of MTTF on interconnects line widths for different thin film metals. The encouraging increase in lifetime occurs below 2 urn is reported due to bamboo grain development [3]. The quality of thin film metal lines is very important in achieving desired properties and subsequent device performance. Thus, knowledge of thin film growth is key issue in device properties. A comprehensive description of thin film fabrication and characterization is given recently in ref. [4].

Progress in microelectronic products not only means the continued production of reliable high-device density chips, but corresponding advances in the way chips are packaged. Over the years, techniques have been developed to increase the number of input-output connections while simultaneously reducing the size of packages. The type and shape of packages are evolved substantially in the last three decades. In addition, other strategies have developed to enhance the reliability of systems composed of multiple chips and packages. One such example is the voice recognition module where multi-chip module package is replaced several separately packaged chips mounted on a circuit board. Furthermore, by flip-chip packaging of the Si chips on a Si substrate, reliability is increased because thermally induced stresses due to thermal expansion coefficient mismatches between materials is minimized. It is precisely these thermal stresses that are responsible for a host of solder failure and reliability issues.

Adherence to Moore's Law in the future will only be possible with the adoption of new materials and correspondingly new processes that enable them to be incorporated into devices. In recent years, an unprecedented number of such materials have been introduced into circuitry (Fig. 8). They include copper metallizations that are now replacing aluminum alloys, and low dielectric

174

constant insulator that will eventually serve as intermetal fillings. Similarly, thin silicon oxynitride films are scheduled to replace the critical Si02 gate oxide. Whereas traditional metal and dielectric materials served for almost 30 years their replacements are being placed into service within much shorter time frames. Furthermore, increasing device densities mean more processing and lithographic patterning steps, all of which enhance the probability of defect formation. Key to the reliability of these materials will be controlled of such defects, which incidentally are becoming ever smaller in size. As seen in Fig. 9 defects not only control the yield or economic viability of products, but they also critically affect their reliability. Extension of Moore's Law into 21st century can thus only occur with reduced defect counts. Therefore, "Less is Moore".

References

1. M. Ohring, Reliability and Failure of Electronic Materials and Devices (Academic Press, Boston, 1998).

2. J.R. Black, IEEE Trans. Electron Devices ED-16, 338 (1969). 3. S. Vaidya, T.T. Sheng, A.K. Sinha, Appl. Phys. Lett. 36, 464 (1980). 4. M. Ohering, Material Science of Thin Films: Deposition and Structure,

Second Edition (Academic Press, Boston, 2002). 5. D.R. Wolters, J.J. Van der Schoot, Philips J. Res. 40, 115 (1985). 6. J. Tao, K.K. Young, N.W. Chung. C. Hu, 30th Annual Proceedings of the

IEEE Reliability Physics Symposium 338 (1992). 7. T. Yamaha, S. Naitou, T. Hotta, 30th Annual Proceedings of the IEEE

Reliability Physics Smposiumy 349 (1992).

175

1G 7

100M -.

f

I 10M -,

1M -

!

100K -

10K 1

1K

16K

4K,

1641

256MA

* 64MJ

. Gate 16M.X Arrav

I ' P e n t i u m y * 80686''

v * - * rifi Mitsubishi

"Gate Array r80486 ' 68040 (

\ L S I Logic 1 ^256rv^8°386 G^e Array

(68020 rS0286 l

'68000 '

'8086

4004

<"# •8085 , . r8080 i ;

* = Microprocessor/Logic

A-Memory (DRAM) I 1 1 1 I 1 1—T 1 1 1 1 f -

70 72 74 76 78 80 82 84 86 88 90 92 94 96 98

Year

Figure 1. Plot of the time-dependent of increase in number of transistors on microprocessor and memory integrated circuit chips. Source: ICE Corp. 1993 Status.

176

2,000

80786?.

70 72 74 76 78 80 82 84 86 88 90 92 94 96 98

Year

Figure 2. Plot of the time-dependent increase in chip area of microprocessor and memory integrated circuit chips. Source: ICE Corp. 1993 Status.

Oxide spacer TiSi2

p+ poly n+ poly Oxide [̂ TiSi2 spacer

i substrate!

Figure 3. Chip cross section for the CMOS structure.

177

Temperature

Figure 4. Schematic model of damage in an interconnect subject to electromigration. Hillocks and voids develop in response to temperature gradients superimposed on the directed electron current flow.

- Al accumulation

e-) Mz

Figure 5. Electromigration damage at Vias between metallization levels Mi and M2: a) void formation in metallization at bottom of via plug (form reference [6]) b) Void formation in metallization at top of via plug (form reference [6]) and c) Thinned metallization at via bottom and sidewalls (from reference [7]).

178

Figure 6. Scanning electron micrograph of dielectric breakdown in a 40 nm thick Si02 film. The crater diameter is - 3 nm (from reference [5]).

- Gun, 450oC/30 min

• Gun/3000A POLY-Si, 450OC/30 min ,

- - Gun/3000A POLY-Si, 450°C/30 min (OVER STEPS)

In - S, 300-C/30 min

4 6 8 10 12 LINE-WIDTH (nm)

Figure 7. Mean time to failure as a function of interconnects line width for evaporated (E-gun) and sputtered (In-S) Al films (from reference [3]).

179

sJ

:

SOI, BST, Ru02, Organics _ . „ ,

SIOF, Ta2Ost CoSh D-Cu, TaSiN

Ti Cu W

1.2 0.8 0.5 0.35 0.25 0.18 0.13

Design Rules (Microns)

Figure 8. Number of new materials used in circuitry.

PROCESSING OPERATING STRESS

MATERIALS FAILURE

DESIGN TIME

Figure 9, A representation of the way yield and reliability concerns are interconnected through defects.

DIFFUSION IN MULTILAYERS

S. LUBY, E. MAJKOVA

Institute of Physics, Slovak Academy of Sciences, 84511 Bratislava, Slovakia

A. LUCHES

INFMand University ofLecce, 73100 Lecce, Italy

An overview on the laser induced diffusion in metallic multilayers (MLs) is given. The method of measurement is illustrated on the studies of Co/Ag bilayers, trilayers and Fe/W and Co/W multilayers for magnetoresistance sensors and of W/Si MLs for X-ray mirrors. Pulsed excimer XeCI and KrF lasers with 30 nm pulses were used. The main advantage of pulsed laser heating is low inertia of the process because the heating is restricted to the certain area of the sample. Moreover, heat is delivered reproducibly in certain quanta. Therefore, as it will be shown, the characteristic diffusion distances in MLs of only a few nm per pulse could be recorded. The laser heating was performed at flounces of 0.05 -0.6 J/cm2 and number of pulses up to 100. They were directed to the same site. In this paper, basic diffusion parameters for the above mentioned structures are summarized. Due to the rapid laser heating and subsequent rapid cooling, some interesting phenomena in layered structures could be observed e.g. the grain boundary diffusion with low activation energy was recorded in Co/Ag system up to the melting point of Ag. This is because the crystallization does not proceed enough fast. Moreover, for mutually immiscible combination of materials the sharpening of interfaces in MLs due to the backdiffusion effect was found.

1. Introduction

Metallic multilayers (MLs) consisting of alternating layers of two different materials are broadly used in soft X-ray optics [1] and in giant magnetoresistance (GMR) sensors [2]. The combination of elements or compounds with high and low Z (e.g. Mo/Si) and of ferromagnetic and non-ferromagnetic elements (e.g. Co/Ag) are the examples of the first and second application, respectively. In this paper we pay attention especially to laser irradiation induced diffusion and intermixing in metallic MLs.

Typical ML periods (bilayer thickness) are d = 2 - 20 nm and a number of periods in the stack may be up to 400. Because of small diffusion distances and high number of interfaces these structures are of great interest for diffusion studies. However, the structure of MLs is complex and all diffusion driving forces and types of diffusion may be present. The most important between them are [3] isothermal diffusion, thermomigration, stress-induced diffusion (baromigration) and electromigration, if structures are electrically loaded. The flux of particles can be expressed as

J = -DVN+ FND/kT (1)

180

181

where D is the diffusion coefficient, N - concentration of particles, F - driving force, k and T have the usual meaning.

Numerous diffusion studies in MLs have shown that kinetics of interdiffusion process is much faster than expected from high temperature data obtained with bulk systems. Due to complex composition and structure of MLs back-diffusion (reverse diffusion, chemical cleaning) with negative interdiffusion coefficient is observed as well, especially in immiscible combination of materials [4].

2. Laser Heating in the Diffusion Studies

In the diffusion studies materials are conventionally heated using standard furnace or RTA (rapid thermal annealing) isothermal methods. Here entire sample is heated. Using laser beam the heating is restricted to certain area of the sample. Therefore, the laser heating is less inertial (Fig. 1) [5]. Both cw and pulsed lasers are used. CW diode pumped doubled YAG was used for diffusion studies of semiconductor heterostructures [6]. An additional advantage of pulsed laser is the reproducible delivery of heat into the samples in certain quanta. Besides our papers [5, 7-11] pulsed excimer laser was employed for diffusion studies by Howari et al.[12]. But lasers are used also for other diffusion studies, i.e. for the study of thermal diffusivity [13]. In laser diffusion studies the diffusion parameters must be calculated by integrating them under the temperature profile evolving during and after laser pulse.

2000

1600

* 1200 a

800

400

I 1 1

0.36 Jem'2

A ,

/ < / Ax

• /£>/• / • / k .

A

i

- A -

Xo'^o^ /0.30 J

--•—— S'1"—

i i i

2

cm

•• i

A

xo 0.20 Jem"2

- • — —•

0.15 Jem2

i i .

10 20 30 40

t[ns]

50 60

Figure 1. Temperature vs. time evolution at the surface of Co(10nm)/Ag(20 nm)/Co(3nm) trilayer on oxidized Si substrate irradiated at flounces 0.15,0.2, 0.3 and 0.36 J/cm2 by one XeCI laser pulse.

182

As an example, it can be mentioned that using standard excimer XeCl laser pulse of 30 ns and flounces of 0.1 - 0.2 Jem"2 the diffusion length in Co/Ag layered structures is 2 - 3 run per pulse and it corresponds to the characteristic diffusion distance in MLs with nm period [5]. In the case of isothermal annealing at lower temperatures (100 - 300°C) similar distances are achieved after approx. one hour heating [14].

3. Composition Measurements

For the diffusion studies in MLs with nm periods information of composition with spatial resolution on nm scale is required. Therefore, if methods like AES, SIMS, EDX and TEM are employed, nanobeam design is preferred [15]. In our works, alternative techniques, especially grazing incidence (GI) XRD, X-ray reflectivity (XRR), GI diffusion scattering (GIDS) and RBS methods are employed. The depth sensitivity of GI XRD depends on the absorption of X-rays in surface layers of MLs e.g. the absorption of CuKoc in 10 nm thick Co layer is approx. 50 % at the grazing angle of 0.5° [5]. This is acceptable for many applications. The XRR technique is combined with simulations of reflecting spectra and from the best fit the thickness of intermixed layers is obtained [9]. In RBS studies lower energy beams are employed to improve the resolution. In our work [16] we have used 300 keV He+ ions.

4. Intermixing Studies in Layered Structures

As examples of laser irradiation induced diffusion in layered structures the results for immiscible system Co/Ag [5,8], miscible system W/Si [7,9] and partly miscible combinations Fe/W and Co/W [10,11] are reported. Heat processing was performed using XeCl and KrF lasers. Duration and repetition rate of pulses were 30 nm and 5 Hz, respectively. All samples were deposited by e-beam evaporation onto oxidized Si wafers at the temperature up to 70°C. Laser flounces were F = 0.05 - 0.60 Jem"2, number of pulses to the same site N = 1, 10, 100. The temperature and depth of melting vs. time in laser-irradiated samples were computed.

Co/Ag system:bilayers and trilayers with Co thickness from 1 nm to 10 nm and Ag thickness from 15 nm to 50 nm were studied. Heating with KrF laser was more effective than with XeCl one because of the Co surface of higher reflectivity for the XeCl wavelength. The XRR spectra (Fig. 2) show considerable changes under laser treatment. At F = 0.10 Jem"2, the interfaces become more sharp due to densification of the structure and the stress release. At higher flounce they deteriorate due to solid-state diffusion. After melting of Ag at F = 0.20 Jem"2 the spectrum shows sharpening (number of peaks increases

183

again) which may be attributed to the backdiffusion process. With further increase of the flounce the spectrum deteriorates again. a) In the case of solid-solid interface the results hint at the grain boundary diffusion. The diffusion coefficient is D = 10'8 c m V at 1210 K and the activation energy of diffusion is Q = 0.4 eV. b) In the case of solid-liquid interface diffusion coefficient with the pre-exponential factor D„ = 3.8xl0'7 c m V and activation energy Q = 1.5 eV were obtained. At 1380 K, i.e. at the top of the flat temperature maximum for F = 0.20 Jem"2 (not shown here), we have found D = 1.3xl0"8 c m V . The values D0 and Q are typical for diffusion in liquid metals and amorphous alloys [15].

0,1

1E-3

3

CO

- , E - 5

(0

c 1 1E-7

1E-9

0,5 1,0 1,5 2,0

0 [deg]

Figure 2. XRR spectra of trilayer Co(2nm)/Ag(40nm)/Col.5nm) in the as-deposited state and after treatment by 1 laser pulse at various flounces in [Jem"2].

W/Si system: W,.xSix/Si MLs with 10 periods and x = 0, 0.33, 0.5 and 0.66 and top Si layer were studied. The nominal thickness of Si and the other layer were 5.5 nm and 2.5 nm, respectively [9]. W layer was mixed with Si to increase the thermal stability of MLs. Then the driving force for the Si diffusion into other layer decreases and for combination of Si with W-disilicide (x = 0.66) the thermal equilibrium is achieved. The samples were treated by XeCl laser. Some results are shown in Figure 3. From simulations of XRR spectra, it follows that

T r

J i L

184

under laser irradiation the ML period fluctuated around as-deposited value. This corresponds to laser fast melting and cooling process (Fig. 1). However, at F = 0.60 Jem"2 and N = 100 the MLs are considerably deteriorated or completely collapsed. The best resistance against thermal damage is obtained at x = 0.5. Surprisingly, at x = 0.66 it decreases again. This may be explained by the decrease of the melting temperature of the mixed layer with increasing Si content.

2 0 [ d e g r e e )

Figure 3. XRR spectra of W,,xSix/Si multilayers irradiated by F=0.6 Jem2, N=100. The spectrum of Wo.sSioVSi multilayer in the as deposited state is also shown.

Fe/W and Co/W systems: the special interest in these MLs is because of refractory nature of non-ferromagnetic spacer in GMR sensors. These systems are partly miscible. Their heat of formation is AH = 0 for Fe/W and AH = -2 kJ/g.at for Co/W. MLs with 5 or 10 bilayers and with nominal thickness of Co and Fe layers 1 and 2 nm and of W layers 1, 2, 5 run were studied. Samples were treated by KrF laser. Also here XRR was used for evaluation, however in combination with GIDS. In Figure 4, some results for Fe/W ML are shown.

In Fe/W and Co/W MLs no intermixing at the interfaces under laser treatment was observed at all used flounces up to 0.25 Jem"2. In Fe/W MLs roughness increased slightly from 0.4 to 0.7 nm with increasing flounce. In Co/W MLs the roughness was lower and the superlattice structure was found. Nevertheless, the magnetoresistance was lower in Co/W stack, which supports

185

the idea that the roughness is of some importance for GMR effect [2]. At W on (Co or Fe) interfaces the roughness was higher than at their (Co or Fe) on W counterparts. This is ascribed to the kinetic mixing - the deposited W atoms have higher kinetic energy.

3

f

0 [deg]

Figure 4. Measured (a) X-ray reflectivity (dots) and (b) GIDS detector scan taken on the Is' ML Bragg maximum and their simulations (full lines) of (1 nm Fe/2 nm W)xlO ML. The data are shifted upwards for clarity.

5. Conclusions

Laser is a useful tool in the study of intermixing in layered structures in the nm range of thickness. The localized heat processing and delivery of heat in certain quanta provide broad variability of experimental conditions. Short laser melting does not deteriorate the layered structure as a rule. On the other hand, attention should be paid to the further improvement of various computation procedures to extract the intermixing characteristics from the experimental profiles, because it is practically impossible to monitor the thermal processes within the duration of ns laser pulses.

186

Acknowledgements

The support by Scientific Grant Agency VEGA, Bratislava under the contract 2/3149/22 and by Slovak state order contract "Submicrometer thin film technologies" No.2003 SO 51/03R 06 00/03R 06 01 is acknowledged.

References

1. E. Spiller, Soft X-ray Optics, SPIE, (Washington, 1994). 2. P. GrUnberg, Acta Mater. 48, 239 (2000). 3. R.P. Feynman, R.B. Leighton, M. Sands, The Feynman Lectures on

Physics, (Addison-Wesley Publ. Co 1966). 4. M. Stroemer, S. Faehler, H.U. Krebs, A. Crespo, P. Schaaf, Nucl. Instr.

Methods B 122, 503 (1977). 5. E. D'Anna, G. Leggieri, A. Luches, M. Martino, G. Majni, G. Barucca, P.

Mengucci, S. Luby, E. Majkova, M. Jergel, Thin Solid Films 343 - 344, 206(1999).

6. S. Balasubramanian, D.D. Nolte, M.R. Melloch, J. Appl. Phys. 88, 4576 (2000).

7. E. D'Anna, S. Luby, A. Luches, E. Majkova, M. Martino, Appl. Phys. A 56,429(1993).

8. S. Luby, E. Majkova, M. Jergel, R. Senderak, A. Anopchenko, E. D'Anna, G. Leggieri, A. Luches, M. Martino, P. Mengucci, G. Majni, A. Di Cristoforo, Mater. Sci. Engn. C19, 145 (2002).

9. E. D'Anna, A. Luches, M. Martino, M. Brunei, E. Majkova, S. Luby, R. Senderak, M. Jergel, F. Hamelmann, U. Kleineberg, U. Heinymann, App. Surf. Sci. 106, 166(1966).

10. E. Majkova, S. Luby, M. Jergel, A. Anopchenko, Y. Chushkin, G. Barucca, A. Di Cristoforo, P. Mengucci, E. D'Anna, A. Luches, M. Martino, Hsin-Yi Lee, Mater. 5c/. Engn. C 19, 139 (2002).

11. E. Majkova, S. Luby, M. Jergel, Y. Chushkin, E. D'Anna, A. Luches, M. Martino, P. Mengucci, G. Majni, Y. Kuwasawa, S. Okayasu, Appl. Surf. Sci. 208-209, 394 (2003).

12. H. Howari, D. Sands, J.E. Nichols, J.H.C. Hogg, T. Stirner, W.E. Hagston, J. Appl. Phys. 88, 1373 (2000).

13. Y. Takata, H. Haneda, T. Mitsuhashi, Y. Wada, Appl. Surf. Sci. 189, 227 (2002).

14. W. Brueckner, S. Baunack, M. Hecker, J. -I. Moench, L. van Loyen, CM. Schneider, Appl. Phys. Lett. 11, 358 (2000).

15. R. T. Huang, F. R. Chen, J. J. Kai, I. F. Tsu, S. Mao, W. Kai, J. Appl. Phys. 89,7625(2001).

16. R. Senderak, M. Jergel, S. Luby, E. Majkova, V. Holy, G. Haindl, F. Hamelmann, U. Kleineberg, U. Heinzmann, J. Appl. Phys. 81, 2229 (1997).

DYNAMICS OF INTERACTING ADATOMS ON COMPLEX SURFACES

Z. CHVOJ

Institute of Physics, Academy of Sciences of the Czech Republic Na Slovance 2, 182 21, Praha 8, Czech Republic

E-mail: chvoj@fzu.cz

The collective behaviour of atoms adsorbed on a surface is determined by the thermodynamic properties of the system of adsorbates. Under certain conditions, systems of atoms with repulsive interaction form ordered surface structures with specific dynamic and thermodynamic features. In the present paper, we have investigated an ensemble of interacting adatoms on surfaces with two non-equivalent adsorption sites, the adsorption energy and the number of which are different in the elementary cell. We have taken an fee surface with hollow sites (higher adsorption energy) and bridge sites as a model lattice. For repulsive interaction, the coverage dependence of the critical temperature is non-monotonic. The adsorption isotherms exhibit features characteristic for the system with simultaneous repulsion and attraction between ad-particles. The complex thermodynamic properties of adatoms result in step like dependence of the diffusion coefficient D on coverage and in non-Arrhenius temperature dependence of D.

1. Introduction

The theoretical description of the various surface phenomena observed in experimental studies presents a considerable challenge, since it not only has to include the influence of the underlying substrate but also the lateral ad-particle-ad-particle interactions in a two-dimensional layer. Substantial insight into surface processes has been achieved by developing simple lattice models based on short-range lateral adsorbate interactions. In the majority of models of diffusion on crystal surfaces the particles occupy equivalent lattice sites formed by the minima of the potential relief. Such simplification differs substantially from real surfaces, where different binding sites have been experimentally identified. To estimate the role of this effect we considered a lattice with fourfold hollow and two-fold bridge adsorption sites on a fcc(lOO) surface (Fig. 1).

In previous papers [1-5] the dynamic properties of this lattice were studied. The resulting chemical diffusion coefficient showed characteristic features as a function of coverage, such as a maximum or step at specific coverage; a characteristic behaviour also observed in various experiments [6-8]. It must be emphasized that the variation in the diffusion coefficient arises from the presence of two different sites, even in systems without interaction between particles. The increasing interaction resulted in the shift of the critical coverage to a step-like increase in the diffusion coefficient. Such behaviour is not clear

187

188

and in [3] it was speculated that it is a result of the formation of ordered phases. Thus, in the present paper we will determine the phase diagram of atoms adsorbed on a fcc(lOO) surface with two non-equivalent adsorption sites and discuss its consequence on the dynamics of adatoms.

Figure 1. Schematic view of a lattice with two kinds of sites. The circles at the middle of the square lattice bonds denote A-sites. Squares denote four-fold c-sites and the filled circles denote substrate atoms forming a fcc(lOO) surface with square symmetry.

2. Model of Adlayer and Hamiltonian

We take an ideal solid surface (with the symmetry of a fcc(lOO) crystal face) as a model lattice. We will assume that the adsorbed particles occupy the minima of the potential relief which forms a two-dimensional lattice shown in Fig. 1. In this case the thermodynamical state of the ad-particle system is completely described by a set of site occupation numbers «/ with

fl, if site is occupied n,={ (1)

[0,if site is empty here the index / = 1,2,...,N labels the i-th site of the lattice. These sites have different symmetry, i.e. four-fold hollow sites in the vertices of the square lattice (c-sites) and two-fold bridge sites in the middle of the square lattice bonds (b-sites). One can easily see that the number of 6-sites Nh is two times greater that the number of c-sites Nc

In principle, the depths of the b- and c-sites can be also different. We will assume that the depths are equal to eb and ec. for b- and c-sites, respectively. In addition, we have considered only lateral repulsive interaction between the nearest-neighbour nn pairs such that

Nb=-N, N=-N " 3 ' 3

189

of ad-particles. Each fe-site has two nn c-sites and each c-site has four nn 6-sites. The strength of the interaction between nn sites is set to K. The system Hamiltonian H can than be written as:

H=-fc& -£bYr>+rL"inJ (2)

here symbols <c>, <b> and <nri> denote summation over all c-sites, 6-sites and the nn pairs, respectively. The thermodynamical properties of the system are described by the free energy per particle/ which we define as follows (in units of kBT)

f = N~l\nQ (3)

e=Zexp M

• N \

kT\^n''H (4)

The grand canonical partition function Q has the form where u is the chemical potential of ad-particles; summation is carried out over all 2N configurations of the system.

3. Phase Diagram

The phase diagram of the system described in the previous paragraph was determined by the Real Space Renormalization Group method [5,9-12]. Phase diagrams for repulsive interaction are plotted in Fig. 2. It should be noted that the repulsion between ad-particles produces quite interesting phase diagrams. Firstly, a phase transition occurs in the narrow interval of surface coverage

1 2 ML <©< M L . The other interesting point is the non-monotonous

dependence of the critical temperature Tc on the surface coverage at the low coverage side of the phase diagram for £c-eb<A, and at the high coverage side for

ec-eb>X (the phase diagrams are symmetrical around © = M L only for £c-

Eb=X). It is possible, in principle, to observe three first order phase transitions, changing the temperature and keeping the surface coverage constant in close

1 2 vicinity of the critical line 0 = M L + A for ec-eb<A. ( © = M L - A

3 3

for £c-£b>X), where A « M L . The behavior shown by our system is the

result of switching between the dominant role of the difference in adsorption energies and the interaction energy at different temperatures. The difference in

190

adsorption energy leads to preferential occupation of c-sites. On the other hand, the repulsive interaction results in the separation of atoms on c and b sites. To characterize the phase transition we divide the whole lattice into two subsystems: particles located at c-sites form the first phase, while particles adsorbed in 6-sites form the second phase. The high temperature produces random distribution of ad-particles over both sublattices. When the temperature is lowered, ad-particles tend to occupy all 'deep' c-sites first. This is true only for

0 < — ML . When the mean surface coverage exceeds ML, some ad-

particles should occupy 'shallow' 6-sites. The interaction forces particles

adsorbed in nn c-sites to leave these sites and occupy further 6-sites. Then, for

1 2̂ 1 the whole surface should separate into regions of ordered c-phase e 3 3

(with a completely occupied c-sublattice and empty 6-sublattice) with local

0 = — ML and a 6-phase with reversed occupation of the sites (local 3

0 = — ML). The ratio of the partial 'volumes' of both phases is determined by 3

the mean surface coverage 0.

0,4 0,5

Surface coverage ®

0,7

Figure 2. Phase diagrams for repulsive interaction. The values of the parameters are as follows: X =4 kJ/mol, Ec-ei, = n kJ/mol, where n is the number of the curve. Inserts denote the detailed course of the phase boundary line at 0 - 1/3-ML and the schematic distribution of atoms on the adlayer under critical temperature. Filled circles or squares denote occupied sites.

191

We calculated the mean surface coverage as well as the individual coverage of c and b sublattices as functions of the chemical potential u at different temperatures. These dependencies give us an insight into the structure of the ordered phases and support the above discussion.

We define the individual coverage as follows:

b 2 B dsk

0C =3kBT (5)

The free energy/is obtained using the RSRG transformation (see[9-12]). The surface coverage <9can then be expressed via the sublattice coverage

6>=26V3 + 6>C73. (6)

The mean coverage & and partial coverage @cb are plotted in Figs. 3 and 4 as functions of chemical potential u. For high temperatures, all isotherms are close to the Langmuir isotherm (the dashed curves in Fig. 3).

1,0-

0,8 CD

SP 2 0,6

I 8 0,4

rP3

oo 0,2

0,0

\ V

i • - i „ _ . — • i . • ^ - j

/

i • i • i • i

y |

i |

. |

i |

.

I \

-4 12 16

(n+ig / I^T

Figure 3. Adsorption isotherm s for repulsive interaction. The values of the parameters are as follows: Si-si, = A = 4 kj/mol, Tc * 157 K. Curve 1— Langmuir gas, 2 — T = 1000 K, 3 — T= 250 K, 4 — T= 170 K, 5 — T= 150 K, 6 — T= 100 K.

192

l.O-i

0 5 10

(M+8h) / kRT 15

<X5° 1

&°^ e j

9-> g <MH o (U

. o 0,4-tt c3

-8 0,2-00

0,0-

l/-h

-

-

1 \

/ \

i • i •

(H+eb)/kBT5 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 c-sites. 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 <9(,/,(u) give us a detailed insight into the distribution of adatoms on the surface. The low temperature behavior of the c sublattice occupation is also unusual (see Figs. 4a,b). At first 0C increases up to 1, meaning that almost all c-sites are occupied. This situation corresponds to a mean surface coverage ©= 1/3-ML. A further increase in the number of ad-particles leads to a decrease in the mean occupation of the c sublattice 0C and an increase in 0b. The repulsive interaction plays a dominant role, decreasing the simultaneous occupation of the adjacent c and 6-sites. The adlayer disintegrates into islands with pure c or b occupied sites. At 0 -> 2/3-ML, 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. Ad-particles behave identically to non-interacting ad-particles. The b and c-sublattices 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 ad-particles. 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 = ec- eb and X) we use in the discussion:

k = exp

f v = exp

f AE^

kj (8)

1 ' (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(<9), Kl&) and I\0).

The calculated behaviour of these quantities, together with D{&) for v = 100 and for three particular values of k, is shown in Fig. 5. For k=l0'2 the difference in binding energy between the b and c sites is much higher than the repulsion energy and the features of D, r and Kd (curves 1 in Figs. 5a,b and c, respectively) are correspondingly similar to non-interacting systems [14]; the diffusion coefficient increases step-like at & =0.33 where the thermodynamic factor exhibits a narrow maximum. Furthermore, the jump rate factor increases at this coverage.

The activation energy for a jump out of a c site is rather high and consequently both the jump diffusion rate and the diffusion coefficient are very small. At <9= 0.33, all additional particles must occupy b sites and diffusion will take place via the less activated b-b hopping. The diffusion coefficient thus

195

3.0

2.5

2.0

1.5

1.0

0.5

0.0

-

1

3

v

k=103

k«0.3 k=1

= 100

,

r

i . _ i —

1

— •>

— . — " • 1

Figure 5. a) Quantitative dependence of D (normalized) on coverage ©for v = 100. Curves 1,2 and 3 correspond to the it-values of 10"3, 0.3, 1, respectively, b) Quantitative dependence of r (normalized) on coverage <9for v= 100. Curves 1,2 and 3 correspond to the A-values of 10"3, 0.3, 1, respectively, c) Quantitative dependence of thermodynamic factor K,, on coverage 0 for v = 100. Curves 1,2 and 3 correspond to the it-values of 10"3, 0.3, 1, respectively.

196

increases by one step to a value corresponding to lower activation energy and the jump diffusion coefficient increases because it shows the average hopping rate of the whole system. The decline of r at higher coverage is due simply to the typical site blocking effect of a lattice gas model. The other end of the parameter range is reached for k=\. The system's behaviour is entirely governed by the repulsive interaction and is illustrated by the curves numbered 3 in Fig. 2. There is also an increase in the diffusion coefficient but with the onset starting at <9=0.5 and smoothening over a wider range of the coverage. .Tand Kd similarly, exhibit mild increases with the latter showing a broadened, less intense maximum at this coverage.

5. Summary

We studied thermodynamic properties and surface diffusion in a simple model of an adlayer occupying the non-equivalent fourfold hollow (c) and bridge (b) adsorption sites on a fcc(lOO) surface. It was shown that the phase diagram for repulsive interaction differs considerably from the phase diagram of an ordinary lattice gas with equivalent sites, illustrating an unusual non-monotonous dependence of the critical temperature on surface coverage. The characteristic feature of the low temperature phase is the separation of the adatom system into regions of complete occupation of the c or b sublattice. The adlayer under the phase boundary line is composed of regions of completely occupied c or b sites.

Diffusion was studied with respect to three parameters: 0, the total coverage including both sites; k, the parameter describing the difference in activation energy for diffusion of both sites; v, the change in activation energy due to the occupation of an adjacent site of the other kind. The diffusion coefficient, and also Kj and r, vary in a characteristic way at a critical coverage 0C, which shifts continuously from 0C =1/3 (in cases where the difference between the two sites dominates the interaction) to 0C =1/2 (in cases where the interaction determines the system's behaviour). In intermediate situations the position of 0C reflects the equilibrium balance between the two tendencies.

Acknowledgments

This work was supported by the Grant Agency of the Academy of Sciences of the Czech Republic under grant No. IAA1010207.

References

1. Z. Chvoj, H. Conrad, V. Chab, M. OndfejCek, A.M. Bradsha, Surf. Sci. 329, 121 (1995).

2. Z. Chvoj, H. Conrad, V. Chab, Surf. Sci. 352-354, 983 (1996).

197

3. Z. Chvoj, V. Chab, H. Conrad, Surf. Sci. 426, 8 (1999). 4. Z. Chvoj, H. Conrad, V. Chab, Surf. Sci. 442, 455 (1999). 5. A.A. Tarasenko, Z. Chvoj, L. Jastrabik, F. Nieto, C. Uebing, Phys. Rev.

B63, 165423 (2001). 6. A.G. Naumovets, M.V. Paliy, Yu.S. Vedula, Phys. Rev. Lett. 71, 105

(1993). 7. A.G. Naumovets, V.V. Poplavski, Yu.S. Vedula, Surf. Sci. 321, 321 (1988). 8. E.D. Westre, D.E. Brown, J. Kutzner, S.M. George, Surf. Sci. 294, 185

(1993). 9. Th. Niemeyer, J.M.J.van Leeuwen, Physica 71, 17 (1974). 10. M. Nauenberg, B. Nienhuis, Phys. Rev. Lett. 33, 1598 (1974). 11. B. Nienhuis, M. Nauenberg, Phys. Rev. Lett. 35, 477 (1975). 12. B. Nienhuis, M. Nauenberg, Phys. Rev. Bl l , 4152 (1975). 13. R. Ferrando, E. Scalas, M. Torri, Phys. Lett. A 186,415 (1994). 14. Z. Chvoj, H. Conrad, V. Chab, Surf. Sci. 376, 205 (1997).

COPPER SURFACE SEGREGATION DURING V2Os THIN FILM DEPOSITION

M. M. AHADIAN, A. IRAJI-ZAD

Department of Physics, Sharif University of Technology,

P. O. Box 11365-9161, Tehran, Iran E-mail: ahadian@mehr.sharif.edu

M. GHORANNEVISS, M. HANTIZADEH

Plasma Physics Research Center, LA. U.

We observed surface segregation of copper onto vanadium pentoxide thin films in V2Os/Cu/Si samples with oxide thickness of 50-200 nm during room temperature deposition. Auger Electron Spectroscopy (AES) and X-ray Photoelectron Spectroscopy (XPS) showed that Cu exists on the surface of the samples. The thickness of accumulated Cu was estimated to be about one monolayar. The amount of detected Cu was lower after Ar ion bombardment. Secondary Ion Mass Spectroscopy (SIMS) was used to measure Cu and V concentration profiles to investigate the mechanism of diffusion. The results indicate the accumulation of Cu on the surface and a small concentration of Cu in the V205 layer. Therefore, Cu surface segregation is confirmed using AES, XPS and SIMS techniques.

1. Introduction

The phenomenon in which some component of an alloy or compound accumulates at a surface is called surface segregation. Surface segregation causes the concentration of materials near the surface to change from that in the bulk. This plays an important role in surface science and engineering. Much research has been focused on this subject during the last decades. In the case of binary metallic alloys systematic experiments and studies have been performed [1]. Although the nature of surface segregation is complex, there are semi-empirical theories and simulations that can be used to explain experimental results in metallic alloys. Sometimes it is possible to use surface segregation to control surfaces properties. In addition to works done on binary alloy, studies of ternary and quaternary have been performed recently [2].

Surface segregation also occurs in thin films but there are some differences compared to bulk materials. As in thin films, atoms often diffuse more rapidly, this phenomenon can be observed at lower temperatures compared to bulk samples [3]. In addition, in thin films surface segregation can be observed even in compound layers that are stoichiometric in bulk form. The simplest situation is the segregation of substrate material during deposition of a thin film. In recent

198

199

years, systematic studies on metallic systems have been performed and there are some models for surface segregation of metallic substrate onto the metallic layer

[!]• Surface segregation often occurs during heat treatments. It may also occur

during thin film deposition where it affects in particular the quality and properties of multi-layers such as those deposited by molecular beam epitaxy (MBE). Although it is expected that this effect can be observed only on ultra thin films, but there are few reports for surface segregation during deposition of rather thick films.

We have reported Cu surface segregation during thin film deposition of V205 by thermal evaporation [4]. This phenomenon was also observed during sputtering deposition [5]. V205 thin films have been used in many applications like catalysts for oxidation, lithium solid-state batteries, and gas sensors. Addition of Cu into a V2Os thin film can modify its physical properties. Therefore, Cu surface segregation during V2Os deposition might be used to fabricate layers with desired properties. Also, metallic Cu layer forms good electrical contacts in device fabrication and for certain application its diffusion should, however, be stopped in order to produce a pure V205 layer.

In this work, we have studied surface segregation of Cu onto V205 during deposition. We performed AES, XPS and SIMS experiments on V205/Cu/Si samples and Sn02/Cu/Si samples for comparison.

2. Experimental

Samples were deposited onto Si wafers in a high vacuum system by thermal evaporation. Initially, a 200 nm Cu layer was evaporated onto the Si substrate. Then a V205 layer with a thickness of 50-200 nm was deposited from an alumina boat onto the Cu layer. The base pressure in the system was 10"5 torr. For the sake of comparison Sn02 (200 nm)/Cu (200nm)/Si samples were prepared under the same conditions. The substrate temperature did not increase by more than 100 °C during deposition. The samples were transferred to an AES/XPS system just after the deposition. A 5 keV electron beam gun and a dual Mg and Al anode X-ray source were used for AES and XPS analyses, respectively. A concentric hemispherical analyzer (Specs model EA10 plus) was employed to measure the energy of emitted electrons from the surface of the sample. The pressure during these experiments was 10"9 torr. An Argon gun was used to clean the surface of the samples with 500 eV Ar ion. SIMS experiment was done using Cameca 6f instrument with Cs ion bombardment.

200

3. Results and Discussion

As we mentioned before, surface segregation changes the composition of the surface. This effect influences usually only the top few monolayers and in some cases only one monolayer [6]. Therefore surface analysis techniques can be employed to investigate this phenomenon. AES and XPS are sensitive to about 3-5 nm in depth and can be used to obtain information about the nature of the elements and their chemical states at the surface.

AES analysis of the V205/Cu/Si samples with oxide thickness of 200,100 and 50 nm indicated the presence of V, O, C and Cu at the surface of the deposited films. The existence of Cu at the surface indicates surface segregation during the deposition of the vanadium pentoxide film. The carbon signal vanished after a short ion bombardment. Hence, the carbon peak is due to contamination. After more bombardment, the intensity of the Cu peaks decreased while the vanadium and oxygen peaks increased. This means that the Cu concentration substantially decreases below the top layers. The amount of accumulated Cu was independent of V205 layer thickness in the range of 50-200 nm. After the Cu peaks had vanished, the ratio of O/V was measured to be 2.5 ± 0.5 [4]. XPS spectra also show Cu peaks indicating Cu accumulation at the V205 surface (Figure 1). Similar to the AES results, Ar bombardment reduced the intensity of the Cu peaks in XPS spectra.

Cu

^y^\ Cu

Cu O Cu

V

25000

20000

•15000 J3 c

10000 8

^ . UA I- 5000

0 1400 1200 1000 800 600 400 200

Binding Energy (eV)

Figure 1. XPS spectra of V205 (50nm)/Cu(200)/Si sample. The experiment was performed using the Al anode.

To understand the mechanism of surface segregation, the chemical state of the materials is important. The segregation energy depends on the chemical state

201

of the elements at the surface i. e., Cu and V. It is possible that vanadium pentoxide reduces to lower oxidation states during evaporation in vacuum. We focused on XPS results to understand the chemical state of vanadium using XPSPEAK software. XPS peaks of vanadium and oxygen are shown in figure 2.

530 525 520 515

Binding Energy ( eV )

Figure 2. XPS result of V2Os (50nm)/Cu(200nm) /Si indicating V205. The anode was Al and 5 min ion bombardment was done to remove C contamination on the surface.

The existence of the V 2p3/2 peak at 517 eV and the spin orbit splitting of 7.3 eV between the V 2p3/2 and V 2pi/2 peaks confirm V205 composition [7]. However, the 2.8 eV FWHM of the V 2p3/2 peak is more than the expected value. This could be related to amorphous state of the deposited layer. The O/V peak area ratio is more than the value reported for V205. This is due to extra oxygen at the surface because of Cu accumulation. Our XPS and AES results together confirm that the chemical state of Cu is Cu20. For the sake of comparison an AES experiment was done on a Sn02 (200nm)/Cu (200nm)/Si sample. The deposition condition was the same as for the V205/Cu/Si samples. The result in figure 3 shows Sn and O peaks but no sign of Cu on the surface. However, after annealing the Sn02 /Cu /Si samples in air at 500 °C for 3 hours, AES spectra indicate Cu on the surface.

An important question is the thickness of the Cu layer. The presence of vanadium peaks in the spectra indicates that the accumulated Cu thickness is below the inelastic mean free path of electrons, which is about 1-5 nanometers. The simplest model for accumulated Cu is a uniform layer on the V205 layer. Using this simple model it is possible to calculate the Cu layer thickness by the ratio of elemental peaks for different binding energies. The basis for this

202

estimation is the dependence of the mean free path on electron energy. We compared the ratio of Cu peaks in the AES spectra of the vanadium pentoxide samples and those of a thick Cu layer. In this way, the Cu thickness was estimated to be about one monolayer.

n 4

3 u 3 Sn02/Cu/si before

annealing Sn02/Cu/Si after annealing

200 400 600 000

Energy (eV)

Figure 3. AES result for SnO2(200 nm)/Cu(200nm)/Si sample before and after annealing at 500 °C for 3 hours in air atmosphere.

Another method for thickness determination is using background electrons in XPS spectra. Inelastic scattered electrons form the background of the peaks. The background carries information on the thickness of the layers, because the depth from which electrons are emitted is important in the probability of inelastic scattering. We used Tougaard's model that uses the ratio of the peak area (Cu in this case) and the increase in the background signal at 30 eV higher binding energy [8]. This method also yields a thickness estimate of the segregated Cu layer of about one monolayer.

To address further the question of surface segregation, the Cu depth profile was measured using SIMS. In SIMS experiments, a fine ion beam sputters the sample and the secondary ions are collected and analyzed by a mass spectrometer. Therefore, SIMS can be employed to obtain information about the concentration of compounds versus depth. The extreme range of detection makes SIMS a powerful technique. To avoid edge effects, which are due the secondary ions from the edge of the crater, the result was obtained from the central area of the crater. The SIMS result for V205 (100nm)/Cu /Si sample (figure 4a) indicates a Cu peak at the surface. At a greater depth the Cu concentration decreases substantially. Our results in these systems based on SIMS showed that the Cu concentration in the V205 layer is small and that Cu accumulates at the surface. SIMS in addition to the surface experiments discussed above indicates that the Cu accumulation is indeed due to surface segregation.

203

Figure 4b shows SIMS result for the V205 (50nm)/Cu /Si sample. In this case Cu accumulation at the surface is observed as well. But in this sample there is a mixing between Cu and V205 layers. Usual diffusion can mix layers but if we compare figures 4a and 4b we see that the thickness of the mixed region does not increase with deposition time. We conclude that at the beginning of the V205

deposition, the surface segregation is different from that occurring after thicker layer deposition. At the beginning, the Cu concentration is high and a V2Os

island is surrounded with Cu, but as the thickness increases the Cu concentration is reduced and only a few monolayers of Cu are present at the surface. Because our SIMS measurements are less sensitive to the surface compare with AES and XPS results we calculate the thickness of the accumulated Cu using AES and XPS experiments. In order to confirm this description more experiments need to be performed.

10000 -, a)

w 1000 • V c r^ v

8 100 - h\ C u

1 0 - M — • — - , , — 0 500 1000 1500 2000

Sputtering "Time(s)

10000 j

1000 •

V)

o 100 •

I 10 +

0 500 1000 1500

Sputtering Time ( s )

V)

Figure 4. a) SIMS result for V20 5 (100nm)/Cu(200nm) /Si sample. and b) Result for V20 s (50nm)/Cu(200nm) /Si sample.

204

4. Conclusions

In this work, we have studied copper surface segregation on V205 and Sn02

layers during deposition. The XPS and AES results show that Cu accumulates during V205 deposition. SIMS results show that the mechanism of accumulation is not ordinary diffusion. There is no sign of Cu segregation when Sn02 is deposited under the same condition.

Acknowledgements

The authors wish to thank Research Council of Sharif University of Technology for financial support and Dr N. Taghavinia for his useful suggestions and reading the manuscript. The authors also wish to thank Mr. Khajeaminian for AES and XPS results and Mr. Haji Hoseini for SIMS experiment.

References

1. M. Yoshitake, et al, J. Vac. Sci. Technol. A. 19, 1432 (2001). 2. G. Bozzolo, et al, Computational Material Science 15, 169 (1999). 3. M. Yoshitake, K. Yoshihara, Applied Surface Science 100-101, 203 (1996). 4. A. Iraji-zad, M. M. Ahadian , Z. Vashaei, J. Phys. D: Appl. Phys.35 1167

(2002). 5. H. Miyazaki, M. Kamei, I. Yasui, Thin Solid Films 343, (1999) 168. 6. Y. H. Chung; Surface Science and Spectroscopy; 2001; Academic Press,

California, USA, p. 119. 7. A. Kohl, E. Taglauer, H. KnOzinger, Phys. Stat. Sol. (a) 173, 85 (1999). 8. D. Briggs, M. P. Seah (eds), Practical Surface Analysis, vo/.l, John Wiley,

Chichester, England, 238 (1994).

THE PREPARATION AND SURFACE STUDIES OF Fe/Pt THIN FILMS

G. VARGHESE

Crystal Physics Centre, St Berchmans' College, Mahatma Gandhi University

Kottayam, Kerala 686101 India

E-mail: gvar@sify.com

Thin film deposition of Fe/Pt over silica substrates using rf/dc magnetron sputtering technique is studied. The microstructure and surface roughness of the coated samples are investigated using grazing incidence x-ray reflectomerty and secondary ion mass spectroscopy. The analysis of specular reflectivity data provides information along the depth of the film

1. Introduction

Two-dimensional material systems play very vital role in present day technology, which include the multitude of uses in electronics, in biotechnology and in pharmacology. Recent developments in film growth techniques and in the theoretical advancements in nucleation processes have made the ground clear for further progress in this field. The preparation of materials with tailored characteristics require a detailed understanding of the atomic interactions, and how these are influenced by factors such as composition and preparation process. For example, in ultra thin films used for magnetic recording, the data storage capability is sensitive to variations in the compositions. Deviations from flatness of surfaces or interfaces may have substantial effect on their coercive-and demagnetizing-fields, domain walls, electrical conductivity, and giant magnetoresistance [1,2]. The predetermined operating parameters would help us to design a system having a desired ability.

Characterization of a film is very important to know whether the desired level of perfection has been achieved during the growth process. Always the quality of the interfaces and surface roughness depends very much on the growth parameters and chemical constituents. Surface characterization using x-rays is widely used owing to their high resolution and penetration power.

Magneto electronics is a rapidly expanding field in basic research as well as in industrial application [3,4]. Surface and interface determined properties of magnetic layers of thin films have attracted much interest in high quality magnetic films. The knowledge about the surface structures and inter layers of such magnetic multilayer systems will enable us to obtain progress in tailoring a

205

206

variety of materials exhibiting exotic properties [5]. Recently metallic multilayers with perpendicular magnetic anisotropy have been reported as alternative to magneto-optic storage materials. Compositionally modulated multilayer films of iron (Fe) and platinum (Pt) have also been reported as potential candidates for high-density magnetic recording because of their ability to achieve small grain sizes [6-8]. The magnetic properties of these films are strongly dependent on the relative as well as the absolute thickness of Fe and Pt layers.

In this article the fabrication of nearly perfect Fe/Pt films on silica substrates and their surface characteristics are reported. DC and RF magnetron sputtering technique under high vacuum conditions grow the bi-layer films. The thickness of Pt was varied while keeping Fe layer thickness unaltered. The film deposited is found to have preferences in the (111) plane geometry for both Fe and Pt under normal growth conditions. The microstructures of the layers were analyzed by X-ray diffraction (XRD), X-ray reflectivity with Cu-Ka radiation source, and by secondary ion mass spectroscopy (SIMS). X-ray scattering technique have an advantage over others since it can probe buried layers with Angstrom resolution as well as being a non-destructive technique.

2. Experimental

2.1. Sample Preparation

The Fe Pt films were deposited on chemically clean bare silica substrates using sputtering technique in cryogenically controlled high vacuum chamber [9]. The silica substrates were pre-cleaned with acid and base solutions and degreased in alcohol vapor. A heat treatment was also carried out on the substrate at 350°C to remove any further contamination. The substrate was then cleaned in situ by ion bombardment in an Ar atmosphere. Prior to sputtering, the base pressure was maintained 5.5x 10"6 Torr, and the substrate was kept 15 cm above the sputtering source. During deposition process the substrate temperature was kept below 350°C. The target materials, Fe and Pt were 99.99% pure, and the argon ion beam, at a pressure of 10"3 Torr, was used to sputter the target. Ultra clean Ar gas with an impurity level less than lppb was used as the process gas. No impurities were introduced into the chamber. Fe was deposited on the silica substrate using the RF plasma at 100W power for 30 min without rotating the substrate. Over the Fe film, a thin layer of Pt was deposited in the DC mode with 25 W power out put for 12 min. To maintain uniform thickness, the carriage containing the substrate was rotated at the rate of 60 rpm. Different samples were prepared by varying the power as well as time of deposition.

207

2.2. X-ray Reflectivity Studies

Thickness measurement of thin films is a basic problem, which can be easily tackled by X-ray scattering methods. The most accepted technique using X-ray scattering was the one developed by LG Parratt in 1954 [10]. The Fresnel reflectivity of layer deposited on a semi-infinite substrate is:

R = {[r, +r2 exp. (-2ikzt)]/[l+r,r2exp (-2ik2t)]}2

where xx and r2 are the Fresnel reflectivity coefficients at the free surface and the substrate interface respectively. The term kz is the vertical component of the wave vector of the beam transmitted through the layer of thickness, t. The intensity of the beam decreases rapidly with angle of incidence. Intensity maximum occurs when argument of exponent reduces to zero and this will generate an oscillatory pattern in intensity peaks [Kiessig fringes], corresponding to different angles of incidence. The condition for this is a modified form of Bragg relation,

2t[sin2ai-sin2ac]1/2 = nX

where ot; is the incident angle, etc is the critical angle for total external reflection and n is the order of the peak.

In most cases the incidence angle is too small and therefore the above equation can be approximated to measure the thickness as:

t = nA/2[a2i-a2c]"

2

In the reciprocal space the equation has a further simple look;

t = 2n/Aqz

where Aqz = 4 n A6/X, is the difference in the scattering vectors inside the crystalline layer.

Grazing incidence X-ray reflectomerty was employed both in specular and off specular modes to find the layer thickness and surface roughness. Figure 1 shows the Kiessig fringes obtained on the specular scan on one of these samples grown under different conditions. The high-resolution reflectivity data with step size of 0.00014 A'1 and resolution 0.0018 A'1 was obtained using an 18 kW rotating Cu-anode (Enraf Nonius FR591). Si (111) crystal was used to choose the Cu-Ka radiation and a triple axis spectrometer was employed to collect the data.

208

1000

c o o

0.001

2000 4000 eooo mdeg

Figure 1. The grazing incidence X-ray reflectivity of the Fe-Pt film deposited over silica substrate [o data, solid line fit].

2.3. XRD Studies

The lattice structure of the films was measured with X-ray diffraction. The XRD pattern was recorded on Philips X-ray photometer for various 20 values from 0 to 80°. The Figure 2 shows the XRD patterns recorded.

80

60

I " 20

Pt(111) i

Fe (111)

20 40 . ...)j.i.ii«1»...iit.i li>liwf I 1 " I

[°20] 6 0

Figure 2. X-ray diffractogram of FePt films deposited over silica substrate.

-t—r-% 80

209

2A. SIMS Analysis

SIMS measurements on the Bragg mirror were performed in CAMECA-ims-4f ion microscope. Oxygen ion beam was used as a Primary ion source, accelerated and rastered across the sample surface over an area of 250x250 urn2. Secondary ions emitted due to the collision of accelerated oxygen ions were collected from the central portion of the rastered area. The depth resolution was above 50 A.

3. Results and Discussion

Nearly perfect FePt films were prepared by vacuum sputtering method on silica (111) substrates. These films were grown non-epitaxially on silica substrates. Since the orientation of these non-epitaxial films is different from that of epitaxial ones, these films are particularly important. The specular scan detail shows that the sample for which Pt thickness was low has a higher degree of the surface roughness. The layer thickness can be controlled by deposition time at each target. Substrate temperature during the deposition was maintained at about 35°C by water-cooling system. The total film thickness was found to lie between 65-70 nm. The peaks of the reflectivity curve indicate the roughness of the surface and interface.

The SIMS studies show expected surface composition of Fe and Pt elements over reasonable depth.

The XRD studies confirm that Fe layer crystallizes in bcc structure (111) and Pt layer grows in (111) direction of fee structure. Only (111) peaks were observed in each profile. Looking for a face-centered cubic structure for Fe at elevated temperatures will be interesting, since only a relaxed fee phase can support the anti ferromagnetic coupling at top layers [11]. Tomoyuki Maeda etal [12,13] have reported that the system attains a more ordered structure upon annealing above 300°C. The absence of (001) and (110) peaks here indicates that ordering has not set in these samples [10,11]. Promoting (001) and (110) orientations will suppress the grain growth. This can be achieved by reducing the film thickness. For binary alloys the driving force for disorder- order transformation will be large.

Detailed structure analysis of the films annealed at different temperatures is also in progress. The electron density profile, film roughness, etc are also to be evaluated from the reflectivity data. Results obtained in XRD measurements have a direct link to the magnetic properties of the samples. Only VSM studies can alone reveal the magnetic properties of these films. The magnetization ratio depends on the layer thickness, impurity concentration and the type of the substrate. Satoshi Miura et al. [14] reported that bare Si substrate has a strong influence on the Giant magneto Resistance for Fe layers. The anti ferromagnetic

210

coupling energy between the layers has a strong correlation with the microstructures of the layers, especially in regards to interfacial roughness.

Acknowledgements

The author is grateful to the Indian Academy of Sciences for the award of summer research fellowship and to Surface Physics Division, SINP, Kolkota, for sample analysis and discussions.

References

1. G. Palasantzas, Y.P. Zhao, G.C. Wang, T.M. Lu, J. Barnas, J.Th.M.De. Hosson, Phys. Rev. B 61, 11109 (2000).

2. Hartmann (Ed), Magnetic Multilayers and Giant Magnetoresistance, (Springer 1999).

3. S.D. Bader, Surf. Science 500, 172 (2002). 4. J.F. Gregg, I. Petej, E. Jouguelet, C. Dennis, J. Phys. D Appl. Phys 35,

R121 (2002). 5. G. Palasantzas, Y.P. Zhao, G.C. Wang, T.M. Lu, J. Barnas, J.Th.M.De

Hosson, Phy. Rev B 61, 11109 (2000). 6. C. Christides, I. Panagiotopooslos, D. Niarchos, T. Tsakalakos, A.F.

Jankowski, J. Phys. Condens. Matter 6, 8187 (1994). 7. CM. Kuo, P.C. Kuo, H.C. Wu, Y.D. Yao, C.H. Lin, J. Appl. Phys. 85,

4886(1999). 8. G.D. Waddil, J.B. Tobin, A.F. Jankowski, J. Appl. Physics 74(11),

6999(1993). 9. S.K. Chen, F.T. Yuan, W.C. Chang, Tsung-Shune Chin, J. Mag. and Mag.

Materials 239, 471(2002). 10. L.G. Parratt: Phys. Rev. 95, 359 (1954). 11. R.E. Camley, D. Li: Phys. Rev. Lett. 84, 1947 (1999). 12. T. Maeda, T. Kai, A. Kikitsu, Toshihiko, J. Ichi Akiyama, Appl. Phy. Lett.

80,2147(2002). 13. H. Zeng, M.L. Yan, N. Powersw, D.J. Sellmeyer, App. Phys. Lettt. 80,

2350 (2002). 14. S. Miura, D. Takahasi, M. Tsunode, M. Takahasi, J. Appl. Physics 91(7),

4461 (2002).

V. NANOMATERIALS

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1D-NANOSTRUCTURES ON SILICON CARBIDE THIN FILMS

P. G. SOUKIASSIAN

Departement de Physique, Universite de Paris-Sud, 91405 Orsay Cedex, France E-mail: psoukiassian@cea.fr

The atomic scale ordering and properties of cubic silicon carbide thin film surfaces are investigated by room and high temperature scanning tunneling microscopy. In this review, I focus on the Si-terminated P-SiC(lOO) surfaces only. Self-formation of Si atomic lines and dimer vacancy chains on the p-SiC(lOO) surface is taking place at the phase transition between the 3x2 (Si rich) and c(4x2) surface reconstructions. Using a rigorous protocol in surface preparation, it is possible to build very long, very straight and defect free Si atomic lines, forming a very large superlattice of massively parallel lines. These self-organized atomic lines are driven by stress. They have unprecedented characteristics with the highest thermal stability ever achieved for nanostructures on a surface (900°C) and the longest atomic lines ever built on a surface (urn scale long). Investigating their dynamics, we learn that their dismantling at high temperature results from collective and individual mechanisms including one-by-one dimer removal. Overall, this is a model system especially suitable in nanophysics and nanotechnologies.

1. Introduction and Historical Background

Silicon carbide (SiC) is certainly not a new material since it is older than the solar system. Indeed, SiC has been discovered in 1895 by Henri Moisan (1904 Chemistry Nobel Prize laureate) on a meteorite located in the Diablo Canyon (Arizona) [1]. Initially, silicon carbide has been established for its excellent mechanical properties as "carborundum" since it was primarily used for many decades as a hard material (the highest hardness after those of diamond and boron nitride). SiC now became very well known as an advanced material having many versatile and promising applications in e.g. matrix composites, biocompatibilty or microelectronics [2-4]. In the latter field, SiC appears to be especially suitable for high-power, high-temperature, high voltage, high frequency and radiation resistant electronic devices and sensors [2-4]. Its average figures of merit scale up to 3 orders of magnitude above those of conventional semiconductors such as Si or III-V compounds, SiC being outclassed only by diamond [5-10]. Figure 1 shows the representative figures of merit of various conventional and novel semiconductors according to the criteria established by Keyes (high speed logic and high integration density electronic devices) [5] and by Johnson (high power, high speed, high temperature and high voltage analogic devices) [6].

213

214

lOOOOf

High Power, High Temperature and High Speed Analogic Devices

00 High Speed Logic Devices

Si GaAs InP GaN SiC Diamond Figure 1. Figures of merit of various semiconductors according to the criteria of Keyes [5] (high speed logic devices) and Johnson [6] (high power, high temperature and high speed analogic devices).

Furthermore, SiC is chemically rather inert which, combined with its excellent ability to resist to radiation damages, makes it a very suitable material for harsh environments [2,9]. Also, SiC is a "refractory" IV-IV compound semiconducting material belonging to the class of wide band gap semiconductors (together with diamond and group III nitrides) and a very high thermal stability [2-4]. This makes it very useful for operations at elevated temperatures (> 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 self-organized 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 [8,12-16,19,33-38].

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 p-SiC(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 STM topograph. 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 STM topograph showing the Si-Si dimers forming atomic lines on the p-SiC(lOO) c(4x2) surface, b) 125 A x 125 A STM topographs (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

Fi l led S t a t e s IMVSJHV Sc : ! f ' 1

Filled States

X X 1

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 ranging from 25 °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 - 25 *€ ^ T-600°C

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.

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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 dimer-dimer 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 results from an 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 results from both individual (one-by-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 ranging from a 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 circuit from the 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).

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53, 13121(1996). 28. P. Kackell, J. FurthmUller, F. Bechtedt, G. Kresse, J. Hafner, Phys. Rev. B

54,10304(1996). 29. A. Catellani, G. Galli, F. Gygi, Phys. Rev. Lett. 11, 5090 (1996). 30. A. Catellani, G. Galli, F. Gygi, F. Pellacini, Phys. Rev. B 57, 12255

(1998).

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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

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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. Aboulet-

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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 - p0) / Po ] x 100

where PH (RH) is the resistivity (resistance) of the sample in presence of the magnetic field, p0 (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 Mn03). 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 random- access 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 magneto-tunneling 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], Cu-FeNi[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

Cu8]Co,9

CU80CO20

Cu70Fe3o Cu8oFeioNi10

Co30Ag7o

Cu80Co2o Ni-Co-Cu/Cu

C o x Cllioo-x

C1190C010

Cu70Co3o Co18Cu72

Cu8 4Co ] 6

Co2oCu80

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

15 7 5 10 0.6 4

4.5 6.2 3

Temperature (K) 10 5 5

4.2 5

4.2 300 300 77

300 4.2 300 300 300

H (Tesla)

2 5 5 7

1.5

14 0.8 1.3 0.9 1.0 1.0 1.5 1.5 1.5

Ref.

[51 [6] [61

[171 [24]

[261 [50] [53]

[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.

Voltmeter or

Oscilloscope

Keithley 224 programmable constant current source

. . • I 1 f

Reference electrode

Cathode

Anode

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/(pm2-p s

2)1 / 2cos2e 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 low-temperature 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 systems-DMS 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

* • >

s< l_ 4>

4-1 4>

E CO L . CO Q. o>

ttic

CO _ l

j o * : -

3.60-

3.58-

3.56-

3.54-

D n -u U

metastable phase

i • i '

n-a

\ Curich \ \ \ \ \ 1 \ \ \ Co rich

M x o -1 - • 1 • 1

D

Cu

Co

1 • — - o

200 400 600 800 Annealing Temperature ( C )

Figure 2. Phase and lattice parameter change in the electrodeposited Cu75Co25 film on annealing [43].

It is to be noted that although the stable phase of Co at temperature lower than 425°C, is HCP, no trace of HCP phase was detected in the present case. The bright field TEM images showed that the average crystallite size in the as deposited film C lies in the range 5-20 run and increases to 40-60 nm on annealing the films.

3.2. Giant Magnetoresistance (GMR) of Electrodeposited Cu-Co Films

The requirement of a conducting substrate for electrodeposition could be a drawback of the process since during MR measurement it leads to current shunting in the current-in-plane (CIP) geometry, which reduces the signal in electroposited thin films. Electrodeposition of GMR thin films directly onto semiconducting substrates, such as suitably doped Si, is a better alternative, as the conductivity of these substrates could be good enough to allow the electrodeposition without significant short circuit during the transport measurements due to the relatively higher resistivity of Si. High quality thin films of composition Cu8oCo2o with excellent metallic lustre were, therefore, obtained on n-Si (100) by electrodeposition (samples S), Besides, Cu75Co25 films are deposited at room temperature onto Cu-coated alumina substrates, at a pH of 6.0 and deposition current density of 2mA/Cm2 (sample C). The resistance of the as-deposited and annealed films was measured at room temperature and upto 10 kOe, and employed for calculating MR. As shown in figure 3 the variation of room temperature MR with applied magnetic field for a sample C annealed at

238

different temperatures revealed that the magnitude of the negative MR increases with both the increase in magnetic field and sintering temperature. Maximum room-temperature MR at a magnetic field of 10 kOe increased from 0.65% in as-deposited sample to 4% in a sample annealed at 450°C for 1 hr. MR values in the present case are comparable to the reported value of 6.7% MR at 15kOe. In another sample C an MR value of-7.8% was estimated at 20K even at a smaller magnetic field of 30 kOe. The MR as well as magnetization studies are found to be supportive of the presence of fine Co particles even at lower temperatures of annealing, though these are observable only at higher annealing temperatures in XRD studies [Figure 2].

*5"

s

0 -

•

-1 -

- 2 -

. 4 -

• *

J . ' 1 ! : **x ******

as-deposited T * • - # . TA = 350aC T r * • * * • TA = 4G0*C * * » » , . T , = 450aC ^

i 1 • 1 • 1 • r

•

•

r

• m

* * • 4

* W

, ! .

a 4 6 H {kOe)

10 12

Figure 3. Magnetoresistance vs. magnetic field, measured at room temperature for electrodeposited Cu75Co25 films [43].

The resistivity of a thin film is known to decrease with the increase in the film thickness. This may lead to larger MR for thicker films, even if there is no significant increase in change in resistance. This indeed has been experimentally observed in the case of samples C. The MR values of three films having different thickness i.e. 150, 210 and 300 nm were estimated to be 3.3, 3.6 and 4%, respectively.

GMR measurement at room temperature is carried out on 350 nm thick Cu8iCoi9 films electrodeposited on n-Si (sample S) when the applied magnetic field is varied upto lOkOe. It can be seen from the figure 4 that the maximum magnetoresistance of the as-deposited film S is nearly 1%. The maximum GMR of a similar film deposited on Cu-coated A1203 substrate is (film c) -0.6%. The higher value of the magnetoresistance of the film on S is attributed to the higher resistivity of Si in comparison to that of the coated Cu layer (~20nm) on A1203.

239

The optimum MR obtained for a film S (on n-Si) annealed at 425°C for 30 min is 2.7%. The value of MR for a film C is 3.6% that is higher than the value of the film S. The lower value of MR for the annealed film S is attributed to the formation of silicide upon annealing [45]. The MR of magnetic granular films is directly proportional to the electron mean free path for the matrix. As the resistance of copper silicide is larger than the copper silicide matrix is less than that for copper matrix. Hence a reduced MR is observed for the annealed films S as compared to films C. Further increase in annealing temperature to 450°C or higher results in lowering of MR owing to the formation of non-ferromagnetic cobalt silicide. The maximum value of GMR at 10k in a magnetic field of 50 kOe is - 5 % . This typical data confirms the presence of ferromagnetic very fine Co (or highly Co-rich) nanograins, which are responsible for the observed GMR behavior even in the as-deposited state [60].

H(kOe)

Figure 4. Magnetoresistance vs. magnetic field of CusuCom films electrodeposited on (a) n-Si: (i) as deposited (ii) annealed at 350°C for 1 h, (iii) annealed at 425°C for 15 min, and on Cu/Al203: (v) as deposited and (vi) annealed at 425°C for 30 min [45].

4. Conclusions

The research activities in the area of GMR materials are gaining momentum owing to their potential applications in magnetoelectronics VLSI, MEMS and magnetic sensors. Various material systems, both in nano-multilayers and nano-granular thin films can be successfully deposited by using electrodeposition, a versatile technique. The technique has successfully been employed to prepare nanowires and for interconnects and packaging.

240

Nanogranular Cu-Co thin films have been gavanostatically electrodeposited on Cu-coated glass or alumina. The as deposited films are found to consist of a single metastable fee phase of Cu-Co alloy with a grain size of 5-20 nm. On annealing at 700°C a total phase separation into Cu (fee) and Co (fee) was detected. High quality Cu-Co films have been directly deposited onto n-Si substrate, thereby eliminating the need of a conducting layer, which is required for depositing films on an insulating substrate. A systematic study of the effect of various processing parameters including bath temperature and deposition current density has led to the conclusion that electrodeposition is a nucleation and growth controlled process like a physical vapour deposition technique.

The value of MR for Cu-Co films was also found to increase many times with the increase in temperature of annealing, which in turn controls the microstructure of the granular thin films. A maximum room temperature MR of 4% was observed at 10 kOe for films grown on conducting alumina and annealed at 400°C. At 20 k and 3 kOe the maximum MR increased to 7.8%. In the case of the films grown on n-Si substrate too, the room temperature MR was observed to increase with annealing, temperature. The annealing induced enhancement in MR values has been attributed to the segregation/separation of Cu and Co phases resulting into magnetic nano-granular structure.

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238, 145 (2002).

SELF-ASSEMBLED QUANTUM DOTS: STRUCTURAL AND OPTICAL PROPERTIES, AND DEVICE APPLICATIONS

M. HENINI

School of Physics and Astronomy, University of Nottingham, Nottingham NG7 2RD, U.K.

E-mail: Mohamed.Henini@Nottingham.ac.uk

Low dimensional structures (LDS) form a major new branch of physics research. They are semiconductor structures, which have such a small scale in one or two spatial dimensions that their electronic properties are significantly different from the same material in bulk form. These properties are changed by quantum effects. Throughout the world mere is increasing interest in me preparation, study and application of LDS. Their investigation has revitalised condensed matter science, in particular semiconductor materials. These complex LDS offer device engineers new design opportunities for tailor-made new generation electronic and photonic devices. New crystal growth techniques such as molecular beam epitaxy (MBE) and metal-organic chemical vapour (MOCVD) deposition have made it possible to produce such LDS in practice. These sophisticated technologies for the growth of high quality epitaxial layers of compound semiconductor materials on single crystal semiconductor substrates are becoming increasingly important for the development of the semiconductor electronics industry. This article is intended to convey the flavour of the subject by focussing on the technology and applications of self-assembled quantum dots and to give an elementary introduction to some of the essential characteristics.

1. Introduction

Crystal growth and post-growth processing technologies have developed to the extent that it has become possible to fabricate semiconductor structures whose dimensions are comparable with inter-atomic distances in solids. These structures are known as LDS. The movement of charge carriers in these structures are constrained by potential barriers. This results in the restriction of the degrees of freedom for motion to two, one or even zero. The system becomes two, one or zero dimensional depending on whether the potential barriers confine the carriers in one (layers), two (wires) or three (dots) dimensions (Figure 1), respectively. Carriers exhibit wave-like characteristics and when the layer thickness is comparable with the carrier wavelength the carrier motion is constrained and exciting new physical properties result.

The study of LDS began in the late 1970's when sufficiently thin epitaxial layers were first produced. The layers used in the early investigations became available following developments in the technology of the epitaxial growth of semiconductors, mainly pioneered in industrial laboratories for device purposes.

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One of the mam directions of contemporary semiconductor physics is the production and study of sttuctures with a dimension less than two: quantum wires and quantum dots, in order to realize novel devices that make use of low-dimensional confinement effects. During the last few years much attention has been devoted to the strain in the grown layer and characterization of self-assembled semiconductor quantum dots (QDs). The strong interest in these semiconductor nanostaictures is motivated by the possibility to use them as active media in future high-speed electronic aid photonic devices.

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2. Fabrication Methods of Quantum Dots

Several methods for the fabrication of QDs have been reported over the last decade including lithography-based technologies. Although this technique is widely used to provide QO predominantly by the combination of high-resolution electron beam lithography and etching, the spatial resolution required for reaching the size regime where significant quantization effects can. be expected tends to be larger than the desirable level In addition, lithographic methods and subsequent processings often produce contamination, defect formation, size non-uniformity, poor interface quality, and even damage to the bulk of the crystal

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Itself. A new attractive method of defect free lOnm scale QD fabrication Is the Stranski-Krastanov (SK) growth in lattice-mismatched systems. In the SK growth mode, the mismatched epitaxy is initially accommodated by biaxial compression in a layer-by-layer (2D) growth region, traditionally called the wetting layer. After deposition of a few monolayers the strain energy Increases and the development of Islands (3D) becomes more favourable than planar growth [1] (Figure 2).

Figure 2. Scanning tunnelling microscope pictures (100x100 ran) of InAs/GaAs QDs grown by MBE on (100), (311)A and (311)B GaAs substrates [1] As can be seen, using substrates with different orientation can control the shape of the QDs.

In the III-V semiconductor material system, SK growth has been used to grow InAs islands on GaAs and it has been shown that the size fluctuation of dots is relatively small (<10%) and the small dots and surrounding host matrix are dislocation-free and strained coherently with GaAs. It has been reported that the InAs p*owth mode changes from 2D to 3D upon the deposition of less than 2 monolayers of InAs, so as to reduce the strain in grown layer, since there is about a 7% lattice mismatch in the GaAs/InAs system. The strained (In?Ga)As/GaAs material system has been the most widely studied for which various quantum effects have been demonstrated. Various combinations of III-V semiconductors based on phosphorus or antimony compounds, and Si/SiGe alloys have also been studied.

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The advantages of this technique of QD fabrication are that no nanotechnology and no further etch or implantation induced process is necessary. Since the dots are grown in-situ a homogeneous surface morphology is maintained and defect creation is avoided. However, the inherent problem associated with this method is the size non-uniformity and the position uncontrollability of the QD. Controlling the dimension and arrangement of the self-organized 3D structures is thought to be very important for obtaining good properties of the structures.

The islands become technologically more interesting if it is possible to manipulate their arrangement laterally and vertically (Figure 3) in order to achieve the three-dimensional arrays. There are already several reports on spontaneous lateral ordering due to the preferential nucleation along surface steps. Kitamura et al. [2] demonstrated successful alignment of InGaAs by using a 2° off (100) GaAs substrate with multi-atomic steps in MOCVD growth process. Pre-patterned substrates have also been used for ordering of QDs in a more direct way. Miu et al [3] grew by MBE on etched GaAs gratings and found islands to form on the sidewalls of ridges running along [1-10] direction. Similar results were obtained by Jeppesen et al. [4] for Chemical Beam Epitaxy (CBE) deposited InAs islands in wet-etched and partially overgrown, trenches and holes on a (100) GaAs surface. They formed chains of InAs islands aligned in trenches along [011]. The chains of islands have 33nm minimum periods.

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Figure 3* Schematic diagram showing ordering of quantum dots due to strain fields effects to form (a) vertical alignment, and (b) lateral alignment on patterned substrates.

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The vertical alignment is expected and the total density can be increased by stacking the QDs with a spacer layer. Vertically aligned and electronically coupled islands has several advantages including the application of the tunnelling process to novel electronic devices such as single electron tunnelling devices, the study of tunnelling dynamics between QDs, and the high QD density for QD lasers. Several groups have recently successfully grown stacked InAs self-assembled QD structures separated by GaAs spacer layers by MBE Sugiyama et al. [5] reported vertically aligned InAs QDS up the ninth layer with 2.5nm monolayers InAs and 1.5nm GaAs spacer layers. Solomon et al. [6] demonstrated arrays of InAs islands, which are vertically stacked, vertically aligned and electronically coupled in the growth direction. They have achieved vertical alignment of up to 10 islanding layers with no associated dislocation generation.

3. Photoluminescence Properties of InAs/GaAs Quantum Dots

In this section I will report on the photoluminescence properties of multiple (InGa)As/(AlGa)As QD layers grown by molecular beam epitaxy under different conditions (i.e., different Al content, number of QD layers, and different spacer thickness between QD layers). We found that by varying the Al content in the (AlGa) As matrix and/or stacking several QD layers, the room temperature dot luminescence is tuned over a wavelength range from 0.8 urn to 1.3 urn.

Three different sets of samples were considered. In the first set, three InAs layers were embedded in an ALGa^As matrix grown at TG=520 °C. The average thickness, L, of each InAs layer is 1.8 monolayer, ML. The three InAs layers are separated from each other by 20 nm-thick ALGaiyYs barriers (y=(0.0 - 0.8)), resulting in uncoupled dots. In the second set, 1 and 10 layers of InAs dots (L=1.8 ML) were embedded in a GaAs matrix. The vertically stacked InAs layers were separated from each other by a distance d = 1.7 nm. This structure was grown at TG = 500 °C. Finally, in the third set of samples, three InAs layers were embedded each of them in a GaAs/Al<uGao.7As quantum well and grown at rG=500 °C. The dot formation was controlled in situ by monitoring the reflection high-energy electron-diffraction pattern. Photoluminescence (PL) measurements were performed from 7M0 K to T ~ 500 K. The optical excitation was provided by the 514.5 nm line of an Ar+ laser. The luminescence was dispersed by a 3/4 m monochromator and detected by a cooled Ge diode.

Figure 4 shows the room temperature QD PL emission for three representatives InAs QD samples. The three samples have different Al content in the AlyGai.yAs matrix surrounding the dots or have different spacing {d) or number (N) of the InAs QD layers. With increasing y, the QD PL band blue-

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shifts from 1.1 urn (sample A: y = 0, d = 20 nm, N = 3) to 0.8 urn (sample B: y = 0.8, d = 20 nm, N = 3). This blue shift is due to the deeper carrier confining potential of the dots at higher values of y. In contrast, with decreasing d and/or increasing N, the PL red-shifts from 1.1 urn (sample A: y = 0, d = 20 nm, N=3) to -1.3 nm, (sample C: y = 0, d = 1.7 nm, N = 10), evidence for electronic coupling between vertically stacked QDs. Therefore by engineering the carrier potential profile of the dots it is possible to cover a broad energy range for the room temperature light emission of QDs. This is of particular interest for extending the optical emission range of QDs to 1.3 (am, the window for signal transmission through silica fibers.

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Samples shown in Figure 4 exhibit a different thermal behaviour. In fact the thermal stability of the dot emission is strongly dependent on the composition of the matrix incorporating the dots. We found that the dots embedded in a (AlGa)As barrier and/or in a GaAs/(AlGa)As QW exhibit the highest thermal stability. This can be attributed to the low level of thermal escape of carriers from the dots towards the high energy levels of the AlGaAs barrier or the GaAs/(AlGa)As QW, which therefore act to prevent carrier depopulation of the dot levels.

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4. Quantum do Teasers

QDs find significant interest especially for application in laser diodes. It is worth noting that the device, which benefited most from the introduction of quantum wells (QW), is the injection laser. The QW laser reached mass production within very few years because of its low cost, high performance and high reliability. QDs are believed to provide a promising way for a new generation of optical light sources such as injection lasers. QD lasers are expected to have superior properties with respect to conventional QW lasers. Theoretical predictions [7] of the intrinsic properties of QD lasers include higher characteristic temperature T0 of threshold current, higher modulation bandwidth, lower threshold currents and narrower linewidth (Figure 5).

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The principal advantage of using size-quantized heterostructures in lasers originates from the increase of the density of states for charge carriers near the band-edges (Figure 1). When used as one active medium of a laser, this results in the concentration of most of the injected nonequilibrium carriers in an increasingly narrow energy range near the bottom of the conduction band and/or top of the valence band. This enhances the maximum material gain and reduces the influence of temperature on the device performance. Figure 6 shows the development of semiconductor diode lasers in terms of threshold current density as a function of time for various heterostructures based on double heterostructure (DHS), QWs and QDs. The recent developed QD lasers [8] based on InAs QD system have already showed record low threshold current density of 24 A/cm2 at room temperature for a wavelength of 1.28 urn. This is

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almost a factor of two lower than what has been achieved for QW lasers [9]. Recently, InAs QD laser diodes with high light output power of 4.7W at 1.135 u.m were reported [10]. Red light emitting QD lasers have also been successfully fabricated with AlInAs/AlGaAs [11] QD, and InP/GalnP [12] QD. It is possible to access new energies by combining materials with different lattice constants and energy gaps. Currently, the spectral range of III—V semiconductor QD lasers extends from near infrared (1.84 jam for InAs-(In,Ga,Al)As QD lasers on InP substrates [13] to the visible red range [11].

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5. Other Applications of QDS

Methods of detection of infrared (IR) radiation have been investigated for almost 200 years since the astronomer William Herschel discovered what is now called the infrared portion of the spectrum in 1800. Infrared detectors have been key components in thermal imaging, guidance, reconnaissance and communication systems.

All important applications of infrared techniques, both for military and civil purposes (sciences, meteorology, medicine, industry, etc) rely on the detection of radiation in the 1-3 urn, 3-5 urn and 8-14 um spectral range (the so-called atmospheric windows). The 8-14 um wavelength region is especially important for imaging since the temperature of the human and environments bodies is around 300K corresponding to a peak wavelength of thermal radiation of about 10 um. The materials, which cover the above wavelength regions, include II-VI, III-V and IV-VI compound semiconductors. With respect to the well-known HgCdTe detectors, GaAs/AlGaAs quantum wells devices have a number of potential advantages

252

including the use of standard manufacturing technology based on advanced GaAs growth and processing techniques, highly uniform and well controlled MBE growth on large GaAs wafers, high yield, greater thermal stability, and intrinsic radiation hardness.

Recently, there has been significant interest in developing novel QD infrared photodetectors (QDIPs). There are two major potential advantages of quantum dots over quantum wells as photodetectors [14], namely: (1) intersubband absorption may be allowed at normal incidence. In quantum well infrared photodetectors (QWIPs) only transitions polarized perpendicular to the growth direction are allowed, due to absorption selection rules. The selection rules in QDIPs are inherently different, and normal incidence absorption is, indeed, observed, (2) thermal generation of electrons is signicantly reduced due to the energy quantization in all three dimensions. Generation by LO phonons is prohibited unless the gap between the discrete energy levels equals exactly to that of the phonon. This prohibition does not apply to quantum wells, since the levels are quantized only in the growth direction and a continuum exists in the other two. Hence thermal-generation or recombination by LO phonons results, with a capture time of few picoseconds. Thus, it is expected that the signal-to-noise ratio in QDIPs will be significantly larger than that of QWIPs.

The operation of photovoltaic QDIP fabricated from (InGa)As/GaAs heterostructures have been demonstrated by Pan et al. [15]. These detectors are sensitive to normal incidence light. At zero bias and low temperature of 78 K, peak detectivity of 2x108 cm.Hz1/2/W, with a responsivity of 1 mA/W at a wavelength of 13 um, was obtained. The authors claim that this is the higest detectivity achieved for a QDIP operating in the photovoltaic mode, and device might be attractive for focal plane array imaging applications.

As reported above QDs can be used to make better optical devices. They can also be used to make devices with a totally new functionality. Another application is an optical data storage medium, in which bits of information are stored as a single or few electrons within the dots. These memory devices based on QDs should require a very low switching energy because the information is stored as a single or few electrons. Optical illumination can be used to write the charge to be stored in the QDs. It is predicted that because each QD would carry a bit of information, the memory devices could potentially achieve ultra dense storage capacities of Terabit/in2.

Sugiyama et al. [16] at Fujitsu Laboratories, Japan, observed memory effect of InAs QDs buried in Schottky barrier diode with memory retention time of 0.48ms for the first time. The band diagram of their structure which contained a single layer of QDs is shown in Figure 7. Two sequential laser pulses with an interval time At irradiated the diode. Electron-hole pairs are generated in the InAs QDs after a pulse

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Eradiation torn the first laser. The electron escape from the QDs by tunnelling or thermo-ionic emission, and the holes stay in the QD. The residual holes decrease the photocurrent when the second laser pulse is irradiated. The retention time of the optical memory was determined from the interval time, At, dependence of the photocttrrent difference !write and Iretd. These are preliminary results. However, this group believe that there is a possibility of using QDs for high density optical memory.

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Figure 7. Device showing memory effect of InAs quantum dots.

Some devices in which electrons are transported in the vertical direction have shown a clear memory effect [17] of InAs QDs. Very recently, Son et al. [18] have demonstrated the memory operation of InAs QDs in a lateral heterojunction field effect transistor (FET) structure. Lateral transport occurred through the parallel conduction of the two-dimensional channel and QD layers. This channel, which is induced by gate bias, is strongly influenced by the charge stored in InAs QD layers. This was shown by the current hysteresis curve and the capacitance-voltage measurements. The authors believe that the memory operation is due to the charge trapping effect of InAs QDs.

QD memoiy devices using optical illumination to store charge in the QDs require a sensitive method of detecting the trapped photo-excited charge within the QDs. Shields et al. [19] demonstrated that using a transistor structure (Figure 8) it is possible to detect the presence of single photo-excited carrier in a single QD. Conventional semiconductor single-photon detectors rely upon the avalanche process. However, in the device of Shields et al , the gain derives from the fact that the conductivity of the FET channel is very sensitive to the photoexcited charge trapped in the QDs. Their FET contains a layer of InAs QDs adjacent to the channel and separated from it by a thin AlGaAs barrier. The capture of a single photo-excited carrier by a QD leads to a sizeable change in the source-drain current through the transistor, allowing the detection of a single photon.

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Figure 8. Cross-sectional view of a QD FET structure with a small area Schottky gate and scanning electron microscope image of the gate region for an FET with a 2 jim wide mesa and 4 pan long gate (Courtesy of Dr AJ.Shields, Toshiba Research Europe Ltd). A positive gate bias charges the QDs with electrons, which limits the mobility of the adjacent electron channel. Single photons liberate a trapped electron via capture of a photoexcited hole. This results in a detectable increase in the conductance of the electron channel.

6» Conclusions

In this article, I have described progress made on some QDs devices. Other work on QD structures is processing in many laboratories worldwide. Some of the studies cited in this article, like many reported in international journals, are still in their infancy, and many challenges remain in the development of high-performance QD devices for optical and electronic applications. For the case of QD lasers it is believed that they are ready for practical applications and their future is extremely promising. However, the replacement of QW lasers with QD lasers depends on the industry and its willingness to switch to new technology. If recent history of DHS and QW lasers is any guide, it is likely that QDs will lead to entirely new classes of materials and devices.

255

References

1. M. Henini et al., Microelectronics Journal 28, 933 (1997). 2. N. Kitamura et al., Appl. Phys. Lett. 66, 3663 (1995). 3. D.S.L. Miu et al., Appl. Phys. Lett. 66, 1620 (1995). 4. S. Jeppesen et al., Appl. Phys. Lett. 68,2228 (1996). 5. Y. Sugiyama et al., JpnJ.Appl.Phys. 35, 1320 (1996). 6. G.S. Solomon et al.,Phys.Rev.Lett. 76, 952 (1996). 7. M. Asada et al., IEEE J. Quantum Elect. 22, 1915 (1986). 8. X. Huang et al., Electron. Lett. 36, 41 (2001). 9. N. Chand et al., Appl. Phys. Lett. 58, 1704 (1991). 10. R.L. Sellin et al., Electron. Lett. 38, 883 (2002). 11. K. Hunzer et al., J. Appl. Phys. 87, 1496 (2000). 12. T. Riedel et al., Jpn. J. Appl. Phys. 38, 597 (1999). 13. V.M. Ustinov et al., Tech. Phys. Lett. 24,49 (1998). 14. E. Finkman et al., Physica E 7, 139 (2000). 15. D. Pan et al., Appl. Phys. Lett. 76, 3301 (2000). 16. Y. Sugiyama et al., JpnJ.Appl.Phys. 35, 1320 (1996). 17. N. Horiguchi et al., JpnJ.Appl.Phys. 36, L1246 (1997) 18. H. Son et al., JpnJ.Appl.Phys. 40,2801 (2001). 19. A.J. Shields et al., JpnJ.Appl.Phys. 40,2058 (2001). 20. Zh.I. Alferov et al., Sov. Phys. Semicond. 4, 1573 (1970); Zh.I. Alferov et

al., Fiz. Tekh. Poluprovodn. 4, 1826 (1970) ; I. Hayashi et al., Appl. Phys. Lett. 17, 109 (1970); R.C. Miller et al., J. Appl. Phys. 47, 4509 (1976); R.D. Dupuis et al., Appl. Phys. Lett. 32, 295 (1978); W.T. Tsang, Appl. Phys. Lett. 39, 786 (1981) ; Zh.I. Alferov et al., Pis'ma v Z.Tekn.Fiz. 14, 1803 (1988.); N. Chand et al., Appl. Phys. Lett. 58, 1704 (1991); N. Kirstaedter et al., Electron. Lett. 30, 1416 (1994); N.N. Ledentsov et al., Phys. Rev. B 54, 8743 (1996); G.T. Liu et al., Electron. Lett. 35, 1163 (1999); R.L. Sellin et al., Appl. Phys. Lett. 78, 1207 (2001).

21. D. Bimberg et al., MRS Bulletin July 2002, p.53\, R.L.Sellin et al., Appl. Phys. Lett. 78, 1207 (2001).

PREPARATION AND CHARACTERIZATION OF ULTRATHIN FILMS AND FILM COATINGS FOR MICROELECTRONICS

Y.A. POGORYELOV

Institute for Magnetism, National Academy of Sciences of Ukraine 36-b Vernadsky Ave., 03142, Kyiv, Ukraine

E-mail: pogor@imag.kiev.ua epogor@ukrpost. net

The work is devoted to the preparation and characterization of new thin metal film media for various types of non-volatile memories, and also to the development of new techniques for characterization of film parameters in microelectronics. We studied the physics of indirect exchange coupling in trilayer thin-film structures based on rear earth (Tb) and transition metals (Fe) with the nonmagnetic spacer (Au), which were prepared by electron-beam evaporation in an ultrahigh vacuum system. Investigations were carried out using magneto-optical and magneto-transport techniques, including know-how based on the Hall-like effect at zero applied external magnetic field. Oscillations of the Hall resistivity with the change of its sign, typical for the Ruderman-Kittel-Kasuya-Yosida model of exchange interactions, were experimentally observed. Also, a new method is presented for determining thermophysical parameters. It can be applied for non-contact and non-destructive investigation and monitoring of the adhesion of coatings or films to the substrate, both during and after deposition.

1. Introduction

Modern progress in micro- and nano-electronics, especially in areas concerned with data storage systems based on non-volatile memories, demand ceaseless search for new ideas, creation of new materials and structures, and development of new techniques for materials characterization with the aim to improve materials properties.

At present new types of nonvolatile random access memory (RAM), e.g., magnetic RAM (MRAM), based on multilayer thin-film structures, are being developed. As is well known, when magnetic films are separated by nonmagnetic spacer, their magnetizations are coupled to each other by an exchange interaction through the conduction electrons of the spacer layer [1^4]. As the thickness of the spacer layer is varied, the coupling can oscillate in sign. As the magnetic moment of most of rear-earth (RE) atoms is 7-10 uB, the use of RE metal as one of the magnetic layers in a multilayer structure can result in considerable increase of the net magnetic moment of the system, when the magnetic sublattices are oriented in parallel. Such multilayer film structures can become the film analogue for permanent magnets, which can be applied in the elements of MRAM, magnetic tunnel junctions, etc. The existence of different types of interfaces in such structures complicates the understanding of the nature

256

257

of the exchange coupling [5-8]. That is why the investigation of the exchange coupling in bi- and trilayer film structures is topical. It is also known that exchange coupling depends on the degree of conduction electrons polarization. Enhancement of spin-polarization due to quenching of the kinetic energy of the system at application of the external magnetic field [9] can distort the intrinsic exchange effects. Therefore in order to get the real picture of magnetic interactions in the investigated film structures it is necessary to use not only the conventional measurement techniques, but also a technique, which will allow avoiding the application of the external magnetic field. Such technique can be based on the Hall-like effect, predicted by Hirsh [10].

At the same time data storage devices based on magnetic and magneto-optical carriers still hold their positions. Therefore not only the magnetic measurements are important when studying characteristics of materials and structures for data storage applications. The specificity of the recording process onto magneto-optical carriers and compact-discs (CD) stipulates not only the search and application of new materials, but also demands the knowledge of thermophysical parameters of the metal film data carrier on CD. Taking all these into account will provide the optimal data density and also reliability and durability of data storage under the condition of long-term usage.

2. Preparations and Characterization of Film Structures for MRAM Applications

2.1. Sample Preparation

Exchange coupling between ultrathin Fe (8 A) and Tb (12 A) layers separated by Au spacer of varied thickness (3 - 20 A) was studied. Film samples were prepared on Si substrate by electron-beam evaporation in an ultrahigh vacuum system with a background pressure of ~ 10"8 Torr. This vacuum system was designed and built especially for these purposes at the Institute for Magnetism of the National Academy of Sciences of Ukraine. Evaporation rates were about 0.2 - 0.5 A/s. Layer thickness and deposition rates were controlled by a quartz crystal monitor. Samples were protected by a 30 A thick layer of SiO.

When Fe is deposited directly on Si substrate the Fe-silicide can be formed at Si/Fe interface. In the case of ultrathin Fe films such silicide formation can significantly affect its magnetic properties. That is why thin nonmagnetic sublayer can be deposited directly on Si substrate prior to deposition of Fe layer to prevent the formation of a magnetically dead layer of Fe [11]. Thus another analogous set of samples was prepared, but with an Au layer of one monolayer (ML) thickness at Si/Fe interface.

258

2.2. Characterization of Film Structures by Conventional Methods

Polar magneto-optical Kerr effect (PMOKE), magnetoresistance and Hall effect measurements, including a new technique based on the Hall-like effect at zero applied external magnetic field, were used to characterize the films magnetically.

PMOKE measurements showed oscillations of the Kerr angle (9K) with increasing thickness (dAu) of the Au spacer. Also, magneto-optical investigations showed that introducing 1 ML of Au at the Si/Fe interface significantly increases the amplitude of the magneto-optic response from the whole structure, and at the same time does not alter the behaviour of magnetic interactions between Fe and Tb layers. This is demonstrated in Figure 1. The upper curve (a) corresponds to samples with an Au sublayer. The lower curve (b) corresponds to PMOKE data obtained on the samples without Au sublayer and prepared by molecular-beam epitaxy [12]. As is clearly seen, the periods of 9K oscillations coincide almost entirely. Oscillations of 0K are connected with the periodical change of the exchange interaction between Fe and Tb layers through the Au spacer with its thickness increase.

Figure 1. PMOKE data for Fe/Au/Tb film structures with (a) and without (b) Au sublayer at the Si/Fe interface.

The existence of exchange coupling oscillations in trilayer Fe/Au/Tb film structure is also supported by results of anomalous Hall effect (AHE) measurements [12]. It was shown that the Hall conductivity oscillates, changing in sign as a function of dAu (see Figure 2(a)). It means that owing to indirect exchange coupling magnetic moments of Tb atoms are aligned parallel or antiparallel to magnetic moments of Fe atoms while the Au spacer thickness is increasing.

259

2.3. Hall-Like Effect at Zero Applied External Magnetic Field

As it was mentioned above the exchange coupling can be affected by the application of an external magnetic field. To avoid such distortions and to evaluate intrinsic magnetic interactions in Fe/Au/Tb structures, Hall-like effect based technique was applied. The idea of this technique is based on appearance of the transverse spin current in ferromagnetic structures [10] at zero magnetic field. Therefore, measuring the transverse current Im, that appear in the film plane at application of the electric potential to the film-substrate system, one can obtain the information about intrinsic exchange processes. This idea was put into know-how for determining the magnetic properties of magnetic films [13]. Results of these studies are presented on Figure 2(b) in comparison with AHE measurements (a) and calculated curve (c), which was obtained in [14] on the basis of Bruno and Chappert theory [15] and adapted for our case.

0,02

s 0,01-o X

s <§ 0,00

a s -°>°i "I

(a) •••©••• Hall resistivity (p ) i ( b ) " * " I n ( H = 0)

(c) Calculations

-0,02 - i — i — i — i — i — | — i — | — i — i — i — | — i — | — i — | — i — | — i — | —

0 2 4 6 8 10 12 14 16 18 20 d , A

Au

Figure 2. Exchange coupling in trilayer Fe/Au/Tb structure as a function of dA„. (a) AHE easurement results [12]; (b) Hall-like effect at zero applied external magnetic field; (c) Theoretical curve [19].

The explanation of this Hall-like effect is based on the fact that there is one more contribution to the Anomalous Hall Effect, besides skew scattering and side jump. It arises simply from the fact that in general a spin current exists in ferromagnetic metals when a charge current exists [10]. Electrons carrying a spin and associated magnetic moment (m) experience a transverse force (F) when they are moving in a longitudinal electric field. When a spin-unpolarized

260

current flows in a metal, the spin-orbit interaction produces asymmetric scattering of the conduction electrons so that electrons with one particular spin direction, e.g., spin-up electrons, have a larger probability to be scattered to the right compared to spin-down electrons [16, 17]. Similarly, spin-down electrons would tend to scatter to the left more than spin-up electrons. If there is a net magnetization in the system there will be a magnetization current (Im) [18, 19] associated with the flow of electric current (I,), and the transverse force will give rise to a charge imbalance in a direction perpendicular to the current flow (Figure 3). As was shown in [20] the direction of such a current in ferromagnet/normal metal/ferromagnet system depends on the mutual orientation of the magnetizations of both ferromagnetic layers.

v f y v f y n

^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ 1 ^^L ^ \ * ^^^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^ ^

Figure 3. Scattering of spin-polarized electrons in ferromagnetic metal.

The current Im, which appeared in the film plane during the application of the electric potential to the film-substrate system, was measured [13]. Results of these investigations are presented on Figure 2(b) in comparison with the theoretical curve (Figure 2(c)) [14].

As is well seen from Figure 2 the possible influence of the external magnetic field is manifested by the difference of the Hall resistivity oscillations behaviour (a) [12] from the calculated one in comparison with data for zero magnetic field case (b). According to the RKKY model of exchange coupling oscillations should have damping behavior with the increase of spacer thickness, but oscillations of Hall resistivity obtained by conventional AHE measurements are almost undamped. Periods of oscillations observed for the Fe/Au/Tb system were about 2.5 and 8.9 ML, and correlate well with the oscillation periods 2.5 and 8.6 ML, which were measured for the Fe/Au/Fe system [21], and also with values of 2.51 and 8.6 ML, extracted [15] from measurements of the Au Fermi surface.

As one can see the presented technique more precisely reveals the character of oscillations of the exchange coupling in the investigated trilayer film structure as compared with the conventional AHE measurements.

261

3. Peculiarities of Characterization of Thermophysical Properties of Film Structures

In the systems which use thermal recording of information, the temperature source originating in the film after laser pulse irradiation can be considered as initial conditions for the substrate. Since the substrate in comparison with the film can be considered as a semi-infinite medium, let us use results of [22] in order to analyze thermostimulated processes in the former. In this work the three-dimensional distribution of the temperature field in a semi-infinite absorptive medium, the surface of which is irradiated by laser pulse, was obtained. From this distribution it is easy to obtain the distribution of thermal stresses in the system "metal film - substrate" by substituting in [22] the layer thickness that absorbs laser irradiation by the metal film thickness h. Volumetric stresses in the substrate and radial stresses, progressing in the substrate (in the central part of the interaction region the material undergoes tension and shrinks to its periphery) can lead to its detachment, corruption and, as a result, to partial loss of data.

The monitoring of the film carrier adhesion to the disc can be achieved by determination and differentiation of thermophysical characteristics of the carrier from one point to another as the laser beam is moving over the CD [23]. For the purpose of monitoring of the data carrier and its adhesion to the substrate during the process of data recording by the laser beam a new technique [24] of express-analysis of thermophysical characteristics of the metal-film data carrier on CD was developed.

The method is based on the analysis of thermal front propagation dynamics in the metal film during the irradiation of the film - dielectric substrate structure by laser pulse. Let thermal diffusivity of the substrate be much smaller than the one of the film (as« aj). Then in the case of a cylindrical film with radius r » h, the central part of which with radius r0« r is irradiated by the axially symmetric flow of the laser radiation

af=1.38r2/t1/2 (1)

here t12 is the time, during which half of the maximal temperature is reached at a distance r from the laser irradiated area, with the condition that the duration of the pilot laser pulse is smaller than the time constant of the film: vt < h2/af. This criterion allows neglecting the heat extraction to the substrate during the laser pulse. After determination of afi the maximal temperature Tm in the controlled point and the power density in the pilot thermal pulse Q it is easy to calculate other thermophysical parameters - heat capacity and thermal conductivity [24]. The heat capacity C is given by

262

C = TtQr2/^ (2)

and the thermal conductivity K is given by

K = afC/Trr2!! (3)

The usefulness of this technique was tested by means of a setup that allows measurements both in atmosphere and vacuum at different temperatures. The experimental setup included a neodymium glass solid-state laser (T,<~ 50 ns, A = 1,06 um), a vacuum chamber with sample holder and a device for sample heating, an optical system providing shaping of the laser beam and irradiation of the samples, and finally a temperature registration circuit. Registration of the temperature change AT < 10 K with an inaccuracy below 8 % was carried out using a photosensor connected via a DC amplifier to the input of an oscillograph. This setup allowed obtaining thermophysical characteristics of materials having phase transitions in a wide temperature. The pilot laser pulse power density Q was determined using a solid calorimeter. As an example of application of this technique Figure 4 shows the results of thermal conductivity measurements of films of 1 um in thickness on polycrystalline corundum, prepared by electron-beam evaporation of nonmagnetic Al, and also of ferromagnetic Fe which is applied for preparation of films for magneto-optic recording [25]. The reliability of this technique is confirmed by comparison of experimental results for Fe with the reference data [26].

1,0

0,8

0,2

0,1

n n

i

i

-

r A -

i

i 6 1

i

* i

1

• A

- Al -Fe

- -Fe [26]

.

-

•

-

T̂

~

-

300 600

T ( K )

900 1200

Figure 4. Thermal diffusivity of the metal films vs. temperature.

263

4. Conclusions

Oscillations of the Hall resistivity, changing its sign as a function of the Au spacer thickness in trilayer Fe/Au/Tb film structures were clearly observed. It means that at certain Au spacer thicknesses the magnetization of the Tb layer is aligned parallel to the magnetization of the Fe layer. The possibility to control the direction of Tb layer magnetization in the investigated trilayer structure lets us to anticipate large values of the net magnetization (equivalent to several Tesla) in multilayer structures (Fe/Au/Tb/Au)N with N ~ 40-50. This is a promising result for the further development of nonvolatile electronic memories based on such multilayer structures. Also by means of the new technique based on Hall-like effect it was shown that magnetic interactions in trilayer Fe/Au/Tb structures are possibly influenced by the external magnetic field.

Besides, the characteristics of optical data carriers with thermal type of recording can be optimized by applying a method developed for the determination of thermophysical parameters of metal films. The method can also be applied where new materials exist in very small quantities only, but are sufficient for deposition of a thin film sample. Due to its high sensitivity the method can be applied the study of other products comprising thin film coatings. The method can be used for non-contact and non-destructive investigation and monitoring of the adhesion of coatings or films to the substrate both during and after their deposition.

References

1. B. Heinrich, J.F. Cochran, Adv. Phys. 42, 523 (1993). 2. S. S. P. Parkin, Annu. Rev. Mater. Sci. 25, 357 (1995). 3. P. Bruno, J. Phys.: Condens. Matter. 11, 9403 (1999). 4. A. T. Costa Jr., J. d'Albuquerque e Castro, and R. B. Muniz, Phys. Rev. B.

56, 13697 (1997). 5. Z. S. Shan and D. J. Sellmyer, Phys. Rev. B. 42, 10433 (1990). 6. A. E. Freitag and A. R. Chowdhury, J. Appl. Phys. 85, 4696 (1999). 7. A. E. Freitag and A. R. Chowdhury, J. Appl. Phys. 85, 5756 (1999). 8. H. Hoffmann, R. Scherschlicht. In Festkorperprobleme ed. by Helbig

(Vieweg, Braunschweig/Wiesbaden, 1998), 275. 9. C. B. Hanna et al., Phys. Rev. B 61, 13882 (2000). 10. J. E. Hirsch, Phys. Rev. B 60, 14787 (1999). 11. F. Zavaliche, W. Wulfhekel, H. Xu, and J. Kirschner, J. Appl. Phys. 88,

5289 (2000). 12. E. Shypil, A. Pogorilyy, Ye. Pogoryelov, T. H. Kim, G. Berera, J. Moodera,

J. Magn. Magn. Mater. 242-245 PI, 532 (2002).

264

13. A.F. Zhuravlev, Ye.A. Pogoryelov, A.N. Pogorilyy, E.V. Shypil. Method to detect magnetic properties of magnetic materials. Ukrainian patent No. 43616A, (2001).

14. Ye. Pogoryelov, Metallofiz. Noveishie Technol. 25, 47 (2003). 15. P. Bruno, L. Chappert, Phys. Rev. Lett. 67, 1602 (1991). 16. J.E. Hirsch, Phys. Rev. Lett. 83, 1834 (1999). 17. The Hall effect and its applications, ed. by C.L. Chien and C.R. Westgate

(Plenum, New York, 1980), 55. 18. M. Johnson and R. H. Silsbee, Phys. Rev. B 37, 5312 (1988), ibid. 37, 5326

(1988). 19. S.T. Chui, J.R. Cullen. Phys. Rev. Lett. 74, 2118 (1995). 20. S. Takahashi and S. Maekawa (unpublished). 21. J. Unguris, R.J. Celotta, D.T. Pierce, J. Appl. Phys. 75, 6437 (1994). 22. M.E. Gurevich, A.F. Zhuravlev, Yu.V. Kornyushin, A.E. Pogorelov,

Metallofizika 7, No. 2, 113 (1985). 23. A. Pogorelov, Ye. Pogoryelov, A. Zhuravlev, J. Magn. Magn. Mater. 249

(3), 428 (2002). 24. A. E. Pogorelov, A. F. Zhuravlev, Ye. A. Pogoryelov. Method to determine

thermo-physical characteristics of thin films. Ukrainian patent No. 43224A, (2001).

25. M. Naoe, N. Kitamura and T. Hirata, J. Appl. Phys. 61, 3337 (1987). 26. H. Gonska, W. Kiershpe and R. Kohlhaas, Z. Naturforsh B 23a, 783 (1968).

NANOCRYSTALLINE FILMS IN THE Ag-Ni SYSTEM

I.K. BDIKIN

Institute of Solid State Physics, Chernogolovka, Moscow distr., 142432, Russia

and

Department of Ceramic and Glass Engineering, CICECO, University ofAveiro 3810-193 Aveiro, Portugal E-mail: bdikin@issp.ac.ru

G.K. STRUKOVA, D.V. MATVEEV, S.A. ZVER'KOV, V.V. KEDROV, G.V. STRUKOV

Institute of Solid State Physics, Chernogolovka, Moscow distr. 142432, Russia

Ags-Ni|.s (x=0.0-1.0) films grown on Cu substrates by electrodeposition were studied. The films were found to be a nanocrystalline mixture of pure silver and nickel. The grain sizes were determined by X-ray diffraction and electron microscopy techniques. The minimal value was 3.3nm for the alloy with 70wt% Ni concentration. The stability of the grown films upon heating in air was examined. An increase in the grain size was found to begin at 150°C.

1. Introduction

The magnetic and nonmagnetic metal two-componental nanostructures are of interest not only in fundamental physics, but also because of their possible use in recording units and magnetoresistive devices [1]. The Ag-Ni system is among the least studied ones. We know only one work that reports obtaining a nanocrystalline Ag-Ni alloy by mechanical alloying [2]. But there are several publications on amorphous and nanocrystalline Ni-Co [3], Ni-W [4], Ni-Si [5], Ni-Fe [6-8], Ag-Co [9] alloys. The phase diagram of the Ag-Ni system has the simplest form, there are no intermediate phases, the elements are mutually insoluble and do not intermix even in the liquid state [10]. This may be due, in particular, to a large atomic misfit (-15%) of the crystal lattices of the elements with the same lattice type (fee).

The electrodeposition method of depositing thin metal films has many advantages. It can be carried out at ambient temperature and pressure, and therefore requires much less complex apparatus than vacuum-based techniques such as molecular beam epitaxy or sputtering. The electrodeposition with nonaqueous solution has additional advantages in comparison with aqueous solution: high level of ecology safety, opportunity of depositing larger number of metals and alloys.

265

266

Our goal here was to study the structure of Ag-Ni films of various compositions grown by co-electrodeposition (with non-aqueous solution containing Ag and Ni ions). This co-electrodeposition method was previously developed by the authors for different alloys of precious metals [11].

2. Experimental

Ag-Ni films on a copper substrate with typical thickness 100-500nm were grown by electrodeposition. For this we used a single electrolyte containing Ag and Ni ions [11]. Deposition was carried out at room temperature in a three-electrode in a standard electrochemical cell using potentiostatic conditions with potentials of 1.5-1.8 V, the current density being 2-5.4 mA/cm2. Potentials are quoted relative to a saturated calomel reference electrode.

Structures of the grown films were studied by X-ray diffraction techniques (SIEMENS D500, CuKcc radiation). The chemical analysis was performed on an X-ray microanalyser YXA-5, provided with an analytic system LINK AN-10000. The grain size was defined by transmission electron microscopy (JEM 100CXII).

3. Results and Discussion

Ag-Ni alloys of different concentration were grown by electrodeposition. The substrate was a ~0.3mm thick rolled copper foil. 0-29 diffraction patterns are shown in Fig. 1. The copper substrate peaks are seen. These peaks have intensity different from the reference copper specimen (Here and below we used PDF-2 Data Base, JCPDS International Centre for Diffraction Data). The 220 peak intensity is increased indicating texture. Comparatively narrow Ag or Ni peaks are seen only in Ni-depleted or Ag-depleted alloys, respectively. Pure Ni, however, exhibits broadened peaks, and pure silver has narrow peaks. The peaks intensity ratio of pure silver also differ from the reference intensity, which is also indicative of the presence of texture. As for the substrate, the 220 peak is enhanced. But if for the substrate this is accounted for by the rolling technology for the copper foil, then the electrodeposited silver film texture is caused by preferential epitaxial film growth. As the concentration of nickel in the alloy is increased, the peaks of silver shift towards the peaks of nickel, broaden and their intensity ratios become close to the reference ones. No peaks of pure nickel are present on the diffraction patterns. At a high nickel concentration (>20wt%) only broadened double film peak is observed.

(53

30,0 40,0 50,0 60,0 70,0 80,0 2 9 ( ° )

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 (<10nm) and the sample is two-phase. One phase corresponds to pure Ag, the other to pure Ni. It has to be noted that there were regions of coarse grains (more often nickel). Based on the electron diffraction data, our X-ray diffraction patterns must be analysed in terms of a nanocrystalline metal mixture rather than their solid solution.

The conclusions drawn from the electron microscopic study of the alloy correlated with the X-ray diffraction data. A detailed analysis of the X-ray diffraction patterns suggests that the peaks are smeared towards smaller diffraction angles (larger lattice parameter) as well as large ones. The peak intensities ratio correlated with peak intensities ratio for texture substrate, which is also indicative on the presence of texture in film. We used the relation for grain sizes (L) and width of diffraction peaks (A29): L~X/(A26*cos0), where l

268

is the wavelength, 8 is the diffraction angle. M Table 1 the estimated grain sizes are listed from the first two X-ray diffraction peaks.

Figure 2. Typical cross-sectional TEM image of the Ag3«iNi70 alloy. (On the Inset: electron diffraction pattern).

Table 1. Grain sizes of the Ag-Ni alloys,

Concentration ! Ni (wt%)

0

10 40

70

100

Peak width A20 (°)

0.36

1.21 2.64

2.89

1.04

Grain size L (nm) \

26.4 |

7.8 |

3.6 |

33 I

9.1

These eanocrystalline alloys are stable which is important for their practical use. We have examined the stability of the alloys grown by us after heating and mechanical treatment (polishing). The X-ray spectra of the polished specimen are identical to those of the unpolished one. So, one can conclude that polishing does not affect the alloy's structure. Also, no changes were detected for room temperature storage for several months. But, already after slight heating (beginning at 150°C) the diffraction pattern of the alloy is changed (Fig.3). The peaks get narrow, the narrowing occurs immediately with the increasing temperature and changes monotonically up to 600°C. The spectrum does not

269

change at the same temperature during time. This result can be regarded as recrystallization and grain coarsening but only grains of a specific size appear unstable at a particular temperature. As the temperature is increased, this range of unstable grain size grows. This suggests that nanocrystalline Ag-Ni alloys cannot be obtained by traditional quenching techniques. So, electrodeposition is the suitable method to grow such alloys.

• • * ' * i i . . . i i

35.0 40.0 45.0 50.0 55.0

2 © ( ° )

Figure 3. X-ray di(Traction pattern of Ag-Ni alloy with 80wt% Ag concentration at different temperatures.

4. Conclusions

Ag-Ni films of a broad compositional range were grown by electrodeposition. X-ray diffraction and transmission electron microscopy techniques show that the films are a nanocrystalline mixture of Ag and Ni metals. As the Ni concentration is increased in the alloy, the grain size is decreased, and with the Ni concentration in excess of 70wt% it goes down to 3.3nm. The nanocrystalline state of the Ag-Ni system is unstable as the temperature is increased. On heating in air the grain size grows and after annealing at 600°C the grain size of such specimens becomes equal to that in the one component films.

270

Acknowledgements

We wish to thank V.F.Degtyareva, A.V.Serebryakov and G.E.Abrosimova for their useful discussion and help in interpreting the results.

References

I. J-Ph Ansermet, J. Phys.: Condens. Matter 10, 6027 (1998). 2. 3. Z.L. Zhao, Y. Zhao, Y. Niu, C.L. Wang, W.T. Wu, J. Alloys Comp. 307,

254 (2000). 4. H. Zhu, S. Yang, G. Ni, D. Yu, Y. Du, Scripta Mater. 44, 2291 (2001). 5. T. Yamasaki, Scripta Mater. 44, 1497 (2001). 6. W. Lee, J. Lee, J.D. Bae, C.S. Byun, D.K. Kim, Scripta Mater. 44, 97

(2001). 7. F. Czerwinski, Electrochim. Acta 44, 667 (1998). 8. F. Czerwinski, H. Li, M. Megret, J.A. Szpunar, Scripta Mater. 37, 1967

(1997). 9. M. Aus, U. Erb, J.A. Szpunar, C. Cheung, B. Szpunar, G. Palumbo, J.

Magn. Mater. 187, 325 (1998). 10. G.K. Burnat, N.P. Fedot'ev, P.M. Vyacheslavov, K.M. Smiridonova, Russ.

J. Appl. Chem. 41, 291 (1968). II . M. Hansen, Constitution of Binary Alloys, (McGraw Hill, New York,

1958). 12. G.V. Strukov, V.V. Kedrov, G.K. Strukova, N.V. Klassen, Slectroplating &

Surface Treatment 7, 24 (1999). 13. V.I. Krysov, S.K. Krysova, I.S. Maksimov, Y.V. Medvedev, Physica B:

Condens. Matter 265, 291 (1999).

FRAGMENTATION OF POSITIVELY CHARGED METAL CLUSTERS IN STABILIZED JELLIUM MODEL WITH

SELF-COMPRESSION

M. PAYAMI Center for Theoretical Physics and Mathematics, Atomic Energy Organization of Iran

P.O. Box 11365-8486, Tehran, Iran

Using the stabilized jellium model with self-compression, we have calculated the dissociation energies and the barrier heights for the binary fragmentations of singly-ionized and doubly-ionized Ag clusters. The results for the singly-ionized clusters show that not all of the clusters are stable against spontaneous decay. However, the decays of doubly-ionized clusters can proceed via two different mechanisms: evaporation and fission. Our results show that fission is the dominant decay process for small clusters whereas for large clusters the evaporation process dominates.

1. Introduction

The study of the stability of nano-structures against fragmentations is of great importance in the nano-structure technology. Metal cluster in a matrix forms the fundamental structure in electronic devices. Charging a cluster would lead to its fragmentation and thus the micro-electronic devices should not be touched with bare hands. On the other hand, knowledge of the fragmentation processes of metal clusters enables us to enhance the efficiency of metallic-cluster catalysts. A catalyst provides the required activation energy in a specific reaction. The more the effective surface of the catalyst, the more the reaction rate is. However, after functioning for a period of time, the efficiency decreases because, those clusters come together and combine to make larger clusters, which in turn, results in the decrease of the effective surface. It is possible to re-increase the efficiency by inducing fragmentations. To initiate fragmentations, the clusters should somehow be ionized.

The properties of metallic clusters have been extensively studied [1-3] using the jellium model (JM). In this model, the discrete ions are replaced by a uniform positive charge background of density n = 3/(4^rs)

3 in which rs is the bulk value of the Wigner-Seitz (WS) radius of the valence electrons of the metal. A refined version of the JM, the stabilized jellium model (SJM), which was introduced [4] by Perdew et al. in 1990, has improved some drawbacks [5,6] of the JM. Using the SJM, the fragmentation of charged metal clusters has been also studied [7] by Vieira et al. However, since the surface effects have a large contribution in the energetics and sizes of small clusters, a more sophisticated use [8] of the SJM is needed to predict the correct energetics of the clusters in the study of the fragmentation processes. This method is called SJM

271

272

with self-compression (SJM-SC) and has been used to predict the equilibrium sizes and energies of neutral as well as the charged [9,10] metal clusters. The SJM-SC has been also used by Sarria et al. [11] to calculate the surface energies and the work functions of metals. In contrast to the JM and the SJM in which the rs value is borrowed from the bulk system, in the SJM-SC, the density parameter rs of the jellium sphere assumes a value such that a cluster with a given number of electrons and specific electronic configuration achieves its equilibrium state. The SJM-SC calculations on neutral metal clusters [8,10] has shown that the equilibrium rs value of the jellium sphere is less than the bulk value and tends to its bulk value for large cluster. This phenomenon is called self-compression which is due to the dominant effect of surface tension in small metal clusters. However, it has been shown that [9] charging a small metal cluster can result in an equilibrium rs value which is larger than the bulk value. This effect is called self-expansion. The self-expansion has been also predicted for highly polarized metal clusters [10,12]. These two effects have different origins. In the former, the repulsive coulomb force dominates the surface tension whereas, in the latter' the Pauli force is responsible for the self-expansion. Comparing the SJM-SC results with the SJM results and the experiment provide information about the possible structural relaxations of the clusters in the fragmentation processes. In this work, using the SJM-SC, we have studied the binary decay processes of

positively charged silver clusters AgN ( Z = + l , + 2 ) containing up to 100 atoms in all possible channels. The possible decay channels for singly-ionized Ag clusters are

Ag+N->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

+^Ag2N

+_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

+^Ag+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

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)

N2,z2\D „ O N1,z1

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) = +-0{rc-r) 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. £(rsArc) = 0. (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

rf(r?) = Hr?y B\in, 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) +

( < S v ) B S ( r I , r ; ) ( i V ^ - r ) [ « W - « t ( r ) ] (12)

where,

EJM ["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. The first and 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, <5v, is defined as the difference between the pseudo-potential of a lattice of ions and the electrostatic potential of the jellium positive background. The effective potential, used in the self-consistent KS equations, is obtained by taking the

277

variational derivative of the SJM energy functional with respect to the spin densities as

^([n^ni,n+];r) = -——(ESJM-Ts) CT=t,4 (15) 8na(r)

By self-consistent solution of the KS equations

'4v 2 +^(r)A <(r) = s,>r('0 =t,4- (16)

with the density

n(r) = Y,V?(r) (17)

the total energy of the cluster with the density parameter rs for the jellium

background is obtained. In our spherical JM, we have

n+(r)= 0(R-r) 4;zr,

(18)

r l / 3 . in which R = N rs is the radius of the jellium sphere, and n(r) denotes the

electron density at point r in space. Using the Eq. (21) of Ref. [4], this average value is given by

fi\2

( * L ('•..'•/) = B, XK)

2r] lOr, (19)

The SJM ground-state energy (Eq. (12)) for a cluster with TV electrons is a

function of N , rs, and rc . The equilibrium density parameter, rs (N) , for a cluster in the ground state electronic configuration, is the solution of the equation

_d_ EuM(N,rs,rc

B) = 0 (20)

r.=f,(.N)

278

where the derivative is taken at fixed values of N and rc . The total energy in

the SJM-SC then becomes a function of the equilibrium density parameter as

EsjM-sc=ESJM(N,rs(N),rf).

3. Results and Discussions

After an extensive self-consistent SJM-SC calculations, we have calculated the

equilibrium rs values and the energies of AgN clusters (Z — 0,1,2) for

different cluster sizes (N < 100). To show the main differences in the

equilibrium rs values of these clusters, which are appreciable for small clusters,

we have plotted the corresponding rs (N) values only up to N = 34 as shown

in Fig. 3. As is obviously seen in the figure, the neutral and singly-ionized clusters are self-compressed for all values of N . However, for doubly-ionized clusters, the rx(N) values cross the bulk border (i.e., r / =3 .02) at

N = l.

4.2

= 3.7

N 3.2

2.7

2.2

-

- ] • \

: I

^ ^ V l ,-wQ-C

- \r

— z=0

•0-z=+1

-Q-2= + 2

- -Bu lk

10 15 20 25 Number of Atoms N

30 35

Figure 3. Equilibrium rs values in atomic units.

Figure 4 shows the SJM-SC energies per atom in electron volts for neutral, singly-ionized, and doubly-ionized silver clusters with iV < 34. For comparison, we have also plotted the bulk value (e = -7 .89eF ) by a dashed line.

279

-2

^ "4

R |LU -6

— z=0

-0-z= + 1

• z = + 2

--Bulk

10 15 20 25 Number of Atoms N

30 35

Figure 4. SJM-SC energies in electron volts.

In Fig. 5, we have plotted the dissociation energies of the most favored

decay channels for the evaporation process in the singly-ionized clusters. We

have shown the most favored value of p by p . The solid small square

symbols indicate the most favored values p on the right vertical axis whereas,

the corresponding dissociation energies, D]+(N,p*), are shown on the left

vertical axis by large open squares. The positivity of the Di+(N,p*) implies

that the parent cluster is stable against spontaneous fragmentation whereas, the clusters with negative dissociation energies undergo the fragmentation. However, the liquid-drop model (LDM) calculations [19], which neglects the shell structure, predicts no fragmentations for the singly-ionized clusters.

The dissociation energies of the most favored evaporation channels of the doubly-ionized clusters are plotted in Fig. 6. Here, also as in the singly-ionized case, some of the dissociation energies are positive and some are negative. The implications are the same as in the singly-ionized case. The fission barrier heights for the most favored channels, Bf(N,p*), as well as the corresponding fragment sizes, p , are plotted in Fig. 7. As is seen, some of them are negative. The negativity of a barrier height means that the parent

280

cluster is not stable against the fission and it would spontaneously decay. However, the most favored channels with positive barrier heights imply that those parent clusters are stable against spontaneous fission and energy must be supplied to induce any fission.

Q.

10 20 30 40 50 60 70 80 90 100 Number of Atoms N

Figure 5. Dissociation energies in electron volts for singly-ionized clusters and the most favored sizes.

0 10 20 30 40 50 60 70 80 90 100 Number of Atoms N

Figure 6. The same as in Figure 5 for doubly-ionized clusters.

281

Since the doubly-ionized clusters can decay via either evaporation or fission, the parent cluster would be stable against any spontaneous fragmentations when both the dissociation energies and the fission barrier heights are positive. Thus, to determine the stability of different doubly-ionized clusters, we have compared D2+(N,p ) and Bf(N,p ) in Fig. 8. It is clearly seen that in a certain size range (N > 34), the fission and evaporation definitely start their competition. Because of the pronounced shell effects in our results here, it is difficult to determine a well-defined size beyond which the evaporation dominates. However, Fig. 8 shows that the evaporation dominates for N > 84 , consistent with our LDM results [19] ( N > 94).

? j"

a.

CD

0 10 20 30 40 50 60 70

Number of Atoms N

Figure 7. Fission barrier heights in electron volts for doubly- for doubly-ionized clusters are compared.

? <D

J ^ Q. ^ Q

;v c >

en

4

2

0

-?

-4

~ l JI JL • M k ' J n h 1 'Tlil/^l " \if\l V Vt'' ' \fmi VT • \\ w « H

- v I l • I *

- D E -""Bf

JI . / jrA

A. _/s LA I \ J ^ W jsr *» i \ Jr H &Jf r IT Y*St y fit*r ty'

i v yv i

0 10 20 30 40 50 60 70 80 90 100

Number of Atoms N

Figure 8. Dissociation energies and the barrier heights ionized clusters and the most favored sizes.

282

4. Conclusion

In this work, we have studied the fragmentations of singly- and doubly-ionized Ag clusters. Using the SJM-SC energies for the clusters, we have calculated the dissociation energies and the fission barrier heights. Our results show that some of the singly-ionized clusters are stable against any decay. The doubly-ionized clusters can decay via two different processes: evaporation and fission. The smaller doubly-ionized clusters prefer to decay via fission. The doubly-ionized clusters for which both the dissociation energies in the evaporation and the fission barrier heights are positive, do not undergo any spontaneous fragmentations and are stable. However, beyond N = 34 the competition between the evaporation and the fission mechanisms is started. The evaporation mechanism dominates for N > 84 which is close to the LDM result ( J V > 9 4 ) .

References

1. W.E. Ekardt, Phys. Rev. B 29, 1558 (1984). 2. W.D. Knight, K. Clemenger, W.A. de Heer, W.A. Saunders, M.Y. Chou,

M.L. Cohen, Phys. Rev. Lett. 52, 2141 (1984). 3. M. Brack, Rev. Mod. Phys. 65, 677 (1993). 4. J.P. Perdew, H.Q. Tran, E.D. Smith, Phys. Rev. B 42, 11627 (1990). 5. N.D. Lang, W. Kohn, Phys. Rev. B 1, 4555 (1970). 6. N.W. Ashcroft, D.C. Langreth, Phys. Rev. 155, 682 (1967). 7. A. Vieira, C. Fiolhais, Phys. Rev. B 57, 7352 (1998). 8. J.P. Perdew, M. Brajczewska, C. Fiolhais, Solid State Commun. 88, 795

(1993). 9. M. Brajczewska, A. Vieira, C. Fiolhais, J.P. Perdew, Prog. Surf. Sci. 53,

305(1996). 10. M. Payami, J. Phys. Condens. Matter 13, 4129 (2001). 11. I. Sarria, C. Henriques, C. Fiolhais, J.M. Pitarke, Phys. Rev. B 62, 1699,

(2000). 12. M. Payami, J. Chem. Phys. I l l , 8344 (1999). 13. U. Naher, S. Bjcrnholm, S. Frauendorf, F. Garcias, C. Guet, Phys. Rep.

285,245(1997). 14. W. Kohn , L.J. Sham, Phys. Rev. 140, Al 133 (1965). 15. P. Hohenberg, W. Kohn, Phys. Rev. B 136, 864 (1964). 16. M. Payami, N. Nafari, J. Chem. Phys. 109, 5730 (1998). 17. J.P. Perdew, Y. Wang, Phys. Rev. B 45, 13244 (1992). 18. N.W. Ashcroft, Phys. Lett. 23, 48 (1966). 19. M. Payami, Proc. 9th Conf. of the Nuclear Physics, Kerman University,

Iran (26-28 Feb 2003).

VI. OPTICAL MATERIALS

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ORGANIC FILMS FOR OPTOELECTRONIC APPLICATIONS

X. LIU, T. MICHELY, M. WUTTIG

/. Physikalisches Institut der RWTH Aachen.Lehrstuhlfur Physik neuer Materialien D-52056 Aachen, Germany

E-mail: wuttig@physik. rwth-aachen. de

A systematic study comparing the growth of perylene films on a Au substrate and on a Au substrate coated additionally with a self-assembled monolayer (SAM) of 1-octadecanethiol is presented. The films have been characterized by atomic force microscopy and x-ray diffraction. Compared to the Au substrate, on the SAM substrate, smaller island sizes and a better c-axis texture result. Possible mechanisms for these changes in the perylene film growth are discussed.

1. Introduction

The advancement of semiconductor technology has created a high-tech industry with annual sales of more than 300 billions $. This success has to be mainly attributed to the constant improvement in packing density, i.e. the ability to create smaller and smaller structures on increasingly larger substrates. Hence devices could be developed with steadily improving functionalities. This trend and the related success of the semiconductor industry is presently challenged. On the one hand, since the decreasing structure size in electronic devices is approaching physical limits, new concepts and materials have to be explored to maintain the present speed of device development. In addition, traditional manufacturing schemes are challenged since they are accompanied by ever increasing costs which show an exponential increase from one to the next generation of chips. Hence alternative routes to semiconductor nanostructures are explored.

Recently, molecular films have emerged as candidates for electronic and optoelectronic applications. They offer alternative and presumably cheaper manufacturing possibilities. Furthermore molecular materials might be used to store information. Organic semiconductors have been successfully employed to fabricate field effect transistors, light emitting diodes and photovoltaic devices [1-4]. Furthermore, molecular materials are expected to play a prominent role in next generation electronic technology due to their versatility in modifying the physical properties by molecular design. For example, single-molecule transistors have been demonstrated [5] and many efforts are directed towards exploring organic magnets [6]. It is therefore not surprising that more and more attention is focused on organic film deposition, since films are a prerequisite to exploit the full potential of molecular materials. In organic devices, carrier

285

286

transport and luminescent behavior are governed by the orientation and packing of molecules. For organic field effect transistors, one always aims at achieving the utmost perfection of molecular order and particular molecule orientations to obtain large carrier mobility. For light emitting devices, on the contrary, one hopes that the active organic material exists in amorphous form to avoid concentration quenching. Thus reaching a complete understanding of the nucleation and growth of molecular films and exploring new methods to tailor the structure of molecular films for a given application are of paramount importance.

In contrast to the strong atomic binding forces in inorganic crystals, in organic, molecular crystals and films the building blocks of the crystal - organic molecules - are bound to each other by the relatively weak van der Waals forces. Furthermore, the organic molecules possess extended, generally anisotropic shapes, which introduce specific steric requirements with respect to molecular orientation and order for the crystal and thin film nucleation and growth. Therefore, the knowledge accumulated during the last several decades on the deposition and growth of inorganic thin films cannot be directly applied to the growth and nucleation of organic films. Compared with what is known about conventional inorganic systems, there is up to now only a limited understanding of the microscopic growth mechanisms of organic systems [7- 9]. Even the dependence of the film structure and morphology upon growth conditions is still missing for many organic films. Nevertheless, it has been realized that the molecule-substrate interaction plays a crucial role in determining the molecular orientation [9], growth mode [7], film morphology [9], and even the crystal structure [10], though the mechanisms by which the molecule substrate interaction affects the film properties still remain unclear. Self-assembled monolayers (SAMs) have already been successfully explored to modify surface properties, for example, to control the surface wetting behavior [11], to modify the electrochemical potential of metal electrodes [12], and to develop novel lithographic methods [13]. Recently, SAMs have also been used as growth templates to define size, morphology, structure and orientation of growing crystal grains in thin films through modifying the interfacial structure. A number of reports have been published but most of them are concerned with the growth of inorganic films predominantly prepared by chemical solution or chemical vapor deposition [14- 19]. In the following section, it will be shown how a SAM can modify the evolution and properties of a vacuum evaporated perylene film on a polycrystalline Au substrate [20]. Before the experimental results are presented, a compact introduction to self-assembled monolayers follows in the next section.

287

2. Self-assembled Monolayers

Self-assembled monolayers (SAMs) consist of densely packed long-chain organic molecules which are chemisorbed on inorganic substrates through a head group which has a specific affinity to the substrate [21]. Three SAMs are popularly investigated. They are organic acids on mica or sapphire, alkyltrichlorosilanes on Si02, and alkanethiols on Au. Here we will focus on the alkanethiols SAM on a Au(ll l) substrate. An alkanethiol molecule consists of an alkane group and a thiol group with a formula of CH3(CH2)«S, just as suggested by its name. The thiol group is chemisorbed on the gold substrate. The S-Au bonding energy is estimated to be greater than 1 eV [23]. While the monolayer/ substrate interaction is determined by the type of bond between the head group and the substrate, the intra-layer interaction, i.e., the intermole-cular cohesive energy comes from Van der Waals (VDW) interactions and is directly proportional to the hydrocarbon chain length. Although the VDW interactions is only 0.07 eV per CH2 unit and thus is much weaker than the S-Au chemisorption bond, the sum of the interactions between a hydrocarbon chain and its neighbors is also considerable. The structure of the SAM is mainly determined by these two interactions. While an epitaxial registry of the SAM on the substrate is strongly preferred by the S-Au bonding, the SAM has a trend to form close-packing structures to minimize the inter-molecule interaction. The actual structure of the SAM results from the energetic equilibrium between these two interactions. Thus it is easy to understand why the tilt angle of the molecule chains changes with the coverage and depends on the length of the alkyl chain (see Fig. 1). The structure for full coverage of alkanethiol on Au(l 11) has been determined by He atom diffraction [24] and confirmed by grazing incidence X-ray diffraction [25], which reveals an orthorhombic primitive unit cell with a dimensions of 3a*2 V3 a, where a - 2.884 A, is the Au lattice constant(see Fig. 2).

3. Experimental

The films were prepared by sublimation in a vacuum system with a base pressure in the low 10'7 mbar range. Perylene evaporation for thin film growth causes typically a pressure rise to 1-2 x 10'6 mbar. The deposition rate was measured by a quartz crystal microbalance. Two kinds of commercial substrates [27] were used. One is a 10 nm Au/2 run Ti/glass substrate, subsequently denoted as Au-substrate. The second type of substrate has additional a

288

monolayer of self-assembled 1-Octadecanethiol molecules adsorbed on the Au layer (SAM-substrate). To enable a direct comparison for each growth experiment, these two kinds of substrates were installed side-by-side in the deposition chamber. Perylene was deposited upon evaporation at room temperature. Afterwards, the films were characterized using scanning force microscopy (AFM) and x-ray diffraction.

f V W

«ii& 0SSB <88888

Figure 1. The following are illustrated; (a) n-alkanethiol molecule, (b) interaction potential between thiol molecule and Au(l l l ) surface, (c) decreasing tilt angle, 0, with increasing coverage, (from ref. [22]).

3JI4A *«?

Figure 2. Structure of alkanethioi monolyer (large circles) on bulk-terminated Au surface (small circles). Diagonal slash in large circles represents azimuthal orientation of plane defined by all-trans hydrocarbon chain, (from ref. [26]).

4. Effect on tie Nieleatlon Mate To obtain a coherent picture of perylene growth on the Au- and the SAM-substrate, we have examined perylene films of different thickness on both

289

substrate types. To unravel the initial growth stage, discrete thin films, in which islands do not coalesce, were prepared. The evaporation rate was kept constant at 0.11 nm/s and the deposition time was varied. By analyzing the island density as a function deposition time, the nucleation behavior of the perylene films on the two kinds of substrates was investigated. Fig. 3 shows the AFM pictures of discrete perylene films with deposition times between 10 s and 40 s. Several features can be deduced from the pictures: 1) For all pairs of samples installed side-by-side in the chamber with exactly the same amount deposited (arranged in the same row in Fig. 3), the films prepared on the SAM-substrate present a larger island density and hence smaller island size compared to the Au-substrates. Please take notice that the scanning size is 5 /an for the films on the SAM-substrate (left column in Fig. 3) and 15 fjm for the films on the Au-substrate (right column in Fig. 3), respectively. 2) With increasing deposition time, the island density rises until coalescence occurs for the films on the SAM-substrates. For films on the Au-substrates the island density stays almost constant (see Fig. 4), whereas the grain size increases with deposition time. The obvious difference in perylene island density and its deposition time dependence between growth on the Au-substrate and on the SAM-substrate indicates that the adsorbed SAM considerably alters the surface properties of the Au-substrate. We envision two different scenarios. First, if the adsorption energies of perylene on both substrates are large enough, nucleation and growth occur in the complete condensation regime. The single molecule density increases initially

until they turn to decrease at time TC= [28], where D is the molecule oDN

diffusion coefficient, N is the island density, and a is the capture number, which is a slowly varying quantity with values between 5 and 10. As nucleation is essentially terminated on the Au-substrate already for the smallest deposition time of 10 s, xc must be considerably smaller than the measurement time scale. On the contrary, TC must be comparable to the measurement scale on the SAM-substrate. As the island density on the Au substrate is much smaller, it follows that if"/lfAM = A^/Af4", i.e. the surface diffusion coefficient on the Au substrate must be much larger than on the SAM substrate. Second, we are unable to rule out incomplete condensation. In this case the observed continuing nucleation on the SAM-substrate would imply a much larger adsorption time ra

(longer increase of single molecule density) and thus a much larger adsorption energy of the molecules on the SAM-substrate. In this scenario, in addition a large molecule diffusion coefficient on the Au-substrate is still needed to explain the much lower island density. We consider this last scenario as less likely, but in order to unambiguously identify the mechanisms of how the SAM modifies nucleation, additional measurements would be needed.

290

mm An

10-

&.*'•

30

* 1

40 i *Lk

¥ rt'5'^-4 j -*. - ^ s ** "

o „ < \ - i ^ r *> '«rx^ * ^ r ^ s, ^ : „ „ -

^>rio; ^ : ' v ^i r ^ K \ - -" --.«v* v* *v *"i- - " -. * *'*^S& t^v; ,-^rf-M-/ |\ ;> "s s> * "" ^ :V ^ -V*.- \ '*

^ ^ *v#^v ; ; r ; v :

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

• w ^ + v̂ j

^ A ; • i v -s v s

!

v ' f -

;; w;

: ;>!* '

l«l*m 15 fan

Figure 3. AFM topographs of 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, the right column depicts films deposited on Au-substrates. The image size for each column is indicated.

291

8-

7-<-C I6' "8 5-(0

^ 4 -

'g 3-s ?2-c

Isia

0-1

—•-- 0 -

o'

u—

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10 15 20 25 30 35 40 45 50 Deposition time (s)

Figure 4. Deposition time dependence of island density. The island density increases with deposition time on the SAM-substrate, while it stays nearly constant on the Au-substrate. Lines to guide the eye.

5. Effects on the Film Roughness and Texture

As we saw, the introduction of a 1-Octadeca-nethiol monolayer at the perylene-Au interface strongly modifies the nucleation behavior of perylene films. The change in nucleation behavior might further result in a modification of the film microstructure such as grain size, roughness, etc. Hence we have also examined to what extent the microstructure of a thick film is modified by the SAM. We prepared several pairs of thick films at various evaporation rates. For all samples, the product of evaporation rate and evaporation time was kept constant, i.e. the same amount of material was deposited. The resulting film thickness is calculated to be 100 nm, using the single crystal density and assuming that the accommodation and sticking coefficients do not vary significantly between the gold covered quartz of the quartz crystal microbalance and the substrates used in the present work. The AFM topographs are presented in Fig. 5, all in the same scale. The left column displays the samples prepared on the SAM-substrates and the right column those deposited on the Au-substrates. The two samples of each row were installed in the chamber side-by-side and deposited at the same time. The evaporation rates labeled on each row vary from 0.027 nm/s to 6.8 nm/s. The film roughness calculated on the basis of the AFM is plotted in Fig. 6 versus the deposition rate. In agreement with our finding discussed above that on a SAM-substrate compared to an Au-substrate a larger maximum island density is reached during the initial deposition stage, all films on the SAM-substrates shown in Fig. 5 exhibit a smaller grain size and accordingly also a smaller roughness than their counterparts on the Au-substrates. The evaporation rate itself greatly affects the film morphology for both types of 5 substrates. With

292

increasing evaporation rate, the grata size decreases and the film becomes smoother. However, at a deposition rate above about 1 nm/s, the dependence of film roughness on deposition rate as well as substrate type becomes weak.

0.027 nm/s 64 mitt

0.11-nm/s 960 s

0.44 amis 240 s

6.8mn/s 15 s

:̂ % ;<w 1^4% y$ gib* i^X^^

th^f

iiPftl l*'-'iJ'ri

- * -4-

Figure 5. AFM topographs comparing the effect of evaporation rate on film microstructure for SAM-substrates (left column) and Au-substrates (right column). Evaporation rates and tie deposition time for each row are indicated. The nominal thickness of all films is about 100 nm.

293

120

100

I »•

f 40

20

0.01

\ - • - Au - c - SAM

\

0.1 1 Deposition rate (nnVs)

10

Figure 6. Dependence of films roughness on the deposition rate. Typical morphologies of the associated samples are shown in Fig. 5.

Another important feature seen from Fig. 5 is that most crystal grains in the film deposited on the SAM-substrate show a flat top facet, in contrast to the films directly grown on gold films. This feature is seen more clearly when the scanning sizes are smaller (see Fig. 7). This observation implies that the perylene films deposited on the SAM possess a better ordering of grain orientation. X-ray diffraction measurements confirm this conclusion. Fig. 8 shows the 0-20 diffraction pattern of a pair of samples, which both were 100 nm thick and which both were prepared with an evaporation rate of 0.11 nm/s. Only three peaks are observed, the (001) and (002) peaks of a-perylene and the (111) peak of Au. To obtain a clear comparison, we normalize all peak intensities to the Au(l 11) peak. Since the thickness of the perylene films and the underlaying Au films are the same for both samples, the higher intensities of the perylene (001) peaks suggest a better c-axis texture. The corresponding rocking curves of the perylene (001) peaks are shown in the inset of Fig. 8. As is expected, the full width at half maximum (FWHM) is much smaller for the perylene films grown on the SAM-substrate. The introduction of the 1-Octadecanethiol monolayer thus introduces a stronger c-axis orientation. A c-axis orientation means that the perylene molecules prefer to stand upright on the molecular monolayer [30]. The formation of c-axis texture in perylene thin films can be understood on the basis of our observations in single crystal growth of perylene by the temperature-gradient sublimation method. Single crystals of perylene fabricated in this way are always thin platelets, with the c-axis perpendicular to the platelet surface. It can thus be assumed that the dense packed (001) surface is the surface of lowest surface free energy. Thus in thin film growth, there is an energetic driving force,

294

which tends to push the average grain orientation towards a c-axis texture. The different degrees of perfection of the c»axis texture in the thin films grown on the Au- and the SAM-substrate point to a second factor, which determines texture formation. Evidently, this second factor must be related to the substrate. It is the interaction of the deposited molecules with the substrate in the nucleation phase. Let us assume two different scenarios. First, a situation of weak molecule substrate interaction (which is equivalent to a high interface energy), imposing no specific constraints on molecule orientation during nucleation. In such a case the nucleating crystal grains will right from the beginning possess the molecule orientation minimizing their surface energy and thus develop a perfect texture. Second, let us consider a situation of strong molecule-substrate interaction (corresponding to a small interface energy). This strong interaction may easily imply a specific orientation of the molecules with respect to the substrate in the nucleation phase, which differs from the molecule orientation favored by the surface energy of the perylene crystallites. In such a case a transition from the substrate dominated molecule orientation to the molecule orientation favored by the surface energy may occur with increasing film thickness, but texture formation will be more difficult. Indeed we believe that our observations are related to these two schematic scenarios. For the SAM-substrate the substrate-molecule interaction can be assumed to be weak and purely of van der Waals type, allowing the molecules to obtain already during the nucleation phase the upright molecule orientation compatible with the energetically preferred c-axis texture. In contrast, for metallic substrates it is known that the first monolayer of aromatic molecules forms /r-type bonding orbitals with the metal surface, implying that the molecules lie flat on the metal surface [29], implying also for perylene on the Au»substrate a flat lying molecule configuration [31]. This molecule orientation is at variance with the energetically preferred c-axis texture and thus makes an inferior c»axis texture on the Au-substrate plausible. We note that most likely only after adsorption of a molecular layer of flat perylene molecules the nucleation of islands on the Au-substrate will take place.

Figure 7. AFM pictures of perylene film deposited on SAM (left) and on Au substrate (right). The scanning size are 2|im x 2jun for the left and 5\an x Sum for the right.

295

C

5

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Figure 8. X-ray diffraction 0-28 patterns. The rocking curves of perylene (001) are shown in the inset. The films on the SAM-substrate possess strong c-axis texture. Lines to guide the eye.

6. Summary

In summary, AFM and x-ray diffraction were used to show that an additional self assembled monolayer of 1-Octadecanethiol molecules on a Au-substrate strongly modifies the properties of the subsequently vacuum deposited perylene films. The grain size becomes smaller and a strong c-axis texture is introduced. Both factors tend to reduce the roughness of the perylene films. These effects are attributed to the changes in the perylene molecule-substrate interaction by the additional self assembled monolayer. It appears very desirable to extend this investigation to other SAMs and organic molecules. Such a systematic study might reveal general trends that would enable structure and hence also property tailoring for organic films.

References

1. C.W. Tang, S.A. Van Slyke, Appl. Phys. Lett. 52, 913 (1987). 2. C.W. Tang, S.A. Van Slyke, C.H. Chen, J. Appl. Phys. 65, 3610 (1989). 3. J. H. Burroughes, D.D.C. Bradley, A.R. Brown, R.N. Marks, K. Machay,

R.H. Friend, P.L. Burn, A.B. Holmes, Nature (London) 347, 539 (1990).

296

4. M. Granstrom, K. Petritsch, A.C. Arias, A. Lux, M.R. Andersson, R.H. Friend, Nature (London) 395, 257 (1998).

5. W. Liang, M.P. Shores, M. Bockrath, J.R. Long, H. Park, Nature (London) 417, 725 (2002).

6. E. Coronado, J.R. Galan-Mascaros, C.J. Gomez-Garcia, V. Laukhin, Nature (London) 408, 447 (2000).

7. S.R. Forrest, P.E. Burrows, Supramolecular Science 4, 127 (1997). 8. P. Fenter, F. Schreiber, L. Zhou, P. Eisenberger, and S.R. Forrest, Phys.

Rev.B 56, 3046 (1997). 9. H. Peisert, T. Schwieger, J. M. Auerhammer, M. Knupfer, M.S. Golden, J.

Fink, P.R. Bressler, M. Mast, J. Appl. Phys. 90, 466 (2001). 10. Q. Chen , T. Rada, A. McDowall, N.V. Richardson, Chemistry of Materials

14, 743 (2002). 11. G. E. Poirier, Chem. Rev. 97, 1117 (1997). 12. S. Wang, D. Du, Q. Zou, Talanta 57, 687 (2002). 13. S.Y. Chou, L. Zhuang, J. Vac. Sci. Technol. B 17, 3197 (1999). 14. A.D. Polli, T. Wanger, A. Fischer, G. Weinberg, F.C. Jentoft, R. Schlgl,

M. Rhle, Thin Solid Films 379, 122 (2000). 15. G. Dahlgren, A. Smith, D.B. Wurm, Synthetic Metals 113, 289 (2000). 16. J. Flath, F.C. Meldrum, W. Knoll, Thin Solid Film 327-329, 506 (1998). 17. H. Shin, M. Agarwal, M.R. De Guire, A.H. Heuer, Acta Materialia 64,

801 (1998). 18. G.C. Herdt, D.R. Jung, A.W. Czanderna, Progress in Surface Science

50, 103 (1995). 19. F.C. Meldrum, J. Flath, W. Knoll, Thin Solid Films, 348, 188 (1999) 20. X. Liu, S.H. Mohamed, J.M. Ngaruiya, M. Wuttig, T. Michely, J. Appl.

Phys. 93, 4852 (2003) 21. L. H. Dubois, R. G. Nuzzo, Annu. Rev. Phys. Chem. 43, 437(1992). 22. Y. Yourdshahyana, A. M. Rappe, J. Chem. Phys. 117, 825 (2002). 23. P. Fenter, A. Eberhardt, K. S. Liang, and P. Eisenberger, J. Chem. Phys.

106,22(1997). 24. N. Camillone, C.E.D. Chidsey, G.-Y. Liu, G. J. Scoles, J. Chem. Phys.

98, 3503(1998). 25. P. Fenter, P. Eisenberger, K.S. Liang, Phys. Rev. Lett. 70, 2447(1993). 26. G.E. Poirer, Chem. Rev. 97, 1117(1997). 27. The substrates are supplied by Georg Albert, PVD-Beschichtungen, Angew.

Physikalische Chemie, Im Neuenheimer Feld 253, D-69120 Heidelberg. 28. J.A. Venables, G.D. Spiller, M.Hanbucken, Rep. Prog. Phys. 47, 399

(1984). 29. F.P. Netzer, M.G. Ramsey, Crit. Rev. Solid. State Mater. Sci. 17,

397(1992) 30. D.M. Donaldson, J.M. Robertson, F.R.S., J.G. White, Proc. Roy. Soc.

A 220, 311(1953). 31. C. Seidel, R. Ellerbrake, L. Gross, H. Fuchs, Phys. Rev. B. 64, 195418

(2001).

DEVELOPMENT OF HIGHLY REACTIVE PHOTO-CATALYTIC Ti02 FILMS

S.H. MOHAMED

/. Physikalisches Institut (IA) der RWTHAachen, Lehrstuhlfur Physik neuer Materialien 52056 Aachen, Germany

and Physics Department, Faculty of Science, South Valley University, Sohag, Egypt

R. DRESE

/. Physikalisches Institut (IA) der RWTH Aachen, Lehrstuhl fur Physik neuer Materialien 52056 Aachen, Germany

M.M. WAKKAD

Physics Department, Faculty of Science, South Valley University, Sohag, Egypt

M. WUTTIG

/. Physikalisches Institut (IA) der RWTH Aachen, Lehrstuhl fur Physik neuer Materialien 52056 Aachen, Germany

E-mail: wuttig@physik. rwth-aachen. de

The investigation and optimization of the parameters affecting photo-catalytic properties of Ti02 films are of scientific and technological interest. In this paper a successful attempt is reported to obtain crystalline, photo-catalytic Ti02 films on unheated substrates. Some of the films were deposited at different total pressures in a mixture of argon and oxygen and at a fixed oxygen gas flow ratio of 27%. Others were sputtered in a mixture of argon, oxygen and nitrogen. Both the effects of total pressure and nitrogen to oxygen ratio on the photo-catalytic properties of the deposited films were examined. The anatase Ti02 films obtained at a total pressure of 2 Pa have the lowest water contact angle and the highest decomposition efficiency for methylene blue while an anatase/rutile mixture of Ti02, obtained at a total pressure of 0.33 Pa, has the highest water contact angle and the lowest decomposition efficiency. The photo-catalytic activities were found to increase with increasing total pressure. This was ascribed to the increase in surface roughness and the decrease in density as revealed by X-ray reflectometry.

1. Introduction

The photo-catalytic behavior and the superhydrophilicity of Ti02 films continues to

motivate scientific investigations. Some of this work has been focused on the

examination of the mechanisms, processes and the parameters, which can affect

and/or enhance the efficiency. The overall catalytic performance has been suggested

to be dependent on quite a number of parameters including particle size, the specific

297

298

surface area, the ratio between anatase and rutile phases, the amorphous content, the illuminating light intensity as well as the material to be degraded [1-4]. Although the photo-catalytic decomposition of organic materials and the superhydrophilicity result from the creation of electrons and holes by the incident light, the mechanisms affecting the efficiency of both phenomena could be different. For example, Takeda et al. [5] showed that there is no pronounced correlation between surface roughness and the decomposition ability of acetaldehyde (CH3CHO) while Uelzen and Miiller [6] have recently shown that the surface roughness is a necessary factor for the enhancement of wettability (superhydro-philicity).

Ti02 exists in three different crystalline phases: anatase, rutile and brookite. So far only rutile and anatase have been observed in thin films. Anatase shows a higher photo-catalytic activity than rutile. This enhancement in photo-activity is attributed by Sumita et al. [7] to the larger band gap of anatase. The efficiency of photo-catalysis strongly depends on the lifetime of electron-hole pairs generated by light absorption. As the recombination probability is inversely proportional to the magnitude of the band gap, the lifetime is prolonged and results in a high activity for anatase [7].

In this work we examine the effect of the rutile content on the photocatalytic activity of Ti02. In addition, we examine the effect of the incorporation of a small amount of nitrogen into the Ti02 films.

2. Experimental

Transparent Ti02 films were deposited from a Ti target on microscopic glass substrates without heating by a reactive dc magnetron sputtering system. The sputter system has already been described elsewhere [8, 9]. In the present study, sputtering was performed at a constant cathode current of 950 mA and at a target-to-substrate distance of 55 mm. Some of the samples were sputtered at different total gas pressures and at fixed gas flow ratio of q0^ i(q0^ +? / l r)of 27 %. Other samples were

deposited at a constant total pressure of 2 Pa. Then, 0 2 at q0 l{qQ + qAr) = 27 % is

partially replaced by N2. The N2 flow was increased from 0 to 12 seem while at the same time the 0 2 flow was decreased from 30 to 18 seem to deposit oxynitride films for a specific oxygen/nitrogen ratio.

To determine the thickness, density and surface roughness of the films, X-ray refiectometry (XRR) was used [10,11], employing a Philips X'pert diffractometer. Computer simulations were performed to reproduce the measured spectra. The thickness of the films was determined from the period of the intensity oscillations and the density was determined from the position of the total reflection edge. The roughness was obtained from the decay of oscillation amplitude and overall intensity [10].

299

To study the crystallographic order and the chemical composition, X-ray diffraction (XRD) and Rutherford Backscattering Spectroscopy (RBS) were used, respectively. A tandetron accelerator was used for RBS measurements with 1.4 MeV 4He+ particles and a current of 14 mA. The backscattered particles are detected at an angle of 170° with respect to the incident beam direction by a semiconductor detector with an energy resolution of 10 keV. The thickness of the titanium oxide layer can be calculated from the energy width of the Ti peak if we take the energy loss of 4He+ per unit depth in the titanium oxide matrix into account. The stoichiometry of the films, on the other hand, can be determined from the height of the Ti, O and N peaks. The XRD patterns of the as deposited titanium oxide films were examined by a Philips X'pert MRD diffractometer. Data were recorded for samples prepared on glass substrates using the grazing incidence geometry. The angle of incidence was 0.7°. The diffraction patterns were obtained using Cu Ka

radiation (k = 1.542 A). The photo-induced hydrophilicity was evaluated by measuring the contact angle

of pure water on Ti02 after UV illumination using a DSA10 contact angle meter. The illumination was carried out using an UV light lamp with a wavelength of 366 nm and an intensity of lmW/cm2. The contact angle measurements were carried out at room temperature in ambient atmosphere with an experimental error of less than 1°.

The photo-catalytic behavior of the Ti02 films was mainly tested by measuring the decomposition of Methylene Blue (MB) (Ci6Hi8N3S.ClxH20 (x = 2-3)). Each Ti02 sample was dipped into an aqueous MB solution with a concentration of 1 mmol/1 for 1 h and subsequently dried in air for 45 min in a dark place. The transmittance of the films before and after UV irradiation was measured employing a Perkin-Elmer UV/VIS spectrophotometer. Using the value of transmittance at 580 nm before illumination T0 and after illumination Tt (in time steps of 5 minutes

each), the change of absorbance AABS characterizing the decomposition of MB T

was calculated: AABS = ln(—) . i

3. Results and Discussion

Figure 1 shows XRD patterns of samples prepared at different total pressures and at a fixed gas flow ratio (q0i l(q0i +qAr))oi 27%. It can be seen that the film prepared at a total pressure ptot of 0.33 Pa contains a mixture of anatase and rutile. With increasing total pressure, the rutile content decreases while the anatase content increases. At p,ot = 2 Pa the rutile phase disappeared and a pure anatase phase is obtained. The higher concentration of the rutile phase at low pressures is related to

300

the higher energy provided to the film by the impinging sputtered particles, such as fast neutrals reflected at the target and negative oxygen ions [12-14]. Fig. 2a shows the change in water contact angle as a function of UV illumination time for Ti02 films deposited at different total pressures and constant oxygen flow ratio of 27%. It is seen that the film deposited at ptot = 2 Pa has an initial water contact angle around 22° and the film deposited at ptot = 0.33 Pa has an initial water contact angle around 60°. The initial water contact angle increases with decreasing total pressure i.e. with increasing rutile content. Upon UV illumination, the water contact angle started to decrease, i.e. the photoinduced hydrophilicity increased. After 20 min of UV illumination, the water contact angle of the films prepared at a Ptot of 2, 1, 0.8 and 0.7 Pa, respectively, was less than 1° while complete wettability was observed after 35 and 40 min of UV illumination, (for films prepared at 0.6 and 0.45 Pa, respectively). On the other hand films prepared at 0.33 Pa have a water contact angle around 10° even after 60 min of UV illumination. Fig. 2b shows the variation in absorption at 580 nm by MB for the Ti02 films deposited at different total pressures as a function of UV illumination time. It is seen that the decomposition efficiency increases with increasing total pressure. The highest decomposition efficiency was achieved for the anatase Ti02 film prepared at 2 Pa. With the conversion from anatase to a mixed anatase/rutile film the decomposition efficiency decreases i.e. the rutile phase has a lower ability to decompose MB than the anatase phase.

— A ~J> -_A

c=> J

•

; 1

-k—,

* - . - i

A.

rj\o£

1

• I

* •

_ P ^ * « _

C M

•

i •

2

1

O

O

o

o.

o

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. 8

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. 6

P a

P a

P a

P a

P a

4 5 P a

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I •

2 0 3 0 4 0 5 0 6 0 7 0 8 0

21) [ ° ]

Figure 1. XRD patterns of Ti02 films prepared at different values of the total pressure and at a constant

gas flow ratio q0l l(q02 +gAr) = 27%.

301

0.00 - r

—•-2 .0 Pa - • - - 1 . 0 Pa -••A-0.8 Pa —T~0.7Pa - 0 ~ 0.6 Pa —A—0.45 Pa --O-0.33 Pa

0 10 20 30 40 50 60

UV illumination time [min]

Figure 2. (a) Change in water contact angle and (b) change in absorption of light as a function of illumination time for Ti02 films deposited at different total pressures and at a constant gas flow ratio

Figure 3 a shows the variation of the water contact angle as a function of UV illumination time for films prepared at ptot = 2 Pa and at different oxygen/nitrogen flow ratios. Before UV illumination, the water contact angle decreased with increasing the ratio of nitrogen up to 20%, and above 20% nitrogen the contact angle increased again. After 15 min of UV illumination, the water contact angle of all films prepared in an oxygen/nitrogen mixture became zero while the water contact angle of the film prepared without nitrogen was still around 7°. This is an indication that the addition of a small amount of nitrogen to the sputter gas improves the hydrophilicity of anatase Ti02 films. Also there is a small improvement in the

302

efficiency of decomposition of MB (Fig. 3b) up to 20% nitrogen, while above 20%

the efficiency decreases. This behavior can be understood by analyzing the

composition and structure upon nitrogen addition. Our RBS analysis could not detect

any nitrogen in films. This is also in agreement with XRD analysis which reveals

crystalline T i 0 2 films. Figure 4 shows the XRD of films prepared at p,0t = 2 Pa and

different oxygen to nitrogen flow ratio. It is seen that there is no obvious change in

the crystalline structure. From Fig. 5a, it is seen that the density initially decreases

and then increases with increasing nitrogen flow and simultaneously decreasing

oxygen flow. Also there is a small increase in surface roughness. The decrease in

density and increase in surface roughness result in a more porous microstructure,

which leads to enhanced hydrophilicity and better photo-catalytic properties.

0.00

CO

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-0.04 _

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25

ftiv

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o Hz o o

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to

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15 -

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UV Illumination time [min]

Figure 3. (a) Change in water contact angle and (b) change in absorption of light as a function of the illumination time for Ti02 films deposited at different oxygen to nitrogen flow ratios and at total pressure of 2 Pa.

303

i—•—r

qN/(qN. + q0 ) = 0.267

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-r 1 1 1 1 1 r

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29 [°]

Figure 4. XRD patterns of Ti02 films deposited at different oxygen to nitrogen flow ratios and at total pressure of 2 Pa.

4. Conclusions

Crystalline Ti02 films with anatase and mixed anatase-rutile phase have been successfully deposited on unheated substrates by reactive dc magnetron sputtering. With increasing total pressure the rutile fraction and density decrease while the surface roughness increases. The dependence of the photo-catalytic properties on the rutile fraction, density, surface roughness and the addition of small amount of

304

nitrogen has been studied. The films prepared at higher total pressures that are characterized by higher surface roughness, lower density and lower rutile fraction have a higher photo-catalytic activity. The films prepared at lower total pressures that are a characterized by; lower surface roughness, higher density and higher rutile fraction have a lower photo-catalytic activity. The enhancement of photo-catalytic activity of films prepared at p,ot = 2 Pa with different nitrogen to oxygen flow ratios could be attributed to the decrease in density and the increase in surface roughness.

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Figure 5. Variation of (a) density and (b) surface roughness of Ti02 films at different oxygen to nitrogen flow ratios at total pressure of 2 Pa.

Acknowledgments

One of the authors (S. H. Mohamed) would like to thank the DAAD for a scholarship to carry out this research work at the I. Physikalisches Institut der RWTH Aachen, Germany. The authors are also thankful to C. Horn, O. Kappertz and S. Ziegler for fruitful discussions and experimental support.

305

References

1. M.R. Hoffmann, S.T. Martin, W. Choi, D.W. Bahnemann, Chem. Rev. 95, 69 (1995).

2. M.A. Fox, M.T. Dulay, Chem. Rev. 93, 341 (1993). 3. M. Anpo, T. Shima, S. Kodama, Y. Kubokawa, J. Phys. Chem. 91, 4305 (1987). 4. L. Gao, Q. Zhang, Scripta mater. 44, 1195 (2001). 5. S. Takeda, S. Suzuki, H. Odaka, H. Hosono, Thin Solid Films 392, 338 (2001). 6. Th. Uelzen, J. MUller, Thin Solid Films 434, 311 (2003). 7. T. Sumita, H. Otsuka, H. Kubota, M. Nagata, Y. Honda, R. Miyagawa, T.

Tsurushima, T. Sadoh, Nucl. Instr., Meth. in Phys. Res. B 148, 758(1999). 8. S.H. Mohamed, O. Kappertz, J.M. Ngaruiya, T.P. Leervad Pedersen, R. Drese,

M. Wuttig, Thin Solid Films 429, 135 (2003). 9. S.H. Mohamed, O. Kappertz, T.P. Leervad Pedersen, R. Drese, M. Wuttig,

Phys. Stat. Sol. (a) 198, 224 (2003). 10. B. Lengeler, M.H. Hii ppauff, Fresenius J. Anal. Chem. USSR 346, 155 (1993). 11. E. Chason , T.M. Mayer, Crit. Rev. Solid State Mater. Sci. 22, 1 (1997). 12. P. LSbl, M. Huppertz , D. Mergel, Thin Solid Films 251, 72 (1994). 13. M. Migliuolo, R.M. Belan , J.A. Brewer, Appl. Phys. Lett. 56, 2572 (1990). 14. K. Ishibashi, K. Hirata, N. Hosokawa, J. Vac. Sci. Technol. A 10, 1718 (1992).

MULTILAYER THIN-FILM OPTICAL FILTERS: DESIGN, FABRICATION, AND APPLICATIONS

S.H. KESHMIRI

Electrical Engineering Department, Faculty of Engineering Ferdowsi University, Mashhad 91775, Iran

and Microelectronics Research Laboratory, Faculty of Sciences

Ferdowsi University, Mashhad 91775, Iran

MM. MIRSALEHI

Microelectronics Research Laboratory, Faculty of Sciences Ferdowsi University, Mashhad 91775, Iran

Thin-film optical filters have many applications in a wide range of instruments and high-technology systems. In this article, a short introduction on the design methods of multilayer optical filters is first given, and the three methods that have been used at Ferdowsi University for the design of various optical filters are explained. A brief description of the common optical materials and deposition techniques are also given. Finally, a few examples of the work done at Ferdowsi University on design and fabrication of multilayer thin-film optical filters are included.

1. Introduction

A multilayer thin-film optical filter is a stack of materials deposited on a substrate. The materials used are usually dielectrics and the substrate is normally an optical glass. Optical filters are divided into several types. Some specific types are: antireflection coatings, high reflectance filters, neutral and polarizing beamsplitters, edge filters, minus filters, wide- and narrow-bandpass filters, heat mirrors, and cold mirrors [1,2].

Optical filters have found a wide range of applications in various fields. They are used as part of many scientific instruments, and also in optical communication systems, laser systems, data storage, computers, display devices, energy conversion and conservation systems, IR detection devices and some architectural applications [3, 4].

In this paper, we describe some of the design and fabrication techniques of multilayer optical filters that have been used at Ferdowsi University. Section two is devoted to the design techniques in which three design methods are briefly described. Fabrication of these filters is described in section three, and some experimental results are presented in section four.

306

307

2. Filter Design

A schematic diagram of a multilayer optical filter is shown in Figure 1. The index of refraction and the thickness of the i* layer (1< i < M) are represented by nj and dj. The indices of refraction of the incident medium and the substrate are shown by n0 and nsub. Iinc, Ir, and I, are the power of the incident light, the reflected light, and the transmitted light, respectively. Also, 0inc is used to represent the angle of incidence.

The percentage of power of an incident light that is transmitted (or reflected) by a filter is considered as its main characteristic. This is normally drawn versus the wavelength over an interval of interest. The indices of refraction of the materials used, the number of layers (N), and

m

d3

MM

dm

Figure 1. Schematic diagram of a multilayer optical filter.

the thickness of each layer (usually chosen to be less than the wavelength of the light) affect the characteristic curve of a multilayer optical filter. Also, the type of polarization and the angle of incidence of the light, the refraction indices of the medium and the substrate, and the absorption coefficients of the materials affect the characteristic of the filter. The absorption coefficients are normally ignored; since most of the materials used have low losses and the thicknesses of the layers are small. The filter design problem is, therefore, determining the number of layers, their thicknesses, and the materials to be used in order to achieve a particular spectral characteristic curve.

Different techniques have been used to design thin film optical filters [5]. There are some classical structures, such as the stack of quarter wavelength layers of low- and high-index materials that can be used to obtain some specific filters. However, the design of a filter with an arbitrary spectral characteristic curve is more complicated. Considering the desired spectral curve, the filter design can be viewed as an optimization problem in which a parameter, such as

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the root mean square deviation between the desired curve and the obtained curve, is minimized. This usually requires some extensive numerical calculations. However, the complexity of the calculations is quite within the range of the capability of the present computers.

The numerical techniques in thin-film optical filter design can be divided into two categories. In the first category, an approximate solution is first obtained and then numerical techniques are used to refine the solution. In the second category, there is no need for an approximate solution and a random pattern can be used as the starting point.

In this paper, we briefly describe three methods that we have used to design optical filters. These are: Fourier transform, genetic algorithm, and flip-flop methods. We have developed software packages for these methods at Ferdowsi University. In the following subsections, these methods are described.

2.1. Fourier Transform Method

The discovery of an analytical relation between the spectral characteristic of an optical filter and its index of refraction was first made by Pegis [6] and Delano [7]. This method was later expanded by Sossi [8] as well as Dobrowolski and Lowe [9]. As described by Delano, the following relation can be used to relate the index of refraction and the spectral characteristic of an optical layer

lnm==j_+jP(v}expi_a;rox)d(7 (1)

where a = 1 A, k being the wavelength of the light, and p is a parameter that is related to the spectral characteristic of the filter. The index of refraction in the above relation is considered as a continuous function of x, where the x-axis is normal to the filter surface and x = 0 is the interface of the filter and the surrounding medium. Several choices for p have been suggested [9]. In our work, we have used the relation p = (R/T) m, where R and T are the power reflection and transmission coefficients of the filter. Due to the analogy of Eq. (1) and the Fourier transform integral, this filter design method is known as the Fourier transform method or inverse Fourier transform method.

The procedure of designing a filter is as follows. For a given T (or R) function, first, the p parameter is found. Then, by solving the integral equation (1), the index of refraction of the filter is determined. The integral equation is usually not easy to solve and its solution is found by numerical techniques. As a result, n will be determined as a continuous function of x. This type of filter is called inhomogeneous. The resultant n can be made discrete to obtain a multilayer filter. In the latter case, n is constant within each layer. Hence, this type of filter is called homogeneous.

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The discrete values of n obtained as described above, might not be physically available. To solve this problem, Herpin's method [10] is used to obtain each desired value of n from a combination of two available materials with indices of refraction nH and nL. It should be mentioned that the Fourier transform method provides an approximate solution. To refine the design, it should be followed by an improving technique. In our case, we used the gradient method to obtain a better solution [11].

2.2. Flip-flop Method

The flip-flop method introduced by Southwell [12] is an example of a design technique that does not need an approximate solution. In this method, two materials with indices of refraction nH and nL are used. The starting point is a stack of layers with equal optical thicknesses. The initial distribution of the index of refraction is not important. All the layers can be considered having nH

or nLor any arbitrary combination of these two distributions. A merit function based on the desired spectral response is considered for

evaluating the designed filter. An example of such function is

f 1 N Y'2

Afl^j-SfaW-GWpj (2) where 0,(Xj) and O(Xj) are the target and the actual values of the characteristic

parameter (usually reflectance or transmittance) of the filter at wavelength X̂ , respectively. N is the number of wavelengths at which the difference between Ot

and O is included. The procedure for designing a filter using the flip-flop method is as follows.

First, a total physical thickness is selected for the coating and it is divided to a large number of "thin" layers. These thin layers should have thicknesses much smaller than (e.g., on hundredth of) the wavelength of the incident light in the material. The layers are arbitrarily assigned values of nH and nL. Next, the spectral response of the present filter and its merit function are calculated. The state of each layer is changed one at a time from low index to high index or from high index to low index. If the merit function is improved, the new state of that layer is kept; otherwise, it is discarded. This procedure is continued until the merit function is kept constant after all the layers are tested. After the procedure is ended, the neighboring layers that have the same index of refraction are combined. Therefore, the final number of layers is usually much smaller than its initial value.

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2.3. Genetic Algorithm Method

Genetic algorithm is a powerful optimization technique. It is based on the survival law of the fittest species in nature. Here, a population of different filters that are obtained by random selection is used as the starting point. A fitness function, such as Eq. (2), is then calculated for each filter. Those filters that have lower values of fitness function are selected and the others are discarded. Applying the genetic algorithm operations of crossover and mutation [13], new filters are created to keep the population at a predetermined level. The above procedure is repeated as many times as required to get close enough to the desired characteristic curve.

The genetic algorithm is capable of obtaining the optimum solution. However, in order to reduce the computational time, one might prefer a local minimum with a relatively close characteristic curve to the global minimum. The genetic algorithm method has been used to design different types of filters [14].

At Ferdowsi University, we have developed a software package for designing optical filters by genetic algorithm method. Using this software, we have designed several types of filters including polarizing beamsplitters [15, 16], as is described in the next section.

3. Optical Materials and Filter Fabrication Techniques

The most common optical materials are mainly oxides, fluorides, and sulfides. For example, the following materials can be used in optical filter fabrication: A1203, A10xNy, Sb203, BeO, Bi203, BiF3, CdS, CdTe, CaF2, Ce02, CeF3, Chiolite (5NaF.3AlF3), Cr203, Cryolite (Na3AlF6), GdF3, Ge, Hf02, HfF4, LaF3, La203, PbCl2, PbF2, PbTe, LiF, LuF3, MgF2, MgO, NdF3, Nd203, Pr60„, SrnF3, Sc203, Si, SiO, Si203, Si02, Si3N4, NaF, SrF2, Ta205, Te, Ti02, TICl, Th02, ThF4, YbF3, Y203, ZnSe, ZnS, and Zr02.

Important parameters in choosing an appropriate material are: optical properties (such as refractive index, region of transparency, and percentage of absorption), suitable deposition method, mechanical properties (such as hardness, resistance to abrasion, and magnitude of built-in stresses), chemical properties (for example, solubility, resistance to the attack by atmosphere, and compatibility with other materials), price, availability, and other properties that might be important in some particular applications (for instance, electrical conductivity, or dielectric constant of the material). Some metal films are also used in optical filter applications (as an absorbing or reflecting layer); some examples are: Ag, Al, Au, Cu, Ni, Cr, and Rb.

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In choosing a suitable material, the following considerations can also be important: Generally, oxides are harder than fluorides, sulfides, and semiconductors. Optical properties of some materials (specially semiconductors) depend on the temperature. Metal films are generally soft and may tarnish when exposed to the air. Adhesion of some materials to the glass substrate is low. Microstructure and properties of a thin film depend on the material, deposition technique, and deposition parameters. Optical properties depend on the microstructure (films can be dense or porous, amorphous or polycrystalline, or have a columnar structure). Films can be under tensile or compressive stress (this can lead to deformation of the substrate or fracture of the film).

Thin film deposition techniques can generally be divided into two main categories: chemical vapor deposition (CVD) and physical vapor deposition (PVD). CVD techniques are faster and less expensive, but they usually do not produce very high quality coatings (compared to PVD methods). Since film uniformity, thickness accuracy, and process repeatability are of high importance in optical coatings, the PVD techniques are commonly used in the fabrication of multilayer optical filters. Plasma-enhanced CVD processes are also used for deposition of some organic polymer films.

The main PVD techniques used for production of optical filters are vacuum (thermal) evaporation and sputtering. Vacuum evaporation is the predominant technique for production of optical coatings. In this technique the material to be deposited is heated in a vacuum chamber to a temperature at which it vaporizes. The vapor then condenses on a substrate as a thin solid film. For heating the material, resistance heating, RF induction, electron-beam heating, or laser heating can be used. The pressure in the chamber is usually of the order of 10"5

mbar (or less). The evaporation process can also be done reactively, by introducing some reactive gas into the chamber (at a suitable partial pressure). The substrate is sometimes heated to few hundred degrees Celsius. Evaporation rate and the total film thickness are monitored during the film-deposition period. Since in most cases the characteristics of the final filter critically depend on the film thicknesses, this monitoring step is an essential part of the fabrication procedure.

Due to the rather low density of the films deposited by conventional thermal-evaporation technique, the humidity and temperature stability of these films is not so good (unless they are specially protected). Film porosity depends on the material, substrate temperature, pressure in the vacuum chamber, evaporation rate, and angle of incidence of the evaporate species. Water vapor can be trapped in the pores, and result in change of index of refraction of the layer. Consequently, spectral characteristics of the filter can be shifted. This effect is somehow reversible. Advantages of the conventional thermal

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evaporation technique are: simplicity, flexibility, and relatively lower cost. On the other hand, porosity (low-density) of the films, and poor humidity and temperature stability of the fabricated filters are its major shortcomings.

To improve the quality of deposited films, more energetic processes such as ion-assisted evaporation, reactive ion plating, and magnetron sputtering techniques have been used. These methods produce films with improved humidity and temperature stabilities. With bulk-like properties, these films have also more reproducible optical constants [17]. Sputtering techniques (reactive, non-reactive, DC, RF, and magnetron) produce more stable layers, with bulklike densities of the layers (no ageing effects), and good film-thickness control. Deposition on large-area substrates is also possible. However, the process is slow, and sputtering targets are usually expensive.

4. Design and Experimental Procedure

Design of several optical filters was done by the software programs described in section two. In most cases, the genetic-algorithm software produced the best results. Thermal-evaporation technique was used for the fabrication process; and an oscillating quartz film-thickness monitor was used for the evaporation-rate and film-thickness measurements. A double-beam optical spectrometer was used for measuring and recording the spectral response of prepared samples.

The following filters were fabricated using the above procedure: antireflection coatings, high-reflectance filters (i.e., all-dielectric mirrors), neutral and polarizing beamsplitters, EDFA (Erbium-Doped Fiber Amplifier) gain-compensating filters, narrow bandpass filters, edge filters, cold mirrors, and heat-reflecting coatings.

Design procedure was as following. The ideal spectrum of the desired filter was given as an input to the genetic algorithm program (along with the number of chosen materials and their refractive indices, maximum number of layers, desired range of film thicknesses, incidence angle, type of polarization, and refractive indices of the substrate and incidence medium). Then, the output of the program was a table of film thicknesses and materials for each layer of the designed filter, as well as a plot of transmittance (or reflectance) of the filter in a specified range of wavelengths.

As an example, the design of an edge filter with a cutoff wavelength of Ac = 1550/7/w , is described here. The ideal filter should reflect all radiations below Xc (i.e., R = 1 for/I <1550ww) and transmit all radiations above Xc

(i.e., R = 0 for/l > 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,

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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 J-

20

10 h

1300 1350 WOO 1450 1500 1550 1600 1650 f. Wavelength (nm)

00

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

1 M : • . - (b) ; ,

• ••" fa) -'

<S~ 60

1400 1500 1«B Wavelength (nm)

ITflO

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).

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As a few examples, designed spectra of an antireflection filter, a high-reflectance 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 front-surface 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 ? r

Wavdeqgh (nm)

Figure 4. Spectrum of an antireflection filter designed for the visible region.

315

100

90

eo

i so ts.

20

10

; • • ' •

VT" T 1

: : :

: i :

: : :

': : :

i i i

i i i

1 i— — r • — " i " V " I

I Li! \ L±.

H ' : : \

• i • ,

I ! | -i i . i f , , 300 350 400 450 500 550 600 650 700 750 800

Wavelength (nm)

Figure 5. Spectrum of a high-reflectance filter designed for the visible region.

100

500 600 700 Wavelength (iim)

800 900

Figure 6. Designed spectrum of a polarizing beam-splitter filter.

316

I - '1 1 "1-1 "1 1 1

j : :

• i ; : : :

1 i [j.. ! ] L ! ! \

: i :

i i i i i i

^-i i i «00 1400 1500 1S00 1700 1800

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. H.A. Macleod, Thin Film Optical Filters, 3rd edition, (Institute of Physics Publishing, London, 2001).

2. J.D. Rancourt, Optical Thin Films User Handbook, (SPIE Optical Engineering Press, Bellingham, 1996).

3. 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. 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).

5. J.A. Dobrowolski, Optics News 6, 24 (1997). 6. R.J. Pegis, Journal of the Optical Society of America 51, 1255 (1961). 7. E. Delano, Journal of the Optical Society of America 57, 1529 (1967). 8. L. Sossi, "A method for the synthesis of multilayer dielectric interference

coatings," Translated by J.A. Dobrowolski, (National Research Council of Canada, 1974).

9. J.A. Dobrowolski, D. Lowe, Applied Optics 17, 3039 (1978). 10. J.A. Dobrowolski, S.H.C. Piotrowski, Applied Optics 21, 1502 (1982). 11. 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.

12. W.H. Southwell, Applied Optics 24, 457 (1985). 13. D.E. Goldberg, Genetic Algorithms in Search, Optimization and Machine

Learning (Addison-Wesley, MA, 1989). 14. E. Michielssen, S. Ranjithan, Mitra, IEE Proceeding Journal 139, 413

(1992). 15. 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)

16. M. Shokooh-Saremi, S.H. Keshmiri, M.M. Mirsalehi, M.M. Bagheri-Mohaghgheghi, 7th Iranian Conference on Electrical Engineering, (Tehran, Iran, Proc. of EM Fields and Applications papers, 139-145, 1999).

17. 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 photo-induced 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, light-induced 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

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

-> atomic mobility

phase transitions, ablation (VOi, SmS, polymers)

photocroms (KC1, glasses )

photorefractive effect (LiNb03:Fe'3,Sn2P2S6)

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

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: a0 =10' - 105 cm"1, Aa/a0 = 10 - 100(photocroms, AsSe for example), n = 1,5-3,0, An=10-4-10"1 (LiNb03, As2S3), Ro = 5 - 20 %, AR/R = 0,1 - 0,5 (V02, 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 +12s — 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

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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, -> RELAXATION -> AMORPHYSATION

CRYSTALLISATION- DIFFUSION

CliG Metal

?^£„ • im lw Charge

Recoriliii|>

i n n n n n n Reamiing

Recording j^sassas 3 Recording

- « Development > Aj4"iw*rrj4wv4"'

Etching Kelief

Figure 4. Optical recording processes in chalcogenide glasses.

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

322

(a) intensity

(b) track

be tore erasing

after recording

(c) detected signal

record in

erasing readout

1

1

o o o

o o o o

TTTT

time

1

1

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><1010 bits per square inch), to realize the holographic recording on the disc and to increase the transformation speed (amorphysation-crystallization).

Intensive investigations are performed in the direction of nano- or even micrometer-high relief formation on the surface, realisation of the self-organization process for data recording. Stimulated interdiffusion in nanolayered structures, sctructural transformations on the nano-scale promise new possibilities for such processes [6]. An example of surface holographic grating, recorded on the amorphous nanolayered structure is presented in Fig.6.

Summarizing the known results and taking into account the physical limits of optical recording processes it is possible to predict the further development in the next few years on the basis of atomic engineering of light-sensitive materials, thin layers and nanostructures, new light sources (blue laser), near-field optics.

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Figure 6. AFM picture (1000x1000x100 mn3) of the surface relief, recorded in a one-step process on the a-Se/As2§3 nanolayered structure.

References 1. Photonics. Bdby M. Balkanski, P. Lallemand. (Gauthier9Vil!ars? Paris,

1975). 2. HJ. Caulfield, Handbook of Optical Holography. (Academic Press, New»

York, 1979). 3. K. Svartz, Physics of optical recording in dielectrics, (semiconductors.

Zinatne, Riga, in Russian 1986). 4. K. Tamaka, Rev. Sol St. Scl 4,641 (1990). 5. A. Kikineshy, Optical Memory, Neural Networks 4, 177 (1995). 6. V. Palyok, M. Mishak, I. Szabo, D. Beke5 A. Kikineshy, Appl. Phys. A68,

489, (1999).

DIFFUSION OF ATOMIC HYDROGEN AND PASSIVATION OF STRUCTURAL DEFECTS IN SILICON AND IN TRANSPARENT-

CONDUCTING THIN FILMS

S.H. KESHMIRI

Microelectronics Research Laboratory, Faculty of Sciences Ferdowsi University, Mashhad 91775-1436, Iran

Dangling bonds associated with various forms of structural defects in semiconducting materials have a negative effect on electrical characteristics, and hence, performance of the devices made of these materials. Atomic hydrogen can deactivate dangling bonds and improve device performance. Structural defects were created in single-crystal silicon wafers by ion-implantation technique. ECR (Electron Cyclotron Resonance) method was used to study the effect of atomic hydrogen in damaged silicon samples. Hot-filament hydrogenation was also employed for polysilicon, indium-tin oxide, and zinc oxide samples. In all cases, the hydrogenation process produced considerable increases in drift mobility and lifetime of charge carriers in the samples.

1. Introduction

Varieties of structural defects (with different concentrations) exist in all forms of solid silicon. In bulk single-crystal silicon these are mainly point defects (vacancies, interstitials, and impurities); which may be originally present in the starting material (e.g., created during crystal growth or wafer-preparation steps), or can be created during some fabrication process (e.g., ion-implantation or plasma-etching). Grain boundaries and point defects are major defects in polycrystalline silicon; and surface states are the main defect in porous silicon samples (due to presence of nanometer-size wires and dots). Point defects are abundantly present in amorphous and thin-film silicon samples. In transparent-conducting thin films, the main defects in amorphous films are point defects, and in polycrystalline samples there are grain boundaries and point defects.

The "dangling bonds" created by these defects act as traps for charge carriers, and hence reduce lifetime and drift mobility of the carriers. This, in return, produces negative effects on the performance of the devices made of these materials.

In the past two decades, there has been a great deal of interest on study of the effects of atomic hydrogen in semiconductors [1-6]. One reason is due to the fact that hydrogen is incorporated into semiconductor samples during several device-fabrication processes. Also it is shown that atomic hydrogen can easily diffuse into the semiconductor material, deactivate the traps (i.e., dangling bonds), and hence, increase charge-carriers lifetimes and drift mobility in the

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325

material. Therefore, it can noticeably improve performance of the devices made of these materials.

For these reasons, several techniques have been developed for hydrogenation of semiconductor samples [1,2]. For example, Kaufman source, ECR (Electron Cyclotron Resonance), and hot-filament techniques have been used for producing atomic hydrogen plasma, and hydrogenation of the samples.

In single-crystal silicon samples, the effects of hydrogenation are: passivation of deep-level defects related to some metal impurities (e.g., Au, Pt, Pd, Mo, Cu, Ni, and Fe), passivation of point defects, passivation of surface and interface states, deactivation of host lattice dopants, and generation of some H-related defects. In polycrystalline silicon specimens, atomic H can result in: passivation of grain boundaries, improving performance in MOSFETs & TFTs (i.e., increasing carrier mobility, increasing on/off ratio, and reducing threshold voltage), and improving conversion efficiency in poly-silicon solar cells. In amorphous silicon films, considerable passivation of dangling bonds has been observed; and in porous silicon samples, the effect of H on the luminescence process has been studied. Also, in silicon grown by the EFG (Edge-defined Film-fed Growth) technique, increased carrier lifetime (due to H incorporation) has been observed [7].

In transparent-conducting films, a considerable increase in carrier drift mobility (and hence, electrical conductivity) has been reported [8-12]. Considering the widespread use of these films (as a transparent electrode) in many optoelectronic devices, this increase in electrical conductivity of theses films can be important in reducing power-consumption and improving device performance.

On the negative side, high-energy hydrogen ions can produce additional defects in the samples [1, 2]. Therefore, a low-energy hydrogenation process (such as the hot-filament method) has a clear advantage over high-energy techniques (such as the Kaufman and ECR). On the other hand, the hydrogen plasma densities produced by ECR are much larger, compared to a single hot-filament source; so the hydrogenation times are shorter when ECR technique is used.

In this research, ECR technique was used for hydrogenation of the damaged single-crystal silicon samples; and hot-filament method was employed for hydrogenation of polysilicon and transparent-conducting films.

326

2. Sample Preparation and Characterization

2.1. Single-Crystal Silicon Samples

In order to produce structural defects in single-crystal silicon samples, silicon wafers were implanted by Si ions. In this way, the defects created in the samples were "pure" structural damages (and not related to impurity atoms). N-type (boron-doped) Si wafers (p = 1-10 ficm, [100] orientation) were implanted using a Varian (model 350D) implanter. Si ions were derived from SiF4 gas, and were accelerated to 20 keV energy. The implantation dose was varied from 1012

to 1015 cm"2; and the implantation process was done at room temperature. The effects of the induced damages were evaluated by Schottky-diode I/V

measurements. Thermally evaporated Au, Ti, and Al films were used for Schottky and back-contact electrodes. On each sample, several Schottky dots (with known areas) were evaporated, and their I/V characteristics were recorded. I/V characteristics of the diodes were recorded by a HP 4145B Semiconductor Parameter Analyzer, at room temperature, and in some cases, at lower temperatures (down to 100 K) as well. Deep Level Transient Spectroscopy (DLTS) technique was used to measure trap concentration; and spreading-resistance profiling was used for studying the depth of the implanted region. ECR technique was used for hydrogenation of these samples.

2.2. Potycrystalline Silicon Samples

In this case, samples were p-type polysilicon wafers taken from a polysilicon solar-cell fabrication line (after n-type-layer diffusion step). So the samples had a p-n junction already built into them. To complete the diode fabrication process, metal electrodes were evaporated onto the top surface of the wafer. The back electrode covered the whole wafer surface; but the front electrode was deposited through a patterned mask. Fig. 1 shows a schematic view of these samples. For these samples, hot-filament method was used for the hydrogenation process.

N-Type Diffused Layer ^ p r

P-Type — 4 — Poly-Si Wafer

Top Electrodes

Back Electrode

Figure 1. Basic configuration of the polysilicon samples.

327

A sample holder with a resistance heater was used for heating the sample during the hydrogenation process. I-V characteristics of each sample were measured before and after the hydrogenation process. A board inside a personal computer (which measures 4100 fixed points for each I/V plot) produced the I/V plots.

2.3. Transparent-Conducting Films

Indium-Tin Oxide (ITO) films were deposited by spray-hydrolysis technique. Zinc oxide samples were prepared by two methods. Evaporated ZnO films were deposited by thermal evaporation of ZnO powder in vacuum, and by spray-pyrolysis technique [10, 12, 13].

Aluminum electrodes were evaporated onto the two ends of each sample (for measuring sheet resistance of the sample, and to make electrical connection to the sample for hydrogenation process). Fig. 2 shows the basic configuration of theses samples.

A

Aluminum Electrode *

Glass Substrate

T Transparent-conduting film

1

Figure 2. Configuration of the ITO and ZnO samples.

Sheet resistance and electrical conductivity of each sample were measured by a setup similar to polysilicon samples. Hot-filament method was used for the hydrogenation process in these samples. Film transparency spectra were recorded before and after the hydrogenation process by a double-beam spectrophotometer (Perkin-Elmer, model Lamda-9). Hall-effect measurement was done for determination of free-carrier concentration in the films.

2.4. Hot-Filament Hydrogenation Set-up

The plasma source consisted of a glass tube with a tungsten filament, heated by a 6 V (ac) power supply, as the electron source. This filament was placed between two metal rings. The lower ring (anode) and the upper ring (cathode) were connected to the accelerating voltage. Another power supply was used for adjusting hydrogen ion energy. Sample was placed on top of the plasma source. The whole setup was placed inside a vacuum coating unit (Edwards E306-A).

328

High purity (99.99%) hydrogen gas was fed into the source through its lower end. H2-gas flow was controlled by the inlet valve of the coating unit. A sample holder with a resistance heater, a thermocouple, and a temperature-controller circuit was used for heating the sample during the hydrogenation process.

Zr . Sustrate Holder/Heater

— Sample

. Glass Tube

t t t H2Gas

Figure 3. Configuration of the plasma source.

3. Results and Discussion

The results obtained for silicon implantation, and hydrogenation of single-crystal silicon, polysilicon, ITO, and ZnO samples are given in this section. Generally, all the tests showed some degree of defect compensation by the atomic hydrogen; however, the amount of passivation depended on several factors, as is described below.

3.1. Silicon-Implantation of Single-Crystal Si Wafers

At the low Si ion energy used in these experiments, nuclear stopping should be dominant, and the damage profile is expected to be similar to (though slightly shallower than) the implanted ion profile. For Si ion energy of 20 keV, TRIM (Transport of Ions in Matter) simulation program predicted a range of about 335 A, and a maximum penetration depth of around 700 A. Thus the extent of the damaged region could not extend much farther than 700 A below the substrate surface. In other words, only near-surface damages were created here. This was actually confirmed by the spreading-resistance results, which gave the extent of the damaged area to be around 0.1 (am for the p-type sample implanted with a

329

dose of 1013 cm"2 Si ions, and slightly more, for the 1015 cm"2 Si-implanted sample.

Silicon implantation at 1012 cm"2 dose had almost no noticeable effect on I/V characteristics of the Schottky diodes made on the implanted samples. Evidently, the damages produced by this low dose (which took only a few seconds of the implanter's time) did not alter the Si surface significantly, and therefore, may be considered as the threshold of the implantation effects. On the other extreme, the 1015 cm"2 dose (which took longer than six hours to complete) produced so much damage that the implanted samples behaved like amorphous specimens and showed an almost-ohmic characteristic.

Figure 4 shows the effect of Si implantation on I/V characteristics of Schottky diodes made on implanted and control n-type samples. As this figure indicates, the implantation process (with a dose greater than 1012 cm"2) greatly lowered the Schottky barrier height.

DLTS results showed an increase of about one to three orders of magnitude in trap concentration in the Si-implanted samples (compared to the control samples).

10"

10-*--1

a 10

b 3

-6

O 10-

10' 10

10" 12

(a)

JM

(c> -W).

' *

•

i

i

•̂ ••J

/

-1.2 -0.48 Reverse Voltage (V)

0.48 1.2 Forward

Figure 4. Schottky barrier height reduction of n-type Si due to Si implantation, with a dose of: (a) 1015 cm-2, (b) 1013 cm"2, and (c) 1012 cm-2, (d) control sample (no implantation). (In the reverse bias, the absolute values of current are plotted versus applied voltage).

3.2. ECR-Hydrogenation of n-Type Single-Crystal Silicon Samples

Effect of H in samples with 1013 cm"2 implantation dose was a considerable recovery of I/V characteristics towards the I/V of control sample (see Fig. 5b).

330

Increasing hydrogenation time (tH) increased the degree of recovery. As Fig. 5c indicates, most of the structural defects (produced by implantation) were evidently passivated at two hours hydrogenation time.

1Q-2-

1 0 " 4 .

s o

io-10-

10-12

(a)

_®_

(o) —-—

\ 'a

1 ' ! fi 1/ / If

,'j

/ / if

^ - ^ " # T " ^ * "

-1,2 -0.48 0 0,48 , 12 Reverse Forward

Voltage (V) Figure 5. Effect of ECR-hydrogenation on I/V characteristics of n-type Si implanted with 1013 cm"2

dose and Ts„b=250 °C: (a) Implanted sample with no hydrogenation, (b) Implanted sample after 30 minutes hydrogenation, (c) Implanted sample after 120 minutes hydrogenation, and (d) original control sample (no implantation, no hydrogenation).

Higher substrate temperature produced mixed results. At 30 minutes hydrogenation time, increasing Tsub to 490 °C increased the Schottky barrier slightly, but at 120 minutes time, Tsub of 490 °C reduced Schottky-barrier height [9].

The effect of hydrogenation on samples with 1015 cm"2 implantation dose was not as prominent. In fact, the density of the structural damages produced by ion implantation at this high dose was so great that the Schottky diodes made on these samples showed an almost ohmic behavior (see Fig. 6a). This was clearly due to an amorphous-like nature of the implanted specimens. Thirty minutes hydrogenation did not produce a significant change in I/V characteristics of these samples. Although both forward and reverse currents were reduced to some extent, there was slightly more reduction in the reverse leakage current (see Fig. 6b). However, 120 minutes hydrogenation time showed more significant change and increased the Schottky-barrier height to some extent (see Fig. 6d).

331

Rwense Voltage (V)

Forward

Figure 6. Effect of ECR-hydrogenation on I/V characteristics of n-type Si implanted with 1013 cm"2

dose: (a) Implanted sample with no hydrogenation, (b) Implanted sample after 30 minutes hydrogenation at 1^=250 °C), (c) Implanted sample after 120 minutes hydrogenation at 1^=490 °C, (d) Implanted sample after 120 minutes hydrogenation at Ts„b=250 °C, and (e) Original control sample (no implantation, no hydrogenation).

Increased substrate temperature helped the recovery process considerably and produced a more Schottky-like I/V characteristics, as shown in Fig. 6c. However, as it is evident from this figure, there must have been many defects still left in this sample [9].

3.3. Hot-Filament Hydrogenation of Polysilicon Samples

A polysilicon diode sample was hydrogenated for 30 and 60 minutes durations. Ionic current was 2.5 mA, and the sample temperature was raised to 200 °C during the hydrogenation process. The results are plotted in Fig. 7.

As this figure indicates, the hydrogenation process produced an increase of about 100% in the forward-bias current of the diode, and a considerable decrease (by a factor of -0.3) in the reverse-bias leakage current [14].

332

E

s b o

Voltage (V) Figure 7.1/V characteristics of polysilicon: (a) reference sample, (b) sample heated to 200 °C and hydrogenated for 15 minutes, and (c) sample heated to 200°C and hydrogenated for 60 minutes.

3.4. Hot-Filament Hydrogenation oflTO Transparent-Conducting Films

An ITO film with a thickness of 400 run and sheet resistance of 59.4 Q/D was hydrogenated at five successive stages. The duration of hydrogenation at each step was 45 minutes, but the substrate temperature was increased 50°C for each stage. The results are plotted in Fig. 8. As this figure shows, Rs was reduced to 21.6 Q/D after hydrogenation at Tsub = 200 °C. This indicates 175 % increase in the electrical conductivity of the sample [12].

Substrate Temperature (°C) Figure 8. Variation of sheet-resistance (Cl/D) of an ITO film with substrate temperature during the hydrogenation process ("x" denotes Rs of the sample before hydrogenation).

333

Another ITO sample with sheet resistance of 31.85 Q/D was hydrogenated at the optimum substrate temperature of 200 °C for different durations. The results are given in Fig. 9. At the optimum hydrogenation time of 45 minutes, sheet resistance of the sample dropped to 7.35 Q/D, which indicates an increase of-330 % in the electrical conductivity of the ITO sample [12].

• 8

•a

J5 CO

10 20 30 40 50 60 Hydrogenation Time (min.)

Figure 9. Dependence of sheet resistance (C2/D) of an ITO sample to hydrogenation time at T„b=200 °C.

The transmission spectra of an ITO sample before and after hydrogenation (to the optimal condition for lowest resistivity) are shown in Fig. 10. As this figure shows, the transparency of the film is improved in the visible region, and the transmission in the NIR part of the spectrum is reduced. This is as expected, because of the increased electrical conductivity of the sample [12].

a

C

e

80-

60-

4 0 -

20 -

0 -

( b ) ^ v ^

1 »

J 1 , , r-

\ \ ( a )

\ ( b )

— \ 1 1 •

500 1000 1500 2000

W a v e l e n g t h (nm)

2500

Figure 10. Transmittance spectra of (a) an as-deposited ITO sample, and (b) after hydrogenation at Ts„b= 200 °C for 45 minutes.

334

3.5. Hot-Filament Hydrogenation ofZnO Transparent-Conducting Films

The results for ZnO samples were generally similar to the ITO case. For example, for a ZnO sample prepared by spray- pyrolysis method, hydrogenation at Tsub = 250 °C for 45 minutes resulted in 200% increase in the electrical conductivity [10]. The hydrogenation process was even more profound for the evaporated ZnO films (which are expected to have a higher concentration of point defects, due to the fact that they were deposited at a lower temperature, compared to the sprayed samples). For example, hydrogenation of an evaporated ZnO sample (at optimal conditions) resulted in -340% increase in electrical conductivity of the sample [10]. Hall-effect measurements showed that the hydrogenation process did not affect free-electron concentration in both ITO and ZnO films considerably. For example, a ZnO film before hydrogenation had 3.00x1018 cm"3 free electrons; but after one hour hydrogenation at Tsub= 250 °C, the concentration rose to 4.34x1018 cm"3 [10].

The results of electrical resistivity, free-electron concentration, and electron mobility variations with hydrogenation times for a sprayed sample are given in Fig. 11.

i — i — i — i — i — • — i — i — i — i — i — i — r 0 10 20 30 40 50 60

Hydrogenation Time (min.)

Figure 11. Variations of free-electron concentration (N), electrical resistivity (p), and electron mobility (n) with hydrogenation time for a ZnO film hydrogenated at Tsull = 250 °C.

Optical transmission spectra of a ZnO sample before and after the hydrogenation process produced results similar to the ITO case [10].

335

4. Conclusions

At low and medium implantation doses, the ECR hydrogenation was quite successful in passivation of the structural defects in the implanted silicon samples. In fact, the Schottky-barrier height of the diodes made of the implanted and hydrogenated samples showed better than %90 recovery of the barrier height (compared to control samples). However, at very high implantation dose, the recovery was not so significant.

Hot-filament hydrogenation produced considerable improvement in both forward-and reverse-bias currents of p-n junction diodes made of poly-silicon samples.

Hot-filament hydrogenation of various transparent-conducting films of indium-tin oxide and zinc oxide showed considerable increase in electron mobility and hence electrical conductivity of the samples. Considering the wide range of applications of these films in a variety of optoelectronic devices, the increase in electrical conductivity of TC films may result in significant improvement of performance of those devices.

References

1. J.I. Pankove, N.M. Johnson, (eds.), "Hydrogen in Semiconductors", Semiconductors and Semimetals 34, Academic Press (1991).

2. N.H. Nickel, "Hydrogen in Semiconductors II", Semiconductors and Semimetals 61, Academic Press (1999).

3. S. Ashok, J. Chevallier, I. Akasaki, M.N. Johnson, B.L. Sopori, "Defect and Impurity Engineered Semiconductors and Devices", Materials Research Society Symposium Proceedings 378 (1995).

4. F. Nuesch, E.W. Forsythe, Q.T. Le, Y. Gao, L.J. Rothberg, Appl. Phys. Lett. 87, 7973 (2000).

5. K. Srikanth, S. Ashok, J. Vac. Sci. Technol. A10, 1118 (1992). 6. J. Poortmans, M. Rosmeulen, A. Kaniava, J. Vanhellemont, H. Elgamel, J.

Nijs, in: S. Ashok, J. Chevallier, I. Akasaki, N.M. Johnson, B.L. Sopori, (eds.), "Defect and Impurity Engineered Semiconductors and Devices", Materials Research Society Symposium Proceeding 378, 399 (1995).

7. J.I. Hanoka, C.H. Seager , D.J. Sharp, J.K.G. Panitz, Appl. Phys. Lett. 42(7), 618 (1983).

8. S.H.H Keshmiri, C.W. Nam, S. Ashok, Proceedings of the "Materials Research Society", Symposium A, (MRS Fall meeting, November, Boston, USA 1993).

9. S.H. Keshmiri, Proceedings of "The 12th. Int. Conf. on Microelectronics (ICM 2000)", Oct. 13-Nov. 2 (2000), Tehran, Iran.

10. S.H. Keshmiri, M. Rezaee, Thin Solid Films 382, 230 (2001).

336

11. S.H. Keshmiri, M. Rezaee, A.R. Barati, The 11th. Int. Workshop on the Physics of Semiconductor Devices Dec. 10, 2001, New Delhi, India (IWPSD2001).

12. S.H. Keshmiri, M. Rezaee, S. Ashok, Thin Solid Films 413, 167 (2002). 13. S.H. Keshmiri, M.M. Bagheri, S. Ojaghi, M. Rezaee, "13th. Int. Conf. on

Surface Modification Technologies", Sept. 1999, Singapore. Also Published by ASM Int., (Mterials Park, Ohio, USA 1999).

14. M. Rezaee, S.H. Keshmiri, A.R. Barati, Proceedings of the 5th Conf. on Condensed Matter Physics, Center for Higher Education in Basic Sciences, (Zanjan Iran, May 1999 in Persian).

THE EFFECT OF PARTICLE SIZE ON OPTICAL PROPERTIES OF CdS FILMS FORMED BY PHOTOCHEMICAL TECHNIQUE

S.M. MAHDAVI, A. IRAJI-ZAD, F. RAZI, M. REZAESMAEILI

Department of Physics, Sharif University of Technology, Tehran, Iran P.O. BOX 11365-9161

E-mail: mahdavi@sharif.edu

Cadmium sulfide thin films were formed on glass and ITO/glass substrates using a photochemical method from an aqueous solution. The samples were placed both near the solution surface and the bottom of that during and after UV irradiation respectively. The experimental results showed that the adhesion of films formed on the ITO/glass substrates were much better than those on glass, furthermore all samples which were placed near the surface solution were more adhesive than others. The SEM experiment showed the granular structure of deposited films. The XRD spectra showed an amorphous structure for the as deposited thin films, whereas those annealed at 300°C and 500°C in the N2 atmosphere had cubic and hexagonal structure respectively. The band gap of the CdS films were calculated from the optical spectra as well as the CdS particles in the irradiated solution. The size of the particles in the solution and films were calculated using the effective mass approximation method. The results confirmed that the band edge and size of the particles depend on the irradiation time. By controlling these parameters, the particle size can be reduced to nanometer range.

1. Introduction

Cadmium sulfide has been widely studied during the last 30 years. Its thin films have been used in CdS-CdSe sollar cells, optical sensors for visible light, lasers and optical waveguides[l-7]. In addition, nanocrystals of CdS exhibit attractive optical properties that are used in luminescence devices Different techniques are used to form CdS thin films, such as thermal evaporation, sputtering, pulsed laser deposition, chemical bath deposition and photochemical deposition(PCD) technique. In PCD method the film formation occurs during the UV irradiation of the solution. In fact photons enhance the chemical reaction .

The aim of this paper is to study the physical properties of CdS thin films formed by PCD and comparision with those that formed after sedimention of the irradiated solution. We used SEM, AES and XRD techniques for structural and compositional analysis. The band gap and the size of CdS particles were calculated using spectral transmision of thin films and effective mass approximation method respectively.

337

338

2. Experimental Methods

Glass and ITO coated cornig glass 7039 substrate were used. Aqueous solution were prepared using 5N CdS04 powder and sodium tiosulfate. The pH was varied by adding H2S04 and the solution was stirred by a magnetic stirrer. The solution is tranparent for wavelengths longer than 300 nm. Then the substrate was held in the solution about 3 mm under the solution surface, and was irradiated by a 80W UV source. Among the strong lines of the mercury lamp, only the line at 254 nm (4.9 eV) is absorbed by sodium tiosulfate molecules in the solution, and S203

2" ions are dissociated to release S atoms. S203

2" + hv-» S + S032"

It is known that S is released from S2032" ions in an acidic solution as the

following reaction: 2H+ + S203

2" -> S + H2S03

The S2032' ions also supply solvated electrons. Thus we can consider the

following reactions [7]: 2S203

2" + hv-> S4062"+2e~

S032 ' + S203

2" + hv ->• S3062" + 2e "

At last CdS molecules are formed by the following reaction: Cd2++S + 2e" ->CdS

The S/Cd ratio in the deposited film is related to the pH of solution. We measured S/Cd ratio using peak hight of S and Cd from AES data and considering their sensitivity factors. The results showed that the best value is pH > 5. The samples were also placed at the bottom of the solution's container for few days to form thin layer CdS films by sedimenting CdS particles which were already formed by irradition of solution.

3. Results and Discussions

3.1. Thin Film Structure

The XRD pattern of as-deposited CdS thin films formed on both glass and ITO substrates showed amorphous structure, whereas the annealed thin films at 350C in the nitrogen environment have both cubic and hexagonal structure. When the samples are annealed at 500°C, their structure turns to hexagonal. Fig.l shows the XRD spectra of annealed films that were formed on glass substrate. It shows different peaks related to (100), (002), (101), (110), diffraction lines of hexagonal structure.

The surface topography observed using SEM experiments. As it is shown in Fig. 2 thin films CdS exhibit a granular structure. The existance of a few cracks on the thin film deposited by sedimentary technique shows stress and less adhesion to the substrate. In fact films on ITO substrate showed better quality and" adhesion than glass substrate[10].

339

0 20

Figure 1. The XRD spectra of annealed films at 500°C on glass substrate.

,8L,

80

Figure 2. The SEM photograph shows granular structure of the CdS films, left: thin film deposited by sedimentary technique, right: the thin film deposited by photochemical technique.

32* Optical Properties

The transmitance of CdS thin films formed by PCD and sedimentary technique shows that their band edge is around 400-500 nm and more transmittance is obtainable at higher wavelengths. Figure 3 shows the optical transmittance spectra of solution and three different samples made by sedimentary technique. Graph a is from CdS solution which is irradiated by UV light for 30minutes, whereas the graphs b9c and d belong to the samples formed from solutions which are irradiated for 205 10 and 5 minutes, respectively. These samples

340

remained in the solution for 12 days after specified irradiation time. As it can be seen , their band edge is related to the irradiating time and they show a red shift for longer irradiation time.

430 600 800 1000

u)3velength(nm)

Figure 3. The optical transmitance spectra of the CdS solution (a) and three different samples (b, c and d ) which are irradiated by UV light for 30,20, 10 and 5 minutes, respectively.

The absorption coefficient a of the films is determined from the following formula a = ln(l/T)/df where df is the thickness of the film. CdS is a direct band gap material and for a direct allowed transition, the absorption coefficient is related to the band gap (Eg) by the relation:

a sA(hv-E g ) ' / '

The optical band gap can be obtained by extrapolating the linear portion of the plot (ahu)2 versus hu to a=0. In here A is a constant and a function of the effective mass. Figure 4 shows the curve for sample d as was shown in Fig.3 . The magnitude of Eg of this sample is 2.65 eV that is higher than 2.45 eV for the bulk CdS.

The relation of Eg versus the irradiating time for sedimentary technique and deposition time for PCD are shown in Fig.5 . It can be seen that the energy gap of CdS thin films or particles in the solution is reduced to the bulk value for long time irradiation. We suggest that this is due to the particle size . This means that as the irradiation time is increased the grain size of CdS particles is also increased, that is why the Eg approaches to that of bulk.

341

3.2 E(ev)

Figure 4. The magnitude of Eg (2.65 eV) for the sample d as was shown in Fig.3.

2J66 •

2JS4 •

"S1 2J62 • •2,

iff 2J6

258

256

254

1 } I

I

I

2.9

2.8

9- 27 -•U

u7 2.6 -

2.5

2.4

i i

10 lime(hour)

• CdS in liquid

• deposited on gtass

A deposited on ITO

i i i i

10 20 30

12

40

Time (min.) Figure 5. The relation of Eg versus the irradiating time for PCD (above) and sedimentary technique (below)

342

The quantum confinement effect can be qualitatively explained using the effective mass approximation. For a spherical particle with radius d, the effective band gap E(d) is given by [12]:

E(d)=2.43+(2.446/d 2)+(0.3031/d) where d is in nm. Fig. 6 shows the variation of the size of CdS particles as a function of time which is related to the band gap. This figure also shows that the size of particles of CdS thin films deposited on glass is bigger than that of ITO/glass at the same time.

30

25

~ 20 E & a 15

10

• CdS h liquid

• dep osted on glass

A deposted on n"0

.

• A •

m A •

» A

•

A

*

0 5 10Time(min.)15 Figure 6. The variation of the CdS particles size as a function of time

20 25

4. Conclusions

The experimental results showed that the adhesion of cadmium sulfide thin films formed on the ITO/glass substrates were much better than those on glass, furthermore all samples which were placed near the surface solution were more adhesive than others. The SEM experiment showed the granular structure of deposited films and the XRD spectra showed different structures depends on the annealing temperature. The band gap and size calculations confirmed that both of them depend on the irradiation time.

Acknowledgment

The authors wish to thank to the Research Council of Sharif University of Technology for financial support.

343

References

1. M.Tsuj i, et al., J. of Crystal Growth 214/215, 1142 (2000). 2. P. Nemec, et al., Phys. Sol. (b) 224 No. 2,481 (2001). 3. S. Kumar, et al., J. of Phys. and Chemistry of Solids 61, 1809 (2000). 4. D. Hariskos, et al., Sol. (a) 173, 253 (1999). 5. F. Goto, et al., Thin Solid Films 387, 179 (2001). 6. Z. Yu, et al., Colloids and Surfaces A 181, 145 (2001). 7. U. Winkler, et al., Phys. Stat. Sol. (a) 173,253 (1999). 8. R. Kumaresan, et al, Jpn. J. Appl. Phys. 40, 3161 (2001). 9. M. Pattabi, J. Uchil, Solar Energy Matreials & Solar Cells 63, 309 (2000). 10. A. Martel, et al., Phys. Stat. Sol. (b) 220, 261 (2000). 11. G. Sasikala, et al., Thin Solid Films Ml, 71 (1997). 12. T.R. Ravindran, et al., Nanostructured Materials 11, 603 (1999).

ENHANCEMENT IN PHYSICAL PROPERTIES OF ZnO TRANSPARENT CONDUCTING COATING BY Al

INCORPORATION

B.N. PAWAR

Bharati Vidyapeeth Deemed University, Y.M, College, Pune 411 038, India E-mail: bnp62@yahoo.com

S.R. JADKAR

LPICM, Ecole Poly technique, Palaiseau Cedex 91128, Paris, France

K.C. MOHITE, M.G. TAKWALE

School of Energy Studies, University of Pune, Pune 411 007, India

Over the past few decades, Zinc Oxide (ZnO) has been emerged as transparent conducting oxide (TCO) materials for photovoltaic applications. The opto-electronic properties of ZnO can be improved by doping some suitable impurity ions their cations sites. We found that the aluminum (Al) is a suitable dopant in ZnO lattice to form ZnO:Al films. We report an enhancement in the physical properties of textured Al doped ZnO films deposited on glass substrate by Chemical Spray Pyrolysis technique. Al concentration in the starting solution was varied from 0 to 5 at. % at the step of 1 at. %. The films were characterized for structural properties by X-ray diffraction and electrical properties by Four Probe method and Hall measurements. The spectral transmittances were recorded with double beam spectrophotometer in UV-VIS-NIR regions. The XRD analysis reveals that the films are polycrystalline over the entire range of Al concentrations studied and clearly show the incorporation of Al in the films. The sheet resistance and average optical transmittance were found to be - 69 fi/D and 85 % respectively in the visible region at an optimized Al concentration. The competent values of the resistivity, carrier density and mobility are obtained at 2 at. % Al concentration.

1. Introduction

Badekar was the first who prepared transparent conducting oxide (TCO)

semiconductor thin films of Cadmium by thermal oxidation in 1907. Since then

due to various applications there has been tremendous interest grown in TCO

semiconductor films. The importance of TCO films is based on combination of

their unique and impressive properties such as low electrical resistivity along

with good optical transmittance in visible region and high infrared reflectivity.

Due to these properties, TCO films have made them an attractive material for

applications in opto-electronic, mechanical and architectural systems [1]. Over

the past few decades, several TCO materials have been developed including

binary and ternary compounds. Among these, thin films of Zinc Oxide highly

344

345

favored because its synthesis is economically viable and do not show any chemical and mechanical aging effect. Furthermore, these films can be grown on glass and polymer substrates. It has been shown that ZnO films have good chemical stability against the hydrogen plasma [2]. The environmental stability and physical properties ZnO film can be further improved by doping with trivalent impurity [3]. Thus, doped ZnO films can be a good candidate for substituting commonly used Indium doped Tin Oxide or Fluorine doped Tin Oxide films for a-Si:H based solar cells [4].

Numerous techniques have been reported for the synthesis of ZnO films on glass substrate [1-6]. Among these, the most preferred technique is the non-vacuum chemical spray Pyrolysis technique because of its extreme simplicity having low capital cost, easy to incorporate the impurities and possibility of large area commercial deposition. In the present paper, an attempt has been made to investigate the effect of Al incorporation on physical properties of ZnO films prepared by spray Pyrolysis technique.

2. Experimental

All films were deposited from a 0.5 M solution of zinc acetate onto chemically and ultrasonically cleaned glass substrates by spray paralyzing process, which involves the thermal decomposition of aqua-alcoholic solution of metallic inorganic compounds at high substrate temperature (350-500° C). The zinc acetate was dissolved in a mixture of methanol to distilled water in the ratio 3:1. A few drops of acetic acid were added to avoid the precipitation of the solution. Doping of Aluminum was achieved by adding solution of aluminum chloride to the starting solution. All films were deposited with fixed substrate temperature (Ts) and fixed airflow rate. The substrate temperature was held constant at 450 ± 5° C using a Cromel-Alumel thermocouple located under the substrate and electronically temperature controller unit. The compressed air was used as a carrier gas. The other deposition parameters are listed in table 1.

Table 1: Deposition conditions for ZnO: Al films

Parameter Value

Solvent volume ratio 1:3 Substrate temperature 450° C Air flow rate 7LPM Concentration of zinc acetate 0.5M Al doping concentration 0-5 at. % Sprayed quantity 20 ml.

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The sheet resistance and hence resistivity are measured using the conventional four-point probe method. The XRD patterns of films were recorded under same physical conditions using the CuKa line with X = 1.5405 A (model Rigaku RU200B). The carrier concentration and hall mobility were determined by Hall voltage. For determination of Hall voltage, the sample was kept symmetrically between the magnetic poles and non-ohmic contacts are made with four conductive aluminum strips deposited by thermal evaporation of aluminum in vacuum not less than 10"5 torr. The spectral transmittance was recorded using double beam spectrophotometer in the range of 300 to 1500 run.

3. Results and Discussion

3.1. Sheet Resistance and Transmittance

Figure 1 show the variation of sheet resistance, average transmittance and figure of merit of Al doped ZnO films as a function of Al doping concentration. As seen from fig. 1(a), initially the sheet resistance decreases with increase in Al doping concentration. The initial decrease in sheet resistance with increase in Al doping can attributed to substitution of Al ion at Zn sites, which supply charge carriers to the conduction band. This indicates that the sheet resistance depends on Al content in the film. Similar dependence of sheet resistance on Al content has been reported by Aktaruzzaman et al. [8]. It is also observed that the average value of the visible transmittance remains constant up to the doping level 2 at. %. Further increase in Al doping level decreases the transmittance. The deterioration of the film properties at higher doping level can attributed to the insufficient oxidation of sprayed material which may cause the poor textured surface thus increase in the sheet resistance and decrease in the visible transmittance.

The figure of merit is the measure of performance of transparent conducting films and defined as the ratio of visible transmittance to the sheet resistance [9]. Figure 1(b) show the variation of figure of merit as a function of Al doping level. The 2 % of Al doping level was found to be optimized Al doping level with sheet resistance 69 Q/D at optical transparency >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

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348

o o S? o oo o

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Figure 2. X-ray diffraction spectra of Al doped ZnO films deposited at various Al doping level by spray Pyrolysis technique.

349

aluminum doping. The X-ray spectra of undoped ZnO film show (100) preferred growth orientation. Other weak orientations (101), (110) and (201) are also observed. Doping below 2 at. %, the intensity of preferred orientation (100) attenuated while (002) orientation becomes prominent which indicate the preferred growth orientation of the films with their crystallites oriented with C-axis normal to the substrate. Increasing the Al doping concentration above 2 at. %, the intensity of peaks at (100), (101) and (110) increases while it decreases

for (002) peak. The growth mechanism of the films can be better understood by comparing

the measured standard deviation (ag) of ZnO film for all TC (hkl) values with ASTM data.

The effective change in standard deviation as a function of Al doping level of the films is shown in Figure 3. As seen from the figure the introduction of small amount of Al in the starting solution increases the standard deviation. However, further increase in Al doping concentration the standard deviation begins to decrease and attain the saturated value. The improvement in standard deviation of doped ZnO films can be governed by preferred oriented overgrowth along with the nucleation during the film growth. The saturated value of standard deviation may due to the preferred nucleation in the initial stage of deposition and the improvement in preferred growth. The nucleation during the film growth affects the standard deviation values [10].

3.0-

Doping level (at. %)

Figure 3. Variation of standard deviation as a function of Al doping level for ZnO films.

350

3.3. Electrical Properties

The dependence of resistivity (p), mobility (u) and carrier concentration (n), of Al doped ZnO films as a function of Al concentration is shown in Figure 4. As seen from the Figure, the optimum doping efficiency is achieved at 2 at. % of Al doping concentration. With increase in Al content in the starting solution to the optimal level, the resistivity of the film decreases to minimal values. However, a further increase in doping concentration result an in increase in resistivity of the films.

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.

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Doping concentration (at. %) Figure 4. Dependence of electrical properties of ZnO:Al films on Al doping level in the precursor solution.

351

The lowest resistivity (4.02x10"3 fi-cm) was obtained for 2 at. % Al doping level as compared to undoped ZnO film (5.52xl0"2 Q-cm). The lowest resistivity can be resulted from the higher carrier concentration and carrier mobility (see fig. 4 (a)). The increase in resistivity at high Al doping level can be due to the electrical inactivity of Al from the extra aluminum atom in the film or due to deficient of oxygen vacancies leading to the non-stoichiometric compositions [5].

The ionized impurity scattering and grain boundary scattering may govern the behavior of carrier mobility of such polycrystalline semiconductor films. The initial increase in mobility may be due to ordered growth in the crystallinity and textured growth of the film. The decreasing behavior of mobility at higher Al doping level can be attributed to the three factors viz; i) The Al atoms occupy interstitial position and act as ionized impurity

scattering centers, ii) Existence of (101) and (110) planes that possess high trap density [11]. iii) Segregation of Al atom at the grain boundaries in the form of non-

conductive oxide, from extra Al atom [12]. Consequently, the electrical resistivity increase, which decreases the charge carrier mobility [13]. For Al doped ZnO films, the carrier concentration shows a strong

dependence on Al doping concentration. As seen in fig. 4 (b), carrier concentration increases sharply from ~ 1.02xl019 cm'3 to ~ 6.92X1019 cm'3 with increase in Al doping level from 0 at. % to 2 at. %. A further increase in Al doping resulted a sharp decrease in carrier concentration to ~ 7.49 X1017 cm"3. The behavior of carrier concentration with Al doping can be explain as: Al is a n-type dopant. It replaces Zn atoms within the host ZnO lattice during the crystallization process. The substitution of Al+3 ion with a Zn+2 ion donates one additional free electron than it require for bonding to the conduction band. The process of donation of electron will continue till Al atoms get hold of the Zn sites. However, after certain level of doping, Al atoms do not contribute any free electron to the conduction band but also exist as impurity scattering centers. This increases the electrical inactivity of Al atoms in the film [14], which decreases the carrier concentration at higher Al doping level.

3.4. Optical Properties

Figure 5 show the typical optical transmittance spectra of ZnO:Al film deposited by spray Pyrolysis on 1.2 mm thick glass substrate measured in the range 300 to 1500 nm.

The transmission spectrum of undoped ZnO film is added for the comparison. The transparent regions of the films are at visible and near the infrared. The transparency in these regions depends on wide optical band gap. The film

352

showed high transmittance (> 85 %) in the visible range of the solar spectrum. The absorption edge falls in the ultraviolet region and shift towards shorter wavelength. This effect can be attributed to the increase in carrier concentration (see figure 4 (b)) [1]. The shift of absorption edges can also be related to the increase in filling up of the lower energy levels in the conduction band by the electrons released from the Al atoms. The oscillations on the spectrum are due to the interference phenomena from the uniform film thickness. A similar feature has been reported for vacuum deposited ZnO:Al films on the glass substrate [6]. A linear trend of absorption coefficient indicate occurrence of direct allowed transition in the film [15]. When the absorption curve extrapolated on energy axis (a2 = 0), gives the direct optical band gap energy (Eg) of ZnO:Al films. The band gap of undoped and ZnO:Al films are found to be 3.221 eV and 3.295 eV respectively. The observed shift in band gap may be due to increase in carrier concentration of the films [16].

100

80-

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| 40 '£ CO

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0 at. % 2 at. %

200 400 600 800 1000 1200 1400 1600

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Figure 5. Spectral dependence transmittance of undoped and doped ZnO films.

353

Ts

10-

15-

12-

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—•— 0 at. % -—°— 2 at. % J

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2.6 2.8 3.8 3.0 3.2 3.4 3.6

Photon energy (hv) Figure 6. Variation of absorption coefficient versus photon energy of undoped and doped ZnO films.

4. Conclusions

Optically transparent and conductive ZnO:Al films were successfully deposited by an in-expensive non-vacuum technique of chemical spraying. Metal organic solution of Zinc acetate was used as source compound. The formation of film was confirmed by X-ray diffraction. The X-ray diffraction analysis indicates that the growth mechanism of films predominantly controlled by aluminum doping level in the starting solution. We found that the electrical properties are sensitive to the Al incorporation in the films. The sharp decrease in the transmission in the ultra-violet region related to the increase in carrier concentration. We also found that aluminum incorporation in the film enhances the physical properties as compatible to those of undoped ZnO film. The optimized values of various opto-electronic quantities are listed in table 2.

Table 2. Opto-electronic values of as deposited films.

Opto-electronic Properties Undoped Doped

Resistivity Average Transmittance Carrier Concentration Mobility Band gap

5.52xl(Tfi-cm -75 % 1.02xl02Ocm-3

11 cm2/ V.s 3.221 eV

4.02x1 (T Q-cn > 85% 6.92xlOl9cm-3

22.5 cm2/ V.s 3.295 eV

354

Acknowledgements

The author is thankful to Honorable Dr. Patangrao Kadam, Chancellor, Bharati Vidyapeeth Deemed University for assisting the financial support to present the paper in the workshop.

References

1. K.L. Chopra, S.M. Major, D.K. Pandya, Thin Solid Films 102, 1(1983). 2. S. Major, S. Kumar, M. Bhatnagar, K.L. Chopra, Appl. Phys. Lett. 49, 394

(1986). 3. Sukuzi, T. Matsushita, Y. Sakamoti, N. Wada, Jpn. J. Appl. Phys. 135, 5457

(1996). 4. L. Bhahadur, M. Hamadani, J. Cornnig, P. Chartier, Solar Energy Materials

14, 3023 (1989). 5. J. Hu, G. Gordon, J Appl. Phys. 71, 880 (1992). 6. A Suzuki, T Matsushita, N Wada, Y Sakamoto, M. Okuda, Jpn. J. Appl.

Phys. 235, L56 (1996). 7. M. J. Alam, D.C. Cameron, J. Vac. Sci. Technol. A(19)4, 1642 (2001). 8. A.F. Aktaruzzaman, G.L. Sharma, L.K. Malhotra, Thin Solid Films 198,

67(1991). 9. G. Haacke, J. Appl. Phys. 4086, 9 (1976). 10. D.J. Goyal, CM. Agashe, B.R. Marathe, M.G. Takwale, V.G. Bhide, J. of

Mat. Sci. Lett.ll, 708 (1992). 11. J. Mck Nobbs, F.C. Gillespie, J. Phys. Chem. Solids 31, 2353 (1970). 12. D. Cossement, J.M. Streydio, J. Cryst. Growth 72, 57 (1985). 13. T. Minami, H. Sato, H. Imamoto, S. Takata, Jpn. J. Appl. Phys. 31,

257(1992). 14. D. Cossement, J.M. Streydio, J. Cryst. Growth 72, 57 (1985). 15. J.C. Manifacier, Thin Solid Films 90,297 (1982). 16. E. Burstein, Phys. Rev. 93, 632 (1952).

OPTICAL ENERGY GAP OF MAGNETICALLY CONFINED ARC DISCHARGE D.C. SPUTTERED HYDROGENATED

AMORPHOUS SILICON

M.C. ABDULRIDA

Department of Physics, College of Education (Ibn Al-Haitham) University of Baghdad, Adamyiah, Baghdad, Iraq

E-mail: moafak50@yahoo.com

H.A. HAMED, B.A. HASSAN

Department of Physics, College of Science, University of Baghdad Jaderiayh, Baghdad, Iraq

The optical energy gap of undoped hydrogenated amorphous silicon (a-Si:H) films prepared by using a magnetically confined arc discharge d.c. sputtering system has been investigated. The detail of the sputtering system used here is presented. High purity silicon polycrystalline target has been used. This work discussed experimental evidence that at different hydrogen partial pressures, during sputtering, and substrate temperature will have quite different values for the optical gap. This, results from the different values and distribution of the density of states in the mobility gap of the material. In this technique, it was found that the optimum hydrogen pressure and substrate temperature for good stability, high optical energy gap and singly hydride samples, obtained from the analysis of infrared transmission spectrum, were deposited at 8 x 10"4 mbar and 523 K respectively. This hydrogenated amorphous silicon is useful for photovoltaic technology (i.e. solar electricity...).

1. Introduction

The potential technological application of hydrogenated amorphous silicon (a-Si:H), such as photovoltaic and optoelectronic devices [1,2] have stimulated many investigation into the technique to obtain high quality material. Hydrogenated a-Si has been widely deposited by the decomposition of silane in a glow-discharge [3,4] and radio frequency (r.f.) diode and magnetron sputtering techniques [5,6]. But, recently magnetically confined arc-discharge d.c. sputtering technique to deposit a-Si:H has received an increasing attention.

As part of our study of a-Si:H thin film photovoltaic, we investigated the optical absorption of the material with the aim of characterizing film quality. This can be achieved by controlling on the main deposition parameters, such as hydrogen partial pressure, PH, and substrate temperature, Ts. Also, we reported the first stage of an investigation into the use of magnetically confined arc-discharge d.c. sputtering for preparation the a-Si:H films and devices.

355

356

2. Experimental Work

The schematic sketch of the magnetically confined arc-discharge sputtering system is shown in Fig. 1. A tungsten filament, 0.25 mm in diameter and 28 mm length and supplied with a molybdenum reflector, was used as electron source. Electrons emitted from the tungsten were accelerated through A slit, 4 x 30 mm2, inside a stainless steel T section CF flanges as a deposition chamber. The acceleration voltage, which was used to accelerate the primary electron beam and designed similar to the Nier type [7], was 50 V. The beam was collimated by a permanent magnet (140 gauss). Thus on the admission of argon gas into the chamber a magnetically confined thermoionically sustained argon arc can be maintained in the region above and wider than the target diameter. The characteristics of filament current, arc current and voltage as well as the rate of deposition will be discussed in a later paper.

Figure 1. Schematic drawing of magnetically confined arc ion source d.c. sputtering system. 1. The target 2. Mica ring 3. Copper cathode 4. Mo reflector 5. W filament 6. Slit 7. Magnet 8. Target voltage 9. S/S rod 10. S/S plate to fix the ion source 11. Al substrate holder 12. Thermocouple 13. Glass substrate 14. Reflector 16. NaCI substrate 17. NaCl holder 18. Heater terminals.

357

The Si polycrystalline target was about 6 mm thick and 40 mm in diameter, and it was bonded to a horizontal copper cathode. A circular hole in a ceramic mask exposing the target to the argon ions was used to prevent any bombardment on the cathode. The target was biased negatively with respect to the sputtering chamber to provide ion energies ranging from 100 to 400 eV. The ion current density bombarding the target was 0.7 to 2.7 mA/cm2. The substrates used in this investigation were NaCl crystal, soda glass and single crystal silicon n+ coated with Si02 (0.25 urn) for infrared transmission, optical absorption and field effect measurements, respectively. The substrates were clamped to an Al backing plate which was electrically grounded to the chamber. The target to substrate was 50 mm. The depositions were carried out at different substrate temperature (373 - 573 K). The background pressure of the vacuum chamber was better than 1.5 x 10'6 mbar. The required hydrogen partial pressure ranged between 0 and 3 x 10"3 mbar, which was mixed with argon pressure 1.1 x 10 "2

mbar, was set and the film deposited.

The amorphous nature and the hydrogenation of the films were verified by x-rays diffraction analysis and infrared transmission spectrum respectively. By using the Goodman and Fritzsche method and analysis via measuring the change in the sheet conductivity of the device channel with and without the electric field applied on the gate terminal [8], the density of localized states of a-Si:H can be determined from the field effect measurements data.

3. Results and Discussion

Infrared absorption studies reveal different species of bounded hydrogen in Si matrix. The I.R. spectra presented in Fig. 2 describes transmittance of various samples grown at different hydrogen partial pressures. Spectrum (a) shows only monohydride mode (2000 cm "'and 630 cm"') in the Si - H Stretch and Wag vibrations, respectively [9]. These films have been deposited at 8 x 10"4 mbar hydrogen partial pressure, PH, and 523 K as a substrate temperature. The other spectra (b, c and d) represent the results of films deposited at different PH. They show a broad absorption centered around 1110 and 950 cm"1 in addition to the usual Si - H modes of SiH2 and SiH3 (2100 and 850 cm "'). Such broad absorption, which is probably associated with Si - N and Si - O modes [10], is not existed in spectrum (a). In particular, the conventional sputtered a-Si:H shows that the 2000 cm"1 and 2100 cm'1 bands usually have intensities comparable to each other even for small hydrogen concentration, whereas our results show that the absorption at 2000 cm'1 dominated over that at 2100 cm'1

for samples grown by magnetron sputtering [11] and present technique. Therefore, the possible involvement of SiH2 and/or SiH3 in film formation will

358

definitely affect the quality of the material and then the performance of its devices.

100

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2000 1800 1600 U00 1200 1QOQ 800

WAVENUMBER (cm "1)

600

Figure 2. The infrared spectra of a-Si:H, deposited at different hydrogen partial pressure (TS=523K.)

An estimation of the optical band gap has been made for the magnetically confined arc-discharge d.c. sputtered samples. Fig. 3 shows the optical

absorption curves, according to Tauc's expression, (ahvy2 ochv since the optical transition in the a-Si is indirect transition [12,13], and from which one can obtain the optical energy gap, as the intercept on the energy axis. It is clear from Fig. 3(a) that the variation in PH at Ts = 523 K reflects large changes in the optical absorption spectrum. While at fixed value of PH = 8 x 10 "4 mbar and a variation in Ts; the optical absorption spectrum, Fig. 3b showed a large stability.

The optical energy gap, E0pt, obtained from Fig. 3, was plotted in Fig. 4. The results indicate that the change in PH at constant Ts = 523 K produces a wide variation in E0pt.. As PH increases, E0pt. increased from approximately 1.52 to 1.85 eV with some scatter in data (about 3%). At PH = 8 x 10 "4 mbar, E0pt. is a maximum and then it decreases sharply as PH increases to 1 x 10 "3 mbar. Once again, at PH = 8 x 10 "4 mbar the variation in the deposition temperature of samples shows slightly change in E0pt. as Ts increases from 360 to 523 K. Since previous studies on the material have shown that the photoconductivity and a-Si:H devices such as solar cells [14], Schottky diode [15], field effect transistor [16] can be degraded by hydrogen bonded configuration at 2100 cm "' (SiH2

and/or SiH3), it is of interest to analyze the infrared spectrum results depending on the density of localized states near Fermi - level, N(EF).

359

1900

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Figure 3. Variation of (ahv) 'A verses hv, indirect absorption, of a-Si:H films deposited (a) at different hydrogen partial pressures, (b) at different substrate temperatures.

360

Substrate Temperature (K)

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1.90

1.85

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Figure 4. The optical energy gap of a-Si:H as a function of PHand Ts.

We have studied the influence of variation in PH and Ts at an argon pressure of 1.1 xlO"2 mbar on the density of states near Fermi-level. It can be seen from Fig. 5 that N(EF) exhibits a minimum of 1.5 x 10 l6 eV"1 cm"3 at PH = 8 x 10"4

mbar (Ts = 523 K). Also when PH is increased from 8 x 10"4 to above 1.0 x 10'3

mbar N(EF) increases by nearly three order of magnitude above the minimum. In the same figure, it shows N(EF) as a function of Ts at constant PH = 8 x 10 "4

mbar. As can be seen, N(EF) is not very sensitive to the variations in Ts compared

with the variations in PH . From this figure, it can be seen that the minimum N(EF) occurs around 523 K. So, it is not surprising to find that N(EF) decreases with increasing PH up to certain value, typically 8 x 10 '4 mbar. This can be assumed to be due to the incorporation of an increasing amount of hydrogen in a-Si network structure resulting from the passivation of a large number of dangling bonds and the dominating of monohydride bonds. Possible explanations for the increase in N(EF) are as follows: firstly, new defects states are created by the hydrogen, secondly, at lower and higher PH = 8 x 10* mbar, the incorporation of impurities such as 0 2 and N2 might be enhanced as depicted in Fig. 2; thirdly, the different PH values can lead to different Si - H configuration such as (SiH2, SiH3 and (SiH2)„), which can lead to additional defects states in the mobility gap of the material.

361

350 10 20

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Figure 5. The density of localized states near Fermi-level as a function of PH and Ts.

4. Conclusions

In this paper, we have been able to design and construct a magnetically confined arc-discharge d.c. sputtering system for thin films deposition applications. This technique can be used to controllably introduce hydrogen to deposit a-Si:H films for optoelectronics devices. We have found that the optimum hydrogen partial pressure and substrate temperature are very close to 8 x 10 "4mbar and 523 K, respectively. The results of infrared spectrum, optical absorption and the density of states near Fermi-level demonstrate that the above deposition parameters are the optimum conditions for growing stable and high quality thin film a-Si:H material.

References

1. Y. Hamakawa, Amorphous Semiconductor Technologies and devices, 22, OHM, Tokyo (1987).

2. W.X. Ni, F. Priolo, D. Grutzmacher, R. Tsu, "Si-Based Optoelectronics: Advances and Future Perspectives", E-MRS Spring Meeting, (Strasbourg France, 18 to be published in Physica E. 2002).

362

3. W.E. Spear, P.G. LeComber, Phil.Mag. B 33, 935 (1976). 4. H. Fritzsche, Solar Energy Matt. 3, 447 (1980). 5. M.C. Abdulrida, J. Allison, Appl. Phys. Lett. 43, 768 (1983). 6. J.I. Pankove, "Semiconductors and Semimetals", 21, Hydrogenated

Amorphous Silicon, Part A (Academic press, New York 1984). 7. A.O. Nier, Rev.Sci.Inst. 18, 398(1947). 8. N.B. Goodman, H. Fritzsche, Phil. Mag. B 42, 149 (1980). 9. B. Buttle, J.B. Adams, Phys.Rev. B 53, 16265 (1997). 10. I.H. Lee, K.J. Chang, Phys.Rev. B 50, 18083 (1994). 11. M.C. Abdulrida, Ph.D. Thesis, (The University of Sheffield U.K. 1984). 12. J. Tauc, Amorphous and Liquid Semiconductors, Ed. By J. Taus (Plenum,

New York, 159 1974). 13. D. Bue, M. Zeman, J.W. Metselaar, 43rd International Scientific

Colloquium, (Technical University of Iimenau, Spt. 21-24 1998). 14. M. Tanaka, M. Taguchi, T. Takahama et.al, Proceeding of Photovoltaics:

Research and Applications 1, 85 (1993). 15. H.L. Fernandes-canque, M.C. Abdulrida, J. Allison, Thin Solid Films 110,

241(1983). 16. J. Robertson, Phyl. Mag. B 63, 47 (1991).

PHOTOCATALYTIC STUDY OF Ti02 THIN FILMS DEPOSITED BY DC REACTIVE MAGNETRON SPUTTERING AND SPRAY

PYROLYSIS METHODS

A.I. MARTINEZ, D. ACOSTA, A. LOPEZ

Institute de Fisica, Universidad National Autonoma de Mexico A.P. 20-364; 01000 Mexico D.F.

E-mail: martinez@eros.pquim. unam. mx

Amorphous and polycrystalline titanium dioxide thin films were deposited by DC reactive magnetron sputtering and spray pyrolysis methods on glass and glass coated with fluorine tin oxide (FTO). The films were characterized by X-ray diffraction, atomic force microscopy (AFM) and UV-visible spectroscopy. The photocatalytic activity of the samples was tested on the degradation of methylene blue. For films deposited by the sputtering technique, we have studied the effect of the total pressure of an Ar-02 gas mixture on phase composition, crystallinity and the photocatalytic properties. Also Ti02

thin films have been prepared by spray pyrolysis using a low concentration of titanium precursor at different substrates temperatures. The best photocatalytic properties in respect with methylene blue degradation were obtained for Ti02 thin films in anatase phase with an open structure. These films were prepared at high total pressure (16 mTorr) and at moderate substrate temperatures (400 °C) by sputtering and spray pyrolysis respectively.

1. Introduction

Chemical wastes from manufacturing processes have too often discharged into environment with little or without treatment; this way of discard have a potential environmental damage, the use of titanium dioxide mediated photocatalysis is an alternative for remediation of waste water and contaminated air [1].

Ti02 thin films can be prepared by a variety of methods; among these, reactive sputtering [2] can be used to prepare good quality films, but they are quite expensive when large-scale production is needed. The spray pyrolysis method is a less expensive alternative [3] for large area films production.

In this work, we compare the properties of pyrolytically sprayed and sputtered Ti02 films deposited on glass. We study the effect of some of the deposition parameters such as substrate temperature on the structure and photocatalytic activity of the sprayed Ti02 films. Also in this paper, the effect of total pressure on structure and photocatalytic properties of the Ti02 films sputtered on unheated substrate was studied. The aim of this work is to compare the photocatalytic properties of Ti02 thin films prepared by DC reactive magnetron sputtering and spray pyrolysis methods.

363

364

2. Experimental

2.1. Sputter Deposition

Ti02 thin films were prepared using the commercial DC Magnetron Sputtering system (Hummer XII from Anatech Ltd). A 5 cm diameter titanium target of 99.995% purity was used; a mixture of 99.5% pure oxygen and 99.997% pure argon, a substrate-target distance of 5 cm and a constant discharge power of 100 W. Before the deposition, the chamber was evacuated to 10"5 Torr. We introduced in the chamber then a gas mixture at different total pressures, 4, 8, 16 and 32 mTorr. We made the gas mixture by the simple method presented in figure 1: The Argon and oxygen gas pressures were regulated by opening both tanks valves with the same flux value in order to obtain a 50/50 ratio.

-x-

Ar

02+Ar To > vacuum

chamber

-x- 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)

4

8 16 32

Thickness (nm)

305

280 280 240

Phase

Rutile + Anatase Anatase Anatase

Amorphous

Grain Size (nm)

28

28 30 -

Eg on glass (eV)

3.12

3.17 3.18 3.29

Eg on FTO (eV)

3.14

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 Ti02 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 Ti02 thin films with decomposition of methylene blue (Ci6H18N3SCl-3H20). We used the same method reported by Zeman and Takabayashi [2]. Ti02 films were immersed in methylene blue solution 1 mM for 1 h and afterwards dried for 30 min in dark room. The surface of Ti02 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 Ti02 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 Ti02 films when were deposited on different substrates for a 16 mTorr gas pressure. The strongest peak corresponds to Sn02 (on FTO substrate). The spectrum in the lower part of the graph corresponds to Ti02 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 Ti02 thin films on tetragonal substrates (FTO) are more crystalline than films on amorphous substrates (glass).

366

Figure 2. XRD patterns of Ti02 thin films on glass prepared by D.C. reactive magnetron sputtering at different total pressures of an equimolar gas mixture Ar/02.

AC101)

>

on FTO substrate

AI200) | AI211) ,

on glass substrate

A(101) A(200) A(Z11)

10 20 50 60 30 40 2H

Figure 3. XRD patterns of Ti02 thin films on different substrates prepared by D.C. reactive magnetron sputtering at 16 mTorr.

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 T02 films; it is possible to calculate the weight percentage of the anatase phase, WA, using the following equation [5]:

WA= l / [ 1+1.265 I R/IA] (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 Ti02 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 Ti02 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 Ti02 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. Ti02

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

' fK0^My^ w ',£' 4 mTorr •li - 8 mTorr Ti 16 mTorr

" J J 32 mTorr

jj

(b

300 500 700 >.inm)

900 1100

Figure 4. UV-visible transmission spectra of Ti02 thin films prepared by D.C. reactive magnetron sputtering at different total pressures of an equimolar gas mixture Ar/02. On different substrates (a) on glass and (b) on glass coated with FTO.

Fig. 6 shows the as measured transmission curves for sprayed Ti02 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 Ti02 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)

Ik

/

•" I ' 1 '

(101)

Ts=450 'C

i '

(200j

\\MJ

4^4^^=350 "C

I . I .

• : 2 l i i

WwWv

i I I I I I I I I ,

10 20 30 40 50 60 2H

Figure 5. XRD patterns of Ti02 thin films on glass prepared by spray pyrolysis method at different substrate temperatures.

Table 2. The correlation of physical parameters of the Ti02 films deposited by spray pyrolysis for different substrate temperature is observed in this table.

TS°C

350 400 450

Phase

Amorphous Anatase Anatase

Thickness (nm)

260 270 285

Grain Size (nm)

-32 30

Eg(eV)

3.20 3.25 3.26

3 7 0

1

0.8

0.6

0.4

0.2

'300 500 700 900 1100

A(nm)

Figure 6. Optical transmission spectra of Ti02 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 Ti02 (-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)

02ad + 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 Ti02 [1]. Fig. 7 shows the degradation of methylene blue film formed on Ti02 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.

1 1 ' 1

1 yS

J , .

• i •

y

\ ' / \ • • - • - . . . / v - - ' ,

/""

300°C

•-—• 350 "C

— 400 "C

— 450"C

371

From Figures 7 and 8 we can see that the highest photocatalytic activity is achieved for the Ti02 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 Ti02

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 Ti02

crystals of the films has a particle size of 15-30 nm, only by spray pyrolysis, the Ti02 thin film present particle size of Ti02 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

<

d

0 10 20 30 40 Pressure (mTorr)

Figure 7. The change of absorbance A ABS of methylene blue film formed on Ti02 surface as a function of different deposition parameters, e.g. temperature for spray and pressure for sputtered deposited films.

372

0 2 4 6 8 10 lrndationnmi(h)

Figure 8. Transmittance at 650 nm of an aqueous methylene blue solution as a function of Irradiation time when the Ti02 films are immersed in the test cell.

(a) 2 p,x2 fix37 nm (b)2 jnx2 fixl33 nm Figure 9. 3D AFM micrographs (a) Ti02 thin film by DC reactive magnetron sputtering at 16 mTorr, (b) Ti02 thin film prepared by spray pyrolysis at 400 °C from a solution of titanium (IV) oxide acetyl acetonate.

4. Conclusions

In this work, we Investigated the effect of different deposition parameters on the photocatalytic activity of the Ti02 thin films prepared by D.C. reactive magnetron sputtering and spray pyrolysis methods. The highest photodegradatlon was found for Ti02 thin films in anatase phase with an open

373

structure. These films were prepared at high pressure (16 mTorr) and at moderate temperatures (400 °C) by sputtering and spray pyrolysis respectively. When were used rutile/anatase structures the photocatalytic degradation of methylene blue is very low due to the absence of surface texture.

A simple spray pyrolysis technique provides the best perspectives for the preparation of Ti02 photocatalyst due to their high deposition rates over large areas and least expensive method.

Acknowledgements

The authors are thankful to Carlos Flores, Carlos Magafia and Manuel Aguilar, for their technical assistance. We also thank the financial support of CONACYT Mexico-Project 34-821 E.

References

1. M.R. Hoffmann, S.T. Martin, W. Choi, D.W. Bahnemann, Chem. Rev. 95,69(1995).

2. P. Zeman, S. Takabayashi,, Surf. Coat. Tech. 153,93 (2002). 3. M.O. Abou-Helal, W.T. Seeber, Appl. Surf. Sci, 195, 53 (2002). 4. L Eckertova, Physics of Solid Films, Plenum-Press, NY (1977). 5. D. Mardare, M. Tasca, M. Delibas, G.I. Rusu, Appl. Surf. Sci. 156, 200

(2000). 6. L.J. Meng, M.P. Dos Santos, Thin Solid Films 226, 22 (1993). 7. H. Tang, K. Prasad, R. Sanjines, P.E. Schmid, F. Levy, J. Appl. Phys. 75(4),

2042 (1994). 8. S.G. Schrank, H.J. Jose, R.F.P.M. Moreira, J. Photochem. Photobiol. A:

Chem. 147, 71 (2002). 9. N. Daude, C. Gout, C. Jouanin, Phys. Rev. B 15, 3229 (1977). 10. M. Radecka, K. Zakrzewska, H. Czternastek, T. Stapinski, Appl. Surf. Sci.

65/66,227(1993). 11. H. Yumoto, S. Matsudo, K. Akashi, Vacuum 65, 509 (2002).

NOVEL TRANSPARENT AND HIGHLY CONDUCTIVE ZnO-BASED COATINGS

B.M. ATAEV, A.M. BAGAMADOVA, I.K. KAMILOV, V.V. MAMEDOV, A.K. OMAEV, S.SH. MAKHMUDOV

Institute of Physics, Daghestan Scientific Center of the Russian Academy of Sciences M. Yaragy, 94, Makhachkala, 367003, Russia

E-mail: crystal@dinet.ru

We report on the comparative study of thermal stability of resistance in in situ doped ZnO:M films (M = Al, In, Ga, Sn) fabricated by Chemical Vapor Deposition (CVD) and Magnetron Sputtering (MS) techniques. It is shown that substantially different concentrations of doping elements are required to produce highly conductive and transparent ZnO films in CVD system while those concentrations are similar in MS system. CVD ZnO films featured unique temperature stability of electrical resistance within the wide temperature range.

1. Introduction

Low-resistive and transparent ZnO films doped with the elements of III and IV groups could replace, in principal, more expensive indium-tin oxide coatings in various electronic devices [1-2]. Along with the good electrical and optical parameters, zinc oxide is lower toxic and relatively cheaper than its competitors. On the other hand, there are well-known such constraints of the magnetron sputtering technique as structural imperfection, non-uniform area and volume distribution of the dopant elements and instability of the electrical properties in the films obtained.

It was shown earlier that Chemical Vapor Deposition (CVD) technique allows one to produce ZnO films with perfect structure in commercial amounts [3]. We also reported recently on a successful gallium doping of growing zinc oxide thin epitaxial films in a low pressure system within a single CVD cycle [4].

We report in this paper the study of in situ doped ZnO films prepared by CVD and magnetron sputtering techniques. Metals of III-IV groups (M = Al, In, Ga, Sn) were studied to be used as dopant elements. It was also made a comparative investigation of basic electrical properties in these films.

2. Experimental Procedure

Single-crystalline ( 1 0 1 2 ) sapphire substrates 20x20 mm2 size were used for ZnO films fabrication in both CVD and magnetron systems. The low-pressure reactor description as well as temperature optimization are described elsewhere

374

375

[3]. The preparation of starting material to be placed into evaporation zone was reported in [4]. Ultra high purity ZnO, Ga203, A1203, ln203 and Sn02 powders were used. Impurity content ranged 0.01-5 wt. per cent. Hydrogen of high purity was used as the working gas. The substrates were placed sequentially, 0.5 cm apart, perpendicular to a hydrogen flow into the reactor tube.

The ZnO:M targets 2-3 mm thick and 40 mm diameter were used in dc-magnetron sputtering system. It should be noted that only continuous preliminary annealing for 10-12 hrs at the temperatures as high as HOOK allowed us to prepare low-resistive targets to secure the steady discharge during dc-magnetron sputtering in Ar:02 (4:1) ambient. Discharge current was about 100 mA, voltage was 350 V, and deposition rate was of the order lxlO'2 um-min"'.

Films perfection was monitored by X-ray (XRD) and electron diffraction. Film resistivity, charge carrier density and mobility were studied by Hall measurements. A VUP4 vacuum chamber were used to study electrical properties of the films in various ambient (argon, nitrogen, air) in the 300-950 K temperature range. The films were held at a given temperature for an hour and then gradually cooled to room temperature following which the film resistivity was measured.

3. Results and Discussion

The experiment showed that ZnO:Ga and ZnO:In films can be successfully produced by CVD method in the implemented conditions (the source and substrate temperatures, gas flow rate, partial pressures of reagents etc.) while the films prepared with ZnO:Al and ZnO:Sn sources had the same electrical and structural properties and morphology as those prepared with pure zinc oxide source material. Apparently, the doping process can not be done by CVD with Al and Sn impurities because of a low volatility of their transient compounds.

It was also revealed in experiment that substantially different concentrations of doping elements were required to produce highly conductive and transparent ZnO films. Optimal Ga203 content was found to be 1-2 wt. per cent (i.e. 0.9-1.8 atomic %) while for indium oxide it was an order lower, 0.1-0.2 wt. per cent (i.e. 0.06-0.12 atomic %). In the last case, 0.3 wt. per cent In concentration led to the blacken films with the metallic type of conductivity. With equal impurity content, the doping level in the films depended on the temperature conditions and the distance from a substrate to a source in the reactor tube. These parameters were also being optimized.

For all CVD films obtained, XRD analysis shows that the ( l l 2 0 ) ZnO plane is parallel to the substrate surface. The crystallite misorientation in the films prepared in the optimum conditions did not exceed 30'. Magnetron-sputtered films feature basis orientation in the whole substrate temperature range

376

300-650 K. XRD and electron diffraction data indicate that all CVD filmsjiave more perfect structure than magnetron films. In particular, half-width of (1 1 2 0 ) plane reflection in CVD films was 2 times lower than that of (0001) plane for the magnetron films. At the same time, the study of surface morphology revealed that the magnetron films had a smoother surface. This may be directly associated with the crystallite size, though one should take into account that surface relief of CVD films is more sensitive to imperfection of transition layer (see also [4]). All above-said about morphology is applicable both for doped and undoped films.

Electrical properties of pure and optimally doped ZnO films are presented in the Table 1. As can be clearly seen, the resistivity in the ZnO:M films (M = In, Ga) was as low as 1.2xl0"4 Qcm in the optimum CVD growth conditions. We think the M impurity affects the electric properties in the films as follows: (a) charge carrier density Nd in heavily degenerated ZnO is one to three orders greater than that in pure ZnO; (b) Hall mobility gradually decreases as impurity content increases.

Table 1. Electrical properties of pure and doped ZnO films.

Sample

Pure ZnO

CVD ZnO + Ga203 (2 wt. %)

CVD ZnO + ln203 (0.2 wt. %)

MS ZnO + Al2O3(0.2wt.%)

p, Qcm

1.5x10''

1.2xl0-4

l.lxlO"4

2.0 xlO"4

Nd, cm"3

2x10"

3xl020

8xl020

5xl020

H, cm2V"'sec"1

70

38

33

3

A higher carrier density Nj ~ 1020 cm"3 is associated with M3+ ions replacing Zn2+ ions in the crystal lattice sites. As M content increases, the formation of such centers as interstitial zinc and oxygen vacancies is believed to be improbable [see 2, 5] and, therefore, one can propose a general formula Zn,.xMxOex.

Just a different case was the films prepared by dc-magnetron sputtering. Conducting films of acceptable quality were produced in a number of sputtering systems [1, 2]. Our comparative study of the films with an equal resistivity shows that the CVD films feature larger mobility while the magnetron sputtered films have a greater charge carrier concentration. We believe this is due to more perfect structure in the CVD films leading to a lower lattice scattering of the carriers.

One problem limiting a wider field of conducting and transparent coatings applications is known to be a temperature instability of the electric properties of the films in the presence of various gases. By now, the best temperature stability

377

in a vacuum or in a gas ambient, up to 700 K, has been observed in ZnO:Al films prepared by magnetron sputtering [1]. It has been concluded that for oxygen chemisorption in the films an Al impurity donor, rather than native defect donors such as an oxygen vacancy and interstitial zinc, is stable. We paid particular attention to studying the instability of electrical parameters in ZnO:M films after heat treatment.

Figure 1 shows the resistivity data in the temperature range from room temperature to 950 K. As can be seen, the resistivity in CVD ZnO:M films is much more stable thermally than in the films prepared by dc-magnetron sputtering. Resistivity curve behavior is the same both for Ga and In impurity. Moreover, it should be noted that undoped CVD films behave the same way taking into account a starting resistivity value.

10 4-.

10 3-

R,Q

io2-J

101-

300 400 500 600 T, K

700 — i —

800 900

Figure 1. ZnO:M resistance after heat treatment in air. CVD films: (l)ZnO:Ga(2%), (2) ZnO:In (0.2%); magnetron sputtered films: (3) ZnO:Sn (2%); (4) ZnO:AI (2%).

378

Polycrystalline textured (0001) ZnO films produced by magnetron sputtering feature relatively large volume of intergranular defect regions where the metal impurity atoms exist preliminary in interstitial sites. With the same impurity concentration, relative density of lattice M3+ ions in the CVD films is obviously higher than that in magnetron sputtered films. When ZnO crystals being annealed in the oxygen, interstitial M defects can not be considered as the main compensating defects because the process of defect formation follows «quasiepitaxial» scheme [5, 6] and foremost interstitial M atoms are «extracted» out off the bulk to be oxidized at the surface. Thus, we suggest that the crystal structure perfection must be among the main factors conditioning the thermal stability of the conducting and transparent films.

ZnO is known to be a nonstoichiometric compound. The physical state of the ZnO surface and electrical properties over it are significantly affected by the oxygen content. The physical nature of variations in oxygen-surface binding energy is determined by a dipole interaction between the «donor impurity» (including intrinsic one, i.e. excessive zinc) - «oxygen admolecules» complexes. The surface density of the oxygen admolecules increases with the donor impurity increase. It is also known that only O2' ions are stable within the 300-480 K temperature range, while at higher temperatures O" ions prevailed [5]. It worth noticing that substantial decrease in oxygen-surface binding energy with the reversible behavior in the 353 K temperature region was observed in indium doped semiconductors with surface density as high as 10l7-10l9cm"2 [7]. This was explained by a formation of metal cluster structures - whiskers over whole surface which were stimulated by the presence of screw dislocations and a stressed state of a surface. The fact that in our study electron effective density and mobility are not dependent on impurity concentration N at the values higher than 1020 cm"3 testifies to a possibility of indium/gallium whiskers formation onto a ZnO surface. It should be noted that whiskers are very stable to oxidation even at high temperatures. For example, the oxidation process of iron whiskers is 50 times longer than that of polycrystalline iron [7]. The last reason is, to our opinion, very important in understanding the unique thermal stability of heavily doped ZnO films prepared by CVD method.

4. Conclusions

Thus, ZnO:M films were prepared by CVD (M = In, Ga) and magnetron sputtering (M = Al, In, Ga, Sn) techniques. The effect of thermal cyclings on electrical properties in the films was investigated. The doped ZnO films prepared in low pressure CVD system feature unique thermal stability of their electric properties within the 300-950 K temperature range in various ambient. The

379

crystallinity of the films is found to be among the main factors conditioning the thermal stability of the films.

Acknowledgements

This work was supported by the Russian Foundation for Basic Research, Grants No. 02-02-17627 and 01 -02-16200.

References

1. T. Minami, H. Nanto, Sh. Takata, Jpn. J. Appl. Phys. 23, L280 (1984). 2. Y. Igasaki, H. Saito, J. Appl. Phys. 70, 3613 (1991). 3. A.Kh. Abduev, B.M. Ataev, A.M. Bagamadova, G.A. Krasulin,

Neorganicheskie Materialy (in Russian) 11, 1928 (1987). 4. B.M. Ataev, A.M. Bagamadova, A.M. Djabrailov, V.V. Mamedov, R.A.

Rabadanov, Tech. Phys. Lett. 21, 129 (1995). 5. A.N. Georgobiani, M.B. Kotlyarevskii, V.N. Mikhalenko, Trudy FIAN (in

Russian) 138,79 (1983). 6. T.V. Butkhuzi, A.N. Georgobiani, E. Zada-uly, B.T. Eltazarov, T.G.

Khurdolava, Trudy FIAN (in Russian) 182, 140 (1987). 7. Yu. A. Vashpanov, Tech. Phys. Lett. 23, 13 (1997).

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, S.SH. MAKHMUDOV

Institute of Physics, Daghestan Scientific Center of the Russian Academy of Sciences M. Yaragy, 94, Makhachkala, 367003, Russia

E-mail: crystal@dinet.ru

It is firstly reported on a low-temperature modification of the CVD growth system with RF-discharge applied during the growth process. This method allowed us to significantly increase the effective pressure of the atomic oxygen during the deposition, to lower the substrate temperature more than 200 °C, and thus to enhance the stoichiometry and crystal perfection of the growing layers.

1. Introduction

Because of its high electron conduction and optical transparency, zinc oxide is widely used nowadays in various optoelectronic applications. The recent investigations conducted to achieve p-type conductivity in ZnO by adding the acceptor impurities (N, P, As) allow one to consider this material as most promising one among other wide band gap A2B6 and A3B5 semiconductors. However, the reproducible fabrication of the blue and near-UV light-emitting diodes and lasers based on ZnO epitaxial layers (EL) is restricted by the existing fabrication techniques. Chemical vapor deposition (CVD) technique is known to be one of the widely used methods of ZnO EL fabrication on crystalline substrates. It is also known such disadvantage of this technique as relatively higher temperatures in the substrate area of the CVD reactor during the deposition, 580-680 °C, which stimulate the chemical side-reactions in the gaseous phase resulting in uncontrolled contamination of ZnO films.

For the first time, we report in this paper on a low-temperature modification of the ZnO EL CVD growth on the cc-sapphire substrates with the RF-discharge plasma generated in CVD reactor during the growth process. It is known that the oxygen deficiency enlarges a number of the donor defects such as interstitial zinc and oxygen vacancies in ZnO. Our method allows us to increase the effective pressure of the atomic oxygen during the deposition by more than 6 orders (see [1]), and to shift the stoichiometry of the growing film to oxygen excess simultaneously with lowering the concentration of donor-type intrinsic defects. This approach was used earlier only for thermal annealing of bulk single crystals and EL of oxides in order to improve their stoichiometry [2].

380

381

2. Experimental

The low-pressure CVD reactor described in [3] was modified by applying the 50 W RF discharge along the both low- and high-temperature zones in the CVD system (Fig. 1). The pressure of the working gas (H2) was as low as 1 Torr and lower, and reagents flow rate being 1-3 m sec"1. It should be noted that with RF activation we could lower the substrate temperature more than 200 °C, down to 420 °C, within the epitaxial growth cycle (Fig. 1 curve 2), while the EL growth rate was lowered by four times and thus allowing us to significantly improve the initial stage of the EL nucleation and growth. The obtained ZnO/a-sapphire films 1-2 urn thick featured high crystallinity, perfect surface morphology and good electrical and optical properties at the expense of both the decrease in the width of the transition layer and amount of the uncontrolled impurities.

5 10 15 20 25 L.CM

Figure 1. RF-modified low pressure CVD system, (a) quartz reactor, (b) temperature profile and (c) growth rate. TCU = temperature control unit; M = manometer; GR = growth rate ((im/min).

382

3. Results and Discussion

Figure 2 shows the characteristic micrographs of (1120)ZnO/(101 2)A1203

surfaces obtained with LEICA DM LP microscope at x600 magnification. One can see that the film deposited with RF activation features less developed surface morphology, and the growth pattern of the film is more than one order less than that prepared without activation. Also the size and number of Zn clusters which usually complete growth figures is sufficiently lower. As a result, we have smoother and mirror-like surface of the film.

Figure 2. Surface patterns of the ( 1 1 2 0 )ZnO/ ( 1 0 1 2 )A1203 films: (1) CVD grown and (2) RF-assisted CVD grown.

Figure 3 represents X-ray diffraction patterns obtained with the DRON diffractometer utilizing CuKa radiation monochromized with pyrolitic graphite. The diffractogram corresponds to a perfect (J 1 2 0 ) ZnO EL. Inset in Fig. 3 shows the rocking curve patterns of the (1120)ZnO peak for both (1) only thermally grown film and (2) the film deposited with RF activation. In the last case, the full width of half maximum (FWHM) of the diffraction peak is two times less, reaching 0.5 degree in the optimum growth conditions. We believe, this value can be reduced by several times when applying a buffer layer technology that can engage the autoepitaxy mechanism.

It should be particularly noted that another distinctive feature of our modified CVD technique is a possibility to activate in situ doping of ZnO films in view to achieve the p-type conduction in this material as well. The preliminary results of RF-stimulated CVD ZnO films growth in 2H2+1N2

ambient show that p-ZnO layers with the resistivity as high as 103 Qcm can be produced on the a-sapphire substrates. We believe, it is possible to significantly lower the resistivity of these films.

383

O <

-1.0 , 0.0 1.0 degree

co O

CM

o CM

O c N

if CM CM

u 20 30 40~ 50 So"

20 , degrees

— ^

70

Figure 3. Diffractogram of the ( 1 1 2 0 )ZnO/ ( 1 0 1 2 )A1203 EL grown by modified RF-assisted

CVD technique. Inset shows rocking curve patterns of the ( 1 1 2 0 ) ZnO peak for films: (1) CVD grown; (2) RF-assisted CVD grown.

RF-stimulated ZnO deposition onto GaAs, GaP, InP substrates into low-pressure CVD system is under our current investigation. In this case, we try to compare the two existing approaches to interpretation of electrical and optical properties in these structures. For example, it was stated in [4] that the ZnO layers were doped in situ by the volatile components of the substrate (e.g. As, P) that stimulated the inversion of ZnO conduction type. In earlier works (e.g. see [5]), the same results on these structures (i-v curves and electroluminescence spectra) were interpreted as Zn, O co-doping of the substrate material with fabrication of p-n junction in GaP substrate. Our experimental results confirm the last point of view.

384

4. Conclusions

Thus^ for the first time, we have successfully fabricated (1120)ZnO layers on ( 1 0 1 2 )A1203 substrate in low-pressure CVD system with RF-stimulation of the growth process, that allowed us to significantly lower the substrate temperature and to improve the initial stage of the ZnO film nucleation and growth. The obtained films features high crystallinity and perfect surface morphology at the expense of both the decrease in the width of the transition layer and amount of the uncontrolled impurities.

Acknowledgements

This work was supported by the Russian Foundation for Basic Research, grants No 01-02-16200 and 02-02-17627.

References

1. A.N. Georgobiani, A.N. Gruzintsev, U.A. Aminov, et al., Fizika i tehnika poluprovodnikov (In Russian) 35, 149 (2001).

2. T.V. Butkhuzi, A.V. Bureyev, A.N. Georgobiani, et al., J. Cryst. Growth 117,366(1992).

3. A.Kh. Abduev, B.M. Ataev, A.M. Bagamadova, et al., Neorganicheskie materialy 23, 1928(1987).

4. Y.R. Ryu, W.J. Kim, H.W. White, J. Cryst. Growth 219, 419 (2000). 5. R.A. Rabadanov, S.A. Semiletov, M.K. Guseikhanov, Kristallografiya 26,

645(1981).

VII. SUPERCONDUCTIVITY

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PHYSICS AND APPLICATIONS OF YBazCi^Oy/LaovSroaMnOj HETEROSTRUCTURES

J.G. LIN

Center for Condensed Matter Sciences/Nanostorage Research Center National Taiwan University, Taipei, Taiwan 10617

E-mail: jglin@ccms. ntu. edu. tw

In this article, we have discussed some physics and the possible applications of YBa2Cu307/Lao.7Sr,uMn03 (YBCO/LSMO) heterostructures based on our recent work. We first described the electroresistance (ER) and magnetoresistance (MR) of LSMO thin films with different thicknesses (t), in order to demonstrate the strain effect on these single layered films. Then the transport properties of YBCO/LSMO heterostructures with different thickness of LSMO is discussed. Based on the data of t-dependent resistivity, magnetoresistance and critical current, the proximity and spin-injection effects play a major role on suppressing the superconductivity. Correspondingly, the future applications of the spin-injector may be controlled by fine-tuning the thickness of LSMO.

1. Introduction

Phenomenon of Colossal magnetoresistance (CMR) in perovskite compounds La).xAxMn03 [1,2] (A= Ca, Sr, Ba, Pb) has attracted considerable attention in recent years due to the related novel physics [3,4] and potential applications [5]. A large number of experiments on polycrystals, single crystals and thin films have been carried out to explore the dependence of the magnetoresistance (MR) on the temperature (T), magnetic field (H), composition as well as on synthesis process. Indisputably, making them into the form of thin film is mostly required for the application of magnetronic devices. The basic behavior of CMR generally the same in both bulk and thin film samples, except for some properties in association with the strain effect induced by the lattice mismatch between film and substrates [6,7]. One of the serious problems for practical application of CMR materials is the insufficient magnetoresistive response at the room temperature (RT) under low field (H < lKOe). Fortunately, some studies showed that the manipulation of resistive states in CMR manganites can be achieved not only by a magnetic field but also by an electric (E) field. It has been reported that electric current (I) could trigger the transformation of the electrically insulating charge-ordered (CO) state to a ferromagnetic (FM) metallic state [8,9]. Furthermore, a correlation [10] between electroresistance (ER) and MR has been established in Lao^Cao|8Mn03 (LCMO) single crystal, and the function of an electric current of 0.3 mA was shown equivalent to 1.5 tesla at temperature below Curie temperature (Tc) and 0.4 tesla at RT. A rough

387

388

estimation [10] shows that a 1 nm wide filamentary path biased with 1 mA could produce a magnetic field of 1 tesla. Since the ER effect is strongly correlated with the MR effect and the MR ratio is very much dependent on the film thickness (t) [11,12], the simultaneous investigation of ER and MR effects with different thicknesses is essential for not only the basic research but also technological applications.

Superconductor/ferromagnet (SC/FM) heterostructures were fabricated early in the eighties of the twentieth century [13]. As an SC layer was grown on top of an FM layer, the proximity effect governs the physical properties of the SC/FM structure [14,15]. Previous studies [16-19] showed that a physical condition for the coexistence of ferromagnetism and superconductivity within one compound requires comparable ferromagnetic exchange energy Eex and superconducting energy gap. However, this condition could not be realized in bulk FM or SC, where Eex is typically at least 2 orders of magnitude larger than A. It was suggested later that this condition could be slightly adjusted if Cooper pairs were injected from SC to FM in a FM/SC structure. The density of states with superconducting characteristic has been observed in the FM region near the PdNi/Nb interface [20], implying a coexistence of weak ferromagnetism and s-wave superconductivity [21,22]. However, the pair-breaking effect due to the injection of spin-polarized electron from FM into SC would be enhanced with increasing the spin polarization of conduction electrons in FM and it was seldom discussed. Recently, the study on the cuprate/manganite heterostructures has received much attention since the discoveries of high temperature superconductivity in Y-Ba-Cu-O [23] and colossal-magnetoresistance (CMR) effect in La-Ca-Mn-0 [24]. An advantage of making such heterostructure is that both materials have the same perovskite-related structure. Particularly, the CMR manganite has been shown to be a half metal with its spin polarization very close to 100% [25], which makes it highly suitable for studying whether or not the superconductivity can coexist with ferromagnetism in the high transition temperature (Tc) superconductors. Some reports have already demonstrated the effect of spin injection in cuprate/manganite heterostructures, [26,27] and emphasized the ultimate usefulness of them in spintronic devices. However, the issues of heating and quasiparticle-injection are always the major problems on current-voltage (I-V) measurements [28]. In the work of Goldman et al. [29], a bilayer La2/3Bai/3Mn03/DyBa2Cu307 heterostructure was used, and the area of La2/3Bai/3Mn03 layer is reduced to limit the heating problem. Another work concerning the effects of spin injection on the mixed state of YBCO revealed that the heating effect occurred after the current being applied for tens of ms [30]. We, thus, try a new approach to investigate the spin-injection effect on YBa2Cu307(YBCO)/Lao.3Sro.7Mn03(LSMO) heterostructure.

389

Another important but rarely discussed issue, in association with cuprate/manganite heterostructures, is the behavior of its normal state resistivity ()). Prior to this work, the only possibility to extend the normal-state resistivity of the high-Tc superconductors well below Tc was by suppressing the superconducting phase either with strong magnetic fields [31-34] or high current desities [35]. Different experimental results for the low temperature (T) normal state of high-Tc materials are summarized as follows: (i) Bi2Si2CuOy remains metallic down to the lowest temperature once superconductivity was suppressed with magnetic field [31], (ii) High field measurements of 61 tesla for under-doped and optimal-doped La2.xSrxCu04 [32] revealed a logarithmic divergence of log(l/T) for p(T), (iii) high field measurements of 50 Tesla for YBa2Cu3Oj (with *<6.8) and (Y0.6Pro.4)Ba2Cu3Ox (with all levels of oxygen content [33]) showed an insulating behavior at low temperature, (iv) medium field measurements up to 20 tesla yielded a change of p vs T from logarithmic to exp[T1/4]] relation [34] and (v) a metallic state for YBa2Cu307.s has been obtained by using the intense pulsed current densities to overcome flux-vortex [35]. In this work, instead of applying magnetic field or high current, we direct the spin-injection from Lao.7Sro.3Mn03(LSMO) to YBa2Cu307 (YBCO) to suppress the superconductivity and an insulating characteristic appears at temperature below Tc. Since the heterostructures involves neither the competition of the scattering length and magnetic length [34], nor the instability of flux-vortices [36], the analysis of our results is much straight forward than previous studies.

From the aspect of spin-injection applications, previous reports on the YBCO-LSMO spin-injection experiments were mostly performed on a trilayer structure with a thin insulating/metallic layer in between to form a junction device. However, the bilayer YBCO/LBMO (B = Ba) could work as well as trilayer. In our work, a different approach was taken in a sense that no extra inject current was applied from the third terminal. Some partial cooper pairs were first polarized via the exchange field in LSMO, then injected to YBCO to break the rest of pairs. Our new results suggest that a proximity effect may serve as a spin-polarizer for the future spintronic devices.

2. Experiments

A series of LSMO films with various thicknesses were fabricated by rf magnetron sputtering [37] under identical deposition conditions using a sintered stoichiometric LSMO target in Ar + 0 2 atmosphere with a pressure and a rf power of 20 mTorr and 3.56 Watt/cm2 respectively. Films were synthesized on LaA103 (LAO) (100) substrate. After deposition the samples were post-annealed

390

at 920 °C under flowing oxygen for 6 hours. YBCO layers with fixed thickness of 150 nm was then grown on top of the LSMO films and post-annealed at 700°C in 0 2 for one hour with a slow cooling to 300°C in 6 hours. The four point contact preparation and the experimental setup for the electrical resistivity (p) and current-voltage (I-V) measurements have already been described in [3,4]. For resistance-field (R-H) measurement a sweep field from -1 to 1 tesla is applied. MR ratio is defined as [{(PH=IT ~PH=O) /PH=O} X 100%] and ER ratio is defined as [{(dV/dI(I) - dV/dI(0)) / dV/dI(0)} x 100%]. The film thickness was measured by stylus method on a Dektek-3030 ST profilometer. Surface morphology and composition of films were determined using a high resolution (JEOL-JSM 6700F) scanning electron microscope (SEM) and by energy dispersive X-ray (EDX) analysis (Hitachi-S570) respectively. The phase purity and structure of film were identified by X-ray diffraction (XRD) method.

3. Results and Discussion

3.1. La0.-jSToiMnO)Single Layer

It is essential to understand the properties of LSMO single layer before one can understand the properties of YBCO/LSMO heterostructure. One unique character of CMR material is the coexistence of metal and insulator phases, which make it an advantage of tuning the electronic state. According to the phase diagram, the ratio of La/Sr determines the magnitude of Curie temperature Tc as well as the metal-insulator transition temperature TMi. For the purpose of application, La/Sr = 0.7 is the most practical one since it yields the maximum Tc

around 360 K. Fig. 1 displays p(T) at zero and 1 tesla for films with different thicknesses. As seen in Fig.l-(a) and l-(c), 60 nm film behaves insulating with an upturn at 150K while 100 nm film shows metallic feature with p decreasing with lowering temperature. The result in these two panels clearly demonstrates that there exists a critical thickness for driving the insulating phase to the metallic phase. For the intermediate thickness t = 80 nm (see Fig.l-(b)), p first decreases with a sharp drop at 200K then behaves as an insulator when T decreases from 200K to 10K, demonstrating a mixed phase of metal/insulaor. This type of low-temperature p upturn has been observed in polycrystalline sample as well as in the thin films [38]. The former was attributed to an intergranular Coulomb gap between grains (20-25 nm), and the later to an coexistance of high-strain/low-strain mixed phases. Since the average grain size of our polycrystalline films is more than 50 nm (based on SEM data), the possible origin for the p upturn behavior in our film of 80 nm thickness is the structural disorder by strain in association with the large lattice mismatch (-2.37%) between LSMO (0.388nm) and LAO. Previous work also showed the

391

existence of strain-induced insulating behavior in LSMO [39,40], (La,Ca)Mn03

[41] and La-Sn-Mn-0 films. The structural disorder can result in spin disorder and enhance the electron localization [42], or may lead to the absence of the characteristic insulator-metal transition in ferromagnetic manganites [43]. Recently, similar p upturn has also been seen in and was described as the localization effect [12,44].

0.16

o •

a Q.

0.08

""n.

I

(b)

i i ^ ^ f 8

20

3» 15

i'° 5 0

1 1

M̂ 100 200

^ T ( K )

. i . i

s. 300

•

50 100 150 200 250 300 350

T(K)

Figure 1. p vs. T under zero and lTesla field for LSMO/LAO films with different thickness : (a) 60 nm, (b) 80 nm and (c) 100 nm.

392

In the percolation picture, which is generally applicable in CMR materials, local electric field perturbs the coexistence of phases of different electronic densities and sets up filamentary currents across nonconductive regions. This filamentary current, in turn, produces magnetic field to induce resistive drop (so call electroresitance, ER). Therefore, the resistive drop induced by magnetic field may be equivalent to that induced by current. As described in [10], for the single crystal, a current of 0.3 mA is equivalent to the effect of applying a magnetic field of 1.5-2 tesla at low temperature but only 0.4 tesla at room temperature.

Figure 2 shows the MR(H) graphs derived from R-H plots at RT. We observed a linear decrease of R with increasing H, which indicates spin related elcetron scattering at grain boundaries (GBs) [45]. In this case MR increases from 2.96% to 3.85% as the film thickness increases from 60 to 100 nm. This marginal increase of MR with thickness is consistent with the report [46] that high field MR remains almost constant for film with t > 20 nm. This low MR at RT is ascribed to the spin disordering of LSMO at high T. Our result of MR (3.85%) for t = 100 nm is very close to the reported value of 3.81% for t = 200 nm LSMO film under 1 tesla field [47].

0 ; (a)

-1 -

-2 -

-3 -<r

. 4 l_ l I I I I ' • < L

0

d .1 o

< -3

-4 0

• 1

-2

-3

-4

-8000 -4000 0 4000 8000

H (Oe )

Figure 2. MR ratio (%) vs. H at RT for LSMO/LAO films with different thicknesses t = (a) 60 nm, (b) 80 nm and (c) 100 nm. MR values at lTesla are displayed on right top corner of each layer.

2.96%

J 1 I 1 1 1 I 1 L

J 1 I 1 I 1 I 1 L

393

Inset of Fig. 3 shows the ER ratio (%) vs. current (I) derived from dV/dl - I plots. With I = 0.3 mA, we obtained 5.6%, 6.4% and 3.5% of ER ratio at RT for t = 60, 80 and 100 run respectively. With I = 0.9 mA ER ratio increases to 11.3% and 4.6% for t = 80 and 100 nm, respectively. It indicates that room temperature ER value decreases 1.6 times when t increases from 60 to 100 nm whereas MR increases to 1.3 times (see Fig. 2) for the same thickness variation. Although there is similarity between MR(H) and dV/dI(I) curves (Figs. 2 & 3), the increasing(decreasing) tendency of MR(dWdl) with increasing thickness suggests that both effects may not have exactly the same origin as normally argued for the magnetic and electric field effects. According to the percolation model, electric field perturbs the coexistence of metallic and insulating regions by creating metallic inclusions [48] within the insulating regions. This metallic inclusion may in turn produce filamentary path where outer layer is insulating and the inner one is metallic. Besides CMR materials, this type of filamentary pattern has also been observed by applying current in amorphous hydrogenated silicon devices [49]. Hence, current flow through space limited within the filamentary regions induces intense local magnetic field which polarizes the FM regions and induce CMR effect. In this case the resistive changes due to magnetic field and applied current are expected to be equivalent as reported for LCMO cryatal [10]. However, this study showed that the thickness dependence of MR and ER effects are in opposite direction. Furthermore, based on the ER and MR studies in single crystal [10], the value of ER ratio should be near equivalent to that of MR ratio at various temperatures from 295 and 65 K. Therefore, both ER and MR increase by 5 times at low temperature. Nevertheless, this model could not explain the results of our films with high-strain (60 and 80 nm films), because the ER value is much higher than the MR ratio at room temperature while they are almost equivalent at 15 K.

A comparison of the resistive changes due to magnetic field and current effects leads to a conclusion that the sample with t = 80 nm is a strongly disordered metallic system and the occurance of M-I tranisition may be caused by a quantum effect- Anderson transition [37]. In the model of Anderson transition [50], an electron moving in a random potential may have either localized or extended eigenstate depending on the energy of electron. Extended states can carry a direct current whereas localised states are bound to certain region and can move only with the assistance of other mechanism (e.g. phonon-assisted hopping) [50]. When one passes a current to the sample, localized electrons gain energy and become conductive via phonon-assisted hopping. Accordingly, the electrical resistance decreases. This explains why MR] of 80 nm film for a current of 0.3 mA (equivalent to 1.5 tesla field, see [10]) is 2.3 times larger than that of than that of MRH at 1 tesla field. Hence, to get a clear

394

picture on the phonon-assisted delocalization effect, we studied the ER effect for t = 80 & 100 nm films at low temperature to minimize the phonon contribution. The evidence that current effect is almost the same as magnetic-field effect at low temperature while it is more influential than magnetic-field effect at room temperature is consistent with the model of phonon-assisted delocalization.

124

120

116

112

108

^ 26

•a > 24 •o

22 33

32 -

31 -

30 -

Figure 3. dV/dl vs. current (I) for LSMO(t)/LAO films with different thicknesses t = (a) 60 nm, (b) 80 nm and (c) 100 nm. Inset is ER ratio (%) vs I.

395

Although the strain-induced metal to insulator transition may not be avoid owing to the mismatch between film and substrate, the localization effect could be reduced by optimizing the annealing condition. In our earlier experiment, the optimal annealing condition is 920 C for 6 hours in flowing oxygen. However, we found later that a fine controlled flowing rate of oxygen (-50 cc/min) can eliminate the mix phase of metal/insulator, and the critical thickness of metallic to insulating phase can be reduced to 50 nm (see Fig. 4). We thus used these optimal conditions to produced YBCO/LSMO heterostructures [52].

100000

10000

I 1000

^ ioo

10

1 0 100 200 300 400

Temperature (K)

Figure 4. p vs. T for slow annealing LSMO/LAO films with t = 20 nm (insulating behavior) and 50 nm (metallic behavior).

3.2. Superconductor-Insulator Transition in YBa2Cu30/Lao.7Sro.3Mn03

Heterostructures

Based on the X-ray data of our YBCO/LSMO heterostructure, the LSMO bottom layer has a monoclinic structure with the c-axis preferred orientation and its lattice parameters are given by a = 5.387, b = 5.425, and c = 8. 026 A; while the YBCO top layer has an orthorombic structure with a = 3.731, b = 3.876, and c = 12.043 A. Figure 5 displays the thickness dependence of p at 5 K for pure LSMO films and YBCO/LSMO heterostructures. The open circles are the data for LSMO films, showing that p(5K) is around 2.5xl03 ohm-cm for dLSMO = 10 nm and gradually decreases with increasing dLSM0, and has a sharp drop at dLSM0

= 50 nm. The reduction of p(5K) between 40 and 50 nm is of the order of 104

indicating a transition from insulator to metal at d = 50 nm. On the other

396

hand, p(5K) of YBCO(150nm)/LSMO(t) heterostructures is zero for t < 40 nm but jumps to 55 ohm-cm at 50 nm, as shown as the solid circles in Fig. 5.

50

--v 40

s V 30 & 20

a. 10

0

.

•

.

•

i

o

•

LSMO • -

~ ^ \ * /

^ x -YBCO/LSMO / \

/ \ i . i . i . J_ .

2500

2000

1500

1000

500

o

10 20 30

,(nm)

40 50

Figure 5. t dependence of resistivity at 5K for pure LMSO(t) films (open circles with the right scale) and YBCO/LSMO heterostructures (solid circles with left scale).

Figure 6 is the p(T) plot for YBCO (150nm)/LSMO(dLSMO) heterostructures with dLSM0 varying from 10 to 50 nm. For t < 40 nm, superconducting transition occurs at Tc ~ 75K as shown in the inset of Fig. 6. The up-turn behavior of p(T) at low temperatures can be further fitted by exp [1/TI/4], reflecting the nature of variable-range hopping for its localized electron state. Our result is consistent with those obtained by the high field experiment of Vanacken et al. [33] and the medium field experiment of Karpinska at al. [34], only the tool we use to destroy superconductivity is different. It can be understood by the fact that the ferromagnetic LSMO is very close to a half-metal and the injection of highly polarized electrons from LSMO to YBCO has strong pair-breaking effects. The route of electric current to inject the spins into YBCO is can be considered as an effective circuit depicted in Fig. 7, in which R,(YBCO) and R,(LSMO) are the resistances along the surface of individual YBCO and LSMO layer, respectively, and R2(YBCO) is the transverse resistance (proportional to its film thickness) of the YBCO layer. Since the distance between the current contacts on the surface of YBCO layer is much larger than the thickness, R,(YBCO) is always much greater than R2(YBCO). In this parallel circuit, the current goes through route 1 for t < 40 nm, because YBCO retains metallic/superconducting and LSMO is insulating. For t > 50 nm, the current may go through either route 1 or 2 since Ri(YBCO) and R^LSMO) are comparable to each other. At low temperature such as 5 K, R,(YBCO) for the insulator-like YBCO film is much greater than R,(LSMO)of the metallic film, so that the resistance measured on the

397

YBCO/LSMO herterostructure is mainly determined by that of route 2, approximately equal to a series connection of double R2(YBCO) plus one R,(LSMO). It explains why the resistivity of YBCO/LSMO herterostructure at t = 50 nm is larger than the corresponding LSMO film resistivity in Fig. 5.

60

P V

CL

150 200 250 300

T(K)

Figure 6. Temperature dependence of the resistivity for YBCO/LSMO(t) heterostructures with various t. Inset is the p-T plot for t < 40 nm.

G-

I R2(YBCO)

Route 2

Route 1

•

R,(YBCO)

R,(LSMO)

-e

R2(YBCO) t

Figure 7. Effective circuit for the spin injection in YBCO/LSMO(t) heterostructures.

398

The result of thickness dependent magnetoresistance (MR) ratio, defined by [p(5 tesla)-p(0 tesla)]/p(0 tela), is shown in Fig. 8, for dLSMQ < 40 nm MR ratio at 50 K is positive and is near zero at 77K and 300K; while for d

LSMO

the MR ratio at all three temperatures is negative, and its absolute value gets higher for higher temperature. These results are consistent with the proposed effective circuit in Fig. 7. Accordingly, for dLSM0 < 40 nm the superconducting

YBCO dominates the transport property of YBCO/LSMO heterostructure and LMSO has a negligible contribution. Therefore, there is no significant MR effect except in the mixed state regime (40K < T < 75K), where an applied magnetic field leads to a thermal activated flux movement, yielding a positive MR ratio at 50K. For t > 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 40

d (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.

xSrxCu04 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 Ic-value 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

0.6

03

0.0

> -03

-0.6

-0.9

40nm

lOhn,

20ml 30tn i

YBOO

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.

0.04

0.02

2 " 0.00

> -0.02

-0.04

j,5T

A4* Mi

w— . 1 , 1 . 1 . 1 . 1 , 1 . 1 . 1 . 1 .

-5 4 - 3 - 2 - 1 0 1 2 3 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

1.5

1.0

YBCO

t=20nm

'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.

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DOMAIN STRUCTURE OF YBa2Cu3Ox FILMS ON NdGa03

SUBSTRATES

I.K. BDIKIN Institute of Sol id State Physics, Chernogolovka, Moscow distr., 142432, Russia

E-mai: bdikin@issp.ac.ru

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 NdGa03 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 Ce02 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 the films with 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 YBCO films occurs in accordance with the {110}/<11_0>

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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 Ce02 seeding layer.

2. Experimental

Ce02 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 Ce02 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 DC-sputtering 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

<1200 m

& o

•£• BOO

400

X-RAY SOURCE

010)NGO (22Z)Ce02 ,(222)Ce02

(120)NGO •£— • =H

(130)NGO

WJ

(010)NGO

- 1 5 . 0 -5 .0 5.0

A en 15.0 25.0

Figure 1. Rocking curve in a wide angular scanning range of a Ce02 film on the (130)NdGaO3

substrate. On the inset: schematics of the technique. The angle between reflections is equal to the angle between the normals to the crystallographic planes.

3. Results and Discussion

Film lattice parameters were determined using (0 0 13), (3 0 10) and (0 3 10) reflections. The following values a = 3.827(1) A, b = 3.889(1) A, c = 11.674(2) A were found for films on (110) NdGaC>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 Ce02 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).

A0 (°)

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||<110>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,

&"

8

(a)

(001)Y,=(110)N

26"

• • • • • 1I 8 - "

11°

n = (110)N

A G ( ° ) 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).

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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}/<1H)> 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 Ce02 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 Ce02 interlayer. The YBCO thin films, deposited by DC sputtering at high oxygen pressure, instead, are oriented along the axes of the Ce02 interlayer, independently on the substrate orientation: [001]Y || <001>c. Cubic symmetry of Ce02 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

138.0

(a

(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 Ce02

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 Ce02 layer into the YBCO thin film results in alignment of the axes of both materials ([001]Y || <001>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 Ce02

seeding layer, the YBCO thin film contains a significant (up to 60%) non-twinned 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 Ce02/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 Ce02 seeding layer.

The presence of a thin Ce02 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 <100>c of the Ce02 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

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5. A. D. Mashtakov, K.Y. Constantinian, G. A. Ovsyannkov, E. A. Stepantsov, Techical Physics Letters 25, 249 (1999).

6. 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).

7. 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).

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9. F. Vassenden, G. Linker, J. Geerk, Physica C 175, 566 (1992). 10. F. Miletto Granozio, M. Saluzzo, U. Scotti di Uccio, I. Maggio-Aprile,

Oe.Fischer, Phys. Rev. B 61, 756 (2000). 11. 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: nawazishalik@yahoo.com

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 cm4, due to the coupling of apical oxygen (0A) 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 SrTi03 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.

SrTiQ3

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 c-axis 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 Cu-1234 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)-0A-Cu(2). Such types of apical oxygen modes have also been observed in Hg|.xTlx-1223 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 Cu-1234 [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 (0A) generated by the presence of the interstitial oxygen Og [14].

10000

eooo

6000

4000

2000

_L-^_W ^SA-LJ 20 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

15 -

(a)-EinWc-axis

(b)-Eiruc-axis

_L _L

527cm-1 598cm-1

_L 200 300 400 500 600 700

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)-0A-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. H. Ihara, K. Tokiwa, H. Ozawa, M. Hirbayashi, A. Negishi, H. Matuhata, Y.S. Song, Jpn. J. Appl. Phys. 33, L503, (1994).

2. N.A. Khan, Y. Sekita, N. Terada, H. Ihara, Supercond. Sci. Technol. 14, 603 (2001).

3. H. Ihara, K. Tokiwa, K. Tanaka, T. Tsukamoto, T. Watanabe, H. Yamamoto, A. Iyo, M. Tokumoto, M. Umeda, Physica C 282-287, 957 (1997).

4. N.A. Khan, Y. Sekita, H. Ihara, A. Maqsood, Physica C 377, 43 (2002). 5. A.D. Kulkarni, F. W. deWette, J. Prade, U. Schroder, W. Kress, Phys. Rev.

B 41, 6049 (1990). 6. 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: lykov@maill.lebedev.ru

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 0.2-10 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.

SMV 0 r

0 L

superconducting ' film 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

f, M H z 448 450

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:

V/(f)/V/(fr) (2)

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)<\ practically in the entire frequency range, and is significantly more than 1 only when f=nfr. Thus the harmonics of the radiation disappear with

425

increasing distance between the film and the resonance coil. The dependence HJ) is explained by the fact that the electromagnetic field emitted provides an ordering in vortex motion and narrowing the hysteresis loop of the magnetisation curve of the superconducting films, because the energy source of the electromagnetic field, both coherent and noise, is the hysteresis of the loop. In case of weak influence of the rf radiation on the vortex system the ratio r(J) should be equal to unity, i.e., the frequency dependencies of the signal are similar for these cases.

2.0 -

1 .5

1 .0 -

0.5

0.0 2 3 4

f,M H z

Figure 4. The signal ratio as a function of the frequency for //io=56.1.6 Oe, mvifexc=554 Hz.

4. Discussion

The fine structure of magnetic field dependence vs external magnetic field in the mixed state should be taken into account to explain the observed generation of rf radiation with frequency up to 600 MHz. If the external magnetic field exceeds the lower critical field Hci, the magnetic flux penetrates the type II superconducting films in the form of Abrikosov vortices. The vortex motion and consequently the voltage induced in a pick up coil depends on the presence of pinning centers in the film. If the pinning forces are overcome, the magnetic flux

426

undergoes a jump. In this regime the magnetic flux on the sample surface is a sum of flux jumps of many different scales. The voltage spectral density measured by the selective voltmeter coupled via a resonant circuit has a maximum at the resonant frequency of the LC-circuit. The origin of the self-radiation harmonics lies on the dependence of the inductance on the oscillating current in the LC-circuit. This is because the coil is located in the vicinity of the superconducting film. The magnetic properties of a superconducting film depend on the field produced by the coil when the amplitude of the current oscillation is large enough. It is known [for example, [5] that the harmonics of the radiation can arise in LC-oscillating systems with non-linear inductance. As usual, equation of the oscillation in the LC systems can be written:

^ + ^ = 0 (3) dt C

where / is a current (/ = oq/dt). To take into account the nonlinear dependence

L(i), we believe:

L(i) = L0(\ + yi2) (4)

where y is a coefficient. In the case y « l , it follows from the equation [5], 3 3

q = a(l YO)2 a2) cos 0)t + — va>2a3 cos 3cot, (5) 32 32

where a and <n=2nfr are the amplitude and frequency of the charge oscillations in the resonance circuit when y=0, respectively. The equation 5 contains the oscillations atfr and "if, in accordance with our experiment. In the other words, the self-radiation harmonics are a result from back action of the resonant circuit oscillation on the superconducting film.

5. Conclusions

Thus, a new method of designing active superconducting elements is developed. The generation, amplification and harmonic mixing of the electromagnetic radiation is demonstrated on the basis of the above approach. The system resembles a maser where the active medium is a system of Abricosov vortices in the superconducting film. Its excitation is attained by a low-frequency oscillating magnetic field. At last, the resonant circuit works as a resonator in the maser. The maximum frequency of oscillations is limited by the recombination time of quasi particles, which form the normal (non-superconducting) core of the vortices. The approach offers new facilities to study the non-equilibrium properties of the vortex system in superconductors, and new possibilities for applying superconductors in electronics.

427

Acknowledgments

The author would like to thank Yu.V.Vishnyakov for technical assistance.This work was supported by the Russian ministry of industry, science and technologies (grant No 40.012.1.1.11.46 - Controlled superconductivity);and Russian-Ukrainian project, "Etalon".

References

1. B.D. Josephson, Phys. Lett., 1,251 (1962). 2. A.N. Lykov, Phys. Lett. ^,281,48 (2001). 3. A.N. Lykov, JETP Letters, 73, 549 (2001). 4. A.N. Lykov, Physics C, 367, 327 (2001). 5. V.V. Migulin, V.I. Medvedev, E.R. Mustel, B.N. Parygin, Foundations of

the theory of vibrations (Nauka, Moskow, 1988), 392.

FABRICATION OF YBCO AND BSCCO THIN FILMS

H. SALAMATI, P. KAMELI

Solid State Lab., Physics Dept., Isfahan University of Technology, Isfahan 84154, Iran

M. AKHAVAN

Magnet Research Laboratory (MRL), Department of Physics, Sharif University of Technology, P.O. Box 11365-9161, Tehran, Iran

In this paper, we review two deposition techniques of high-Tc superconductor materials namely sputtering and pulsed laser deposition. Some basic physical concepts and technologies are discussed and compared. The results of structural and physical properties of YBCO and BSCCO thin films made by PLD and sputtering techniques will be reviewed.

1. Introduction

Since the discovery of high-Tc superconductors [1], significant effort has been put into the research and realization of textured and epitaxial thin films [2]. This effort is motivated largely by the potential applications of thin films in a number of cryoelectronic devices and by the possibility of using epitaxial single or multilayer high-Tc superconductor thin films to study new physical properties of these unique layered materials. Different techniques such as molecular beam epitaxy (MBE) [3,4], metal organic chemical vapor deposition (MOCVD) [5,6], sputtering [7,8], liquid phase epitaxy (LPE) [9], and pulsed laser deposition (PLD) [10,11] have been employed.

Sputter deposition has the deserved reputation of being the technique for preparing thin films of alloys and complex materials. So, due to the structural complexity of high-Tc superconductor materials, it was natural to use this technique for deposition of these materials. However, the problem of resputtering (due to oxygen ion generated in the plasma) was recognized in this class of materials some 18 years before the discovery of YBCO for Ba(PbBi)03

deposition and was thoroughly investigated latter [12,13]. A method to avoid this problem is to thermalize the energetic species either by working at very high pressure (which also provides a particularly rich oxygen environment) or by "off-axis" sputtering at a modestly high sputtering pressure which forces all atoms originating from the target to undergo a few energy-reducing collisions before approaching the substrate.

Pulsed laser deposition is a well-known method for the deposition of high-Tc superconductors. High-Tc superconductors films deposition by PLD system is

428

429

carried out by irradiation of a single target by a focused excimer laser beam with the following characteristics: Wavelength (X=308 nm and 248 nm or Nd:YAG, A.=355nm), energy density 1-3 Jem'2 per shot and pulse duration of 10 ns. An intense laser pulse passes through an optical window of a vacuum chamber and is focused on to a target, where it is partially observed.

In this paper, the two techniques of in situ and ex situ processes are implemented in sputtering and PLD deposition. Report will be given on the preparation of the targets of YBCO and BSCCO for sputtering and PLD deposition. The results of structural and physical properties of targets will be presented. Substrate requirements will be mentioned. The results of structural and physical properties of YBCO and BSCCO films made by PLD and sputtering techniques will be discussed.

2. Experimental

2.1. Fabrication of Targets

A series of targets (different sizes and different compositions) were prepared by the conventional solid-state reaction technique. The superconductivity and structural properties of the targets were tested by resistivity and susceptibility measurements, and characterized by SEM and XRD.

2.1.1. YBCO Targets

Four sizes of disk shaped targets with diameter d=100, 75, 55, 25 mm and thickness of 5 mm were made using high purity powders of Y203, BaC03, Ag20 and CuO. The powder were mixed in stoichiometric proportions, calcined at 920°C for 24h and sintered at 930°C in an oxygen atmosphere for 24h. Calcinations and grinding procedures were repeated three times.

2.1.2. BSCCO Targets

Three sizes of disk type shaped targets with d = 100, 55, 25 mm and thickness of 5 m were made using high purity powders of BJ203 SrC03, PbO, CaC03, and CuO. The powder were mixed in stoichiometric proportions, calcined at 830 °C for 24h and sintered at 860 °C in air for 170 h. The calcinations and grinding procedures were repeated three times.

430

2.2. Fabrication of Thin Films

2.2.1. Sputtering Deposition ofYBCO andBSCCO Thin Films

AH the deposition experiments were performed in a non-baked stainless steel vacuum chamber with the base pressure of about 4xl0~7 torr. The YBCO thin films were fabricated on cleaned MgO (100) and LaA103 (100) substrates, and the BSCCO thin films were fabricated on MgO (100) substrates, utilizing DC sputtering technique. Ultra high purity Ar gas was used as discharge. Tables 1 and 2 represent parameters for the sputtering deposition of YBCO and BSCCO thin films.

Table 1. Deposition parameters for the growth ofYBCO thin films by sputtering.

Parameter

Substrates

Total Ar pressure

Power

Post annealing treatment

Condition

MgO, LaA103

100-300 mTorr

40-100 W

Oxygen atm,(800-850°C 3hr)

Table 2. Deposition parameters for the growth of BSCCO thin films by sputtering.

Sample

Thinl

Thin2

Substrate temperature (°C)

750

800

Total gas pressure (mTorr)

600

600

Sputtering time (h)

2

2

2.2.2. PLD Deposition ofYBCO Thin Films

Tables 3 represent the parameters for PLD deposition of YBCO thin films. A high quality c-axis-oriented YBCO epitaxial thin film was prepared by pulsed laser deposition of stoichiometric YBCO target on a SrTi03 (100) substrate. A Lambda Physik LPX 305i excimer laser (KrF, 248 nm wavelength, 25 ns pulse length) was used in combination with an optical train and Neocera vacuum chamber. The SrTi03 (100) substrate was mounted onto a heater block by silver

431

paint. The heater block was positioned 8 cm away from the target. Typically, we used an oxygen pressure of 200 mTorr, a substrate temperature of 780°C, laser pulse repetition rate of 5 Hz and laser energy of 450 mJ/pulse. After the deposition, one of the samples (Thin 3) was cooled down in oxygen atmosphere (760 Torr) to room temperature. The second sample (Thin 4) was cooled down to 480°C, and annealed at that temperature for about 30 minutes.

Table 3. Deposition parameters for the growth of YBCO thin films by PLD.

Sample

Thin3

Thin4

Substrate temperature(°C)

780

780

Oxygen pressure (mTorr)

200

200

Post annealing treatment

No post annealed

Annealed at 480 °C for 1 h in oxygen atmosphere.

3. Results and Discussion

Figure 1 shows the resistive transition behavior for YBCO target. It can be seen that the superconducting transition is very sharp with Tc~92 K. Figure 2 shows the XRD patterns for YBCO target. The XRD analysis shows that the target is almost pure without any detectable impurity.

0.014

0.012 -

0.010

- . 0.008 H

^ 0.006

0.004

0.002 H

0.000

30 60 90 I I I I I I

120 150 180 210 240 270 300

T(k)

Figure 1. Temperature dependence of resistivity for YBCO target.

432

200

180

160

140

120

100

80

60

40

20

0

S

A„_

20

1—

25

- A _ J

o o

I 30 35

UL K T

40

20

I

45

- J —

50 55 60

Figure 2. The XRD pattern for YBCO target.

Figure 3 shows the resistive transition behavior for BSCCO target. It can be seen that the superconducting transition is sharp with Tc~108 K. Figure 4 shows the XRD pattern for BSCCO target. The XRD diagram for BSCCO sample indicates the presence of large amount of high-Tc (Bi2223) phase.

Figure 5 shows the temperature dependence of the normalized resistivity for two representative YBCO thin films deposited on MgO (100) substrates utilizing DC sputtering technique. This Figure shows that the higher onset Tc of about 86 K is obtained for the samples grown on unbiased and off-axis geometry. Figure 6 shows the resistive transition behavior for three sets of unbiased geometry grown on MgO (100) and LaA103 (100) substrates. As it can be seen for the films deposited on MgO (100) substrate, a better metallic resistivity was observed (Fig. 6a) as compared with the samples grown on LaA103 (100) substrate with similar conditions [14].

Figure 7 shows the XRD patterns of the BSCCO films prepared at different substrate temperatures, keeping the Ar gas pressure constant. In this figure, all of the strong diffraction lines are assigned to (001) of 2212 and 2223 phases, but as it can be seen clearly from the pattern, the proportions of 2223 phase will increase as the substrate temperature increases.

Figure 8 shows the temperature dependence of resistivity for two samples. Both samples show superconducting transition, and they have metallic behavior in their normal state. As the temperature of substrate increases, the slop of resistivity in normal state increases, the transition temperature of the film increases, and the normal resistivity of the film decreased.

433

a:

0.12

0.10 -

0.08 -

0.06 -

0.04 -

0.02 -

0.00

50 100 150

T(K)

200 250 300

Figure 3. Temperature dependence of resistivity for BSCCO target.

2500

+ 2223

2000 -\ * 2212

£ 1500 -•51

"c

O O

1000 -

500 -s

+

10

00 00 J - CO to o o o •*" T -

BSCCO

20

Figure 4. The XRD pattern for BSCCO target.

434

0.0 T "i i m" l i 1 i | 0 20 40 60 SO 100 120 140160

T[K]

Figure 5. Temperature dependence of resistivity of YBCO films, (a) biased, 930 °C, lh. (b) off-axis unbiased, 800 °C, 3h.

T[K]

Figure 6. Temperature dependence of resistivity of YBCO films, (a) MgO(100), 800 °C, 3h. (b) LaAlO3(100), 800 °C, 3h. (c) LaAlO3(100), 850 °C, 3h.

435

35

30 H *<2223) +(2212)

S 2 5 - I

| 20H "ST

o o 10 H

5 -

0

Thin2

Thinl

10 15

28 Figure 7. The XRD patterns for Thinl and Thin2 samples.

300

Figure 8. Temperature dependence of resistivity for Thinl and Thin2 samples.

436

Figure 9 shows the XRD patterns for two YBCO thin film samples, which are made by PLD technique. The XRD pattern indicates that the films are epitaxially aligned with the c-axis, normal to the substrate purface. Figure 10 shows the temperature dependence of the resistivity for two different samples. The Thin 4 sample has a better metallic resistivity at its normal state and higher transition temperature. So, for the fabrication of high quality YBCO thin films, a post-deposition heat treatment in oxygen atmosphere is necessary.

15000

12000 -

.ri 9000

*->

§ 6000 H o

3000 -

(005)

(001) (002)

JL

i 10

T

20

(003)

L JL.

(004)

IN

(006)

I 30

1L J

40

(007)

Thin4

I

50

.A Thin3

60

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

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

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. J. Bednors, K. Muller, Z Phys. B 64, 189 (1986). 2. R. Wordenweber, Suprcond. Sci. Technol. 12, R86 (1999). 3. D.J. Rogers, et al., Supercond. Sci. Technol. 12, R75 (1999). 4. H. Ota, et al., Physica C 311, 42 (1999). 5. S.J. Golden, F.F. Lange, D.R. Clarke, L.D. Chang, C.T. Necker, Appl.

Phys. Lett., 61,351(1992). 6. T. Sugimoto, et al., Appl. Phys. Lett. 63, 2697 (1993). 7. Z. Mori, E. Minamizono, S. Koba, T. Doi, S. Higo, Y. Hakuraku, Physica

C339, 161(2000). 8. S. I. Karimoto, S. Kubo, K. Tsuru, M. Suzuki, Jpn. J. Appl. Phys. 36, 84

(1997). 9. G. Balestrino, et al., J. Appl. Phys. 70, 6939 (1991). 10. L. Ranno, et al., Phys. Rev. B 48, 13945 (1993). 11. H. Salamati, P. Kameli, Physica B 321337 (2002). 12. R. Gilbert, et al., J. Vac. Sci. Technol. 17, 389 (1980). 13. S. Rossnagel, J. Cuomo, Thin Film Processing and Characterization of

High-Temperature Superconductors, (Am. Inst. Phys. Conf. Proc. 165 New York, 106, 1988).

14. 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).

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VIII. MAGNETIC THIN FILMS

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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 magneto-resistive angle and position sensitive sensors in the automotive industry , for example. The technological important aspects of exchange biased [10], spin-valve [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 <100>,

441

442

<110>, <111> and for fee Ni <111>, <110>, <100> 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 (uca :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:

1 1 H = D, withD,, =

* 4 * ^

and

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 Mx

2 + Ny My2 + Nz Mz

2) 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.

Fe

Co

Co

Ni

Gd

Afi/ntot

1.7xl0'4

4.5X10"4

1.8xl0'4

~io-3

Easy axis

[100] bec

[0001] hep

[111] fee

[111] fee

[tilted] hep

MAE (0 K) (jieV/atom)

+1.4

+65

1.8

-2.7

+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| + Fmagei. For a cubic crystal Fan is written in the form:Ean = K0 + K4' (ax

20y2 + Oy2az2 + az

2ax2) + K6'

ax2ay

2az2+ 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 K2

V 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 K2

V. 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 10-20 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<i(T=0 K) for Ni, Co, and Fe layers is on the order of 11, 90 and 140 ueV/atom (Ni-Ny = 1). These values are much larger than the intrinsic anisotropy K of bulk 3d elements due to spin-orbit coupling. One can estimate the critical thickness dc for a disk of diameter D above which a perpendicular magnetization is found by dc /D = l/(37i3) (1 -K/Fd). N|| of a thin disk is approximately given by Ny « 7i/4 d/D. For a 10 mm diameter disk of Co at T=0 K one calculates dc » 20um. dc is temperature dependent due to the different temperature dependencies of Fd and K.

In the limit of a few atomic layers it was argued that a film should not be considered as a magnetic continuum. The discreteness of the lattice becomes important. Calculations of the sum of the dipolar interaction for a discrete assembly of magnetic point dipoles have shown that the average demagnetization factor Nx for a film containing n atomic layers is reduced to N± = 1-A/n with A = 0.4245 (n > 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 + 2K2

s/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 Tc 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 Tc. 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 Tc. As the magnetic moment above T c is the same (except for some additional polarization of the conduction electrons) as the one measured below Tc (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 Tc, but the magnetic anisotropy energy vanishes above Tc, 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 Tc that spatially and timely averaged moment vanishes. Hence, also the macroscopically measurable MAE vanishes, it averages out above Tc. 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) + • The coefficients are related to the magnetization by

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 Tc. 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.

K f / arb.units

— o — K

K 4 1

' ^

K2(T) = 1.47 k3- 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 Tc 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/Tc, 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, C02) 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 N6el-stracture, 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 <fi (bottom, 0=-9O°) of the polarization P of the electron beam with respect to the sample normal. Top: Sharp magnetic contrast with the polar angle in the film plane (left), no contrast with 0 perpendicular to the surface (right), i. e. the Ni film is fully in-plane magnetized. Bottom: At ^=0° the MC is maximum, the domains are parallel and antiparallel with respect to the polarization direction. The magnetic contrast disappears at ^ -90° . Two chiralities exist in the Nfel-wall, as seen at ^=90°. Due to (a) parallel and (b) antiparallel orientations the N&l-wall appears white and black, respectively.

A line scan across this MM shows the typical profile consisting of a narrow core and a long tail as discussed in ref. [17, page 244]. The core width has been determined to be about 400 nm). We find good agreement with calculated

452

profiles of Neel-walls for which the width can be estimated by s = xy]2A/Ktff

using the exchange constant A = 0.75 x 10"" J/m, that is the average of values given for Ni thin films with 157 to 250 nm thickness, and an effective magnetic anisotropy parameter Keff = K2 + Kshape = 0.9 x 103 J/m3. The parameter Keff which includes shape (Kshape) and second-order magnetocrystalline anisotropy (K2) is about one order of magnitude smaller than the experimentally determined Keff of a 8 ML Ni film on Cu(001) at room temperature . Such a difference by orders of magnitude was also observed for Co monolayers on Cu (100) [18] where a Neel-wall width of about 500 and 300 nm was measured for 5.5 and 9 ML. One should note here that a calculation for the Neel-walls of "negative anisotropy materials" favoring <111> directions like Ni should be corrected by taking magnetostriction effects into account yielding a wall width £ = 6.5. ..7.2 J A/ K (P- 234 of [17]) in better agreement to our experimental observation. In difference to the expected decrease of the N6el-wall width for thicker films the domain wall width in Ni/Cu(l 10) increases from about 330 nm at 5 ML to 450 nm at 9 ML and dramatically broadens to 1400 nm (yielding a very small Keff < 102 J/m3) at the start of the spin-reorientation transition (SRT) from in-plane to out-of-plane near 9.5 ML. This enormous increase of the domain wall width within a few tens of monolayers can be explained by the decrease of the effective magnetic anisotropy (proportional to the reciprocal thickness). Eventually Keff disappears at the SRT as the result of the compensation of the in-plane shape anisotropy and the out-of-plane, spin-orbit induced (magnetocrystalline) anisotropy. To quantitatively check the theoretical predicted shape and width of domain walls in magnetic monolayers one needs an accurate knowledge of the magnetic anisotropy Keff and the exchange constant, A which turns out to be not available in many systems. For example, to obtain a quantitative agreement between the experimentally measured wall width of in-plane magnetized 1 ML Fe/W(110) [19] and the calculated one according to g K \A~IK > 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.

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13. M. Farle, Magnetic Thin Films, in NANOSCALE MATERIALS eds. L.M. Liz-Marzan, P. V. Kamat, (Kluwer Academic, 395, 2003)

14. H.B. Callen et al. , Phys. Rev. 139 (1965) A455 ; W.J. Car, Jr.; Phys. Rev. 109, 1971 (1958).

15. R. Ramchal, A.K. Schmid, M. Farle, H. Poppa, Phys. Rev. B 68, 054418 (2003).

16. K. Baberschke, M. Farle, J. Appl. Phys. 81, 5038 (1997). 17. A. Hubert,R. Schafer, Magnetic Domains, Springer (1998). 18. A. Berger, H. P. Oepen, Phys. Rev. B 45, 12596 (1992). 19. 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: aip_2001@mail.ru

The microscopic mechanisms of magnetooptical activity in dielectric compounds with f-and 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 low-dimensional 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\ d'egdJge.r +

digdge.r 1 . (i) heg

s[G)eg-(o-ireg <Deg+6>+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 « h\d) - Q)ov\ we can develop

(1) as a power series [AEe - AEg ) « ti{p) - coov)

456

*"i£* "g&g

OH-C0OV

4e4 •g(reg

c*-%i 1+ ^-+..

ti(a)-%v) (3)

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]

<*y = ao#v + iajtjk < L K > +a2 < QV(L) > , (4)

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 =-(2o)Lvr1o>C1(L),P2L =C2(L),

C1(L0+1) = (L0+1)-},C1(L0) = [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 <LK > 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 spin-orbital interaction, but in this case the diamagnetic contribution can be of the same order as or exceed the "unfrozen" contribution from <L >.

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, (5) mc

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

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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 aF 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), (8) where

459

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

J(") = y y y y J

(coo(H)-co, -r2) + 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]-H2)(H]-H)

f(H) = A ' 2 n ] '— (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. The first example is the Faraday effect in KMnF3 [10].

There are three contributions of the Faraday rotation of KMnF3

aF =ad+ap+amat,

where a^ 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 resonance frequency

hcox « ( 1 3 - 1 7 ) x l 0 3 cm1 and the line width fcr«(4-5)xl03 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 F0=13 deg/(cm Moe).

amat ' 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 cm"1, i.e. near transition 6S^>4Tlg(G) • 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

6 0 •

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 the field dependence 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 R3M50i2 (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

200

V* \ . TbGG

TbAO

\ . \ 0 20 * 0 T ( 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| | [ l l l ] ; 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!

15

0.5^

. i6 i

f2i S~

/ /

/

0 2 HIT)

DyGG / /

Ufl ' H ' T I

300

. - 2 0 0

100

\ OvAG, DVGG

(c)

0 20 « T | K |

\ DyGG

20 40 T(K)

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

466

o o o -

SOO -

BOO -

4 0 0 -

2 0 0 -

o ••

2 0 0 -

4 0 0 -1

• k

\

4 \

"\

+S. 4 ^

*̂ "' i 1 1

H11 [HI]

"**•.

T. 1 1 [- - i I I i

100 1 5 0 ZOO 2 5 0 3 0 0

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.

St,

4 0 0

2 0 0 -

- Z O O -

- • • O O -

- BOO -

^

H >

,.

HIIC1003

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].

<$ 12 n

a

E 0.6 08

*•> & 0.3

a

••• r i

%s^

*/

/ • oo ° o ...

•• I • • 1 - i 1

~ Rotation

-

Ellipticity

20 40 60

Iron thickness (nm)

80

- i 1 p 1

a - ••

Ellipticity

Rotation

0 01*-20 40 60

Iron thickness (nm)

80

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. A.K. Zvezdin, V.A. Kotov, Modern Magnetooptics and Magnetooptical Materials (IOP Publishing, UK, 1997).

6. N.P. Kolmakova, A.I. Popov, Physica B, 179, 19 (1992). 7. G.S. Krinchik, M.V. Chetkin, Sov. Phys. JETP, 13, 509 (1960). 8. M.I. Dolotenko, A.K. Zvezdin, G.G. Musaev et al., Phys. Solid State, 42,

726 (2000). 9. A.I. Pavlovskil, V.V. Druzhinin, O.M. Tatsenko, et al., JETP Lett. 31, 622

(1980). 10. A.A. Mukhin, V.V. Platonov, V.I. Plis, A.I. Popov, O.M. Tatsenko, A.K.

Zvezdin, Physica B, 246 -247, 195 (1998). 11. A.K. Zvezdin, S.V. Koptsik, G.S. Krinchik et al., Pis. Zh. Eksp.Teor. Fiz.

37,331(1983). 12. N.P. Kolmakova, R.Z. Levitin, A.I. Popov, N.P. Vedernikov, A.K. Zvezdin,

V. Nekvasil, Phys. Rev. B 40, 6170 (1990). 13. 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 low-noise 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=<Dn2/2k2L (1)

where <I>n is the spectral density of flux noise, k is the coupling coefficient and L is SQUID inductance. For the SQUID sensor to operate without hysteresis, it is essential that the hysteresis parameter (3 should be less than 1

P = 27iI0Rs2C/0>o<l (2)

where I0j Rs and C are the critical current, shunting resistor and capacitance for each Josephson junction respectively, and <I>0 is a quantum of magnetic flux. From theory of DC-SQUID [4], it is known that the minimum noise of the sensor can be achieved if its modulation parameter p, is close to 1:

P, = 2LVOOS 1 (3)

The same theory predicts that the coupled energy sensitivity in this case can be given by

Ec=5kbTAD0/k2I0Rs (4)

From the above equations, it follows that the optimization of SQUID parameters can be reduced to maximizing k2I0 Rs with both condition (2) and (3) fulfilled.

The coupling coefficient for a square Ketchen-type washer and a spiral coupling coil is given by

k2=l/[l+3.2SD(l/d+l/b)] (5)

where S is a ratio of turn step distance by the turn width, D is the height of the coil plane above the SQUID washer plane, b is the outer size of the SQUID washer and d the inner SQUID washer hole size [5]. For optimum coupling with a practical device, the input inductance should be about 1 u,H and therefore the number of turns of input coil should exceed 100, thus b » d . The inductance of the square washer is given by

L = 1 . 2 5 M (6)

where u0 is the magnetic permeability of free space [6]. From (5) and (6) we have the following expression:

472

k2 = l/(l+4Su0 D/L] (7)

Using (2), (3) and (7) with (3,=1 and S = 2, the following is true

k2I0Rs*l/{VL(l+8uoD/L)} (8)

Thus the maximum value of k2I0 Rs can be obtained when L= 8u0D or optimum SQUID inductance is determined by the input coil insulator thickness. With D=10 urn, the optimum SQUID inductance will be L=0.1 nH. From (3) optimum critical current can be calculated as I0 = 10 uA. For a typical Josephson junction capacitance of about 0.5 pF using (2) it is possible to obtain the value of the shunt resistor which will be Rs = 8 Ohm. This completes the theoretical optimization of all the SQUID parameters.

4. Fabrication of Multilayer Thin Film SQUID

All SQUID layers are fabricated using all-refractory materials and traditional Nb-AlOx-Nb Josephson thin film technology [7]. First the protection layer of 200 run of A1203 is sputtered to cover the whole substrate. Next, a 150-200 nm film of RF-sputtered Ti shunt resistor is applied. The configuration of the shunt resistors is formed by photolithographic masking and chemical etching. All subsequent layers of the SQUID structure are formed by lift-off technique. 200 nm of the base superconductor Nb electrode is deposited by DC-magnetron sputtering. After that, two insulation layers (each 270 nm thick) of SiO are thermally evaporated. The critical Josephson junction area of 3x5 um is defined by selective niobium etching and anodization process (SNEAP), where after plasma etching Might' anodization up to 10-25 V is carried out. Tunnel barrier is formed using thermal oxidation of RF-sputtered pure Al film (1.5 nm). The critical current of the SQUID can be controlled at this stage by varying oxygen pressure. In order to decrease the stray capacitance, one more layer of SiO (250 nm) is evaporated and the SNEAP photoresist pattern is used as a mask for this SiO layer lift-off (self-alignment contact). After RF sputter cleaning, the top Nb connector electrode is sputtered to complete the superconductive SQUID loop.

5. SQUID Magnetometry

Many types of magnetic investigations can be carried out using SQUID magnetometers and the optimal configuration of the measurement system depends on the nature of the signal to be measured and the environment in which measurements are being carried out. In order to fully exploit the extremely high measurement sensitivity offered by the SQUID, effective strategies need to be implemented for the reduction or cancellation of ambient magnetic "noise". There are numerous sources of naturally occurring or man-made magnetic fields

473

such as solar and ionospheric activity, thunder storms, trains, cars, transformers, electricity distribution cables etc. The combined amplitude of these magnetic fields is often many orders of magnitude larger than the magnetic signals of interest.

An effective solution for the ambient noise problem is magnetic shielding. Magnetically shielded rooms consisting of layers of high-conductivity materials (for eddy-current shielding of high-frequency fields) and high-permeability materials (for DC and quasi-DC shielding) are commercially available and are widely used in conjunction with large multi-channel SQUID systems.

Gradiometric detection i.e. the measurement of the spatial gradients of the magnetic fields in order to discriminate against distant sources is another commonly used technique. Gradiometers can be implemented by hardware (by winding pick-up coils which sense first, second or higher derivatives of the field), by software (by the appropriate linking of the signal from two or more sensors) or by a combination of these. Electronic noise cancellation, in which the ambient magnetic noise is independently monitored and subtracted from the measured signals, is now a standard feature in many SQUID systems [8]. A magnetic noise spectrum, measured in the laboratory with a second-order gradiometer SQUID, is shown in Fig. 1. As expected, the largest contribution is from mains (50 Hz) and its harmonics. In addition, the two features at 25 and 70 Hz are due to laboratory sources.

pT/Hz'

150 200 250

F requency (Hz)

Figure 1. Magnetic noise spectrum measured with a 2nd-order SQUID gradiometer (15 mm diameter, 50 mm baseline) in a typical urban environment.

474

SQUID sensors fabricated from high-temperature superconductors with transition temperatures exceeding 100 K can operate in liquid nitrogen or even in cryogen-free refrigerated systems. These HTc SQUIDs bring a new level of practicality but, because of their inferior performance, it is unlikely that they will ever totally replace conventional LTc SQUIDs. For example, typical coupled energy resolution of LTc thin-film DC SQUID sensors is of the order of 10"31

J/Hz which in practical applications translates to a field sensitivity of a few fT/VHz. The corresponding figures for the best currently available HTc SQUDs are approximately an order of magnitude larger, with the added disadvantage of much poorer low-frequency (below 1 Hz) performance. This is illustrated in Fig. 2 which shows the noise power spectrum of the measured difference in outputs from two identical HTc sensors during field tests. It can be seen that the 1 Hz noise level is of the order of 1 pT Hz""2 dropping to a white noise of around 42 fT Hz'"2. Figure 2 shows magnetic noise spectrum, measured in an open field, of two identical HTc SQUID sensors connected in series opposition.

Our research group at the University of Wales Swansea has been involved in SQUID-based instrumentation for a wide range of applications. We have developed miniaturized high-sensitivity SQUID systems capable of operation in typical hospital or laboratory environments as well as portable battery-driven SQUID magnetometers for fieldwork [9-13].

1 ( J U F ' 1 I I I I I I ) 1 1 | 1 1 I I I I l l | 1 | 1 1 TTTTTT

tE 1 0 -~r \

0.1 -

0.01 -

' • i • M ••! ' i ' i • i n l 1—i i,,i m i l i—i—i- i - i m J 1—i—i i i i ii

0.01 0.1 1 10 100 1000 Frequency (Hz)

Figure 2. Magnetic noise spectrum measured in an open field of two identical HTc SQUID sensors connected in series opposition.

475

6. Biomagnetism

Many physiological processes, such as neural activity or depolarization of the heart muscle, produce magnetic fields which can be measured over the body surface. These so-called biomagnetic fields (typically ranging from femtoteslas to picoteslas) are generally too weak to be detectable by conventional magnetometers, but are well within the capability of the SQUID. The diagnostic value of biomagnetic measurements is very similar (and sometimes superior) to that of the corresponding electrophysiological measurements which have been routinely used in clinical practice for decades through procedures such as electrocardiography (ECG) and electroencephalography (EEG).

At present, magnetocardiography (MCG) and magnetoencephalography (MEG), the magnetic counterparts of ECG and EEG respectively, are being actively researched in many centers worldwide and the results of this research are beginning to enter clinical practice. Diagnosis of heart block and transplant rejection, localization of arrhythmogenic tissue, management of focal epilepsy and treatment planning in brain surgery are just a few examples of the current medical applications of SQUID magnetometry. Other biomagnetic studies include non-invasive quantification of hepatic iron levels, studies of gastric function and assessment of lung contamination. In Fig. 3, typical recordings of the adult MCG measured with our miniature SQUID system are shown. The frequency bandwidth of 0.1 to 120 Hz was used and the signal-to-noise ratio was better than 60.

=-4—--4 iU-—I——.

JU—U—1~—U—J K""-—i

0 1 2 3 4 5 Time (s)

Figure 3. Adult MCG measured with a SQUID biomagnetometer in an unshielded environment in the frequency bandwidth 0.1-120 Hz.

476

Another interesting and unique application of biomagnetism which has been rapidly gaining popularity is fetal monitoring. Specially designed SQUID systems have been used to study the signals from the fetal heart and the fetal brain. The technique is totally safe and provides information, which cannot be obtained by any other method. Magnetic cardiac signals from a 35-week old fetus, measured with the same instrument as was used to obtain the recordings in the previous figure, are shown in Fig. 4.

-10

-15

0.5 1 1.5 2 Time (s)

Figure 4. Fetal MCG measured with a SQUID biomagnetometer in an unshielded environment in the frequency bandwidth 0.1-120 Hz.

The latest commercial SQUID systems contain hundreds of SQUID sensors making simultaneous measurements of the magnetic field at many locations. This arrangement allows the mapping of magnetic fields over extended areas in order to reconstruct the sources of these magnetic fields. This type of "magnetic source imaging" has been widely used in cardiac and brain function studies and the same approach, with suitably modified instrumentation, can be applied to other problems such as non-destructive testing and mineral prospecting.

7. Geophysics

Geophysicists have long been interested in magnetic surveys, both onshore and of the ocean floor, in order to identify magnetic anomalies which provide information about geological structures. At present most such surveys are carried out using fluxgate or proton magnetometers. However, the accuracy of such magnetic maps can be much improved by using the SQUID. In particular, recent advances in SQUID-based systems for electromagnetic prospecting (EMP) offer new possibilities in mineral exploration.

477

EMP is a technique for probing deep underground structures, mainly in the search for oil and other minerals. The technique is based on the reconstruction of the subterranean conductivity distribution from the electric and magnetic fields measured on the earth's surface. The measured electromagnetic fields are either of magnetotelluric (naturally occurring) origin, or generated by specially designed transmitters and in the latter case measurements can be carried out in the DC (steady-state) or transient response modes. Thus there exist a number of alternative strategies for the field application of this technique.

A major limitation of EMP is the severe attenuation of electromagnetic waves in the earth's crust which restricts the practical frequency range to a narrow bandwidth near DC. There are many sources of 'noise' in this bandwidth which result in a low signal-to-noise ratio. Hence the sensitivity of the detector becomes a critical parameter in the viability of this technique. Although SQUID magnetometers offer the ultimate sensitivity, because of the practical difficulties of operating them in the open field few attempts have been made so far to exploit their extreme sensitivity to improve the results achievable by conventional EMP instrumentation. Our group recently completed an EMP study at an onshore oil field using a specially designed SQUID magnetometer. We were able to map a hydrocarbon reservoir of only 100 m average thickness at a depth of 1200 m.

8. Other Applications and Concluding Remarks

The SQUID can measure magnetic field distributions with high spatial (down to microns) and temporal (down to microseconds) resolution. From these measured magnetic field patterns, the underlying magnetization or current density distributions (the "magnetic image") can be reconstructed. The size of the pickup coil and its distance from the source determine the spatial resolution of the image. In the case of a magnetized object, the amplitude of the measured field can be related to the magnetic susceptibility distribution within the object.

The varied capabilities of SQUID measurements are attracting a host of new applications. In non-destructive testing, for example of aircraft wings, eddy currents are induced in the metal structures to be imaged and the magnetic field distribution generated by these currents is analyzed to detect defects such as subsurface cracks. SQUIDs can also be used to detect the small currents flowing in electronic circuits and components thus enabling quantitative diagnostics, trouble-shooting and design improvements. Recently a scanning SQUID microscope, based on HTc sensors, has been described which can image magnetic fields of down to 10"10 T with a spatial resolution o f - 10 microns. Among other potential and proposed applications are archaeology, charting of underground water reservoirs and detection of buried ordnance.

478

The SQUID is arguably the most sensitive transducer of any type known at present. This quantum-scale sensitivity poses a real challenge for its practical use in everyday applications, particularly in hostile environments such as industrial sites. We are slowly but surely responding to this challenge and SQUID measurements will one day be as commonplace as the use of sensitive amplifiers.

Acknowledgments

The authors gratefully acknowledge the support of the Royal Society, EPSRC, Welsh Assembly Government and the University of Wales Swansea.

References

1. H. Seppa, T.Rahanen, IEEE Trans. Magn. MAG-23, 1083 (1987). 2. C. Carelli, V.Foglietti, IEEE Trans. Magn. MAG-19, 299 (1983). 3. M.B. Ketchen, IEEE Trans. Magn. MAG-17, 397 (1981). 4. V.V. Danilov, K.K. Likharev, O.V. Snegirev, in SQUID '80, edited by

H.D.Halbohm, 473, (1980). 5. S.S. Romanovitch, V.N. Sosnitski, report of Institute of Cybernetics,

Ukrainian Academy of Sciences 87-10, 1 (1987). 6. M.B. Ketchen, J.M. Jaycox, Appl. Phys. Lett. 40, 736 (1982). 7. V. Koshelets, A. Matlashov, I. Serpuchenko, L. Fillipenko, Yu. Zhuravlev,

IEEE Trans. Magn. MAG-25, 1182 (1989). 8. A.N. Matlashov, Y.E. Zhuravlev, A.Y. Lipovich, A.L. Alexandrov, E.

Mazaev, V.N. Slobodchikov, O. Vazhievski, in Advances in Biomagnetism, edited by S. J. Williamson et a!., 725, (1989).

9. A. Matlashov, Y. Zhuravlev, V. Slobodchikov, N. Bondarenko, A. Bakharev, D. Rassi, in Biomagnetism: Fundamental Research and Clinical Applications, edited by C. Baumgartner et al, 526, (1995).

10. J.A. Crowe, J.M. Herbert, X.B. Huang, N. Reed, M.S. Woolfson, D. Rassi, Y.E. Zhuravlev, S.J. Emery, Physiol. Meas. 16, 43 (1995).

11. Y.E. Zhuravlev, D. Rassi, S.J. Emery, in Biomagnetism: Fundamental Research and Clinical Applications, edited by C. Baumgartner et al., 700 (1995).

12. D. Rassi, Y.E. Zhuravlev, M.J. Lewis, S.J. Emery, in Biomag96: Proceedings of the Tenth International Conference on Biomagnetism, edited by C. J. Aine et al., 533, (2000).

13. G. Macmillan, Y.E. Zhuravlev, D. Rassi, P. A. J. de Groot, in Sensors and their Applications, edited by N. M. White and A. T. Augousti, 43, (1999).

DEPOSITION AND CHARACTERIZATION OF Fe/Si MULTILAYERS

S. KHARRAZI, S. ASHTAPUTRE, S. KULKARNI

Department of Physics, University ofPune, INDIA E-mail: skk@physics. unipune. ernet. in

R. CHOUDHARY, S. SHINDE, S. OGALE

Center for Superconductivity Research, University of Maryland, Maryland, USA

Multilayers of Fe/Si with different bilayer thicknesses have been deposited in Ultra High Vacuum (UHV) using electron beam evaporation technique. A number of techniques like X-ray Reflectivity (XRR), X-ray Diffraction (XRD), Vibrating Sample Magnetometer (VSM) and Magnetoresistance (MR) have been used to characterize the structural and magnetic properties of these multilayer samples. It was found that we could deposit reasonably coherent multilayer thin films of Fe/Si. However silicide formation at interfaces was detected which plays an important role in the resistivity of the samples. An analysis of the multilayer samples in terms of coupling through the silicide spacer layers is presented.

1. Introduction

Numerous studies have been carried out on observation of interlayer coupling through non-ferromagnetic spacer layers [1-9]. Fe/Si multilayer system is considered to be quite interesting. Silicon being a semiconductor is very sensitive to doping, defects and temperature. The coupling mechanism also may be different in metal/semiconductor multilayers than that observed in metal/metal multilayers. However the Fe/Si interface is very reactive and the silicide formation takes place in the early stages of interface formation [10-16]. Besides, there is a large variety of Fe-Si silicides like oc-FeSi2, (3-FeSi2 FeSi, Fe2Si, Fe3Si that may be formed at the interface. Some of these silicides may be metallic and some may not be. For example a-FeSi2 is metallic whereas (J-FeSi2

is a semiconductor with band gap of 0.84 - 0.87 eV [17]. y-FeSi2 on the other hand is a metastable phase. e-FeSi (with B20 structure) is a small - gap semiconductor (Eg ~ 50 meV) whereas FeSi with the CsCl (B2) structure [18-23] is metallic [17]. Fe2Si and Fe3Si are also metallic. They too are metastable and do not exist in the bulk form. However metastable phases can exist in thin films. Therefore Fe/Si multilayer system can be metal/metal or metal/semiconductor multilayer system depending upon the deposition conditions and interlayer thickness of silicon and iron.

479

480

The work on Fe/Si multilayers started with paper published by Toscano et al [24] who first observed the Anti-Ferromagnetic (AF) coupling in evaporated Fe/Si/Fe trilayers at low temperature. Subsequently many research groups have tried to investigate different aspects of Fe/Si system like observation of non-oscillatory AF coupling in sputtered Fe/Si multilayers and the photo-induced AF coupling at low temperature in Fe/Fe-Si samplesf 18,25-27]. Increase in magnetoresistance with decrease in temperature for samples with Si thickness of more than 15 A has been reported [28-31] and also observation of increase in coupling strength with decrease in temperature[32-36]. In refrences [19-23] the dependence of interlayer coupling on crystallinity of iron and iron silicide layer has been discussed and shown. Endo et al [37] reported oscillatory AF coupling of Fe layers in Fe/Si/Fe trilayers.

In general the nature of spacer layer inducing AF coupling depends on the method and deposition parameters as can be seen from very different results on the Fe/Si multilayer system. Broeder's group [32-36] found that the spacer layer does not show a clear semiconducting behavior but rather a metallic behavior, where as Mattson et al [26] reported semiconducting behavior of the spacer layer. Thus there has been a lot of work on Fe/Si multilayers but dependence of AF coupling on silicide formation and the role of deposition technique in particular phase formation is not well-known. We found therefore interesting to deposit the Fe/Si multilayers using electron beam evaporation technique and investigate the same.

2. Experimental

Si substrates were degreased in acetone and then dipped in 10% HF solution for 2 minutes to remove the native oxide layer on the surface. Substrates were transferred in an Ultra High vacuum (UHV) chamber. After obtaining the base pressure of 1 x 10"9 Torr,

Electron beam evaporation method was used to deposit the multilayer samples. Different samples of varying thickness of silicon were prepared. The thickness of the Fe layer was fixed to 30 A for all the samples. The rate of deposition was kept at 0.1 A0 / sec for both Fe and Si layers in all the depositions. All the samples were deposited with substrates at room temperature and no heat treatment was used during or after the deposition. Thus following samples viz. Fe (30 A) / [Si (t) / Fe (30 A)] x 20 / Si (111) where t = 8, 10, 12, 15, 18 and 20 A0 were prepared.

We used different structure and magnetic measurement techniques to investigate the Fe/Si multilayers. For structural characterization we performed X-ray Diffraction (XRD) and X-ray Reflectivity Measurements (XRR) on

481

SIEMENS - D500 X-ray diffractometer. Magnetoresistance (MR) measurements were carried out using IT Multimag from Magnetic Solutions Ltd. and cooled in Janis research Co. Inc. liquid He cryostat. Vibrating Sample Magnetometry (VSM) was performed using EG&G PAR 4500, with a magnetic field varying between 0 to 1.5 T and temperature varying between 80 to 300 K.

3. Results and Discussion

X-ray Reflectivity (XRR) and X-ray Diffraction were used for structural characterization of the multilayer samples. Fig. 1 shows the XRR measurements of samples with Si spacer layer of 12, 15, 18 and 20 A thickness. As can be seen in the figure, only the first Bragg peak appeared in the XRR, except in case of the sample with 15 A0 Si layer for which a second peak has been observed. Observation of only first Bragg peak indicates highly rough interface due to interlayer diffusion of Fe and Si. Using the Bragg formula we could calculate the bilayer thickness of different samples. Table 1 shows the comparison between experimental (calculated) and as deposited bilayer thickness, for different samples. Using software based on Parratt [38] formalism to fit the XRR, we found out that the spacer has a density, more than Si, which possibly is due to formation of Iron Silicide.

3 <

2 St c

F«30Si20

F*3QSi1S

F*30Si15

l W * 1 F*30Si12

0.0 0.5 1.0 1.5 2.0 2.5 3.0

0 (degree)

Figure 1. XRR of the samples with Si thickness of 12, 15, 18 and 20 A.

482

Table 1: Comparison of as deposited bilayer thickness with calculated one from XRR.

Bilayer Thickness(A) As Deposited

42

45

48

50

Bilayer Thickness(A) As calculated from XRR data

37

40

42

44

To further explore the interfacial details we carried out X-ray Diffraction (XRD) of the samples. Fig.2 shows the XRD of samples with Si layer thickness of 12, 15, 18 and 20 A. The diffraction peaks are present at 28.68 °, 45.62 °, 59.05 ° and 65.05 °. The peaks present at 45.62 ° and 65.05 °, in Fig 2 indicate the presence of Fe2Si (110) and (210) planes, at the interface, respectively. Unfortunately even though substrates were cleaned in HF, Si oxide peak could be observed. To confirm this, XRD of Si wafer, as it is and treated with HF are done. Presence of oxide on the surface is due to exposure of substrate to atmosphere (before deposition) for sometime, as it happened due to some technical problem.

Vibrating Sample Magnetometer (VSM) measurements performed at room temperature and low temperature (85 K.) for the samples with Si thickness of 12 and 18 A are shown in Fig.3 (a) and Fig.3 (b) respectively. As can be seen in the figure the remanent magnetization has decreased at low temperature for the samples with 12 and 18 A of Si thicknesses, indicating the stronger Anti-Ferromagnetic (AF) coupling of Fe layers at low temperature in these samples. For rest of the samples there is almost no difference between the magnetization loops at room temperature and low temperature.

Using the formula J = H sM stFe/4 erg/cm2 (1)

we can calculate the coupling strength. Here Hs is the saturation field, Ms is the saturation magnetization and tFe is the thickness of Fe layer. From this we find that sample with Si thickness of 12 A has the strongest AF coupling.

483

I

(0 <o„ w a

FO0SI20

FeWSIIS

Fe30SI12

SK111) wafer (HF treated)

; £1(111) wafer 1 (a* >t Is)

10 20 30 40 50 E0

2e (degree) 70 80 90 100

Figure 2. XRD of the samples with Si thickness of 12, 15, 18 and 20 A.

Further we compare the ratio of remanent to saturation magnetization. Fig.4 shows the change in Mr / Ms ratio with Si layer thickness. We observe that sample with 12 A Si layer has the minimum Mr/ Ms ratio.

We also carried out Magneto-Resistance (MR) measurements of the samples. Figure 5 shows the change in MR percentage with magnetic field at room temperature for the samples with Si layer thickness of 8, 10, 12, 15, 18 and 20 A. It can be seen that MR percentage for samples with Si thickness of 12 A and below is negative whereas it is positive for samples with Si thickness of 15 A and more. This implies that Fe layers are Anti-Ferromagnetically coupled when the Si thickness is 12 A or less where as they are not coupled when the Si thickness is 15 A or more and therefore the shape of magnetization loop for Fe30Sil8 sample is not actually due to AF coupling.

Out of the samples showing negative MR%, the sample with 12 A shows the largest value. Endo et al [39] observed a similar value and behavior of GMR in their Fe/Si multilayer samples.

Fig.6 (a) and Fig. 6(b) show the MR% at room temperature and low temperature (4.5 K) for samples with Si thickness of 8 and 12 A, respectively. As can be seen in the figure the MR% is larger at low temperature, which is in agreement with the VSM data, where remanent magnetization is smaller at low temperature. Both suggest the more contribution of AF coupling at low temperature.

484

0.008

0.004

A 0.002 •

E

0.000 -

-0.002 -

-0.004

-3000 -2000 -1000 1 • 1 • I

1000 2000 3000

H(Oe)

0.008

0.004 -

M 0.002 -

E

0.000

-0.002 -

F»(30A")M Si(iaA°)/F»(30A°»x20/Si(111) b)

-0.004 H 1 r I ' 1 ' 1 ' 1 " 1 •—1—

3000 -2000 -1000 0 1000 2000 3000

H(Oer)

Figure 3. The magnetization loop for samples with Si thickness (a)12 A and (b)18 A.

485

0.90-j.

0.85-

0.80-

g " 0.75 -

S" 0.70 -

0.B5-

0.B0-

0.55-

o.50-| • , • 1 1 • 1 • r 10 12 14 16 18 20

Si layer Thickness (A °) Figure 4. The M, / Ms ratio for samples with different Si layer thickness.

0.45-

-0.20 J 1 • | 1 1 • 1 • 1 • 1 — 0 2000 4000 6000 8000 10000

Field (Oe)

Figure S. The MR% of all different samples measured at room temperature.

486

0.00-

-0.05 -

^ -0.10.

-0.15.

-0.20 -

-0.25.

Fe(30 A") /{ Si(8 A0) / Fe(30 A°» x 20 / Si(111)

a)

• at a T. • at 4.5 °K

T «" 2000

— I — 4000

— I ' 1 • I 6000 8000 10000

H(Oe)

v?

0.00-

-0.05-

-0.10-

-0.15-

-0.20 -

-0.25

F«(30 A") / { Si(12 A°) / Fe(30 A°)} x 20 / Si (111)

b)

• at R. T.

•at4.5°K

— I 1 1 1 1 1 1 • 1 2000 4000 5000 8000 10000

H(Oe) Figure 6. The comparison of MR% at Room Temperature and at 4.5 K. for samples with Si thickness of (a) 8 A and (b) 12 A.

487

4. Conclusions

Fe/Si multilayers exhibit rough interfaces due to interlayer diffusion and silicide formation at the interface. Fe2Si phase has been detected at the interface. This phase is a metallic iron silicide. The sample with Si thickness of 12 A exhibits the largest MR%, which may be attributed to the strongest AF coupling between Fe layers for this spacer layer thickness and interface phase. The increase in the coupling strength with decrease in temperature was also noticed.

Acknowledgements

Part of this work was carried out at "Inter University Consortium for Department of Atomic Energy Facilities" (IUC-DAEF), Indore, India, for which we are thankful to Dr. S. M. Chaudhari and Mr. Satish Potdar. S. K. Kulkarni thanks UGC, India and S. Ashtaputre thanks DST, India for the financial support.

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27. E. E. Fullerton, S. D. Bader Phys. Rev. B 53, 5112 (1996). 28. K. Inomata, K. Yusu, Y. Saito, Phys. Rev. Lett. 74, 1863 (1995). 29. K. Inomata, K. Yusu, Y. Saito, Mat. Sci. Eng. B 31, 41(1995). 30. R. J. Highmore, K. Yusu, S. N. Okuno, Y. Saito, K. Inomata, J. M. M. M.

151,95(1995). 31. K. Inomata, S. N. Okuno, Y. Saito, K. Yusu, J. M. M. M. 156, 219(1996). 32. F. J. A. den Broeder, J. Kohlhepp, Phys. Rev. Lett. 75, 3026 (1995). 33. J. Kohlhepp, F. J. A. den Broeder, J. M. M. M. 156, 261(1996). 34. J. Kohlhepp, M. Valkier, A. van der Graaf, F. J. A. den Broeder,

Phys.Rev.B 55, R696 (1997). 35. H. Fredrikze, A. van der Graaf, M. Valkier, J. Kohlhepp, F. J. A. den

Broeder, Physica B 234-236,498(1997). 36. A. van der Graaf, M. Valkier, J. Kohlhepp, F. J. A. den Broeder, J. M. M.

M. 165, 157(1997). 37. Y. Endo, O. Kitakami, Y. Shimada, J. Appl. Phys. 85, 5741(1999). 38. L. G. Parratt, Phys. Rev. 95, 359 (1954). 39. Yasushi Endo, Osamu Kitakami, Yutaka Shimada, Appl. Phys. Lett. 72,

495 (1998).

SURFACE NONLINEAR MAGNETO-OPTICAL EFFECT IN ANTIFERROMAGNETICS

A.K. ZVEZDIN

Institute of General Physics, Russian Academy of Science, Vavilova St., 38, Moscow 119991, Russia

A.P. PYATAKOV, V.I. BELOTELOV

M.V. Lomonosov Moscow State University, Leninskie gori., MSU, Physics Department, Moscow 119899, Russia

V.A. KOTOV

Institute of General Physics, Russian Academy of Science, Vavilova St., 38, Moscow 119991, Russia

Based on the group theoretical analysis of the nonlinear surface electric susceptibility, we propose a new magneto-optical effect that can be used to detect antiferromagnetism at surface and in thin films. With the use of irreducible representation technique for the crystal space groups, the formulas for surface polarization on the second harmonic induced by the incident electromagnetic wave were derived. The principle possibility of antiferromagnetic domain visualization in zero external magnetic field was shown. As an example, the polarization dependence and azimuth angle dependence of the magneto-optical contrast for (001) surface of YFe03 were calculated.

1. Introduction

Nonlinear magneto-optical effects have received a lot of attention over the past years [1-14]. Surface nonlinear Kerr effects are of particular interest as they permit direct probing of the magnetization on the surface and buried interfaces with femtosecond time resolution. Surface nonlinear Kerr effects were considered in great number of papers both experimentally and theoretically [15-37]. So far the basic formulas for nonlinear Kerr effects were derived on the assumption that the optical properties of the adjacent media are isotropic and gyrotropy is described with the use of only magneto-optical parameter Q [32-34]. This standard assumption is good enough for metallic ferromagnetics (Fe, Co, Ni and their alloys) but not sufficient for anisotropic materials. In this context, the antiferromagnetic material whose magnetic structure and physical properties are very sensitive to the crystallographic symmetry are of great interest. In [3,4], volume nonlinear magneto-optical effects in rhombic antiferromagnets are considered. In these works the effects of second harmonic generation (SHG) were predicted in the crystals whose symmetry group has

489

490

space inversion operation but magnetic structure is odd relative to the space inversion. As for surface nonlinear Kerr effects in antiferromagnetics there was one paper devoted to NiO [38].

In this work, the nonlinear Kerr effects in antiferromagnetics with orthorhombic symmetry are considered to illustrate the role of crystal symmetry.

In regards to surface nonlinear magneto-optical effects, the most interesting are centrosymmetric materials whose crystallographic and magnetic structure are even relative to the space inversion. Orthoferrites (YFe03) and ortochromites (YFe03) can be regarded as centrosymmetric. Their class of symmetry and space group are mmm and D2h [39]. Rare earth orthoferrites and ortochromites (RFe03, RCr03) can also be related to this class of materials provided that the temperature is considerably high (T > 1 - 4 K) when rare earth subsystem is paramagnetic.

2. Theory

The surface second harmonic polarization psurJ = %^2)(M)E Ek can be presented as a sum of nonmagnetic

Pr=°iklX$Nm (la) and antiferromagnetic part

P , A F M = i b l k l m n K E < N ^ (lb) where Gk is antiferromagnetic vector, the unit vector TV is normal to the surface, Ek are the components of fundamental field that is solution of homogeneous equation for the medium occupying half-space z<0.

2

rotrotE{r) r eE(r) = 0 (2) c

where 0) - fundamental frequency, c — the velocity of light in free space, £ -dielectric tensor: £ = £0, z > 0 in vacuum, £ = £(J•, z < 0 - in medium.

This surface polarization is related with the surface layer that is thin enough to describe polarization in terms of delta function P{r) - Ps"r/(K)S(Z) but much greater than the lattice constant to make the notion of volume vector G meaningful.

To determine the structure of the tensors aMm and bMmn, the irreducible representations technique for space group D2h is used [39]. Table 1 shows all eight one-dimensional irreducible representations for this space group with their elements for three independent symmetry operations: / -space inversion, 2X, 2 - spiral axes of the 2-nd order. Components of polar vectors Pt and JV, and axial vector G, stay in the rows corresponding to the irreducible representation according to which they are transformed. In the next column stay the products of

491

these components EjEJNk, their position in the table rows is determined by the rules of multiplication for irreducible representations (Table 2).

Table 1. The table of irreducible representations of space group Lf2tl

r, r2 r3 r4 r5 r6 r7 r8

i l I l l

-l -l -l -l

2X

l l

-l -l I l

-l -l

2Y

-1

-1

-1

-1

vector components G, Gz

Gs

PX;EX;NX

*v> £*v > J * v

PZ,EZ;NZ

ExEyNz

EZEXNZ

EyEzNz

N:Ex:;NzEy

!;NzEz2

E,2NzGx;ExEyNzGy; EXEZNZGZ;

EyEzNzGx; EzExNzGy; ExEyNzGz;

EZEXNZG,; EvEzNzGv;EfNzGz

ExEyNzGx; E'NZ Gy; EyEzNzGz;

Table 2. Multiplication table for irreducible representations

r, r2 r3 r4 r5 r6 r7 r8

r, r, r2 r3 r4 r5 r6 r7 r8

r2 r2 r, r4 r3 r6 r5 r8 r7

r3 r3 r4 r, r2 r7 r8 r5 r6

r4 r4 r3 r2 r, r8 r7 r6 r5

r5 r5 r6 r7 r8 r, r2 r3 r4

r6 r6 r5 r8 r7 r2 r, r4 r3

r7 r7 r8 r5 r6 r3 r4 r, r2

r8 r8 r7 r6 r5 r4 r3 r2 r,

For example, the relation T5 = T6 x T7 x T8 leads to the fact that the product of polar vectors ExEyN: is transformed in accordance with T 5 . In the last column the components EjEjNfi, are arranged according to the multiplication table for irreducible representations.

The values that are related to the same irreducible representation are proportional to each other. This fact determines the structure of tensors alklm

and blklmn for polarization in (1) as will be shown later in Results. Having completed now the symmetry analysis of the susceptibility tensors it

is necessary to calculate the second harmonic response of the antiferromagnetic material. The second harmonic field E(2(Q) can be found from the inhomogeneous equation:

rotrotEir) - ^ - f J ( r ) = — ^ - Psurf5(z) (3) c £0

c

where S is nonlinear dielectric tensor, P — nonlinear surface polarization that described by formulas (la, lb).

492

The solution of the equation (3) can be found with the use of Green function technique [33, 34] and it can be represented as a sum of two terms, nonmagnetic

EM and antiferromagnetic E*FM :

E^{r) = --^QxV(2ik^d^K^(o,z,Q-)PvNM (4a)

£0 C

EfM (f) = - - %- exp( 2ikfv )dflv (2k, ,2co,z, 0-)PvAFM (4 b)

£Q C

where k,=(kx,ky,0) in-plane part of the wave vector, 7jj = (rx,ry,0) in-

plane part of the radius vector, d (2k^, 2co, z, z') - the Fourier transform of

the Green function (see Appendix). The magneto-optical contrast is determined by

s = IG+(2(O)-IG_(2(D)

IG+(2o)) + IG_(2co)

where IG+ and IG_ are intensities of the reflected light for opposite directions

of antiferromagnetic vector G . It can be easily shown that magneto-optical contrast can be found as

_ 2 Re(E™ (EfM+)' + E™ (EA/M*)' + {E™)' EA/M+ + (E™ )* EfM+) ( 6 )

\vNM\2 , \T?AFM+\2 , Ic-WMl2 . \T?AFM+\2 , \T?NM\2 , \T?AFM+\2

\hx | +\K I +\Ey | + | ^ I +\Kz | +\hz |

where EfFM+, EfFM~ - antiferromagnetic parts of the field for opposite

directions of antiferromagnetic vector G, the value with asterisk * means complex conjugate.

From (6) follows that magneto-optical contrast is nonzero when at least there is

one nonvanishing product E{ Ef in the numerator of the fraction. That

means that for existence nonzero magneto-optical contrast the presence of

nonvanishing EfFM is not sufficient, the corresponding nonmagnetic part

Et should be nonzero as well.

3. Results and Discussions

In the case considered JV = [001] the components of polarization are given by:

493

PxNU=alExEt,

PyNM = a2EyEz,

PZNM =aiE

2+a4E2

y+a5E2,

PfM = buEyEzGx+bnExEzGy+bnExEyG2,

PyAFM = b2lExEzGx + b22EyE:Gy + (b'23E

2x + b'2\E) + V£E)) Gz,

P2AFM =b3,ExEyGx+(bl2E

2x+b;2El+b-E2)Gy+bi3EyEzGz. (7)

These equations follow from table 1 where Nx=0, Ny=0, N^O. The components of polarization are proportional only to that combination of electric field and antiferromagnetic vector components that stay in the same row of table 1. For example, P™,, can be proportional only to the EXE, combination, and no polarization component correspond to ExEy combination.

In the yttrium orthoferrite the structure G = (+1,0,0) is realized. This leads to following formulas for the part of polarization sensitive to the antiferromagnetic structure:

PxAFM±=±buEyEz,

PyAFM±=±b2]ExE:,

PzAFM±=±b3]ExEy. (8)

It follows immediately from these formulas that for s-polarized light ( E. = 0,

see Fig. 1) the only nonvanishing antiferromagnetic part of the polarization

isPz . In nonmagnetic part only Pz =a3Ex +a^Ey is nonzero as

follows from (7). Therefore the SHG response of the media will be p-polarized

and contain both nonmagnetic E^ ( r ) and antiferromagnetic E^ (r) part

of the field. When the plane of incidence of s-polarized light is along OX or OY axis (Fig. 1) there will be no antiferromagnetic response at all.

For p-polarized light all components are nonzero unless the OX or OY axis belongs to the plane of incidence (Ey = 0 or Ex = 0, respectively). In the

case Ey = 0 the only nonvanishing antiferromagnetic component is P .It

is interesting to note that nonmagnetic component P = 0 . Therefore the s-

polarized SHG response is entirely antiferromagnetic. Nevertheless magneto-

494

optical contrast 8 (see Eq. (6)) is zero due to the fact that P = 0 and thus

ENyM = 0 .

Magneto-optical contrast 8 is nonvanishing only when the angle of polarization y/ is different from zero (p-polarization) and 90° (s-polarization) or when the azimuth angle <p (the angle where x-axis makes with the plane of incidence) is different from 0° and 90°.

This fact illustrated in Figures 2a and 2b. In Figure 2a the calculated polarization dependence of magneto-optical contrast 8 is shown. The ratio for

b,, nonlinear magneto-optical coefficient ——«0.01 was taken, angle of

a, i. j

incidence 45°. Figure 2b illustrates the calculated dependence of magneto-optical contrast on the azimuth angle cp for the case of p-polarized incident

light. It can be seen from Fig. 2b that the maximum of the absolute value of 8 is at cp - 45° and the contrast reverses (sign of 8 changes) when the plane of incidence forms right angle with x-axis.

/ / / / / / / / / / / / /

/ / / / / / / / / /

I I I I I |G| | | | \\\N\ I I I I I I I I I I I I I I I I I I I I I I I I I I I

Figure 1. The geometry under consideration.

495

8

7

6

s ? 5

•JO

4

3

2

1

0 0

• / \

"I \

10 20 30 40 50

V. deg 60 70

a)

80 90

8

6

4

2

1

•2

-4

-6

-8

•

\ 20 40 60 60 / 100 120 140

b)

160 180

<p, deg

Figure. 2. a) Polarization dependence for magneto-optical contrast in YFe03 (lj/ - angle of

polarization, ^ = 0° - p-polarization, ^ = 90" - s-polarization) and b) The dependence of magneto-

optical contrast on azimuth angle (p.

Nonzero magnetic contrast allows the optical detection of antiferromagnetism in nonlinear magneto-optical observation that is interesting in context of antiferromagnetic domain visualization. It should be remembered that in linear magneto-optical observation the presence of external magnetic field was necessary [40].

496

4. Conclusions

The symmetry analysis based on irreducible representation technique is applied for theoretical investigation of the nonlinear magneto-optical properties of antiferromagnetics with rhombic symmetry. The possibility of antiferromagnetic domain imaging in second harmonic observation is demonstrated. It is a new effect as in linear magneto-optical observation the presence of external magnetic field was necessary factor for visualization of the antiferromagnetism.

Acknowledgements

The work is supported by the RFBR grants (01-02-16595, 02-02-17389) and Federal Program "Physics of solids state nanostructures"

Appendix

Fourier transforms of Green functions for the plane of incidence lying in plane (XOZ):

-k B gu(kpk0,z,z') = , zH , expOV)exp(-/ygz'),

K[P + nK) gu(kll,k0,z,z') = Q,

k g13(Aj,,k0,z,z') = k{{ , : , exp(^.z)exp(-//?z'),

k^/l + nk.)

g2](kpk0,z,z') = 0,

g22 (\ ,k0,z,z') = - Qxp(ikzz) exp(-//?z'),

g22(kvk0,z,z') = 0,

k B g^k^z^') = -—r-L—— exp(/̂ z)exp(-̂ z'),

k0[P + nkz) gi2(k]l,k0,z,z') = 0,

-k2

g33 (A:,,, k0, z, z') = . " v exp(^.z) exp(-2*/?z'), kl[P + nkz)

497

where k0 = wave vector in vacuum, k^ = {kx,ky)- in-plane part of the 00

c

wave vector kQ, k,- normal component of the wave vector kQ,

I ") •? y P = Jn kQ - |̂| - in-plane part of the wave vector in a medium;

Fourier transforms of Green functions for general case (plane of incidence makes with X-axis angle (p)

. _k*8u+k2yg22 _kxky(gu-gn) _kxgu

" l l ~~ , 2 ' "12 - , 2 ' "13 ~ k2 U P li lr •>

^-^^-W^-^:' , , -^ , _ kyg3l

" 3 2 — , » " 3 3 — ^ 3 3 '

where kx = £0 sin 6 COS <p, ky = k0 sin 0 sin (p , 8 - angle of incidence.

References

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Weber, Phys. Rev. Lett. 87, 13 (2001). 15. P. Guyot-Sionnest, W. Chen, Y.R. Shen, Phys. Rev. B33, 8254 (1986).

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(1994). 21. H.A. Wierenga, W.De Jong, M.W.J. Prins, et al, Phys. Rev. Lett. 74, 1462

(1995). 22. B. Koopmans, M. Koerkamp Groot, Th. Rasing, H. Van den Berg, Phys.

Rev. Lett, 74, 3692(1995). 23. R. Vollmer, A. Kirilyuk, H. Schwabe et al., J. Magn. Magn. Mater.l4$,

295 (1995). 24. H.A. Wierenga, M.W.J. Prins, Th. Rasing, Physica, B204, 281 (1995). 25. T.A. Luce, W. Hiibner, K.H. Bennemann, Phys. Rev. Lett, 77, 2810

(1996). 26. T.M. Crawford, C.T. Rogers, T.J. Silva, Y.K. Kim, J. Appl. Phys. 81,

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FABRICATION AND CHARACTERIZATION OF Co/Cu/Co/NiO/Si(100) MAGNETIC MULTILAYER

A.Z. MOSHFEGH, P. SANGPOUR, O. AKHAVAN

Dept. of Physics, Sharif University of Technology, P.O. Box 11365-9161, Tehran, Iran E-mail: moshfegh@sharif.edu

G. KAVEI

Materials and Energy Research Center, P.O. Box 31787-316, Kara), Iran

A. IRAJI-ZAD

Dept. of Physics, Sharif University of Technology, P.O. Box 11365-9161, Tehran, Iran

In this investigation, we have fabricated Co/Cu/Co/NiO/Si(100) magnetic multilayer structure using a combinative sputtering-evaporation technique. Nickel Oxide with a thickness of about 30 nm was deposited on Si(100) substrate by thermal evaporation method. Following the oxide deposition, cobalt and copper films with a thickness of about 3 and 2 nm were grown by employing DC conventional and DC magnetron sputtering methods without breaking the vacuum, respectively. To understand the role of interfaces in this structure, we have studied surface of the Co/NiO, Cu/Co/NiO and Co/Cu/Co/NiO systems by atomic force microscopy (AFM) and sheet resistance (Rs) measurements. According to our data analysis the surface roughness of the top Co, Cu spacer and bottom Co were measured as 1.57, 0.56 and 1.52 nm, respectively. Based on potential dynamic corrosion measurements, the rate of corrosion for the top Co layer in H2O electrolyte was 0.22 mpy.

1. Introduction

With the discovery of the giant magnetoresistance (GMR) effect [1,2], many researchers have been focused their attention on maximizing GMR value for practical applications. Spin valves are important magnetic devices that provide higher GMR values with relatively low saturation fields [3]. A typical spin valve structure consists of two ferromagnetic layers, which are separated with a very thin nonmagnetic layer where one of those layers is pinned to an antiferromagnetic layer and the other is free.

Among the various antiferromagnetic materials used to pin the magnetization of an adjacent ferromagnetic layer, NiO is a possible candidate that offers the advantages of no current shunting, [4], high thermal stability and excellent corrosion resistance [5]. Thus, active GMR region (e.g. Co/Cu/Co) deposited on antiferromagnetic NiO layer has a great potential application for the use in magnetoreistance (MR), sensors and other devices [6-8]. Among many

499

500

multilayers investigated, the largest GMR effect has been observed in Co/Cu based structure [9]. However, this Co/Cu/Co structure has been extensively studied in the last few years, but surface properties of this structure needs further investigation.

Therefore, in this context, growth and surface characterization of each layer in the Co/Cu/Co/NiO/Si(100) structure have been investigated using atomic force microscopy (AFM) and electrical sheet resistance (Rs) techniques.

2. Experimental

The substrate used for this experiment was n-type Si(100) wafers with a typical resistivity of about 5-8 Q-cm and the dimension of 5x11 mm2. After a standard RCA cleaning procedure and a short time dip in a diluted HF solution, the substrates were dried in high purity N2 (99.999%) environment and then loaded into deposition system. The chamber with three sputtering targets and two evaporation boats was evacuated to a base pressure of about 4x 10"7 Torr prior to each run. We have used a combinative sputtering-evaporation technique to deposit the Co/Cu/Co/NiO/Si(100) multilayer structure. A rotating quartz crystal oscillator was used to monitor the thickness of the desired material in the structure. Schematic diagram of the deposition system is shown in Figure 1.

Negative Bias Voltage

• Water In/Out

Biased Substrate Holdei

Shuttei

i l l i l i r * Insulatoi r+ Grounded Substiate Shield

I—• Substrate Holder Si Substrates

Waiter in/out

W boat coated byAI2Oj

Figure 1. Relative position of the deposition system used to grow the Co/Cu/Co/NiO/Si(100) magnetic multilayer structure.

Nickel oxide (NiO) with a thickness of about 30 nm was deposited on Si(100) substrate using thermal evaporation method. Following the NiO deposition, cobalt layer with a thickness of about 3 nm and then copper films

501

with a thickness of about 2 nm was grown without breaking the vacuum by utilizing DC sputtering and DC magnetron sputtering methods, respectively.

To deposit nickel oxide layer, first we have used high purity NiO powder as a starting material. It was pressed and baked over night at 1400 C in an atmospheric oven. This process yielded a green solid disk suitable for evaporation deposition as examined and reported earlier [10]. A special W boat coated by A1203 was used as working evaporation boat to prevent reaction between NiO and W. The first cobalt layer was deposited on NiO layer by using DC conventional sputtering technique and then a copper thin film and second cobalt layer were deposited on the substrate using DC magnetron and DC sputtering technique, respectively. Before deposition of each layer, a pre-evaporation and a pre-sputtering process were performed for about 3 to 10 min. The other growth parameters for the fabrication of the Co/Cu/Co/NiO/Si(100) structure are summarized in Table 1.

Table 1. The growth parameters of the Co/Cu/Co/NiO/Si(100) structure.

Parameter

Deposition Method

Ar pressure (mTorr) Applied power (W)

Deposition rate (nm/sec)

Thickness (nm)

NiO

Thermal Evaporation

— 355 0.03

30

Co

DC Sputtering

70 41

0.01

3

Cu

DC Magnetron Sputtering

5 50 1.3

2

The deposited Cu, Co and NiO films were characterized by necessary techniques in order to study the property and quality of each layer in the Co/Cu/Co/NiO/Si(100) structure. The electrical properties of the films were determined by four-point probe sheet resistance (Rs) method at room temperature. The surface topography of the deposited layers was studied by AFM analysis using a Park Scientific Instruments (Auto probe CP-contact mode). Stability of samples against corrosion was measured by polarization method in H20 and NaCl(3%) media for Co and NiO layers, respectively by Pontentiostat/Galvano technique.

3. Results and Discussion

To investigate Co/Cu/Co/NiO/Si(100) structure, initially we have grown and characterized different subsystems involved in the multilayers structure namely

502

NiO/Si(100), Co/NiO/Si(100), Cu/Co/NiO/Si(100) and Co/Cu/Co/NiO/Si(100). In the following section, we will present and discuss our experimental results for each of the subsystem, separately.

To examine the degree of corrosion resistivity of the deposited NiO thin film, we have studied the corrosion properties of the NiO layer in NaCl (3%) electrolyte by using potential dynamic method. The corrosion rate of this layer in the NiO/Si(100) system was measured and compared with the Si(100) substrate under NaCl corrosion environment.

Figure 2 shows corrosion behavior of the NiO/Si(100) system. In this type of figure, the horizontal axis refers to applied current and the vertical axis represents the measured voltage. Based on our data analysis, the corrosion rate for the NiO layer and the bare Si(100) substrate was measured 7.05 and 8.71 mpy, respectively. These measurements support the chemical stability of NiO layer under our experimental conditions.

Log I/Area (A/Cm2)

Figure 2. The variation of voltage with surface current density for the NiO/Si(100) thin film system.

To study surface topography of cobalt and copper layers, we have utilized AFM method for both Co/NiO/Si(100) and Cu/Co/NiO/Si(100) systems. Figure 3 shows 3D-AFM micrograph of the Cu/Co/NiO/Si(100) multilayer structure in 2><2 um scales. A nearly smooth surface is observed for this system.

503

To quantify the surface roughness, we have used the following expression:

*&>*)=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

0

-0.5, , , , r— , , -3 -7 -5 -3

Log I/Area (A/cm2) 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)

Sheet Resistance (Q/D)

433 131 124

Surface Roughness (nm)

1.57 0.56 1.52

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.

* *~w\*w.r.

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. B. Dai, J.N. Coi, W. Lai, J. Magn. Magn. Mat. 27, 19 (2003). 2. H.W. Jiang, M.H. Li, G. H. Yu, J. Magn. Magn. Mat. 242-245, 341 (2002). 3. B. Dieny, J. Magn. Magn. Mat. 136, 335 (1994) 4. 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: pg@innovent-jena.de

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 single-and/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 (ferro-or 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 Random-Access 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.

50-

40-

30

20-

10-

0-

i 1

•

i 1 i

, __ ., -

-r 1 ,

j ! !

..., : — T r - — , • 1

-20 0 20

field (Om)

soft ferromagnet

hard ferromagnet

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 = ^ext ""• t-an "*" texch "*" E<)eIn ( 1 )

with Eext being the energy in an external field, Ean - magnetocrystalline anisotropy energy, Ecxch - exchange energy, and Edem - magnetodipolar interaction energy (stray field or demagnetising energy). These energy contributions can be written in finite-difference approximations as

E e x t = - ? M i H i X t A V / (2)

with Mj as magnetisation in the cell i, H° x - external magnetic field over the cell i, AVj - cell volume given by the product of lateral dimensions of the discretisation cell (in the order of 5 to 10 run) and the film thickness,

pun _ v ^ u r ban - I N l - ( m i n i )

AV, (3a)

for uniaxial anisotropy with n, as unit vector of the anisotropy axis and mj = Mi/Mi,

E ^ I K - f e y +P'xPi +PjyPijAY (3b)

for cubic anisotropy with pxyz as unit magnetisation vector in a local coordination system,

Eexch = I J i j ^Vj .AVj .K^J ( 4 )

with Jy as exchange coefficient, Ky - exchange weakening, ay - angle between Mj and Mj, and

510

E . =- l -SM.W.# m M i (5) dera 2 j i y J

with Wy as interaction coefficients between cells i and j .

4 Switching Behaviour of Single-Layers and MEAM Elements

At first we studied the switching behaviour of a soft magnetic element and found no qualitative difference between the switching in homogeneous and khomogeneoes magnetic fields generated with crossed conductors shown in Fig. 2. For this reason the most simulations were performed with Permalloy to homogeneous magnetic fields and typical lateral dimensions of 200 to 800 nms

film thicknesses of several nm5 and mean grain size of 20 nm. The axes of the cubic magnetocrystaliine anisotropy of Permalloy grains were randomly distributed in the film plane.

4.1. Single-Layers

4 A A. Magnetic Domains in Remanent State

Sometimes people believe that magnetic thin films with lateral dimensions below ~ 500 nm and thicknesses of several nm exist only in the single domain state. Computer simulations of such thin Permalloy (Ni80Fe2o) films demonstrate, however? the existence of different types of magnetic domains. After saturation of the films in positive z-direction so-called C-direct, C-inverse, S-direct, and S-inverse states occurred due to the formation of closure domains as illustrated in Fig. 3. All these four structure types appear for various geometries: rectangular with and without rounded corners or elliptical

C - dir C - IHV S - dir S - inv Figure 3. Different remanent states for Permalloy with lateral dimensions 150 x 300 nm2 and thickness 2 nm.

4'.l.,2. Magnetisation Processes

To study magnetisation processes we assumed a small and remanent positive x-component Hx of the magnetic field. The z-component was varied between +Hmax and ~-Hmax. As expected, the magnetisation process qualitatively depends

511

on the remanent state of the soft magnetic film. In the case of the S-state (c.f. Figs. 3 and 4) the switching process of the magnetic element is relatively homogeneous. The switching fields are much less compared with the C-states in Fig. 5. This Figure clearly shows a 360°-doraain wall and a "strange" hysteresis loop. Therefore9 S-states should be preferred in MRAM's, where switching fields less than ~ 100 Oe are necessary.

The comparison of rectangular and elliptical Permalloy dots results in the conclusion that the elliptical shape is the best one. Accordingly, we deal in the next section only with magnetisation processes in elliptical MRAM elements.

Figure 4. Magnetisation of Permalloy with lateral dimension 300 x 150 nm2 and thickness 2 nm starting from S - stite ! ! „ * 450 Oe.

Figure 5. As Fig. 4, but for C - state Hmax * 900 Oe.

512

4.2. Magnetisation Processes of Elliptical MRAM Elements

Figure 6 shows the switching behaviour of an elliptical MRAM element ax b = 400 x 200 nm2 consisting of 2 nm thick Permalloy and 1 nm thick hard magnetic film with uniaxial anisotropy constant Ku = 105 erg/cm3 and the easy axis parallel to a-axis. The distance between the films is 4 nm and it is assumed that there is no ferromagnetic or antiferromagnetic interlayer coupling. Therefore, the magnetic interaction between the layers is caused only by stray fields.

-0.8 J

Figure 6. z-component of reduced magnetisation mz as a function of the magnetic field Hz (see text above for the parameters)

Figure 6 shows a switching field of roughly 20 Oe, which corresponds to the technical demands. It is interesting to note that the simulations reveal a dependence of the switching field on the uniaxial anisotropy constant of the hard layer - increasing anisotropy leads to decreasing switching field. This surprising effect can be easily explained by the appearance of closure domains at lower anisotropy.

513

References

1. D.V. Berkov, S.V. Meshkov, Sov. Phys. JETP 67, 2255 (1988). 2. D.V. Berkov, S.V. Meshkov, Hysteresis, IEEE Trans. Magn. MAG-23,

1804(1990). 3. D.V. Berkov, Phys. Rev. B B53, 731 (1996). 4. D.V. Berkov, J. Magn. Magn. Mat. 99, L7-11 (1991). 5. D.V. Berkov, K. Ramstock, A. Hubert, Phys. Stat. Sol. (a) 137, 207

(1993). 6. D.V. Berkov, K. RamstSck, T. Leibl, A. Hubert, IEEE Trans. Magn.

MAG-29, 2248 (1993). 7. K. Ramstock, A. Hubert, D. Berkov, IEEE Trans. Magn. MAG-32, 4228

(1996). 8. D.V. Berkov, J. Magn. Magn. Mat. 161, 337 (1996). 9. D.V. Berkov, N.L. Gorn, Phys. Stat. Sol. (a) 161, R7-8 (1997). 10. D.V. Berkov, N.L. Gorn, Phys. Rev. B, B57, 14332 (1998). 11. D.V. Berkov, J. Magn. Magn. Mat. 111,327(1992). 12. D.V. Berkov, J. Magn. Magn. Mat. 117, 431 (1992). 13. D.V. Berkov, N.L. Gorn, Micro Magus: package for micromagnetic

simulations: http://www.micromagus.de. 14. H. Brilckl, Vortrag AK, Magnetische Schichten fur technische

Anwendungen, (Kahl/Main, 14.06.2001).

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CONFERENCE PHOTOS

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;*v

: ^ ^

huiicipants of the IWTF2003 in IVonl of Physics Department building, Sharif University of Technology, Tehran, Iran.

IWTF2003 Workshop view

517

518

IWjr.'OO.* Workshop view . * > %

• • • • • • • i i i i

Iwri'.VW).̂ Woikshop view

519

S. Sohrabpour ai llur opening A.Z. Moshiegh a! the opening,

M. C)hiin«i chainnv the opening session S.C. Kasliyyp

M. Fade (chair) & H.v. Kane! (lecturing) M. Wuttig

520

A.I. Popov (chair) & M. l;arle (lecturing) M. Ohrin"

S.C. Kashyap (chair) & J.G. Lin (lecturing) S.K. Kulkarni

X Rassi (chair) & A. Iraji-/ad (lecturing) B. Rashidian

521

D. Acosta (chair) & N. Radic (lecturing) M. Vesaghi

f j ' U i i M ^ v a ^ N,A. Kh«m (k\ iuriii") M.C\ Ahdulria-i

S.K. Kulkami (chair) & D. Acosta (lecturing) A.N. Lykov

522

Participants at the thin film laser deposition lab., Physics Dept, SUT.

Participants at the thin film sensors lab., Physics Dept., SUT.

ESC A/A ES surface analysis lab. Physics Dept, SUT.

Thin film sputtering lab., PhysicsDept.. SUT

Participants at the thin film Participants"at the thin film chemical sputtering lab., Physics Dept., SUT. deposition lab., Physics Dept., SUT.

523

• 3

Participants at the exhibition It c

Participants at the exhibition

Participants in front of conference hall Participants at the lunch

Participants at the lunch Participant at the goodbye Party

524

At. the 'Mtvsins? session (from left to ri»hi: A.Z. Miv;h»coh. M. Ohriiig. u.v. KlincL E G MiiiKii-ssion and A.L Popov).

Participants ai !h«, luuch nmin:1

Tehran si^hise<*nh? SOU:.

Puneipanis a! si«hiv:cinti lour in the ea>l of Tehran.

Participants in the Tehran sightseeing tour, the Garden, Saadabad Palace.

Participants in front of Tehran's Anthropology museum.

Participants in the post workshop tour, in front of Alighapou Palace (from left to right: M. Hartmanova, A.I. Popov, K.C. Clarissa, Z. Chvoj, J.G. Lin and H.v. KMnel).

526

P^r iillil f§§

1 - . ^ l l l l i l ^ <

Sio-Seh Pol view, Isfahan. /¥

Participants in iheposl workshop tour in front of Zaynndeh-roud River, Isfahan.

527

K>si workshop lour al Zoroaslriaii Ziggurat (from lolt io ri«hi: Z. Thvoj, M. Ohring, A.Z. Moshfegh, P.G. Soukiassian, A.I. Popov and i l.v. KJind).

H / i

S i i JiL^^^ft1

Participants in the post workshop tour, Persepolis, Shira/.

528

Participants at Takhte-Rostam, Shiraz. The entrance of Persepolis, Shiraz.

Persepolis, Shiraz Persepolis, Shiraz

tie* £W<^Y*%>*^*§vr *

^ ^ ^ ^ ^ ^ ^ *****

Participants in Lhc post workshop tour at Drain Garden.- Shiraz.

^ ^pi^

AUTHOR INDEX

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531

Abdulrida, M.C.

Acosta, D.

Ahadian, M.M.

Akhavan, M.

Akhavan, 0 .

Ashtaputre, S.

Ataev, B.M.

Bagamadova.A

Bdikin, I.K.

Belotelov, V.I.

Berkov, D.V.

Car, T.

Chaudhary, R.

Chromik, S.

Chvqj, Z.

Drese, R.

Farle, M.

Ghoranneviss, M.

Gmucova, K.

Gondarilla, F.C.

Gorn, N.L.

Gornert, P.

Hamed, H.A.

Hantizadeh, M.

Hartmanova, M.

Hassan, B.A.

Henini, M.

Jadkar, S.R.

355

363

198

428

499

479

374, 380

374, 380

265, 405

489

507

101

149

158

187

297

441

198

158

158

507

507

355

198

158

355

244

344

Jergel, M.

Ihara, H.

Iraji-zad, A.

Iykov, J.

Kameli, P.

Kamilov.I.K

Kashyap, S.C.

Kanel, H.v.

Kavei, G.

Kedrov, V.V.

Keshmiri, S.H.

Khan, N.A.

Kharrazi. S.

Kikineshi, A.

Komissinski, P.V.

Kotelyanskii, I. M.

Kotov, V.A.

Kulkarni, S.K.

Kundracik, F.

Kuruvilla, B.A.

Lin, J.G.

Liu, X.

Lopez, A.

Luby, S.

Luches, A.

Lykov, A.N.

Mahdavi, S.M.

Majkova, E.

158

414

198, 337, 499

101

428

374,380

228

54,70

499

265

306, 324

414

479

318

405

405

489

129,149

158

149

387

285

363

180

180

420

337

180

532

Makhmudov, S.Sh.

Mamedov, V.V.

Martinez, A.I.

Matveev, D.V.

374, 380

374, 380

363

265

Metikos-Hukovic, A. 101

Michely, T.

Minkina, V.G.

Mirsalehi.M.M

Mohamed, S.H.

Mohite, K.C.

Moshfegh, A.Z.

Mozhaev, P.B.

Navratil, K.

Navratil, V.

Ohring, M.

Omaev, A.K.

Ovsyannikov, G.A.

Pawar, B.N.

Payami, M.

Phase, D. M.

Pogoryeov, Y.A.

Popov, A.I.

Purandare, R.

Pyatakov, A.P.

Radic, N.

Rassi, D.

Rezaesmaeili, M.

Razi, F.

285

119

306

297

344

11,28,499

405

158

158

171

374,380

405

344

271

149

256

454

149

489

101

469

337

337

Salamati, H.

Sangpour, P.

Soukiassian, P.G.

Stubicar, M.

Strukova, G.K.

Strukov, G.V.

Takwale, M.G.

Tonejc, A.

Varghese, G.

Wakkad, M.M.

Wuttig. M

Zemek, J.

Zever'kov, S.A.

Zhuravlev, Y.E.

Zvezdin, A.K.

428

499

85,213

101

265

265

344

101

205

297

285, 297

158

265

469

489