Scanning Probe Microscopy: Characterization, Nanofabrication and Device Application of Functional Materials: Proceedings of the NATO Advanced Study Institute ... II: Mathematics, Physics and Chemistry)

Scanning Probe Microscopy: Characterization, Nanofabrication and Device Application of Functional Materials

NATO Science Series A Series presenting the results of scientific meetings supported under the NATO Science Programme. The Series is published by IOS Press, Amsterdam, and Kluwer Academic Publishers in conjunction with the NATO Scientific Affairs Division Sub-Series I. II. III. IV. V.

Life and Behavioural Sciences Mathematics, Physics and Chemistry Computer and Systems Science Earth and Environmental Sciences Science and Technology Policy

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The NATO Science Series continues the series of books published formerly as the NATO ASI Series. The NATO Science Programme offers support for collaboration in civil science between scientists of countries of the Euro-Atlantic Partnership Council. The types of scientific meeting generally supported are “Advanced Study Institutes” and “Advanced Research Workshops”, although other types of meeting are supported from time to time. The NATO Science Series collects together the results of these meetings. The meetings are co-organized bij scientists from NATO countries and scientists from NATO’s Partner countries – countries of the CIS and Central and Eastern Europe. Advanced Study Institutes are high-level tutorial courses offering in-depth study of latest advances in a field. Advanced Research Workshops are expert meetings aimed at critical assessment of a field, and identification of directions for future action. As a consequence of the restructuring of the NATO Science Programme in 1999, the NATO Science Series has been re-organised and there are currently Five Sub-series as noted above. Please consult the following web sites for information on previous volumes published in the Series, as well as details of earlier Sub-series. http://www.nato.int/science http://www.wkap.nl http://www.iospress.nl http://www.wtv-books.de/nato-pco.htm

Series II: Mathematics, Physics and Chemistry – Vol. 186

Scanning Probe Microscopy: Characterization, Nanofabrication and Device Application of Functional Materials edited by

Paula Maria Vilarinho University of Aveiro, Portugal

Yossi Rosenwaks Tel Aviv University, Ramat - Aviv, Israel and

Angus Kingon NCSU, Raleigh, NC, U.S.A.

Kluwer Academic Publishers Dordrecht / Boston / London Published in cooperation with NATO Scientific Affairs Division

Proceedings of the NATO Advanced Study Institute on Scanning Probe Microscopy: Characterization, Nanofabrication and Device Application of Functional Materials Algarve, Portugal 1–13 October 2002 A C.I.P. Catalogue record for this book is available from the Library of Congress.

ISBN 1-4020-3018-5 (PB) ISBN 1-4020-3017-7 (HB) ISBN 1-4020-3019-3 (e-book)

Published by Kluwer Academic Publishers, P.O. Box 17, 3300 AA Dordrecht, The Netherlands. Sold and distributed in North, Central and South America by Kluwer Academic Publishers, 101 Philip Drive, Norwell, MA 02061, U.S.A. In all other countries, sold and distributed by Kluwer Academic Publishers, P.O. Box 322, 3300 AH Dordrecht, The Netherlands.

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All Rights Reserved © 2005 Kluwer Academic Publishers No part of this work may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, microfilming, recording or otherwise, without written permission from the Publisher, with the exception of any material supplied specifically for the purpose of being entered and executed on a computer system, for exclusive use by the purchaser of the work. Printed in the Netherlands.

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TABLE OF CONTENTS

PREFACE ON THE NATO ASI ON THE BOOK ACKNOWLEDGEMENTS PHOTO OF THE GROUP ADDRESS LIST OF THE AUTHORS LIST OF PARTICIPANTS

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Part I – Fundamentals of Functional Materials FUNCTIONAL MATERIALS: PROPERTIES, PROCESSING AND APPLICATIONS P.M. Vilarinho 3 SCALING OF SILICON-BASED DEVICES TO SUBMICRON DIMENSIONS A.I. Kingon

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UNSOLVED PROBLEMS IN FERROELECTRICS FOR SCANNING PROBE MICROSCOPY J.F. Scott 51

Part II – Fundamentals of Scanning Probe Techniques PRINCIPLES OF BASIC AND ADVANCED SCANNING PROBE MICROSCOPY 77 D.A. Bonnell and R. Shao NANOSCALE PROBING OF PHYSICAL AND CHEMICAL FUNCTIONALITY WITH NEAR-FIELD OPTICAL MICROSCOPY L.M. Eng 103 NANOSCALE ELECTRONIC MEASUREMENTS OF SEMICONDUCTORS USING KELVIN PROBE FORCE MICROSCOPY Y. Rosenwaks and R. Shikler 119 EXPANDING THE CAPABILITIES OF THE SCANNING MICROSCOPE K.F. Kelly, Z.J. Donhauser, B.A. Mantooth and P.S. Weiss

TUNNELING 153

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FUNCTIONS OF NC – AFM ON ATOMIC SCALE S. Morita, N. Oyabu, T. Nishimoto, R. Nishi, O. Custance, I. Yi and Y. Sugawara 173

Part III – Application of Scanning Techniques to Functional Materials SCANNING PROBE MICROSCOPY OF PIEZOELECTRIC AND TRANSPORT PHENOMENA IN ELECTROCERAMIC MATERIALS S.V. Kalinin and D.A. Bonnell 199 SFM-BASED METHODS FOR FERROELECTRIC STUDIES A. Gruverman

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SCANNING TUNNELING SPECTROSCOPY: LOCAL DENSITY OF STATES AND SPIN DISTRIBUTION OF INTERACTING ELECTRON SYSTEMS M. Morgenstern 251 NANOINSPECTION OF DIELECTRIC AND POLARIZATION PROPERTIES AT INNER AND OUTER INTERFACES IN FUNCTIONAL FERROELECTRIC PZT THIN FILMS 275 L.M. Eng MICROSCALE CONTACT CHARGING ON A SILICON OXIDE S. Morita, T. Uchihashi, K. Okamoto, M. Abe and Y. Sugawara

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CONSTRUCTIVE NANOLITHOGRAPHY S.R. Cohen, R. Maoz and J. Sagiv

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NANOMETER-SCALE ELECTRONICS AND STORAGE K.F. Kelly, Z.J. Donhauser, P.A. Lewis, R.K. Smith and P.S. Weiss

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Part IV – Contributed papers STM TIPS FABRICATION FOR CRITICAL DIMENSION MEASUREMENTS A. Pasquini, G.B. Picotto and M. Pisani

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SCANNING PROBE MICROSCOPY CHARACTERIZATION OF FERROELECTRICS DOMAINS AND DOMAINS WALLS IN KTiOPO4 363 C. Canalias, R. Clemens, J. Hellström, F. Laurell, J. Wittborn and H. Karlsson IMAGING LOCAL DIELECTRIC AND MECHANICAL RESPONSES WITH DYNAMIC HETERODYNED ELECTROSTATIC FORCE MICROSCOPY D.R. Oliver, K.M. Cheng, A. Pu, D.J. Thomson and G.E. Bridges 371

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AFM PATTERNING OF SrTiO3-δ THIN FILMS AND DEVICE APPLICATIONS L. Pellegrino 387 NANOSCALE INVESTIGATION OF A RAYLEIGH WAVE ON LiNbO3 J. Yang and R. K Koch

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SCANNING CAPACITANCE FORCE MICROSCOPY AND KELVIN PROBE FORCE MICROSCOPY OF NANOSTRUCTURES EMBEDDED IN SiO2 G. Tallarida, S. Spiga and M. Fanciulli 405 ELECTRICAL CHARACTERISATION OF III-V BURIED HETEROSTRUCTURE LASERS BY SCANNING CAPACITANCE MICROSCOPY 413 O. Douhéret, K. Maknys and S. Anand PROBING THE DENSITY OF STATES OF HIGH TEMPERATURE SUPERCONDUCTORS WITH POINT CONTACT TUNNELING SPECTROSCOPY L. Ozyuzer, J.F. Zasadzinski, N. Miyakawa and K.E. Gray 425 ANNEALING INFLUENCE ON CO ULTRATHIN FILM MORPHOLOGY IN MBE GROWN Co/Au BILAYERS A. Wawro, L.T. Baczewski, P. Pankowski, P. Aleszkiewicz, M. Kisielewski, I. Sveklo and A. Maziewski 435 CORRELATION BETWEEN THE SURFACE RELIEF AND INTERFACES STRUCTURE OF Fe/Cr SUPERLATTICES AND ELECTROMAGNETIC WAVES PENETRATION A.Rinkevich, L.Romashev and V.Ustinov 443 MAGNETORESISTANCE AND MICROSTRUCTURE OF MAGNETIC THIN FILM MULTILAYERS J. Neamtu, M. Volmer 449 SPM INVESTIGATION OF THIOLATED GOLD NANOPARTICLE PATTERNS DEPOSITED ON DIFFERENT SELF-ASSEMBLED SUBSTRATES 457 F. Sbrana, M. T. Parodi, D. Ricci and E. Di Zitti AFM OF GUANINE ADSORBED ON HOPG UNDER ELECTROCHEMICAL CONTROL A. M. Chiorcea and A.M. Oliveira Brett 467 DYNAMICS IN MODEL MEMBRANES AND DNA-MEMBRANE COMPLEXES USING TEMPERATURE CONTROLLED ATOMIC FORCE MICROSCOPY Z.V. Leonenko and D.T. Cramb 475

INDEX

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PREFACE Today, a wide range of analytical techniques can be used for materials research. The most commonly used high-resolution surface analysis techniques are Scanning Electron Microscopy (SEM) and Scanning Probe Microscopy (SPM). Although both techniques resolve surface structure down to the nanometer scale, the different image formation mechanisms result in different types of information about the structure of the surface, making these techniques complementary. In SEM, an electron beam, guided by a complex array of lenses, interacts with the sample and electrons are emitted from the sample, either as back-scattered or secondary electrons. Secondary electron emission is commonly used for surface morphology analysis. The first SEM equipment was constructed between 1938 and 1942. Improvements in resolution and the development of several types of detectors for local compositional analysis (X-rays, Auger electrons, backscattered electrons, cathodoluminesce) took place since then. Compared to electron microscopy, SPM techniques are quite new. Scanning Tunneling Microscopy (STM), the earliest of the SPM techniques, was invented in 1981 at IBM Zurich Research Laboratory by G. Binnig and H. Roher. STM was the first instrument to generate three dimensional real-space images of surfaces with atomic resolution. This invention earned them the Nobel Prize in Physics in 1986. In STM a sharp conducting tip, with a bias voltage applied between the tip and the sample, scans the surface of the sample. The scanning motion is at the angstrom level and the tip does not contact the sample. The resulting tunnelling current, exponentially dependent on the tip – sample spacing, is the signal used to create the STM image. Since 1981, a large family of SPM related techniques, based on various types of interactions between the tip and the sample, has been developed. It has been demonstrated that the SPM approach allows manipulation of single atoms and molecules. Various SPM techniques such as atomic force microscopy (AFM), magnetic force microscopy (MFM), electrostatic force microscopy (EFM), scanning capacitance microscopy (SCM), near-field scanning optical microscopy (NSOM) and others were proved to be capable of measuring the local physical properties of materials with nanoscale resolution. As a consequence, presently there is an explosion in the application of SPM techniques in a wide spectrum of fields of science, ranging from condensed matter physics, chemistry and materials science, to medicine and biology. Currently, SPM is widely used for nanoscale characterisation of materials by using mechanical, electrical, magnetic, optical and chemical interactions between the probing tip and the surface. Compared to electron microscopy techniques (SEM, TEM, HRTEM, etc), SPM is a low cost analytical method in terms of equipment, maintenance, accessories and sample preparation. In addition, it can operate in ambient m environment forbidden to electron microscopy. Consequently, it is expected that a growing number of university research groups and R&D industry divisions will acquire such equipment for research, quality control and fabrication. At the same time, the role of SPM in the field of nanotechnology is also growing in importance. The continuous need for the miniaturisation of electronic devices, with improved speed and functionality, has been a ix

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constant driving force pushing towards the development and manipulation of nanoscale devices. However further scaling of digital electronic devices necessitates fabrication and application of materials with nanoscale features. In this sense, SPM is becoming an indispensable tool, playing a key role in nanoscience and nanotechnology. One of most rapidly evolving, yet relatively unknown, fields of material science and of functional materials is the field of ferroelectric thin films. These materials possess a unique set of physical properties, such as switchable polarization, piezo- and pyroelectricity and high nonlinear optical activity, which make them extremely attractive for a number of applications. Over the last 10 years there have been serious efforts to develop ferroelectric memories, which combine nonvolatility with high-speed access, almost unlimited endurance and extreme radiation hardness. Due to these significant advantages of ferroelectric memories, it is expected that they will continue to replace other types of nonvolatile memory systems in many applications. In addition, ferroelectric materials can be used in a variety of other devices that exploit their unique properties, such as piezoelectric transducers and actuators, t infrared sensors, optical switches and computer displays. Recent advances in the processing of high quality ferroelectric films resulted in development of 4 Mb nonvolatile ferroelectric random access memories (NVFRAMs) at Samsung and Matsushita. However, the tremendous potential of ferroelectric films is far from being realized, as further developments in this area are hindered not only by the integration issues related to the present state of the NVFRAM technology, but also by a lack of fundamental knowledge related to reliability, performance and scaling of ferroelectric devices. Integration of ferroelectric thin films into Gigabit memory devices requires substantial improvement in the understanding of the properties and device physics of these materials, and this in turn requires the implementation of new tools suitable for in situ testing of ferroelectric nanostructures. One of the most promising approaches is based on using scanning probe techniques. Recently, SFM has been successfully applied for nanoscale characterization of ferroelectric thin films. Several qualitative experiments demonstrating the capabilities of SFM in controlling domains as small as 20-50 nm in diameter have already been performed. SFM was also used for nanoscale studies of degradation effects, such as ferroelectric fatigue and retention loss. Another very important branch of applications is related to SPM-based electrical characterization. As the characteristic dimensions of electronic devices continue to shrink, the ability to characterize their electronic properties at the nanometer scale has come to be of outstanding importance. Scanning probe microscopy has opened new opportunities to measure semiconductor electronic properties with unprecedented spatial resolution. For example, scanning spreading resistance microscopy (SSRM), and scanning capacitance microscopy (SCM) have already demonstrated device measurements with ~2 and 10-20 nm spatial resolution respectively. Kelvin probe force microscopy (KPFM) has been used for measuring electrostatic forces and electric potential distribution and has found many diverse applications in recent years. In the area of functional molecular materials, SPM is also proving to be an invaluable tool. It is being used as a probe to contact molecular structures in order to characterize their electrical properties, as a manipulator to assemble nanoparticles and nanotubes into simple devices, and as a tool to pattern molecular nanostructures.

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To contribute to the development in these technological and scientific fields, there is a need to “collect” and disseminate this new and growing knowledge at the cross disciplinary level. The NATO Advanced Study Institute (ASI) on “Scanning Probe Micrsocopy: Characterization, nanofabrication and device application of functional materials,” held in Albufeira, Algarve, Portugal from the 1stt to the 13th October 2002, by bringing together highly expertise researchers from the nanoscale techniques and functional materials disciplines, by presenting the fundamentals, technological advances and the needs for further developments in their respective fields and by stimulating active discussions, contributed to foster new scientific contacts and to develop new ideas in the field.

Paula Maria Vilarinho Aveiro, Portugal, June 2004 Yossi Rosenwaks Ramat – Aviv, Israel, June 2004 Angus I. Kingon Raleigh, USA, June 2004

ON THE NATO ASI The main objective of this ASI was to disseminate knowledge concerning the new and emerging applications of SPM to the field of material science, especially in the areas of characterisation, device application and nanofabrication of functional materials. Timing of the proposed meeting was extremely appropriate, as the subject reflects the growing importance of SPM as a key tool for further development of nanoscale science and technology. Rapid progress in the field of SPM and functional materials demands energetic efforts on organising meetings to contribute to the scientific education of a new generation of engineers and scientists with new expertise and deeper understanding of nanoscale device physics. The ASI was attended by 83 researchers, postt doctoral and students from 24 countries, including: Austria, Belgium, Canada, Czech Republic, France, Germany, Greece, Israel, Italy, Japan, Korea, Latvia, Portugal, Poland, Romania, Russia, Spain, Slovenia, Sweden, Switzerland, Turkey, Ukraine, United Kingdom and United States of America. To achieve the proposed objectives the scientific programme of the ASI comprised three main parts: Part I - SPM Techniques and Functional Materials, Part II – SPM in Functional Materials: Characterization and Part III – SPM in Functional Materials: Nanofabrication and Device Application. Introduction, development and the expanding capabilities of SPM, as a powerful nanoscience technique, were addressed in Part I. To complete Part I, the scientific and technological importance of the fundamental knowledge of structure / properties / applications relationships of functional materials were presented and discussed. In Part II recent progress in nanoscale SPM characterization of advanced functional materials, namely semiconductors, magnetics, dielectrics, ferroelectrics, were covered. The newest advances on fabrication of nanostructures and the links between nanofabrication and nanoscale characterization were discussed in Part III. During 10 working days these topics were systematically presented and treated in depth by an interdisciplinary team of leading scientists in lecture format. This tutorial activity was complemented by rump sessions on related subjects, namely: (i) Comparision and direction of SPM methods and (ii) Fabrication: future of the bottom up approach. In addition, to stimulate scientific contacts and discussion, poster sessions and short presentations by the participants were held on their own scientific activities in the field. At these discussions the participants were motivated to expand themselves, to be innovative in the field of characterisation and fabrication with SPM, and to present their own innovative ideas on the topic. The theme for this ASI has its own scientific value. Its uniqueness is in the combination of the fundamental nanoscale research with the progress in fabrication of realistic nanodevices. In addition, it developed new educational advances. By bringing together leading researchers from the material science and SPM communities, relevant information and experience was conveyed that allowed scientists to learn more about the actual developments and future trends of each field. Contacts among the scientists xiii

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were fostered and in this way contributed to the development of this new and technological important interdisciplinary field of science. For PhD students and postdoc scientists, participation in this meeting ledd to significant improvement in their knowledge of the basic properties of functional materials, as well as in application of SPM techniques. With SPM becoming a ‘must-know’ technique in many scientific disciplines, this meeting helped to improve the qualification level of university graduates from different countries and provided manpower with new expertise in the field of nanotechnology and SPM.

ON THE BOOK This book is the output of the ASI NATO meeting on Scanning Probe Micrsocopy: Characterization, nanofabrication and device application of functional materials. The book content reflects the scientific content of the school by itself presenting the main lectures that were given, and some of the participants works that were also presented and discussed during the school. The book is organised in four parts. Part I, Fundamentals of Functional Materials, is an introductory chapter that addresses the general properties of functional materials, highlights some of the unsolved problems of functional materials and reports the progress in silicon technology from the perspective of scaling to submicron devices and the expected performance at the end of the silicon scaling era. Part II, Fundamentals of Scanning Probe Techniques, presents the principles and basics of various SPM techniques, such as near-field optical microscopy, kelvin probe force microscopy and non-contac – AFM, showing the capacity of the techniques to measure the local physical properties of materials with nanoscale resolution. The application of SPM techniques to the characterization of specified functional materials such as piezoelectric ceramics and ferroelectric materials is discussed in Part III, Application of Scanning Tecnhiques to Functional Materials, which also presents the utilization of such techniques to the fabrication of some nano electronic devices. Part IV, Contributed papers, includes some of the R&D work related to the utilization of SPM techniques to functional materials, as presented by the participants.

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ACKNOWLEDGMENTS

The Directors of the ASI (P. M. Vilarinho and Y. Rosenwaks) acknowledge the financial support of NATO, through the Scientific f and Environmental Affairs Division, of the Instituto de Cooperação Científica e Tecnológica (ICCTI) of the Portuguese Ministry of Science and Tecnology, of the Luso – American Foundation (FLAD), of the Fundação para a Ciência e Tecnologia (FCT), of the Portuguese Ministry of Science and Technology and of the University of Aveiro. The travel support of some of the participants by the Portuguese, Greek and Turkish NATO Agencies and National Science Foundation (NSF), USA is also acknowledged. The Directors are grateful to Isabel Salvado, Aiying Wu, Javier Perez de la Cruz and Gerardo Gonzalez for their help with all the organizational details before and during the school. The Directors wish to acknowledge the assistance of Alexander Tkach in the text preparation process. The Directors are particularly thankful to the authors for their contribution to this book. One of the directors (P. M. Vilarinho) expresses her gratitude to Angus Kingon. The idea of organising such an event came from m the discussions had with Angus Kingon during her stay at North Caroline State University (NCSU), USA, on a sabbatical leave. The recognition that Scanning Probe Microscopy was an emerging technique, namely in the field of Functional Materials and that was not addressed in a systematic way, combining expertises coming from different fields of materials science, such as physics, processing and device construction, was the embryo m of this ASI. Alexei Gruverman is also thanked for his valuable contribution and help in finding the lecturers and for the critical analysis of the proposal for such ASI.

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Participants in the Group 38. Markys CAIN 39. Jenica NEAMTU 40. Angus KINGON 41. Cristina ROTARU 42. Hiroshi TOKUMOTO 43. Karlis KUNDZINS 44. Nassima KHALFAOUI 45. Lutfi OZYUZER 46. Alexei GRUVERMAN 47. Luca PELLEGRINO 48. Viktor BOVTUN 49. Paulo MOREIRA 50. Derek OLIVER 51. Lukas ENG 52. Vladimir SHVARTSMAN 53. Claudia RITTER 54. Bert STEGEMANN 55. Maria SHVEBELMAN 56. Oren TAL 57. Isabel SALVADO 58. Waldemar NAWROCKI 59. Gabriella LEO 60. Valeria FERRANDO 61. Grazia TALLARIDA 62. Jianshu YANG 63. Márcia C. NEVES 64. Carlota CANALIAS 65. Dawn A. BONNELL 66. Ana-Maria CHIORCEA 67. Zoia LEONENKO 68. Anton MAIDYKOVSKI 69. Esra OZKAN 70. Ana VIANA 71. Metan TANOGLU 72. Kevin F. KELLY 73. Pascal MARCHET 74. Paula Maria VILARINHO

1. Andrey GAL 2. Olivier DOUHERET 3. Dionizy CZEKAJ 4. Agata Lisiska CZEKAJ 5. Krzysztof PIELICHOWSKI 6. Javier PÉREZ 7. Ioannis TSIAOUSSSIS 8. Marian LEHOCKY 9. Martin PUSTKA 10. Fabiano ASSI 11. Stefan KUYPERS 12. Marlies VAN BAEL 13. Vismants ZAULS 14. Sascha KREMMER 15. Domenique WEINER 16. Sangmin SHIN 17. James SCOTT 18. Gerardo GONZALEZ 19. ALBERTO PASQUINI 20. Stefano GARIGLIO 21. Francesca SBRANA 22. Adrian Mihail MOTOC 23. Cristiano ALBONETTI 24. Maria José MATOS 25. Yossi ROSENWAKS 26. Andrei KHOLKIN 27. Teresa SIERRA GARCIA 28. Harvey AMORIN 29. Jesús RICOTE 30. Elena SOLERA CARLAVILLA 31. Alexander TITKOV 32. Monika KRISTKOVA 33. Maciej WOJTAĝ 34. Silvia KARTHÄUSER 35. Markus MORGENSTERN 36. Lech Tomasz BACZEWSKI 37. Sidney COHEN

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ADDRESS LIST OF THE AUTHORS Germany Phone: 49 351 463 3427 Fax: 49 351 463 7065 Email: [email protected]

Paula Maria VILARINHO Department of Ceramics and Glass Engineering University of Aveiro 3810-193 Aveiro Portugal Phone: 351 234 370354/259 Fax: 351 234 425300 Email: [email protected]

Yossi ROSENWAKS Department of Physical Electronics Faculty of Engineering Tel-Aviv University Ramat-Aviv 69978 Israel Tel: 972-3-6406248/7974 Fax: 972-3-6423508 Email: [email protected]

Angus KINGON North Carolina State University Materials Research Center, NCSU 1001 Capability Drive, Centennial Campus Raleigh NC 27695-7919 USA Phone: (voice Email) 1 919 515 8636 Fax: 1 919 515 3419 Email: [email protected]

Kevin F. KELLY ECE Dept., MS-366 Rice University PO Box 1892, Houston, TX 770051892 USA Phone: 1 713 348-3565 Fax: 1 713 348-5686 Email: [email protected]

James SCOTT Department of Earth Science University of Cambridge Cambridge CB2 3EQ England Phone: 44 1223 333461 Fax: 44 1223 333450 Email [email protected]

Seizo MORITA Department of Electronic Engineering Graduate School of Engineering, Osaka University 2-1 Yamada-Oka, Suita, Osaka 5650871 Japan Phone: 81 6 6879 7761 Fax: +81 6 6879 7764 Email:[email protected]

Dawn A. BONNELL Materials Science and Engineering Director, Center for Science and Engineering of Nanoscale Systems The University of Pennsylvania 3231 Walnut Street Philadelphia, PA,19104 USA Phone: 215 898 6231 Fax: 215 573 2128 Email: [email protected]

Alexei GRUVERMAN North Carolina State University Department of Materials Science and Engineering 1010 Main Campus Drive, EGRC, Rm 339, Campus Box 7920 Raleigh, NC 27695 USA Phone: 1 919 513-3319 Fax: 1 919 515-3027

Lukas ENG Institut für Angewandte Photophysik Technical University of Dresden 01062 Dresden xxi

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Email: [email protected]

Email: [email protected]

Markus MORGENSTERN Institute of Applied Physics Hamburg University Jungiusstrasse 11 D-20355 Hamburg Germany Phone: 49 40 42838 3282 Fax: 49 40 42838 2944 Email: [email protected]

Luca PELLEGRINO Dipartimento di Fisica Università di Genova via Dodecaneso 33 16146 Genova Italy Phone: 39 010 353 6323 Fax: 39 010 311066 Email:[email protected]

Sidney COHEN Department of Materials and Interfaces The Weizmann Institute of Science Rehovot 76100, Israel Phone: 972 8-934 2703 or 3422 Fax: 972 8 934 4137 Email: [email protected] Alberto PASQUINI Strada Delle Cacce 73 10135 Turin Italy Phone: 39 011 3977471 Fax :39 011 3977459 Email :[email protected] Carlota CANALIAS Department of Physics Royal Institute of Technology Roslagsvägen 30b, S-11347 Stockholm Sweden Phone: 46 8 55378192 Fax: 46 8 55378216 Email: [email protected] Derek OLIVER Electrial & computer Eng. University of Manitoba Winnipeg MB R3T 5V6 Canada Phone: 1 204 4749563 Fax: 1 204 2614639

Grazia TALLARIDA Laboratorio MDM-INFM Via Olivetti 2 20041 Alrate Brianza, Milan Italy Phone: 39 039 6036540 Fax: 39 039 6881175 Email: [email protected] Olivier DOUHERET Department of Microelectronics & Information Technology KTH, Royal Institute of Technology Isafjordsgatan 22, P.O. Electrum 229 S-164 40 KISTA, Sweden Phone: 46 8 752 1166 Fax: 46 8 750 5173 Email: [email protected] Lutfi OZYUZER Department of Physics Izmir Institute of Technology Gulbahce Campus, Urla TR-35437 Izmir Turkey Phone: 90 232 4987518 Fax: 90 232 4987509 Email: [email protected] Lech Tomasz BACZEWSKI Institute of Physics Polish Academy of Sciences Al. Lotnikow 32/46, 02-668 Warszawa,

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Poland Phone: 48 22 8431331 Fax:48 22 8430926 Email:[email protected] Anatoly RINKEVICH Institute of Metal Physics Ural Division of Russian Academy of Sciences 18, S.Kovalevskaya St, GSP-170 Ekaterinburg 620219 Russia Phone: 007 3432 499395 Fax: 007 3432 745244 Email: [email protected] Jenica NEAMTU ICPE-CA, Advanced Research &Development Institute for Electrical Engineering, SplaiulUnirii 313, Bucharest 030138 Romania Fax. 40 21 346 82 99 Email:[email protected]; [email protected] Francesca SBRANA DIBE-Università di Genova via All’Opera Pia 11a-16145 GenovaItaly Phone: 39 010 3532167 Fax: 39 010 3532290 Email: [email protected] Ana-Maria CHIORCEA Instituto Pedro Nunes Quinta da Nora, Rua Pedro Nunes 3030-199 Coimbra Portugal Phone: 351 239 700 978 Fax: 351 239 700 965 Email: [email protected] Zoia LEONENKO Department of Chemistry

University of Calgary 2500 University Drive NW Calgary, AB, T2N 1 N4, Canada Phone: 1 403 220 6248 Fax: 1 403 289 9488 Email: [email protected]

List of participants Wagistrasse 2 CH-8952 Schlieren Phone: 41(1)6336168 Fax. 4181)6331048 Email: [email protected]

Cristiano ALBONETTI Institute for Study of Nanostructured Material (I.S.M.N.) C.N.R. Bologna Via Gobetti 101, 40129 Bologna ITALY Phone: +390516398523 Fax: +390516398539 Email: [email protected]

Lech Tomasz BACZEWSKI Institute of Physics Polish Academy of Sciences Al. Lotnikow 32/46, 02-668 Warszawa, POLAND Phone: 48-22-8431331 Fax: 48-22-8430926 Email:[email protected]

Marin ALEXE Max-Planck Institute Microstructural Physics Weinberg, 2, D – 06120, Halle Germany Phone: 49-351-463 33427 Fax: Email: [email protected]

Andreja BENCAN Institute Jozef Stefan Jamova 39, 1000 Ljubljana Slovenia Phone: 38614263126 Fax: 38614263126 e-mail: [email protected]

Harvey AMORIN Ceramics and Glass Engineering Department University of Aveiro 3810-193 Aveiro PORTUGAL Phone: 351 234 370354 Fax: 234 351 425300 Email: [email protected]

Dawn A. BONNELL Materials Science and Engineering Director, Center for Science and Engineering of Nanoscale Systems The University of Pennsylvania 3231 Walnut Street Philadelphia, PA, 19104 Phone: 215 898 6231 Fax: 215 573 2128

Fabiano ASSI ETH Zurich

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Email: [email protected] Viktor BOVTUN Microelectronics Department National Technical University Ukraine “Kiev Polytechnic Institute” Peremogy ave. 37, 252056 Kiev UKRAINE Phone: 38 (044) 236 96 76 Fax: 38 (044) 236 96 76 Email: [email protected] Markys CAIN National Physical Laboratory Queens Road, Teddington Middlesex, Tw11 0lw UK Phone: 44 (0) 208 943 6599 Fax: 44 (0) 208 943 2989 Email:[email protected] Carlota CANALIAS Department of Physics Royal Institute of Technology Roslagsvägen 30b, S-11347 Stockholm SWEDEN Phone: +46 8 55378192 Fax: +46 8 55378216 Email: [email protected] Elena SOLERA CARLAVILLA Electroceramic Department

INSTITUTO DE CERAMICA Y VIDRIO CSIC -CAMPUS DE CANTOBLANCO Camino de Valdelatas, s/n 28049 -MADRID SPAIN Phone: (34) 91 735 5840 fax (34) 91 735 5843/5 EMAIL: [email protected] Ana-Maria CHIORCEA Instituto Pedro Nunes Quinta da Nora, Rua Pedro Nunes 3030-199 Coimbra PORTUGAL Phone: 351 239 700 978 Fax: 351 239 700 965 Email: [email protected] Sidney COHEN Department of Materials and Interfaces The Weizmann Institute of Science Rehovot 76100, Israel, Phone: 972 -8-934 -2703 or 3422 Fax: 972-8-934-4137 Email: [email protected] Agata Lisiska CZEKAJ Department of Materials Science University of Silesia

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2, ĝnieĪna St. 41-200 Sosnowiec POLAND Phone: (+4832) 291-82-43 Fax: (+4832) 291-82-43 Email: [email protected] Dionizy CZEKAJ University of Silesia Department of Materials Science 2, ĝnieĪna St. 41-200 Sosnowiec POLAND Phone: (+4832) 291-82-43 Fax: (+4832) 291-82-43 Email: [email protected] Elena DIMITRIU National institute for Materials Physics RO76900, tr. Atomistilor 105bis Magurele-Ilfov, P.O.Box:MG-7 ROMANIA Phone: 00 4021 4930195 Fax: 00 4021 4930267 Email: [email protected] Olivier DOUHERET Department of Microelectronics & Information Technology KTH, Royal Institute of Technology Isafjordsgatan 22, P.O. Electrum 229 S-164 40 KISTA, SWEDEN

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Email: [email protected]

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Fax. 492518333602 Email. [email protected] Maciej WOJTAĝ Faculty of Chemistry University of Wrocáaw F. Joliot-Curie 14, 50-383 Wrocáaw POLAND Phone: 48 71 375 72 36 Fax: 48 71 328 23 48 Email: [email protected] Aiying WU Departement of Ceramics and Glass Engineering University of Aveiro 3810-193, Aveiro PORTUGAL Phone: (351) 234370354 Fax: (351) 234425300 Email: [email protected] Jianshu YANG Paul-Drude-Institute für Festkörperelektronik Hausvogteiplatz 5-7 D-10117 Berlin, GERMANY Phone: +49-30-20377393 Fax: 49-30-20377257 Email: [email protected] Vismants Zauls Institute of Solid State Physics,

University of Latvia Kengaraga 8, LV-1063, Riga, LATVIA Phone: +371-7260803 Fax: +371-7132778 Email: [email protected]

Part I – Fundamentals of Functional Materials

FUNCTIONAL MATERIALS: PROPERTIES, PROCESSING AND APPLICATIONS

P.M. VILARINHO Department of Ceramics and Glass Engineering, CICECO, University of Aveiro, 3810 – 193, Aveiro, Portugal

Contents 1. 2. 3. 4. 5. 6.

Introduction Fundaments of Properties of Functional Materials: General Concepts Functional Materials Functional Materials Processing Technologies Applications of Functional Materials Future Trends in Functional Materials

“On September 7, 2001, doctors in the United States performed the first long distance operation, with surgeons in New York performing a laparoscopic cholecystectomy on a patient in Strasbourg, France. The surgery was successful, and the patient was discharged from the hospital with no complications 48 hours later. The ability to perform complex surgical manipulations from remote locations will eliminate geographical constraints and make surgical expertise available throughout the world, improving patient treatment and surgical training. Needless to say, the potential uses of such remote, robotic technology are boundless, not only in medical settings but in search and rescue missions, scientific discovery missions, and countless other arenas…” – in http://www.stanford.edu/telemedicine [1]. Keywords: Piezoelectrics, Pyroelectrics, Ferroelectrics, Incipient Ferroelectrics, Ferroelectric Relaxors, Applications, Fabrication

1. Introduction It is extremely difficult for anybody not to recognise the importance of the so called Information Technology. Even for those not aware of it, Information Technology dominates our life today. There are numerous examples m with the “first long distance medical operation” being just one among them. The introduction of robotic and computer technology into surgical operations allows surgical procedures to be carried out from a distance (telesurgery). Besides the expert surgeons, these telesurgeries 3 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 3-33. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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require a secure, reliable and fast network connecting the two points (surgeon to patient) and a robotic system capable of translating the surgeon's hand movements in place A (New York, in this case) to the instruments inside the patient in place B (Strasbourg, France) [1]. From a narrower perspective - materials science and technology - this is possible due to the sustained research and development of materials, namely functional materials. As the name suggests, the definition of functional materials reflects the ability of a material to perform a certain “function” under a determined stimulus. In this general definition a wide spectrum of materials can be included together with an ample range of material properties and applications. However, the classification off functional materials is usually related to materials whose “function” f is associated with their electric, magnetic, and or optical properties. This group of functional materials mainly includes dielectrics, pyroelectrics, piezoelectrics, ferroelectrics, ferroelectric relaxors, incipient ferroelectrics, semiconductors, ionic conductors, superconductors, electro-optics and magnetic materials. It is common to identify the class of functional materials with applications such as Materials for Information Technology, Materials for Electrical Energy Conversion, Materials for Biologic Applications, Materials for Space Technology, among others. The application of functional materials, typified by electroactive materials including piezoelectrics, pyroelectrics and ferroelectrics, for sensing and actuation spans most if not all industrial sectors. This includes medical diagnostics such as ultrasonic imaging, aerospace such as accelerometers and micropositioners, automotive such as solid state piezoelectric fuel injectors, and chemical and process control, which requires the use of thermal, strain and force sensors. Although the discovery of certain properties off functional materials goes back to the nineteenth century some of the materials that exhibit such properties became useful only during the Second World War and others very recently. (see references [2, 3] for more details on the history of such functional materials). The utility off functional materials in these applications reflects their unique properties, such as spontaneous polarization, piezoelectricity, superconductivity and magnetoresistance. All these properties are directly dependent on the chemical composition, singularities of the crystallographic structure and manufacture process. The knowledge of the relationships between composition, structure, processing and properties allows the production of improved materials for known applications as well as for new uses and in an economic way. Materials scientists, chemists and physicists have been searching for these relationships for years. The characterization of the bulk properties of functional materials and their description by theoretical models were their main activities in the recent past. Nowadays a considerable knowledge on the relationships between composition, microstructure, processing and macroscopic properties can be established for a wide variety of functional materials. But researchers developing and producing materials try to optimize materials to meet tomorrow’s needs and this “materials bulk approach” has recently changed. If the progressive miniaturisation of electronic components is to proceed with the speed of the past decades devices of nanometric dimensions will be needed soon (see part I of this book by A.I. Kingon). Figure 1 shows the scales involved in some of the actual microelectronic devices and those that will be involved in the future. Starting from m the semiconductor of microelectronics and

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following the Moore's law, which predicts the exponential decrease of the size of technological components [4], a progression into the field of nanoelectronics is happening, and because of the approaching of a limit to the size of current electronic systems molecular electronics is gaining importance. Molecular electronics looks for the fabrication with electronic devices of molecular dimensions and also from molecular components so as to perform functions in electronic circuitry now performed by semiconductor devices. Individual molecules are very small and electronic devices constructed from molecules will be very small (100 x smaller than semiconductor-based counterparts). Besides the dramatic reduction in the size, the ability to produce billions of identical molecules is important benefits of molecular electronics [5, 6]. For more details on molecular electronics t see part III of this book.

Figure 1. Dimensions of microelectronic devices: microsensors, microelectricalmechanical systems (MEMS), nanoelectricalmechanical systems (NEMS), micromachines, integrated circuits (IC) [7].

Consequently, the drive for smaller size, greater functionality and the replacement of products by services emphasise on the less mature study of nanostructural attributes of materials: electrical, optical, magnetic and biological. Many of these attributes are properties of films, surfaces or interfaces. The use of conventional techniques is inadequate to evaluate properties at a nanoscale level, to design nanofeatures and to

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hold a position of functional molecules. Nanoscale tools are required. Due to its nanometer lateral and subangstrom vertical resolution Scanning Probe Microscopy (SPM), with its various techniques, is now a fundamental tool for the nanoelectronics era. The following text is mainly dedicated to some functional materials, technologically important in the present day. The text is not intended to treat all the functional materials and their peculiarities. In view of this and in relation to the final applications emphasis is given to a few electrical properties, such as piezoelectricity, pyroelectricity and ferroelectricity, the family of materials, for example perovskite type and materials formulations and their fabrication. The limitations of the materials and processes will be highlighted from the point of view of the new materials requirements and industry demands, and related with the future “materials nanostructure t approach”. The aim of this text is to introduce the reader to the following texts in which the Scanning Probe Techniques devoted to the characterization (part II), fabrication and device application of functional materials (part III) are presented and discussed. Along the text, references to extended literature or text books on the reported topics will be given.

2. Fundaments of Properties of Functional Materials: General Concepts In this section, the most representative properties of some functional materials, such as dielectrics, pyroelectrics, piezoelectrics, ferroelectrics, ferroelectric relaxors and incipient ferroelectrics are briefly presented in order to justify their actual applications. Exhaustive descriptions on these properties can be found in several text books and reviews [8-15]. Dielectrics are a class of materials with high electrical resistivities. Dielectrics do not conduct electricity due to the very low density of free charge carriers and, because of this, they can perform insulating functions. However, dielectrics exhibit a number of unique assets when placed under the effect of an electrical field. The ability of the dielectric to store charge in a capacitor is related to the polarization of the dielectric under the electric field and is the simplest practical application of a dielectric. Nonetheless, the various polarization responses of the dielectric under an electric field are being increasingly used in micro and nanoelectronic devices and open a wide range of new devices. When an electric field E is applied to an ideal insulator, there is no long-range charge transport, as in a conductor, solely a short-range dislocation of the positive and negative charge centre which causes the appearance of electric dipole moments in the material. The material is called a dielectric. The dielectric is said to be polarised. Several polarization mechanisms were identified: atomic, ionic, dipolar and space charge; each is related to the nature of the charged entities which suffer charge displacement or to the nature of the displacement. The charge displacement process in dielectrics is systematically described in references [11, 15]. If there is a linear relationship between the applied field E and the induced polarization P and P disappears when E is removed the material is called a linear dielectric. If at zero field conditions a mechanical stress provokes the development of electric charges (polarization) these materials are called piezoelectric. Piezoelectricity is

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the ability of certain crystalline materials to develop an electrical charge proportional to a mechanical stress or vice versa. Of the 32 classes of symmetry, 11 possess a centre of symmetry and are non-polar. For these an applied stress t results in symmetrical ionic displacements so that there is no net change in the dipole moment. The remaining 21 point groups do not have a centre of symmetry (i.e. non-centrosymmetric) and possess one or more crystallographically unique directional axes; all non-centrosymmetric point groups, except the point group 432, exhibit piezoelectric effect along unique directional axes. In a piezoelectric the relationship between the applied deformation and the induced polarization is linear and reversible. This effect is different from the electrostriction effect. All the materials suffer a small change in dimensions when subject to an electric field. However, if the resultant strain is proportional to the square of the field, the effect is called electrostriction (equation 1). x = ξE 2 , (1) where ζ is the electrostrictive coefficient. In a piezoelectric the magnitude of P depends on the magnitude of the stress and the sign of the produced charge depends on the type of applied stress (tensile or compressive). The polarization generated from a mechanical stress is called the direct or generator effect, while the converse or motor effect is associated with the mechanical deformation derived from an applied electric field. The relationships between the strain x, stress σ, electric field strength E, electric polarisation P in a piezoelectric material are: P = dσ (direct effect), (2) x = dE (converse effect), (3) where d is the piezoelectric coefficient or strain constant (dij relates a field along the i axis to the strain in the j direction). The d33 coefficient is the most commonly cited of these coefficients and it is the corresponding coefficient for both strain and field along the polar axis. Another important parameter to evaluate the performance of a piezoelectric is the effective coupling coefficient (kefff) which is a measure of the amount of electrical energy that is converted into strain and defined as: electrical energy converted in mechanical energy (4) Keff 2 = (direct effect) input electrical energy mechanical energy converted in electrical energy 2 Keff = (converse effect) (5) input mechanical energy For the fundamentals of piezoelectricity references [8, 9, 14] are suggested. If, in zero field conditions, there are dipolar moments due to a non-symmetric structure, the materials will have spontaneous polarization and they are called pyroelectrics. Ten of the crystallographic groups contain a unique polar axis and are, therefore, spontaneously polarised. In a pyroelectric the change in temperature produces a change in spontaneous polarisation. The pyroelectric effect can be described in terms of the pyroelectric coefficient p that relates to the change in spontaneous polarisation dPs as a function of temperature (dT) as: p = dPs / dT (6) As for piezoelectrics, since polarisation is a vector, the value of the pyroelectric coefficient is different depending on the direction in which it is measured: pi = ¨Psi / ¨T (7)

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where i = 1,2,3. Two figures of merit are defined for the pyroeletric performance of a material FV and FD in the form of: FV = p / 'c İ , (8) where İ is the dielectric permittivity and 'c is the volume specific heat, which describes the effectiveness of a pyroelectric element in terms of materials properties and usually has units of m2/C. In an imaging array, where noise is of primary concern, the figure of merit FD, defined as: (9) FD = p / ('c İ1/2 tan1/2į) , where tanį is the loss and typically expressed in (µm3/J)1⁄2 is more useful. Some pyroelectric materials have an additional property; the direction of spontaneous polarisation can be switched by an applied electric field. These materials are called ferroelectrics. If the spontaneous polarization direction is changed by mechanical stresses the materials are called ferroelastic. Ferroelectricity and ferroelasticity, in opposition to the other two polar phenomena, cannot be predicted from the crystalline symmetry but need to be experimentally verified. All ferroelectrics are piezoelectric and pyroelectric but a piezoelectric or a pyroelectric is not inevitably a ferroelectric. For example, materials such as ZnO and tourmaline ((Na,Ca)(Li,Mg,Al)3(Al,Fe,Mn)6(BO3)3(Si6O18)(OH)4) have a non-reversible dipole but are still pyroelectric. For a material that contains electric dipoles, the local electric field will promote the dipole alignment in a certain region, contributing to the increase of the polarization, which, by itself, will promote the increase of the local field. These co-operation phenomena will align the dipoles along the same direction, resulting in the spontaneous polarization of the material. Consequently, the electric polarization in pyroelectrics and ferroelectrics does not vary linearly with the applied field and hence they are called non-linear dielectrics. In ferroelectrics the relationship between the applied field and the polarization is described by a hysteresis loop (figure 2) similar to the one exhibited by the ferromagnetic materials. The name ferroelectrics refers to the analogy with ferromagnetic behaviour although, for ferroelectrics, the phenomena is not related with the presence of iron. The application of a low electric t field to a non-polarized ferroelectric provokes a linear and reversible increase of the polarization as the field increases. The slope of this variation corresponds to the initial dielectric permittivity of the material. As the field increases, the further increase of the polarization is non-linear and, for high field values, the variation of P with E is small and approaches to a saturation value. The polarization value extrapolated for zero field (E=0) gives the saturation polarization Ps. When the external field is removed, the polarization does not fall to zero, keeping a remnant value designated as remnant polarization Pr. To cancel this value, a field in the opposite direction and of magnitude Ec should be applied. This field Ec required to reduce the polarization to zero is called the coercive field. Further increasing of the field in the negative direction will cause the switching of the polarization. Reversing the field direction once again will complete the hysteresis cycle.

9

Figure 2. Typical hysteresis loop of a ferroelectric (adapted from [11]).

The hysteretic behaviour of ferroelectrics is related to their domain structure. A ferroelectric possesses regions with uniform polarization called domains. Within a domain all the dipoles are aligned in the same direction, differing from the direction of the neighbour domain in such a way that for zero applied field the material macroscopic polarization is null. The several existing domains are separated by interfaces called domain walls which typically have the dimensions of 1 to 2 lattice spacing. For materials with a tetragonal symmetry the polarization direction between domains can form 180˚ or 90˚ angles. For a low applied electric field (region of linear relationship between P and E) the field is not large enough to switch any domain, therefore the ferroelectric will behave as a linear dielectric. As the applied electric field increases, a number of domains, which have a polarization opposite to the direction of the field, will be switched in the direction of the applied field and the polarization will increase rapidly until all domains are aligned in the field direction. Eventually, for a high applied field, the sample will only be a mono domain. As the field strength decreases, the polarization will generally decrease but not return to zero. When the field is reduced to zero the majority of the domains will remain aligned in the applied field direction and the ferroelectric will exhibit a remnant polarization Pr. Usually, the value of Pr is inferior to the value of the saturation polarization Ps since some of the domains will reassume their original orientation. The process of switching all the domains under a single orientation is called poling. The remnant polarization Pr in a ferroelectric cannot be removed until the applied field in the opposite (negative) direction reaches the value of the coercive field Ec. Further increasing of the field in the negative direction will cause a complete alignment of the dipoles in this direction. Reversing the field direction once again will complete the hysteresis cycle. A ferroelectric material can undergo a phase transition adopting a non-polar centrosymmetric structure at a temperature called TC (the temperature of the phase transition). Above TC, with the loss of the polar structure, t the material does not exhibit

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spontaneous polarization and it is said to be paraelectric. Below TC, due to the appearance of the spontaneous polarization and to the mutual interaction between the dipoles which causes a significant increase of the local field, the material exhibits ferroelectricity. The structural phase transition from the paraelectric to the ferroelectric phase is reversible. Near TC due to a distortion in the crystalline lattice as the phase structure changes, the thermodynamic properties, including dielectric, elastic, optical, and thermal constants show an anomalous behaviour; the permittivity raises, reaching a maximum at TC (figure 3). In the ferroelectric region the increase of the thermal agitation as the temperature approaches TC facilitates the growing of the domains oriented along the field. Above TC the permittivity of the material decreases and at the same time a sudden reduction of the resistivity and a marked increase in the losses is observed.

Figure 3. Dielectric permittivity as a function of temperature for BaTiO3 [17].

In most ferroelectrics, the temperature dependence of the dielectric constant above TC can be described by the Curie-Weiss law: İr = İ0 (1 + C/(T í Θ) (10) where İr is the dielectric permittivity of the materials, İ0 is the dielectric permittivity of vacuum, C is the Curie-Weiss constant, T is the temperature and Θ is the Curie-Weiss temperature, which is in general smaller than TC. For first order transitions Θ < TC, while for second order phase transitions Θ = TC. In ferroelectrics the non-linear variation of P as a function of E is also detected in the optical properties. Electro optical effects such as square birefringence (Kerr effect) and linear birefringence (Pockles effect) are some of the optical effects whose magnitude depends on the intensity and direction of the applied field. By analogy with magnetic dipoles, the relative orientation of electric dipoles may create different polarization patterns; electric dipoles can orient in a parallel or antiparallel way. If the electric dipoles are oriented in an antiparallel way, visualised as upward and downward directions, the total polar momentt is null and the material is said to be antiferroelectric. Antiferroelectrics (antipolar materials) can revert to a ferroelectric state when subjected to a sufficiently high electric field. The polarization

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dependence on the electric field of an antiferroelectric is displayed in figure 4. For low electric fields the induced electric polarization is quite low. Only after the application of a high enough electric field, able to break the antiferroelectric order, the polarization increases. For high fields a hysteric behaviour, similar to the one observed in ferroelectrics, is visible [16, 17].

Figure 4. Antiferroelectric hysteresis loops [18].

For additional information on pyroeletricity and ferroeletricity the references [9, 10, 13] should be consulted. Although the majority of the ferroelectrics, such as BaTiO3 (BT), exhibit a transition from a low symmetry and temperature ferroelectric phase to a high symmetry and temperature paraelectric phase, this is not a necessary condition to possess ferroelectricity. Some ferroelectrics, such as BaMgF4 (BMF), melt at temperatures lower than the transition temperature. On the other hand, incipient ferroelectrics (or quantum paraelectrics) such as SrTiO3 (ST), that exhibit structural phase transitions do not possess ferroelectric behaviour. ST falls into the unique category of ferroelectrics known as incipient ferroelectrics along with KTaO3, CaTiO3 and TiO2 [19]. As the temperature decreases, the cubic ST undergoes a transition to the tetragonal phase at 105 K [20] and at 65 K a transition from the tetragonal to the orthorhombic structure [21]. However, no ferroelectric long range order is established. The first transition is purely structural with almost no influence on the dielectric response. For T > 50 K the temperature dependence of dielectric permittivity of ST obeys Curie-Weiss law, although, at lower temperatures the permittivity versus temperature dependence deviates from Curie-Weiss law and saturates reaching high values (~24000 for single crystals [22]) as the temperature approaches 0 K. Quantum fluctuations of the atomic positions are thought to suppress the ferroelectric transition and lead to a stabilization of paraelectric state characteristic for a quantum paraelectric limit [22]. Consequently, the ferroelectric anomaly can be easily induced in incipient ferroelectrics by the application of high enough electric field [20], uniaxial stress [23], chemical substitutions in the lattice [24-26] or oxygen isotope exchange [27]. A certain class of ferroelectrics, designated as relaxors, differenciate themselves from conventional ferroelectrics by the following typical characteristics [28], depicted in figure 5:

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a) A broad phase transition, in which the sharp peak in the dielectric permittivity versus temperature curve related to the phase transition of a classical ferroelectric at Curie temperature at TC is substituted in a ferroeletric relaxorr by a broad peak that occurs over a range of temperatures. This particular feature can also be seen in other properties related to the phase transition. b) A strong frequency dependence of the dielectric properties. While the properties of classic ferroelectrics do not vary intensely with the frequency in the radio frequency range, the dielectric properties of relaxors are strongly dependent on the frequency in rf frequency range. The peak of the dielectric permittivity decreases in value and shifts to higher temperatures as the frequency increases. c) The absence of a macroscopic polarization below the temperature of the permittivity maximum. At temperatures well below the permittivity maximum, no evidence of optical anisotropy or of X-ray line splitting is found. This characterizes the absence of a macro-volume changing to a polar phase, usually observed in classic ferroelectrics. No macroscopic domain state can be observed except when strong electric fields are applied (during hysteresis or under bias). There is not a uniform atomic shift throughout the entire crystal, therefore the net polarization remains nil, but microscopic polar domains develop. In the absence of a field, relaxors contain very small nanodomains that differ from each other in some detail of structure or composition

Figure 5. Schematic representation of (a) dielectric permittivity and (b) spontaneous polarization versus temperature for classic ferroelectrics (curve 1) and ferroelectric relaxors [29].

A common feature of relaxors, in both perovskite and tungsten bronze structures, is that more than one type of ions occupies equivalent crystallographic lattice sites [28]. Although this is a necessary condition, it is not a sufficient one, for there are numerous solid solutions where sharp behaviour is preserved throughout the whole composition range. However, it is generally accepted that ferroelectric relaxors are highly heterogeneous materials and that the compositional fluctuations markedly affect the macroscopic response. There is strong evidence that the physics of relaxorr behaviour is intimately linked to nanoscale chemical heterogeneity. The ferroelectric relaxor behaviour has been the subject of research for many years; several models have been proposed to explain it and have been reviewed by Ye [30]. Smolenskii et al. [31] originally proposed that the dielectric “diffuse phase transition” of relaxors was due to the local compositional fluctuations. In A(B’B’’)O3 perovskites the random distribution

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of B cations creates a range of Curie temperatures that results in the broadened macroscopic variation of permittivity versus temperature. Though this model fails to explain why the “transition temperature” (temperature of the maximum) shifts to higher temperatures as the frequency increases. Later Cross [28] considering the regions of short-range chemical order as nanoscale polar clusters, proposed that the dipole moment of the clusters thermally switches between equivalent directions so macroscopic polar domains never form after the structural transition, as in classical ferroelectrics (superparaelectric model). By analogy with superparamagnetic behaviour it was proposed that there is no interaction between clusters. In this model the frequency dependence of the temperature of the dielectric maximum should obey a simple Debye relationship. However, physical unrealistic values for both activation energy and characteristic frequency were obtained. The dipolar glass-like behaviour, suggested by Viehland et al. [32], is an extension of the superparaelectric model, in which the interactions between polar clusters are considered. Accordingly, the interactions between the polar microregions control the kinetics of the polarization fluctuations and the development of a frustration state near the freezing temperature (Tf) leading to a broadening of the relaxation-time distribution and strong deviations from the Curie Weiss behaviour near Tf. Above Tf the ferroelectric clusters are superparaelectric with dipole moments fluctuating between identical orientations and, as the temperature decreases, the superparaelectric moments freeze into a glassy state due to correlations between dipole moments. The polar nanosized regions are claimed to appear at a temperature above the dielectric maximum, at the so-called Burns temperature [33, 34], and to grow during the cooling till the freezing temperature. Another theory invokes the contributions from a quenched random field which is again due to the compositional heterogeneity [35]. In this model the relaxor state is a ferroelectric state broken up into nanodomains under the constraint of quenched random fields. Charged compositional fluctuations are considered as sources of random fields. In a more recent theory, Spherical random-bond - random-field (SRBRF) model [36], the relaxor is considered as a new type of dipolar glasses, namely, the spherical vector glass, in which the polar clusters are formed when two or more cations, moving in a multisite potential, create a m fluctuations – single reorientable polar unit. Accordingly the regions of compositional “chemical clusters” are essentially static. The above mentioned models generally agree thatt the physical origin of the dielectric peak in relaxors is then different from the one in classical ferroelectrics; in relaxors it is related to the thermal slowing of dynamic polar nanoregions rather than a paraelectric-ferroelectric phase change as in classical ferroelectrics. However, the relationships between the local chemical heterogeneities, the local polar clusters and their dynamics are not yet clearly established. Some of the models consider the regions of short-range chemical order as nanoscale polar clusters [28, 32, 37] and others, the charged compositional fluctuations as sources of random fields. A universal model for ferroelectric relaxorr behaviour is still missing. The full understanding of the relaxor behaviour, due to its nanometric scale, requires the investigation of the polar cluster dynamics at a local level. Until now the majority a of the results on the polar nanocluster structure of relaxors have been acquired by indirect approaches. The classical visualization of ferroelectric domains by optical methods, which have a limited resolution, is not suitable for this purpose. Methods with higher resolution are required.

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The high spatial resolution, down to a few nanometers, and the high sensitivity to local polarization make Piezo Force Microscopy (PFM) a well suited technique for these studies. Indeed, Scanning Force Microscopy (SFM) has been recently successfully applied for nanoscale characterization of ferroelectric thin films [38]. Several qualitative experiments demonstrating the capabilities off SFM in controlling domains as small as 20-50 nm in diameter have already been performed (see A. Gruverman, part II of this book). However, few results have been reported on the observation of the polar structure of relaxors via PFM [39-41]. 3. Functional Materials Among the functional materials which exhibit piezoelectric, pyroelectric, ferroelectric and ferroelectric related properties, four mainn groups of materials have been considered: hydrogen bonded systems, ionic crystals, narrow gap semiconductors [17] and organic polymers. One of the most important is the group of the ionic crystals. Within this group a main type of structure should be considered the corner sharing oxygen octahedral, which includes the four following important families (structures) of materials: bronze tungsten (A2B2O6), perovskite (ABO3), pyrochlore (A2B2O7) and bismuth-layer (Bi4Ti3O12) structures (Figure 6). Among these, the perovksite group is particularly significant, from the point of view of applications. This is because such materials undergo a phase transition on cooling from a high symmetry temperature phase (cubic paraelectric phase) to a non-centrosymmetric ferroelectric phase. The materials with the highest piezoelectric coefficients belong to the lead based perovskite family. Materials with a high ferroelectric transition temperature show piezoelectricity at room temperature, whereas those with transition near or below the room temperature exhibit an important electrostrictive effect. For these, due to the large anharmonicity of the ionic potential, the electrostriction is extraordinarily large [14]. Besides the ability to design the physical properties required for certain applications, by formation of solid solutions, the possibility of fabrication as single crystals, ceramics, textured ceramics and thin and thick films adds value to this family of materials. The ideal structure of the perovskite is cubic with a spacial group Pm3m. The general crystal structure can be thought of as either (i) a body centred cubic (BCC) lattice with A ion (+2) at the centre surrounded by four B ions (+4), each at a corner, and twelve O ions (-2), each of which occupies the midpoint of the edge, or (ii) as a face centred cubic (FCC) lattice with a B ion at the centre surrounded by six O ions, each at a face centre, and four A ions, each at a corner. In the non-polar state the geometric centres of A, B and O coincide giving rise to a non-polar lattice. In the polar state the A and B ions are displaced from their geometrical centres with respect to O2- ions, giving a net polarity to the lattice. These ionic displacements are concomitant with the structural phase transitions that take place as the temperature of the crystal changes. The lattice constant of perovskites is close to 4 Å due to the rigidity of the oxygen octahedra network and the well defined oxygen ionic radius of 1.35 Å. A large number of cations can be accommodated in the cages formed by the oxygen anions, giving rise to wide variety of materials. Complete solid solutions are easily formed between many cations. Moreover, many different cations can be substituted in both A and B sites of the

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perovskite lattice without drastically changing the overall structure. The occupancy of A and B sites of the perovskite structure by more than one type of ions creates complex structures, (A’A”)BO3 or A(B’B”)O3, in which the A or B ions distribution varies, originating different degrees of cation order either as randomly, completely ordered or partially ordered structures. The dependence of the macroscopic properties on the degree of order is well known for different types y of materials such as lead based relaxors [42] and titanates for microwave applications [43]. Moreover, depending on the materials stoichiometry different ratios of order can be defined, such as: 1:1, 1:2 or 2:1. It is then possible to manipulate the material’s properties such as TC, piezoelectric constant or lattice constant with only a small substitution of a given cation, which is very valuable from the technological perspective. Barium Titanate (BaTiO3 - BT), Lead Titanate (PbTiO3 - PT), Lead Zirconate Titanate (PZT), Lead Lanthanum Zirconate Titanate (PLZT), Lead Magnesium Niobate (PMN), Strontium Titanate (SrTiO3), Potassium Niobate (KNbO3 - KN), Potassium Sodium Niobate (K KxNa1-xxNbO3 – KNN) and Potassium Tantalate Niobate (K(TaxNb1-x)O3 - KTN) are some of the technological important functional materials that crystallize with a perovskite type structure.

a

b

c

d

Figure 6. Corner sharing octahedral structures: (a) perovskite, (b) pyrochlore, (c) tungsten bronze and (d) bismuth layer (adapted from [44])

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According to Haertling [18] the discovery of non-linear dielectric properties in BT associated with very high values of the dielectric permittivity, was crucial for the development of the new generation of electronic t materials, which have occurred since the Second World War up to the current day. Hence, BT is among the most systematically studied and widely used ferroelectrics and considered as the prototype ferroelectric ceramic. The temperature dependence of the spontaneous polarization and dielectric constants of BT is shown in figure 3. Three anomalies can be observed: Firstly, above 130 °C BT possess a paraelectric cubic phase and the net polarization is null. The centre of positive charges (Ba2+ and Ti4+ ions) coincides with the centre of negative charge (O2-). At 130 °C, the Curie point, the discontinuity observed is related with the phase transition from a cubic (non-ferroelectric) phase to a tetragonal (ferroelectric) phase which develops on cooling through TC. The centre of Ba2+ and Ti4+ ions are displaced in relation to the O2- ions leading to the formation of electric dipoles. The vector of spontaneous polarization has the [001] direction. The other two discontinuities are associated with transitions from a ferroelectric phase to another one. Between 0 °C and -90 °C, the ferroelectric orthorhombic phase is stable with the polarization direction along the [011] direction. On decreasing the temperature below 90 °C the phase transition from the orthorhombic to ferroelectric rhombohedral phase leads to polarization in the [111] direction. The peaks observed in the dielectric permittivity curve accompany the polarization discontinuities. BT was firstly used as piezoelectric ceramic transducers. However, due to the discovery of better piezoelectric properties in other materials, namely the solid solution between lead titatante and lead zirconate (PZT), BT found use mainly as the high permittivity dielectric in ceramic multilayer capacitors. To optimize the properties for certain applications, BT has been combined with different additives. Various A and B site substitutions in different concentrations have been tried to tailor the dielectric and ferroelectric properties of BT. It was observed thatt solid solutions with isovalent substitutions do not significantly alter the dielectric permittivity versus temperature response. Their main effect is the alteration of the transition temperatures. Ba, Pb and Sr ions can be mixed in any proportions forming complete solid solutions while the solubility of Ca in BT lattice is limited. A site substitution with Srr2+ was found to reduce the Curie point linearly towards room temperature, m whereas differently the substitution of Pb2+ for Ba2+ raises the Curie point. The effect of various isovalent substitutions on the transition temperatures of BaTiO3 ceramic is described in detail in [11]. For multilayer capacitors the modifications of BT are done to maximize the dielectric permittivity value; accordingly the dielectric permittivity maximum should occur near the operation temperature of the device and the dependence of the dielectric permittivity on the temperature should be minimized. Following these modifications, X7R and Z5U ceramic capacitors are fabricated [12]. The formation of complete solid solution in the BaTiO3-SrTiO3 system allows the transition temperature of BT to be shifted from 120 ºC to room temperature. TC and the magnitude of the permittivity can be modulated by the ST content, which makes it very attractive for practical applications. The high dielectric constant, charge storage capacity, dielectric non-linearity, breakdown field, low leakage and microwave loss make the solid solution BST suitable for applications such as Gigabit Dynamic Random Access Memories (DRAMs) and high frequency devices. Moreover the volatilities of

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BST components are lower than the Pb-based ferroelectrics t materials thereby making it relatively easier to introduce into fabrication facilities. For DRAMs applications the material should be in the paraelectric phase in which the hysteretic effects are absent. The basic parameters for applying capacitor thin films on DRAMs are the dielectric constant, leakage current density and reliability. More information on BST for DRAMs is given in the reviews of Ezhilvalavan [45] and H. Schroeder and A. Kingon [46]. In recent times a strong interest appears to exploit the non-linear dielectric response of ferroelectrics and incipientt ferroelectrics above the TC to fabricate tunable microwave devices. Tunable capacitors and resonators devices are based on the tunability effect, i.e., on a dc-electric-field dependence of permittivity that is strengthened on approaching from above the paraelectric-ferroelectric phase transition, at which the dielectric permittivity is maximized, but the dielectric losses remain still sufficiently low. SrTiO3 reveal a high tunability at 80 K and is therefore suitable for planar tunable High Temperature Superconducting (HTS) devices. On the other hand BST solid solution with the maximum of tunability shifted to higher temperatures is ideal for room temperature applications. [47]. The solid solution between lead titanate (PbTiO3) and led zirconate (PbZrO3) is presently the most commonly used compositional system for piezoelectric applications. PT is a ferroelectric material with a transition temperature at 420 °C and PZ is antiferroelectric material with a transition to the paraelectric cubic state at 230 °C [9]. The solid solution between PT and PZ is complete forming the (Pb(ZrrxTi1-x)O3) system known as PZT. PZT has a perovskite type structure with Ti4+ and Zr4+ ions occupying randomly the B site. The PZT phase diagram is represented in figure 7.

Figure 7. PT – PZ phase diagram [8]

The richness of the PbZrO3-PbTiO3 system comes from the existence of several ferroelectric, antiferroelectric and paraelectric phases with various symmetries: tetragonal, rhombohedral, orthorhombic and cubic [8]. The transition temperature TC

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varies from 230 °C to 490 °C depending on the Zr/Ti ratio. Above TC the solid solutions exhibit a cubic structure and are paraelectric. There is no atomic displacement of Zr/Ti ions and no spontaneous polarization is established. Below TC the adopted structure depends on the Zr/Ti content. For high Ti content the ferroelectric phase is tetragonal and the polarisation vectors are along the directions of the pseudocube; therefore, there are six possible orientations. The Zr/Ti ions are displaced along the tetragonal caxis. Because the centres of the positive and negative charge are no longer coincident, a dipole is created originating a spontaneous polarization. For higher Zr concentrations the ferroelectric phase is rhombohedral and the polarisation is along or the body diagonals of the pseudocube, which are eight. In this case the lattice distortion is accompanied by the movement of Zr/Ti ions towards the face centre of the oxygen octahedral. Similarly, the centres of positive and negative charges are displaced and a spontaneous polarization builds up. Below the Zr/Ti ratio of 95/5 the solid solution is antiferroelectric with an orthorhombic phase. Another significant feature of the PZ-PT phase diagram is the existence of an abrupt structural change, with composition at constant temperature the so-called morphotropic phase boundary (MPB). It occurs close to the composition where PZ:PT is 1:1, namely Pb(Zr0.52Ti0.48)O3 - PZT 52/48. MPB compositions exhibit enhanced dielectric and piezoelectric properties [8]. It is believed that due to the 14 possible directions of polarisation (eight [111] directions for the rhombohedral phase and six [001] directions for the tetragonal phase) in MPB compositions the reorientation of the polar axis is facilitated and the electrical properties enhanced [8]. Figure 8 depicts the coupling coefficient kp and the dielectric permittivity İr across the PZ – PT solid solution showing the maximization of the properties in the MPB region. Hence, these compositions are technologically very important and PZT MPB compositions are the most exploited for piezoelectric applications.

Figure 8. Coupling coefficient kp and dielectric permittivity İr values across the PZT compositional range [11].

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The reasons behind the increase of the dielectric properties at the MPB are not yet completely explained. Several justifications have been proposed for the origin of the maximum of electric properties at MPB [48, 49]. As in other materials the PZT properties may be tailored by doping (additives added in concentrations ≤ 3 %) [18]. Doping PZT with acceptor ions such as K+, Na+ (for A site) and Fe3+, Al3+, Mn3+ (for B site) creates vacancies at the oxygen sublattice and usually have limited solubility in the lattice. Such materials are named hard PZT. The highly mobile oxygen vacancies can be re-oriented under an applied field. The aligned dipoles will provide a field, stabilising the domain structure, reducing most of the properties that are enhanced by domain wall motion. Therefore, hard PZT's are characterized by lower permittivities, smaller dielectric losses and lower piezoelectric coefficients and are more difficult to pole and depole (higher coercive fields, poorly developed hysteresis loops) [18]. Hard PZT are then suitable for applications in which it is required to transmit as much power as possible. In contrast, doping PZT with donor ions such as La3+ (for A site) and Nb5+, Sb5+ (for B site) originates A site vacancies in the lattice and the materials are termed soft PZT. In donor-doped PZT the content of oxygen vacancies is minimised reducing the number of domain-stabilising pairs and enhancing domain reorientation. Therefore, soft PZT's possess higher permittivities, higher losses, higher piezoelectric coefficient, maximum coupling factors and are easy to pole and depole (low coercive fields, square hysteresis loops, high remnant polarization). The donor dopants counteract the natural p-type conductivity of PZT and, thus, increase the electrical resistivity of the materials. They can be used for applications requiring very high piezoelectric properties such as sensors [18]. PZT containing 3 to 12 mol% of La forms a relevant class of functional materials, lead lanthanum zirconate titanate (PLZT), with important dielectric, piezoelectric and electro-optic properties. The solubility of La in the PZT lattice is a function of the composition, namely on the PT content. The PLZT phase diagram is well described in [18]. The incorporation of La considerably improves the properties of PZT, explicitly increases the squareness of the hysteresis loop, decreases the coercive field, increases the dielectric permittivity, maximum coupling coefficients and mechanical compliance and enhances the optical transparency. It is believed that the observed transparency is related on one hand to the lowering of the cell distortion caused by La introduction that reduces the optical anisotropy and on the other hand to the uniform grain growth and densification of as a pore free microstructure for La doped materials [18]. Thus, PLZT is very important as an electroptical functional material. Perovskite lead-based relaxor ferroelectrics (Pb(B´1-xB´´x)O3 exhibit excellent dielectric properties, a broad dielectric maximum and large piezoelectric and electrostrictive coefficients. Together with the possibility of designing the properties by solid solution formation with other ABO3 members, relaxor ferroelectric materials are very attractive for multilayer capacitors, piezoelectric transducers and electrostrictive actuator applications to operate under different frequency and temperature conditions [50]. It is well know that the fabrication of monophasic lead based relaxors is difficult. A pyrochlore type phase (A2B2O7-δ), with low dielectric permittivity, precedes the formation of the perovskite phase. Due to its high stability the pyrochlore is difficult to eliminate and, depending on the systems, it remains as a second phase, degrading the

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final properties. Hence, from the technological point of view, is crucial to obtain a single perovskite phase material in these systems [29]. Niobate relaxor compounds, Pb(B1-xNbx)O3, typified by PMN are the most studied. In opposition lead based tantalate (Pb(B1-xTax)O3) and tungstate (Pb(B1-xWx)O3), in which B is Zn2+, Mg2+, and Ni2+, with a low temperature of the dielectric maximum are considerably less studied. However, these are important candidates for utilization in cryogenic conditions such as low temperature capacitors and actuators for space applications [51]. One of them, Pb(Fe2/3W1/3)O3 (PFW) gained higher attention in recent years due to its high dielectric permittivity and low sintering temperature (~950 ºC). These characteristics make it a good material for multilayer capacitors with inexpensive low temperature melting electrodes (such as Ag-Pd alloys). The solid solution formation with PT modifies the relaxor properties of PFW and a dielectric response closer to a classical ferroelectric behaviour is observed [52]. A composition-induced transition from pseudocubic relaxor to a tetragonal ferroelectric state was found in (1-x)PbFe2/3W1/3O3 – xPbTiO3 (PFW–PT) with increasing x. A MPB was reported in the range x = 0.20-0.37 at T = 300 K and x = 0.25-0.35 below 280 K and the phase diagram of this system was proposed [53, 54]. Another interesting feature of PFW is related to the multiferroic magnetoelectric behaviour, i.e. the coexistence of electric and magnetic polarization observed in this material [55]. The coupling between the ferroelectric and magnetic activity provides the possibility to manipulate the magnetic properties through electric fields and vice versa, giving to these materials a wide potentiality for applications in spintronics, multiple state memory elements, or memory devices which use electric and/or magnetic fields for read/write operations [55]. However, the magnetodielectric multiferroism is very rare, being restricted to only a few materials like ferroelectric-ferromagnetic BiMnO3, YMnO3 and Pb2(CoW)O6 and ferroelectric-antiferromagnetic PbFe2/3W1/3O3, PbNi1/3Nb2/3O3 and Pb2(FeTa)O6 [55]. This phenomenon opens a wide range of new possible applications for PFW based materials, although not well studied yet. Moreover, recent studies showed the existence of a core shell structure in PFW-PT system and its effect on the dielectric response [56]. The dependence of the core-shell on the processing parameters was established allowing to tailor the dielectric response of PFW [57]. Bi-layered perovskites are at the present being considered as substitutes for lead based perovskites. The ferroelectric nature of Bi-containing layered perovskite was reported by Smoslenskii and co-workers in the sixties [58]. Nevertheless only recently they become technological important materials due to their almost fatigue free behaviour [59]. Bi-layered perovskites belong to the multilayered Aurivillius family of compounds with the general chemical formula of (Bi2O2)2+ (Am-1BmO3m+1)2- consisting of m perovskite units sandwiched between oxide layers. In SrBi2Ta2O9 (SBT) the perovskitetype groups (SrTa2O7)2- and (Bi2O2)2+ layers are stacked alternately along the pseudotetragonal c axis (Figure 6). The Bi2O2 layers and TaO6 octahedra are considerably distorted and atomic displacements along the a axis give rise to spontaneous polarization. The ferroelectric phase crystallizes with the orthorhombic structure and transforms to the paraelectric phase (with a tetragonal structure) near 600 K [60]. The bismuth layer structure can have perovskite blocks with different thickness depending on the stoichiometry and it was recently shown that there is a dependence of the polarization direction (a-b plane or c axis) on the number of octahedral thickness [17]. The few allowed directions for the spontaneous polarization seems to be the reason for the lower polarization values when compared with PZT. Consequently the

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maximization of the ferroelectric properties and the utilization of bismuth-layered materials in memory devices depend on the ability of producing highly oriented materials. Moreover the structural and physical properties of Bi-layered materials are not completely understood.

4. Functional Materials Processing Technologies The technology to fabricate functional materials depends mainly on their final applications. Current applications require functional materials either in a bulk (polycrystalline ceramics, texturized ceramics and single crystals) or in a film form (thin and thick films). Moreover among each technology the specific application will also define the details of the integration of the functional material into the device. The general processing of polycrystalline ceramic functional materials includes the following steps [11]: (1) synthesis of the powder; (2) milling, usually with additive mixing (lubricant, plasticizers, binders); (3) drying; (4) forming (with the simultaneous application of metallic electrodes for multilayer structures); (5) firing; (6) finishing (including slicing, lapping, polishing, electroding, encapsulation and poling) and (7) evaluation. The critical steps, those determine the microstructure and the final properties of the ceramics are the synthesis of the powders, in which the preparation of monophasic powders, with fine and homogeneous particle size distribution is required, and the sintering, at which the reduction of the porosity is crucial. Details on the preparation technologies of functional ceramics can be found in [11, 12]. There are now innumerable methods to produce synthetic oxide and non-oxide powders with the required high quality characteristics in terms of size, shape and purity. Powder precursors for the fabrication of functional materials can be synthesised by solid, liquid or vapour reactions [61]. Solid-solid reactions commonly involve the mixture of the precursors (carbonates, nitrates, oxalates, etc) followed by a thermal treatment (800-900 ºC) at atmospheric pressure to obtain the desired product. Due to its simplicity and low costs solid state reaction methods are often used and numerous functional materials compositions have been prepared, such as BT and lead based perovskite. For the case of lead perovskite it is well known that the preparation through this methodology usually originates a multiphasic material. The perovksite phase formation occurs via a sequential phase formation process in which intermediate phases, such as pyrochlore type ones, are formed and transformed into the perovskite phase for higher temperatures [29]. Due to the high temperatures required for the nucleation and growth of the new phase, the obtained powders are usually aggregated and require a milling step after. Several comminution techniques, which allow milling until fine particle sizes (tens of nanometers) in shortt periods of time and thus avoiding the undesirable contamination by the milling step, are now available. Fluid impact mills are one of such examples [62]. Recently, comminution techniques have also been used for synthesis of complex oxides. The high energy liberated by the impact of hard milling media can be used to promote the solid-solid reaction between precursors, nucleating new solid phases. The synthesis of complex perovskite oxides has been reported by high energy ball milling, also known as “mechanochemical synthesis” [63-66]. Another modern solid-state reaction technique is the microwave-assisted synthesis (MAS).

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Microwave dielectric heating is related to the ability of substances to absorb microwave energy and to convert it into heat so as to obtain the required energy for the reaction. The attractive features of this technology include the very short time required to complete the reaction (sometimes chemical reaction times can be reduced from hours to minutes with all the related benefits); the reduction of side reactions; the yield increase and the improvement of reproducibility. R&D activities on MAS synthesis of materials were reviewed by Mingos [67], Rao [68] and Adam [69]. If powders with high purity, homogeneity and reactivity are required then other methodologies, although more expensive, should be used. Solution synthesis involves the intimate mixture of the liquid precursors (inorganic, organic and alcoholic solutions) at a molecular level in order to obtain a homogeneous solution at the atomic level which by precipitation will originate a solid precursor. Solution techniques include among others the precipitation and co-precipitation, hydrothermal synthesis, sol-gel, emulsion process, molten salt synthesis and spray pyrolisis. In each of these techniques the nucleation and growth of the solid phase occurs through a reaction in the liquid phase. Due to the high level of atomic and/or molecular homogeneity attained in solution these techniques give rise to high chemical homogeneous solid products, with a high purity and controlled particle morphology and size ( 10-5 Torr) to the substrate, in which the nucleation and growth of the film occur atomistically. Depending on how the particles (atoms or ions) are removed from the target, the following PVD techniques are considered: rf sputtering, ion beam sputtering, electron beam evaporation and laser ablation, among others. The former allows for careful control of film thickness and orientation and compatibility with the semiconductor integrated circuit processing.

24 TABLE I. Thin and thick film deposition techniques (adapted form Haertling [18].)

Thin Film Techniques Physical Vapour Deposition Sputtering (rf magnetron, dc, ion beam) Evaporation (e-beam, resistance, molecular beam epitaxy) Laser Ablation

Thick Films Techniques Tape casting Screen printing Electrophoretic deposition Hybrid sol-gel technique

Chemical Deposition Chemical Vapour Deposition MOCVD (Metal-organic CVD) PECVD (plasma-enhanced CVD) LPCVD (low pressure CVD) Chemical Solution Deposition Sol-gel (solution gelation) MOD (metallorganic deposition)

The difficulty in controlling the stoichiometry of multicomponent films, the slow rates of deposition (normally around 1 Å/s), the need for high-temperature postdeposition crystallization annealing and the high cost related with equipment acquisition and maintenance are the main disadvantages of these methods. For more details on the PVD methods the reader is referred to several text books [71]. Chemical methods allow higher deposition rates, good stoichiometry control, and the production of large area defect-free films and have lower equipment-related costs. Chemical vapour deposition (CVD) is very attractive for industrial manufacturing of functional films. However, the limited availability and toxicity of sources off precursors for functional materials restricts the use of this technology. A comparison between PVD and metalorganic chemical vapour deposition (MOCVD) for the processing of ferroelectric thin films and heterostructures can be found in [75]. On the other hand, chemical solution deposition methods, specially sol-gel, have been increasingly used for the preparation of films of functional materials. Chemical techniques does not require vacuum ambience, are cheaper and faster, allow for a good stoichiometry control and production of large area defect-free films and often produce films with better properties, although the texture degree of the film is inferior. Wet chemical methods entail the preparation of the solution, the deposition of the solution onto the substrate by dip- or spin-coating and the subsequent thermal treatment of the deposited layer to remove the organics and to achieve crystallization and densification of the coatings. Wet processes comprise solgel, metalorganic decomposition (MOD), electrochemical reaction and hydrothermal routes. Comprehensive review texts on solution deposition of piezo- and ferroelecrtic materials can be found in the literature [76, 77-81]. The described methods are mostly suitable for thin films (1 to 5 µm) preparation. The preparation of thicker films (>5 µm) using these techniques is possible although the process becomes time consuming and cost-ineffective. Also the probability of creating defects (cracks, pores, and inclusions) in the films increases with the number of layers. Hence, the preparation of thicker films requires different methodologies, as indicated in table 1. Thick films technologies are mostly based on the densification of powder films. Films with thickness in the range 5-500 µm can be prepared by these methods [82, 83]. A suspension of small powder particles (slurry) is deposited by tape casting, screen printing, jet printing or electrophoretic deposition onto the substrate and dense films are

25

attained after sintering. As the liquid phase is removed by evaporation the particles rearrange and the film shrinks. The final density is attained after sintering which is a critical step in the slurry-based technologies. The temperatures t required to densify lead based functional materials, above 850 ºC, can cause losses of volatile elements (Pb), densification problems and chemical partition of the composition. Usually sintering additives (lead based glass frit or lead borosilicate) are necessary although the final properties of the film might be deteriorated [84-87]. Engineered processes to produce dense thick films at low temperatures are then needed. Tape casting is an ordinary method to fabricate laminated thick layers (in the range of 10-500 µm) in multilayer structures, such as multilayer capacitors and actuators [88-91]. However, the method is not easily compatible with the deposition onto rigid substrates and the suspension and firing without deflection are very difficult. The capability of screen-printing functional films on a variety of flexible or rigid substrates (ceramic, metallic, polymeric) is an important advantage over the tape casting method. Screen-printing method is very well suited to prepare functional thick films with thickness ranging from 10 to 30 µm [9294]. In this technique the interactions between adhesion and sintering promoters, film material and substrate are important, especially for Si-based devices, and might limit the final properties of the device. Electrophoretic deposition (EPD) is another slurry-based coating process to fabricate dense thick films. In EPD, colloidal, charged particles deposit from a stable suspension onto an oppositely charged electrode upon the application of a dc electric t field. Comprehensive reviews on EPD can be found in [9597]. An important advantage of EPD is the possibility of using nanosized powders well dispersed in the suspension, which allow low sintering temperature processes and compatibility with several substrates, making EPD an appropriate process to fabricate lead based functional materials thick layers on silicon based and other substrates. Several works report the EPD preparation of BaTiO3 [98, 99] and PZT thick films [100103]. Almost no investigations have been reported yet on alternative lead free compositions. A further alternative methodology to prepare dense thick coatings compatible with Si technology has emerged as a low- temperature process, named composite (or hybrid) sol-gel technique [104]. It combines the sol-gel processing with fine powder particles in which the amorphous gel is used as the “cement” to glue the sintered piezoelectric powder. The composite suspension is then coated onto the substrate using spin-, dip- or spray-coating techniques. The main advantages over the classical slurry-based approach include [83]: (i) the reduction of the film m shrinkage during firing and crystallization which reduces the tendency to crack formation; (ii) the increase of the viscosity of the deposits which allows relatively thick individual layers to be deposited; (iii) the promotion of the crystallization of the sol-gel matrix due to the nucleation aids role of the nanopowder particles. This allows thick films to be fired at reduced temperatures. The drawback of this technique is the inevitable a porosity. Hence modifications of the hydrid sol gel process have been undertaken. The inclusion of infiltration steps in which the sol is infiltrated between the layers was reported for the preparation of PZT thick coatings on Si platinised substrates [105]. A different approach refers the sedimentation of the powder layer by centrifugation combined with the infiltration of the sol-gel solution. With this technique PZT and PMN-PT thick films with enhanced piezo- and ferroelectric properties were fabricate [106, 107].

26

5. Functional Materials Applications The most important applications for the above mentioned functional materials are summarized in table II. The ample range of applications is related to their varied properties (ferro, piro- and piezoelectric, electrostrictive and electro-optical) and to the possibility of designing their properties and producing reliable devices. Some of the applications are more suited for bulk ceramics, while others require films. The increasing trend for device miniaturization has been accompanied by a growing demand of functional films, either thin or thick. Besides the reduction of the device dimension and compatibility with integrated circuit technology, the use of films offers additional advantages: lower operational voltages, higher velocities, possibility of fabrication of complex structures, low processing temperatures and compatibility with a varied range of materials (metals, ceramics, polymers). The final application of a material is directly related to its properties. Some important properties of selected functional materials were summarized by Haertling [18]. Ferroelectric materials are generally characterized by: i) high dielectric permittivity values (200 – 10000), ii) low dielectric loss (0,1-7%); iii) high resistivity (>1013 Ωcm); iv) high to moderate dielectric breakdown (100-120 kVcm-1 for ceramics and 500-800 kVcm-1 for thin films) and v) non-linear electric, elecromecanhical and electro-optic behaviour. Capacitors with great capacitance density as the multilayer ceramic capacitors (MLC) are based on ferroelectric materials with high εr. In MLC thin (~1 µm) layers (>100) of a high εr material are interspersed with thin layers of metallic electrodes, creating minimised associations of parallel capacitors of high capacitance. MLC are made by tape-casting. Firstly, a slurry of the dielectric powder with suitable additives is prepared. Thin green sheets of the ceramic are then tape-casted. On the top of each dielectric layer a metallic electrode (Pd, Ag-Pd, etc.) is screen printed. These steps are repeated until hundreds of sheets are stacked. After lamination the green sheets are diced with the right size. MLC are then sintered. To finish the fabrication cycle terminations for the internal electrodes are applied [12]. Alkaline earth titanates (BT based) were traditional utilised as the dielectric material, but due to the high sintering temperatures noble metals should be used as electrodes, increasing considerably the price of the capacitor. Alternatively relaxor based materials (such as PMN) are now being employed. Besides the high dielectric permittivity (>30000), a broad dielectric permittivity maxima, low dependence of the dielectric properties on the field lead based relaxors can be sintered at T < 1000 °C, thus allowing the use of cheaper metallic alloys as electrodes. The drawbacks of the relaxors formulations are associated with the higher dielectric losses, the higher frequency dependence of the dielectric permittivity and the difficulty in processing monophasic materials. One recent application also based on high εr values of ferroelectric materials are the Dynamic Random Access Memories (DRAMs). A conventional DRAM is composed of a SiO2 (εr~4-7) capacitor in association with transistors and resistors [46]. To increase the storage capacity and decrease the device size ferroelectric thin films with high εr are required. Ta2O5 (εr~22) has been used to fabricate 256 Mb DRAM. However the Gigabit generation entails higher dielectric permittivity materials. Using a high-permittivity

27

material such as BST (εr~200) the capacitor of the 256 Mb DRAM would be realized in a planar structure [46]. Besides BST other high permittivity materials should be considered. An example is the Ba(ZrTi)O3 (BZT), that with εr similar to BST one exhibits improved dielectric loss, leakage and resistance degradation [46]. Further studies on this family of materials are needed. A comprehensive review on DRAMs was recently published [46]. TABLE II. Applications of functional materials [adapted from 18]

BULK

BULK AND FILMS

FILMS

APPLICATIONS

APPLICATIONS

APPLICATIONS

ML Capacitors

Dielectric Capacitors

Non Volatile Memories

Piezo Generators

IR Sensors

Buffer Layers

Pizo Motors

Piezo Sensors and Actuators

Integrated Optics

Piezo Actuators

Electrooptic Shutters

AR Coatings

Electrostrictive Actuators

Electrooptic Displays

PTC Sensors

The above-mentioned applications are based on the dielectric permittivity and not on the ferroelectric response. For such uses a high dielectric permittivity value stable with the temperature is required. However the ferroelectric hysteresis is a unique characteristic that can also be explored from the practical point of view. A ferroelectric material with a square hysteresis loop has stable Pr values for small changes of E and the switching of polarization occurs for high applied field. These are ideal features for the binary code and non-volatile characteristics required for non-volatile ferroelectric random access memories (NVFeRAMs). In the fifties due to the raising computer industry memories with high storage capacity were required. At that time ferroelectrics were already considered as promising candidates, however lack of reliability, fatigue of the switching cycles, imprint, high operation voltages and leakage currents limited their practical implementation. Magnetic and later semiconductor memories were used [108]. More recently due to the ability to prepare high quality films (epitaxially grown, defect free and with controlled stoichiometry) a renewed interest in ferroelectric memories appear. A FeRAM incorporates a ferroelectric thin film as a capacitor to store the data in a non-volatile state, allowing the date to be rewritten fast and frequently [109]. PZT and SrBi2Ta2O7 (SBT) are the materials under consideration for FeRAM applications. For PZT films the oxygen vacancies and the charge injection at the interface ferroelectric/electrode, considered to be responsible for the fatigue of the switching

28

cycles, can be minimized if oxide electrodes are used. For SBT films the degradation is controlled by the (micro-) structure t (specially the Bi layer). From the fatigue point of view SBT has been considered the most appropriated material. However, PZT can be processed at lower temperatures (500 ºC) than SBT, being more suited for Si based technology. Applications of non-volatile memories include a panoply of highly consumable multimedia equipment (digital cameras, video cameras and digital audio) and portable products (mobile phones, notebooks, PalmPCs, among others) [109]. A recent FeRAM application is the rf-operated Smart Card. A comprehensive review on FeRAMS can be found in [109]. Pyroelectricity is the polarization produced due to a small change in temperature as defined in section 3 of this paper. The pyroelectric response of ferroelectric materials is technologically important for thermal detectors, namely infrared detectors, used for the detection and prevention of fires, night vision, security systems, etc. [110]. Single crystals of triglycine sulfate (TGS), LiTaO3, and (Sr,Ba)Nb2O6 are widely used for heat sensing applications. The use of thin films of ferroelectric materials in pyroelectric devices allow, besides the miniaturization and the compatibility with the IC technology already indicated for FeRAMS, the operation at room temperature that constitutes an important technological step further. The use of ferroelectric thin films for pyroelectric devices is also advantageous due to the high cost of growing single crystals. However, the sensibility of the devices should be improved through the utilization of highly oriented films. PbTiO3, (Pb,La)TiO3 and PZT thin films have been used for pyroelectric sensing applications [110]. The high values of kp, d33 and d31 piezoelectric coefficients off ferroelectrics, namely PZT and PLZT account for the enormous number of piezoelectric applications for these materials, such as: nano- , micro- and macro-actuators, micropositioners, ultrasonic motors, dampers, microphones, gas igniters, accelerometers, pressure sensors, among several others. Piezoelectrics are also involved in the fabrication of microelectromechanical (MEMS) and nanoelectromechanical (NEMS) systems, that have an important impact on medicine and bioengineering (DNA and genetic code analysis and synthesis, drug delivery, diagnostics and imaging), bio- and information technologies, avionics and aerospace (nano- and microscaled actuators and sensors, smart reconfigurable geometry wings), automotive systems and transportation (transducers and accelerometers), etc. [111]. On the other hand, the actual interest in electrostrictive devices acting as tunable piezoelectric components is related to the high values of electrostriction characteristic of some ferroelectric relaxors (PMN-PT and PLZT).

6. Future trends in Functional Materials “Nothing is harder to predict than the future t of science and technology. Remember former IBM president Thomas J. Watson's 1943 remark that there was a worldwide market for "about five computers"? And then there was Bill Gates's 1981 prediction that 640 kilobytes of computer memory "ought to be enough for anybody." Why do even brilliant technologists like Watson and Gates have such trouble reading the future? New technologies don't always arise in a linear fashion. When cutting-edge fields of

29

knowledge come into contact, new disciplines can be spawned, and progress can go zooming off in unexpected directions” – in Untangling the Future - By Paul Saffo, Business 2.0 [112]. If history is any guide, the miniaturization and greater functionality will continue to drive the development of functional materials. New functional materials are expected to be developed trough Molecular Engineering with far above performance in comparison with actual ones (see part I of this book by A.I. Kingon). However for traditional functional materials, as those described in this paper, the future developments will look for: (i) tunable devices, implying the use of materials with very low loss at the microwave frequency range; (ii) deep understanding of the size effects in functional materials; (iii) establishment of the relationships between the nanoscale structure with the nano- and macroscale properties; (iv) very low processing temperatures for easy integration with different materials, (v) development of environmetal friendly materials with optimised properties and (vi) development of heterostructures with maximized properties.

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32 77. Schwartz, R.W., Boyle, T.J., Lockwood, S.J., Sinclair, M.B., Dimos, D., and Buchheit, C.D. (1995) SolGel Processing of PZT Thin-Films – A Review of the State-of-the-Art and Process Optimization Strategies, Integr. Ferroelectr. 7, 259-277. 78. Brinker C.J., Hurd A.J., Schunk P.R., Frye G.C., Ashley C.S., (1992) Review of Sol-Gel Thin-Film Formation, J. Non-Cryst. Solids 147, 424-436. 79. Reaney, I.M., Taylor, D.V., and Brooks, K.G. (1998) Ferroelectric PZT thin films by sol-gel deposition, J. Sol-Gel Sci. Techn. 13, 813-820. 80. Brinker, C.J. and Scherer, G.W. (1990) Sol – gel science: the physics and chemistry of sol-gel processing, Academic Press, New York. 81. Tuttle, B.A. and Schwartz, R.W. (1996) Solution deposition of ferroelectric thin films, MRS Bulletin 21, 49-54. 82. Prudenziati, M. (1991) Thick film technology, Sensor Actuat. A-Phys. 25-27, 227-234. 83. Kholkin, A.L., Wu, A. and Vilarinho, P.M. (2004) Piezoelectric Thick Film Composites: Processing and Applications, in Recent Research Developments in Materials Science, Research Sign Post, 5, pp. 1-24. 84. Akiyama, Y., Yamanaka, K., Fujisawa, E., and Kowata, Y. (1999) Development of lead zirconate titanate family thick films on various substrates, Jpn. J. Appl. Phys. 38, 5524-5527. 85. Jeon, Y., Kim, D.G., No, K., Kim, S.J., and Chung, J., (2000) Residual stress analysis of Pt bottom electrodes on ZrO2/SiO2/Si and SiO2/Si substrates for Pb(ZrTi)O3 thick films, Jpn. J. Appl. Phys. 39, 2705-2709. 86. Kubota, T., Tanaka, K., Sakabe, Y. (1999) Formation of Pb(Zr,Ti)O3-Pb(Zn,Nb)O3 system piezoelectric thick films in low-temperature firing process, Jpn. J. Appl. Phys. 38, 5535-5538. 87. Glynne-Jones, P., Beeby, S.P., Dargie, P., Papakostas, T., and White, N.M., (2000) An investigation into the effect of modified firing profiles on the piezoelectric properties of thick-film PZT layers on silicon, Meas. Sci. Technol. 11, 526-531. 88. Lubitz, K., Schuh, C., Steinkopff, T., and Wolff, A. (2002) Materials aspects for reliability and life time of PZT multilayer actuators, Piezoelectric Materials for the end user, Conference notes, Polecer Meeting, Interlaken, February 2002. 89. Nieto, E., Fernandez, J.F., Moure, C., and Duran, P. (1996) Multilayer piezoelectric devices based on PZT, J. Mater. Sci: Mater. Electron. 7, 55-60. 90. Galassi, C., Roncari, E., Capiani, C., and Pinasco, P. (1997) PZT-based suspensions for tape casting, J. Eur. Ceram. Soc. 17, 367-371. 91. Raeder, H., Simon, C., Chartier, T., and Toftegaard, H.L. (1994) Tape casting of zirconia for ionconducting membranes: a study of dispersants, J. Eur. Ceram. Soc. 13, 485-491. 92. Zhang, H.Z., Leppavuori, S., Uusimaki, A., Karjaleinen, P., and Rautioaho, R. (1994) Compositional and structural behaviour of screen-printed PZT thick films during rapid sintering, Ferroelectrics 154, 277282. 93. Fernandez, J.F., Nieto, E., Moure, C., Duran, P., and Newnham, R.E. (1995) Processing and microstructure of porous and dense PZT thick films on Al2O3, J. Mater. Sci. 30, 5399-5404. 94. Tanaka, K., Kubota, T., Sakabe, Y. (2002) Preparation of piezoelectric Pb(Zr,Ti)O3-Pb(Zn1/3Nb2/3)O3 thick films on ZrO2 substrates using low-temperature firing, Sensor Actuat. A-Phys. 96, 179-183. 95. Sarkar, P. and Nicholson, P.S., (1996) Electrophoretic deposition (EPD): mechanisms, kinetics, and applications to ceramics, J. Am. Ceram. Soc. 79, 1987-2002. 96. Van de Biest, O.O. and Vandeperre, L.J. (1999) Electrophoretic deposition of materials, Annu. Rev. Mater. Sci. 29, 327-352. 97. Boccaccini A.R. and Zhitomirsky, I. (2002) Application of electrophoretic and electrolytic deposition techniques in ceramics processing, Curr. Opin. Solid St. M. 6, 251-260. 98. Ngai, M., Yamashita, K., Umegaki, T., and Takuma, Y. (1993) Electrophoretic deposition of ferroelectric barium titanate thick films and their dielectrical properties, J. Am. Ceram. Soc. 76, 253-255. 99. Zhang, J. and Lee, B.I. (2000) Electrophoretic deposition and characterization of micrometer-scale BaTiO3 based X/R dielectric thick films, J. Am. Ceram. Soc. 83, 2417-2422. 100.Van Tassel, J. and Randall, C.A. (1999) Electrophoretic deposition and sintering of thin/thick PZT films, J. Eur. Ceram. Soc. 19, 955-958. 101.Su, B., Ponton C.B. and Button, T.W. (2001) Hydrothermal and electrophoretic deposition of lead zirconate tinate (PZT) films, J. Eur. Ceram. Soc. 21, 1539-1542. 102.Zhang, R.F., Ma, J., and Kong, L.B. (2002) Lead zirconate titanate thick film prepared by electrophoretic deposition from oxide mixture, J. Mat. Res. 17, 933-935. 103.Wu, A., Vilarinho, P.M., and Kingon, A.I. (2004) Electrophoretic deposition of lead zirconate titanate films on metal foils for embedded components, submitted to J. Am. Ceram. Soc.

33 104.Barrow, D.A., Petroff, T.E., and Sayer, M., (1996) US Patent 5,585,136. 105.Kholkine, A.L., Yarmarkin, V., Wu, A., Vilarinho, P.M., and Baptista, J.L. (2000) Thick piezoelectric coatings via modified sol-gel technique, Integr. Ferroelectr. 30, 245-259. 106. Kholkin, A.L., Yarmarkin, V.K., Wu, A., Avdeev, M., Vilarinho, P.M., and Baptista, J.L. (2001) PZT based thik film composites via a modified sol-gel route, J. Europ. Ceram. Soc. 21, 1535-1538. 107.Vilarinho, P.M., Wu, A., Kholkin, A. (2003) Method for the production of ceramic composites thick films by sedimentation and infiltration of sol-gel solutions, Portuguese patent pending n.º 102 909. 108.Auciello, O., Scott, J.F., and Ramesh, R. (1998) The Physics Of Ferroelectric Memories, Physics Today 51, 22-27. 109.Bottger, U., and Summerfelt, S.R. (2003) Ferroelectric Random Access Memories, in R. Waser (ed.) Nanoelectronics and Information Technology: Advanced Electronic Materials and Novel Devices, Willey-VCH, Weinheim, pp. 567-588. 110.Whatmore, R.W., Patel, A., Shorrocks, N.M., and Ainger, F.W. (1990) Ferroelectric Materials for Thermal IR Sensors: State of Art and Perspectives, Ferroelectrics 104, 269-275. 111.Lyshevski, S.E. (2002) MEMS and NEMS: systems, devices, and structures, CRC Press, Boca Raton, USA. 112.Saffo, P. (2002) Untangling the Future, Business 2.0.

SCALING OF SILICON-BASED DEVICES TO SUBMICRON DIMENSIONS

A.I. KINGON Department of Materials Science and Engineering North Carolina State University, Raleigh, NC 27695-7919, USA

Contents 1. 2. 3. 4. 5. 6.

Introduction Scaling of silicon-based devices 2.1. Trends 2.2. Roadblocks at the end of the silicon scaling era Integration of new materials with silicon 3.1. High permittivity gate oxides 3.2. Advanced dielectrics for DRAMs Integration of new functionality to silicon devices 4.1. Ferroelectric random access memories 4.2. Ferroelectric field effect transistors Where silicon meets molecular electronics Conclusions

1. Introduction The purpose of this paper, within the context of the NATO ASI proceedings, is twofold: a) To describe the competition that advanced functional materials, such as molecular devices, face from existing silicon technology. The new materials and devices must compete with devices fabricated using entrenched t semiconductor technology, especially silicon technology. The major problem is that silicon technology is not stationary, but progresses with a relentless momentum. This paper describes the progress in silicon technology from the perspective of scaling to submicron devices, and the expected performance at the end of the ‘silicon scaling era.’ It is at the end of this scaling era that new molecular electronic devices, such as those described by Kelly, and Sagiv and Cohen, are expected to become commercially viable. These molecular devices must outperform the silicon devices of thatt era, not the present generation. b) To describe difficulties and obstacles being faced by silicon electronics as the end of scaling is approached, and to show how new, functional materials are being considered for integration with silicon. These new materials are being incorporated in 35 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 35-50. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

36

order overcome some of the scaling obstacles, or alternatively to provide new functionality to silicon-based devices. The objectives of the paper are ambitious, and in reality the paper cannot cover all of the developments. As a result, some topics are mentioned only briefly, with references provided to recommended papers that represent either good review material, or represent important breakthroughs. Keywords: Scaling, silicon-based microelectronics, nanoelectronics, functional electronics, new materials. 2. Scaling of Silicon-based Devices 2.1. TRENDS The well-known trend of increasing number of devices per silicon chip is known as Moore’s Law [1]. As shown in Figure 1, the density has been increasing at an exponential rate (doubling every 18 months) for an amazing length of time, more than 30 years. The increasing device density has resulted from a regular decrease in the minimum feature size (along with a small increase in the overall chip dimensions). The decreasing feature size is also the origin of a number of related effects. It is the origin of the decreasing cost per function (storage bit or computational operation), the increasing microprocessor clock rate, and the decreasing power per logic or storage operation (leading to longer battery lifetimes in mobile devices). Scaling is also the origin of the dramatically decreasing size of electronic devices.

Figure 1. Diagram showing the exponential growth of the number of cells (logic or memory) per chip, as well as the decrease in the minimum feature size. (Original data from reference 2, figure adapted from reference 3).

This is the historical perspective. But what lies ahead? Since 1992, the Semiconductor Industry Association has been publishing ‘roadmaps,’ predicting the

37

main trends in the semiconductor industry 15 years into the future [2]. These roadmaps have provided a valuable source of guidance for the industry over the years, providing guidance for equipment and materials suppliers, and targets for researchers to provide solutions for upcoming years. An interesting side-note is that the competition bred by the roadmaps has resulted in semiconductor manufacturers frequently ‘beating’ the roadmap predictions, with individual firms working under the assumption that achieving the targets ahead of the roadmap would result in industry leadership. Table 1 compares important characteristics of past, present and future devices, showing the impact of incremental process innovation. The term ‘technology node’ relates to the minimum feature length allowed by the lithographic processes. Of particular note are the goals projected to be achievable by 2013, in particular the characteristic logic device gate lengths (13 nm), processor clock speeds, the sub-onevolt drive voltages, and the 3.1 billion transistors per chip. These can be contrasted with the 1993 values, showing the dramatic projected progress in a short 20 years. Also interesting is the fact that the year 2003 is the year in which the semiconductor industry entered the ‘nanoscale’ technology era, at 100 nm or below. Modern electronics is now nanoelectronics! TABLE I. Projections of the International Roadmap for Semiconductors (2003) Year of Production

2003

2005

2007

2009

2011

2013

2015

2017

100 nm 80 nm 65 nm 50 nm 40 nm 32 nm 25 nm 20 nm Physical Gate Length MPU/ASIC (nm)

45

32

25

20

16

13

10

8.0

DRAM Pitch (nm)

100

80

65

50

40

32

25

20

Equivalent physical Oxide Thickness Tox(nm) m2) High Performance

1.3

1.1

0.9

0.8

0.7

0.6

0.6

0.5

220

520

930

1200

2100

7700 10000 21000

307

487

773

1227

1948

3092

4908

7791

5.0 65

8.6 45

14.9 32

17.2 25

20

109 16

118 13

10

1.6

1.4

1.2

1.0

0.9

0.8

0.8

0.7

transistors) Unmanaged Gate Leakage Power (W/chip) Physical Gate Length low operating power (nm) EOT for low operating power Tox (nm)

0.51 2.22 5.21 6.07 11.7 20.8 25.6 100.0 m2) LOP Note: There are two types of transistors shown - the high performance logic, and the low operating power devices which are used for mobile applications.

The major message of this paper is to researchers involved in the development of new functional devices – both electronics and other functional devices. This message is that if the new devices are competing with Si-enabled devices, then they need to offer entirely new functionality, orr alternatively need to outperform the Si devices of the same generation. Thus, if one is developing molecular electronics, (for example, using the building blocks being developed by Sagiv and Cohen, see Chapter III of these proceedings), then the new molecular electronic devices must outperform the Si-based logic or memory in terms all the critical system performance parameters. Alternatively it must offer dramatically reduced cost or processing advantages, or offer new functionality currently not offered by scaled Si devices. These requirements provide

38

tremendous challenges for those of us involved in the development of new functional materials. The scaling of Si logic devices is mirrored in the scaling of Si-based memory devices, with dynamic random access memories (DRAMs) being the technology leader. Roadmap projections for DRAMs are also shown in Table 1. 2.2. ROADBLOCKS AT THE END OF THE SILICON SCALING ERA

Figure 2. Schematic cross-section of a field effect transistor (FET) drawn to scale. Note the very small thickness dimension of the gate oxide. The channel corresponds to the region in which the electronic carrier concentration is modulated by charging or discharging of the gate oxide capacitor. By modulating this concentration, the conductivity across this region (i.e., between the source and drain) is changed, hence the 'valve' type action. The source and drain regions correspond to the highly doped sections of the wafer at which two of the contacts are made (the third contact to this three-terminal device corresponds to the gate electrode). (The figure is adapted from reference 3).

Having pointed out the powerful capabilities projected for Si-based devices, we must hasten to add that there are major technical obstacles in the way of the achievement of the projected performance goals [3, 4]. Some of these are readily apparent in Table 1. Besides the difficulty in achieving the required feature resolution through lithographic techniques, there are a number of issues associated with the scaling of the gate dimensions, in particular the gate dielectric thickness. Figure 2 shows a schematic diagram of the CMOS transistor, drawn approximately to scale. As the areal dimensions are scaled, it is necessary to simultaneously scale the thickness of the gate dielectric in order to achieve a constant capacitance density to modulate the carriers in the channel to perform the transistor action. The traditional gate dielectric is a thermally grown amorphous silicon dioxide. However, the thickness dimension of the gate dielectric has dramatically decreased, and it is currently down to a few atomic layers thick (1.3 nm in 2003, Table 1). At these thicknesses, the leakage currents increase concomitantly, through an electron tunneling mechanism (Figure 3). In 2003, the gate leakage currents for logic devices can be as high as 220 A/cm2! By the 32 nm technology node of 2013, the leakage currents are projected to be as high as 7.7 x 103 A/cm2. As these represent

39

dissipative losses, the major implication of the high leakage current lies in the large amount of thermal power that must be dissipated (see Table 1). By the year 2013, logic chips are projected to dissipate about 450 W per chip under static conditions! Thermal management is going to become a major issue. Itt should be noted that at the 32 nm node of 2013, the physical thickness of the dielectric, if SiO2, is expected to be less than 0.6 nm. At this thickness, the SiO2 properties are no longer represented by bulk SiO2, as recently described by Muller et al [6].

Figure 3. Plots of the current densities flowing between gate contact and channel through an SiO2 dielectric, for various thicknesses of SiO2 dielectric, as a function of applied voltage (From reference 4). Although calculated and measured for an FET, the same results are relevant for DRAMs. The horizontal lines indicate maximum allowable current densities for the two cases.

Due to the gate leakage issues, over the past few years the silicon technology roadmaps have distinguished three types of product: high performance logic; low power; and low standby power devices. The last two are demanded by mobile consumer applications. These latter devices require lower leakage currents, and as a result the gate dielectric thickness will not scale as fast as the high performance devices. However, despite the lower scaling rate, it is for these low power devices that the issue of leakage current first becomes critical. The severe constraints imposed upon the gate dielectric and the gate dielectric stack (i.e. including the metal contact, Figure 2), has led to the expectation that the SiO2 dielectric will need to be replaced by an alternative, higher permittivity material. This higher permittivity will allow a thicker physical thickness for the same capacitance density. If a number of other requirements are met, then the alternative dielectric may have a lower corresponding leakage current density. This potential has led to a flurry of

40

research over the past few years into alternative materials for the gate stack of mainstream CMOS devices. In the next section we discuss briefly the progress in the development of alternative dielectrics for Si devices. However, the significance of this work should be noted in the context that the strength of the Si-based microelectronics industry has been based upon the very limited number of materials at the heart of the devices. Even the transition from Al wiring at the ‘back-end’ to Cu-based connections required a vast effort over many years. The materials at the ‘front-end,’ i.e. those associated with the heart of the transistor, have been sacrosanct for half a century. Replacing the SiO2 dielectric truly represents a paradigm shift within the industry. Similar roadblocks exist for DRAMs. Once again, the aggressive thickness scaling of the SiO2 dielectric leads to difficulties in meeting the required leakage current requirements, and there is an aggressive search for alternative dielectrics [3, 7-10].

3. Integration of new materials with silicon 3.1. HIGH PERMITTIVITY GATE OXIDES The substantially increasing leakage currents that result from the thickness scaling of SiO2 have two major negative outcomes [11]. First, it directly impacts the on-off characteristics of the transistor. Secondly, the large leakage currents result in extremely large power dissipation, with thermal management of the logic chips being of increasing concern. The need to solve the problem of the high leakage currents in the Si logic generations from approximately 2005 has resulted in a huge R&D effort in developing and incorporating alternative dielectrics with higher permittivities than SiO2. Simply put, the approach of utilizing a higher permittivity material will allow a dielectric with the same capacitance density as SiO2, but which is physically thicker. If certain rules are followed (see below) the higher physical thickness can reduce the leakage currents associated with direct electron tunneling. However, as one can imagine, the challenge of finding a suitable dielectric to replace thermal SiO2 is enormous. The quest may be divided into the following parts: • Selecting candidate dielectrics that are thermodynamically stable in contact with the doped Si channel, and under the subsequent processing conditions • Subsequently selecting candidates which have appropriate dielectric permittivities, band gaps, and band offsets from the Si conduction and valence bands • Selecting candidates that additionally result in acceptable transistor performance. We will discuss these issues in turn. Firstly, the need to ensure that there is no reaction between the alternative gate dielectric and silicon greatly limits the choice of materials, as discussed first by Hubbard and Schlom [12]. Simple materials with reasonable dielectric constants that were initially considered, such as TiO2 and Ta2O5, suffer from a thermodynamically favorable reduction, through the formation of the metal silicide and SiO2. This has resulted in investigations of simple oxides such as ZrO2, HfO2, Al2O3, La2O3, Y2O3, and Gd2O3, although their permittivities are lower than the two mentioned previously. Secondly, the candidate dielectrics t must have low leakage currents at the operating voltages. While there are other factors that may increase the leakage currents, a

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fundamental requirement is that there is an offset of greater than the operating voltages between the conduction and valence bands of the doped Si of the channel and those bands of the dielectric. Initially, it was assumed that if the optical band-gaps were sufficiently large, for example > 3 eV, then the conduction and valence band offsets would be appropriate. Calculations and experiments by Robertson [13], and subsequently others [14], showed that this simple assumption is inadequate, and that the number of candidates is further limited by this requirement. Furthermore, it has become clear that the band offsets are often not controlled t by the intrinsic properties of the materials, but can easily be modified by extrinsic factors that give charge transfer across the Si-dielectric interface. These processing-related issues are currently being investigated. Thirdly, and possibly most critically, there are many factors that can influence the carrier mobility in the Si channel, and thus the transistor performance, including charge at the interfaces, dipoles at the interface, and roughness at the interfaces, all of which can scatter the carriers. Additionally, it has recently emerged that phonon modes in the dielectric, which are responsible for the high permittivities, may also be reducing carrier mobility in the channel through remote phonon scattering [15, 16]. We now briefly present the current materials status. The preferred material at the present time is HfO2 for implementation at approximately the 65 nm technology node. Results for the closely related material, ZrO2, have been similar, but the community has leaned towards the former. HfO2 and ZrO2 have permittivities of approximately 20 - 25, in other words more than 5 times that of SiO2. The dielectrics are typically deposited in amorphous form. However, during post-deposition processing, two important phenomena occur. Firstly, crystallization typically begins as low as 400 ºC, and even very rapid annealing protocols are typically not successful in preventing crystallization. Secondly, if post processing in oxygen partial pressures greater than approximately 10–5 torr, oxygen from the atmosphere diffuses through the dielectric to the Si interface, and forms SiO2. The rate of SiO2 formation, or at least up to the first 0.5 to 1.0 nm, is extremely rapid. It occurs faster than Si oxidation in the absence of the dielectric, as shown by Garfunkel et al. The phenomenon is shown in Figure 4 [17]. The figure also shows the impact of processing at too low an oxygen partial pressure, namely the reduction of the dielectric to form a silicide [17, 18]. It is essential to avoid the reduction, and thus processing under some controlled t oxygen partial pressure is the norm. As a result, the gate stack does not consistt of a simple layer of high K dielectric (e.g. HfO2) directly on Si. Instead, the dielectric stack consists of Si with a thin interfacial layer of SiO2 (e.g. 0.6 nm thick), and with the HfO2 deposited upon that SiO2 layer. Typically, the HfO2 would be of 2-3 nm thickness. The stack therefore ends up with an effective permittivity corresponding to that of approximately 0.8 to 1.3 nm of SiO2. Figure 5 shows the status of various gate dielectrics, with the Equivalent Oxide Thickness (EOT, corresponding to the equivalent thickness of SiO2) plotted against the important parameter, gate leakage current density. It can be seen that the best alternative gate dielectrics are 2-3 orders of magnitude lower in leakage current density than the equivalent SiO2 gate stack.

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as-deposited Si

~ 26 Å

Ig (A/cm m2)

1000 °C - 10-55 Torr

10

3

10

1

-1

10

-3

10

-5

10

~ 28 Å

10-4

-7

10

10-3

-9

10

-3

-2

-1

0

1

2

3

Vg (V)

as-dep 1.6 10-5

10-4 1.2 10-3

1000 °C - 10-4 Torr

SiO2 ZrO 2

0.8 0.4 0 -3

~ 80 Å Si

1000 °C - 10-66 Torr conversion to ZrSi2

~ 55 Å -2

-1

0 1 2 3 extensive interface growth Vg (V) Figure 4. Composite of electrical and microstructural data for ZrO2-based dielectrics processed and different temperatures and oxygen partial pressures. The TEM of the as-deposited film shows 2.6nm physical thickness -5 of ZrO2, with little or no SiO2 apparent at the interface. High temperature annealing at 10 torr causes some SiO2 to form between ZrO2 dielectric and Si, by oxygen diffusion through the dielectric. Annealing at higher oxygen partial pressures results in more SiO2, and lower capacitance densities. On the other hand, annealing at too low an oxygen partial pressure causes reduction to form a silicide, with a resulting large increase in leakage currents [17, 18].

It must be mentioned that the interfacial SiO2 layer (around 0.5 to 1.0 nm in thickness) appears to play an extremely important role in the achievement of acceptable properties. The HfO2 is normally grown on the thin thermally grown SiO2 interface, and this implies that the interface closest to the channel can be processed to be of high quality, with low interface state defect densities. This approach relies upon the prior knowledge of the Si community to achieve the high quality interface. There are several complications that we have to mention. The first is that of crystallization. The gate stack is typically rapidly annealed to temperatures as high as 1100 ºC for a short period (for example – 5 seconds) in order to “activate” the dopants in the silicon. As mentioned previously, the crystallization occurs at temperatures substantially lower than this value, despite the rapid heating and cooling rates. At this point it is not yet known whether the presence of nanocrystalline grains of HfO2 are deleterious to the properties. Certainly, it does not appear that the grain boundaries impact the leakage currents. However, the crystallization is accompanied by some increase in roughness of the second (SiO2-HfO2) interface, and defects associated with the grain boundaries could easily be expected to act as charge traps, with either filling or emptying of the traps resulting in a negative impact upon transistor properties. To date, the identification of defects, and correlation with transistor properties, has not been accomplished. However, it does appear that attempts to thin or remove the SiO2

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interface results in poor transistor properties. This presents a major problem for scaling this system to equivalent oxide thicknesses of 0.5 nm or below, as predicted in the SIA Roadmap.

Figure 5. J vs EOT for some candidate dielectrics

The problem of crystallization may be addressed by alloying HfO2 with an oxide that retards crystallization, such as SiO2 or Al2O3. In the case of SiO2 alloy additions, the deposition of amorphous HfO2-SiO2 (or ZrO2-SiO2) does retard crystallization, but at the expense of permittivity [3]. Additionally, when crystallization does occur, the crystallization products are not the equilibrium products predicted by the binary oxide phase diagram, but instead they are the products of spinodal decomposition. Once again this is HfO2, surrounded by a SiO2-rich amorphous phase. It is nott yet clear whether the presence of the nanocrystalline phase is deleterious for transistor performance. Alloying with Al2O3 has successfully retarded crystallization, with a smaller penalty in permittivity than the SiO2 case [20], but unfortunately transistor properties have been unacceptable. Secondly, replacement of the dielectric also requires replacement of the gate electrode metal (see the transistor diagram in Figure 2), which is currently polysilicon (p-Si). This is primarily due to the loss off capacitance density associated with the additional SiO2 formation at this interface. Replacement of the p-Si must be by one or more metals with (once again) appropriate band offsets to the high permittivity dielectric, as well as thermodynamic stability in contact with the dielectric during

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subsequent processing. This presents an almost impossibly difficult problem, for which a limited number of solutions may be emerging [21, 22]. In summary, it appears that a high permittivity dielectric based upon HfO2 may be feasible, at least for one or two generations around the 65 nm technology node. However, no solutions have yet emerged for the technology nodes of 32 nm and below. Some research groups are investigating the M2O3 oxides, and M2O3-based alloys as potential gate dielectrics for the last technology nodes. Somewhat encouraging results, at least in terms of capacitor properties, have been shown for La2O3, Gd2O3, La2O3-SiO2, and LaAlO3, with some capacitor properties showing EOT values down to the 0.6 to 0.7 nm range, with leakage current densities 2-3 orders of magnitude lower than SiO2 at the same EOT [11, 17, 23]. The major problem faced is that of the interfacial SiO2 layer. In order to achieve the required EOT values around 0.5 nm, it is essential remove, or replace, this interfacial layer that is contributing about 0.5 nm to the total EOT. However, it is also this same interfacial layer that is allowing the best transistor properties to be achieved as the high K materials are incorporated. There is no obvious solution to this conundrum. It is important to note that the thermodynamic behavior of these M2O3 materials is different from that of HfO2 and ZrO2. Instead of phase separation, the M2O3 materials tend to react with SiO2 at the interface, consuming this material and forming an M2O3-SiO2 amorphous solid solution [17, 24, 25]. The case of amorphous LaAlO3 should also be mentioned. It has recently been shown that amorphous LaAlO3 can be deposited on bare Si, and that complete oxidation can be achieved without oxidation of the Si surface. Capacitor properties are encouraging in terms of EOT and critical current densities [11], although transistor properties have not yet been characterized. Finally, it should be noted that the Semiconductor Industry Association Roadmap suggests that the dielectrics required at the end of the silicon scaling era may well be multicomponent epitaxial oxides, rather than amorphous oxides. This represents a major challenge for the materials community. 3.2. ADVANCED DIELECTRICS FOR DRAMS There is a requirement for high permittivity dielectrics for DRAMs, analogous to that for gate stacks in logic devices. The 1T-IC DRAM cell contains a capacitor which acts as a simple charge storage device, with the presence or absence of the stored charge representing the binary information “1” or “0”. The transistor is simply a switch, which allows the information to be written or read. The capacitor has to be continuously rewritten, or ‘refreshed’, as the charge is lost through leakage. Once again, as DRAMs are scaled to smaller dimensions, the required capacitance remains approximately constant, set by the sense amplifier requirements. Thus, scaling requires an increase in either capacitance density, or alternatively an increase in capacitance dielectric area (as the footprint or area available for the DRAM cell inexorably reduces). Traditionally, the capacitor material has been SiO2, or nitrogen substituted SiO2 (SiON), and the dielectric area increased either by etching a large surface area ‘trench’ in the Si substrate, or by creating large surface area ‘built-up’ structures above the substrate (such as ‘crowns,’ ‘fins’ or multiple ‘discs’). At some point, the difficulty in creating these large area structures becomes sufficiently great that the incorporation of alternative, high permittivity dielectrics becomes a tenable alternative. References 7-10 represent comprehensive discussions of the problems, along with the materials solutions.

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All the DRAM manufacturers have extensive R&D programs investigating the intermediate permittivity dielectrics, such as Ta2O5, Al2O3 and HfO2. Ta2O5 has been studied as a DRAM and gate dielectric for over 10 years. It is envisaged that introduction will first occur in the form of metal/dielectric/silicon (MIS) capacitors, with the Ta2O5 deposited directly on Si, followed by later introduction of metal/insulator/metal (MIM) capacitors [26]. Introduction of MIS Ta2O5 capacitors has been much slower than expected, for the following reasons. Ta is generally deposited at low temperatures, with the resulting dielectric being amorphous or nanocrystalline. This microstructure needs to be annealed in an oxidizing atmosphere in order to increase the dielectric constant and reduce the leakage currents. As this is undertaken, oxygen diffusion through the dielectric results in the oxidation of the Si surface. The resultant low permittivity, series-connected SiO2 dielectric reduces the capacitance density and effective permittivity of the stack. Process integration studies have therefore had to find compromise solutions, including: rapid thermal annealing of the Si surface in N2 (RTN at 800 ºC) prior to dielectric deposition in order to increase the oxidation resistance of the surface; and annealing the Ta2O5 dielectric in UV-O3, followed by a rapid thermal anneal in O2. More effort has recently been devoted to Ta2O5-based MIM capacitors for DRAMs [26-32]. The MIM structure has the direct advantage of eliminating depletion effects associated with the Si interface, and also the problem of SiO2 formation at the interfaces. Attractive structures include TiN or Ru electrodes, or combinations thereof. A consortium of Fujitsu and Toshiba has developed a structure in which the electrode forms a high surface area, high aspect ratio cylinder, upon which the dielectric (in particular Ta2O5) can be deposited by chemical vapor deposition. A study of the mechanical stability of the electrode cylinder by the Fujitsu/Toshiba consortium suggests that this Ru/Ta2O5/Ru technology could be scaled beyond the 100 nm technology node. For example, the existing achievable cylinder aspect ratio of 8 (internal diameter to heightt ratio) would allow the 65 nm technology node (4 Gb) to be achieved [28]. Another dielectric that has received attention over the past few years is Al2O5 [3335]. This material may appear a surprising candidate, as it has a lower permittivity than Ta2O5, around 9 to 10. However, there are two factors favoring the material, namely its low leakage current density, and also the availability of high reliability layer-by-layer MOCVD growth (termed “atomic layer deposition” or ALD), along with the ALD equipment suitable for a manufacturing environment. Samsung have demonstrated a fully functional MIS 1 Gb DRAM using Al2O5 deposited by ALD on ‘rough’ Si with TiN plate (top) electrodes [34]. The low process temperature of 350 °C is attractive to minimize the formation of a low permittivity SiO2 layer at the Si interface. No predictions are yet being made regarding the scalability to further generations. However, it should be noted that the high quality conformal coverage makes the material attractive for high aspect ratio geometries, in particular deep trenches. This is demonstrated in a recent publication from Infineon, which demonstrates the use of ALD Al2O5 in trenches for sub-100 nm technology [36]. This work additionally uses rough polysilicon (HSG) in the trench, and uses a “bottle” geometry to maximize the surface area. Recently, DRAM manufacturers have been developing layered or ‘laminate’ dielectrics, in order to gain maximum advantage from each of two different dielectric materials. For example, Al2O5/HfO2 laminate structures have been demonstrated by Samsung [37]. The apparent advantage is that Al can be deposited with minimal

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accompanying growth of a deleterious SiO2 interfacial layer, while HfO2 has a higher dielectric constant (around 22). Looking beyond these simple oxides, towards the end of the Si scaling era, the silicon industry Roadmap suggests that DRAMs will need to incorporate the very high permittivity perovskite solid solution (Ba,Sr)TiO3 (or ‘BST’). This material was extensively studied during the decade of the 90’s [8-10]. 4. Integration of new functionality to silicon devices 4.1. FERROELECTRIC RANDOM ACCESS MEMORIES One of the important technological fields in which advanced materials have been integrated with Si integrated circuits in order to provide additional functionality is that of ferroelectric random access memories (FERAM) [38]. This memory device is analogous to a DRAM cell in that each cell consists of an access transistor and a capacitor. In this case, the capacitor dielectric is also a ferroelectric material (see Chapter 1). The access transistor allows a short pulse to be uniquely applied to the ferroelectric capacitor of the cell, writing (i.e. ‘poling’) the ferroelectric capacitor in one of two possible net polar orientations, corresponding to bit states ‘0’ or ‘1.’ The cell is read by again applying a voltage pulse to the capacitor, and monitoring the resultant current pulse through a sense amplifier. The current pulse corresponds to two possible cases. If no switching occurs, because the pulse is the same sign as the pulse that wrote the bit, then the current pulse corresponds to capacitor charging only. If, however, the read voltage pulse is of opposite sign to the write pulse, the resultant current pulse will contain the additional current corresponding to the polarization switching charge. The two cases can therefore be discriminated, i.e. the state of the cell can be read. It should be noted that this is a ‘destructive’ readout memory device. This is important, as both read and write cycles must therefore be included in determining the number of switching cycles that the memory must endure to achieve reliability and lifetime goals. The reason for the ongoing interest in ferroelectric memory lies in the unique combination of properties that it offers. In particular, the memory can in principle be as fast and as dense as the DRAM, but has the advantage of non-volatility. In comparison with other nonvolatile memories such as EEPROM or Flash, FERAM has shorter write times, and lower write voltages. The primary ferroelectric materials which are utilized in ferroelectric memories are the Pb(Zr,Ti)O3 (PZT) solid solution system, and modified SrBi2Ta2O9. These materials have been discussed in Chapter 1. The importance of scanning probe techniques for the characterization of the FERAM devices is discussed in the Chapter by Gruverman. The FERAM devices represent a case where complex, multifunctional oxides have been successfully integrated into a silicon integrated circuit. However, it should be noted that it has taken many years, almost 20 years, to get FERAM into commercial production. Furthermore, there are two important issues to point out. Firstly, these ferroelectric memories can be considered backend devices, rather than frontend. This means that the access transistors are processed first, followed by a relatively thick dielectric isolation layer, which protects the silicon frontend from potential contamination by elements from the ferroelectric. The ferroelectric layer is then deposited, and the ferroelectric capacitor stack defined. The ferroelectric stack can even

47

be processed in a portion of the facility separated from the main portion of the silicon fabrication process, in order to minimize contamination. Secondly, the properties of the ferroelectric memories are sufficiently unique that there is a niche market for low density devices, memories which a many generations behind in terms of technology node. If, as in the case of logic and DRAM devices, the oxides would have had to be integrated at the then-current generation, it is certain that the complexity of the materials and integration challenges would have doomed the project to failure. This provides a useful lesson for the integration of the new functional molecular materials with silicon substrates: the community should be aiming at devices that have substantially different, but useful properties, to those of current silicon semiconductor devices. These devices that will address unique applications, may have a market even at low density or low overall performance. 4.2. FERROELECTRIC FIELD EFFECT TRANSISTORS Another memory concept that has been discussed and developed for over 45 years is the ferroelectric field effect transistor (abbreviated FEFET or FEMFET). The concept is simple, replacing the gate dielectric of a field effect transistor by a ferroelectric material. The direction of polarization controls the flatband voltage of the capacitor, i.e. it gives two characteristic voltages at which the transistor can be turned on. This means the transistor can be used as a simple nonvolatile memory, with the added advantage that it has a non-destructive readout (NDRO), i.e. it does not need to be switched in order to read the bit state. The other major advantage is that it is a 1-transistor (1T) memory, and in principal it therefore is much smaller than any other memory type. The early work utilized bismuth titanate films, and uncovered a significant difficulty, namely the fact that the turn-on voltage, particularly in one orientation, changed with time. The origin of the phenomenon is clear – the polarization must be screened, and there are several competing sources of screening charge other than the desired carriers from the transistor channel, viz., charge which has leaked through the ferroelectric gate, or mobile charge from within the ferroelectric gate. The result is that the screening process occurs over a long time period (days and months), with a concomitant change in the turn-on voltage (i.e. the characteristic memory window). This is known as a retention problem. The problem is usually worse for one polarization, i.e. that direction for which screening requires minority carriers from the channel. Research by Westinghouse in the ‘80s and early nineties focused upon the problem of retention by addressing dielectric and interface quality. They moved from Bi4Ti3O12 to PZT and then to MBE-grown ferroelectric BaMgF4, but were unsuccessful in solving the retention problem [39]. More recently, research on the device type has renewed, particularly in Japan [40]. Approaches have included the incorporation of an additional insulator layer between ferroelectric and Si channel (MFIS-type), or even a floating gate (MFMIS-type); the matching of charge requirements at all interfaces; and minimization of leakage currents [40]. Retention of a few days has been achieved. The 1T ferroelectric memory thus remains tantalizing: the obstacles continue to loom large, but success would promise in a revolution in the microelectronics industry. A recent and interesting variation of a single transistor memory is the ferroelectric resistance tunnel junction that has been discussed by the group from Aachen in Germany [41].

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5. Where silicon meets molecular electronics An important question remains, which is whether the incorporation of the new materials will open up silicon technology to the further incorporation of new molecular functional materials directly into the Si chip, or subsequently to the adoption of Si-molecular material hybrid circuits. The major point of this paper is that it is important not to attempt to develop molecular materials which are the direct analogues of existing Si logic or memory devices – this is a monumental task which will not yield competitive devices without enormous resource allocation. This is particularly true for the case of three-terminal devices required for logic, and which also require some ‘gain’, i.e. the output power should be larger than the control power. This is proving extremely difficult to achieve with organic-based materials. Instead, in the short and medium term, the primary attention should be on investigating alternative molecular devices that offer new functionality, and open up new applications. By avoiding direct competition with the incumbents, the future of molecular materials will be bright. 6. Conclusions In this chapter we have emphasized the fact that Si-based devices continue to scale to smaller and smaller dimensions, in accordance with the semiconductor industry roadmap. With this scaling, chips become faster, cheaper, and more powerful, with more devices packed onto each chip. The scaling is at the point that, during 2003, the most advanced logic devices transitioned into the nanoelectronics era, with characteristic length dimensions less than 100 nm. However, obstacles associated with scaling have resulted in the unprecedented consideration of the replacement of the fundamental materials, in particular the replacement of the SiO2 gate oxide within the CMOS transistor. This is a trend that may result in many more multifunctional materials being integrated into mainstream silicon devices. Similarly, molecular materials may be integrated with Si semiconductors to yield hybrid devices, with new and unique functionality for as-yet un-thought-of applications. Acknowledgements The author would like to thank NSF and SRC for financial support for research in the topics described in this chapter. References 1. 2. 3.

Moore, G.E. (1975) Progress in digital integrated electronics, International Electron Devices Meeting 1975, Technical digest, pp. 11-13. International Technology Roadmap for Semiconductors, Semiconductor Industry Association, 1992, 1995, 1997, 1999, 2001, and 2003 editions. (url for 2001 edition is http://public.itrs.net) Kingon, A.I., Maria J.-P., and Streiffer, S.K. (2000) Alternative dielectrics to silicon dioxide for memory and logic devices, Nature, 406, 1032-1038.

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

12. 13. 14. 15. 16. 17.

18.

19. 20. 21. 22.

23. 24. 25. 26. 27.

28.

Lo, S.-H., Buchanan, D.A., Taur, Y., and Wang, W. (1997) Quantum-mechanical modeling of electron tunneling current from the inversion layer of ultra-thin-oxide nMOSFET’s, IEEE Electron Device Letters 18, 209-211. Data from the 2003 SIA Roadmap for Semiconductors. Muller, D.A., Sorsch, T., Moccio, S., Baumann, F.H., Evans-Lutterodt, K., and Timp, G. (1999) The electronic structure of the atomic scale of ultra-thin gate oxides, Nature 399, 758-761. Schroeder, H. and Kingon, A.I. (2003) High-Permittivity Materials for DRAMs, in R. Waser (ed.), Nanoelectronics and Information Technology, Wiley-VCH Verlag GmbH & Co., pp. 539-563. Summerfelt, S.R. (1997) (Ba,Sr)TiO3 Thin Films for DRAM's, in R. Ramesh (ed.), Thin Film Ferroelectric Materials and Devices, Kluwer Academic Publishers, Boston, pp. 1-42. Kotecki, D.E., (1997) A review of high dielectric materials for DRAM applications, Integr. Ferroel. 16, 1-19. Kotecki, D.E., Baniecki, J.D., Shen, H., et al. (1999) (Ba,Sr)TiO3 dielectrics for future stacked-capacitor DRAM, IBM J. Res. Develop. 43, 367-382. Osburn, C.M., Campbell, S.A., Eisenbraun, E., Garfunkel, E., Gustafson, T., Kingon, A., Kwong, D.-L., Lee, J., Lucovsky, G., Ma, T.P., Maria, J.P., Misra, V., Parsons, G., Schlom, D., and Stemmer, S. (2004) Materials and processes for high K gate stacks, to be published in IFST. Hubbard, K.J. and Schlom, D.G., (1996) Thermodynamic stability of binary oxides in contact with silicon, J. Mater. Res. 11, 2757–2776. Robertson, J. (2002) Electronic structure and band offsets of high-dielectric-constant gate oxides, Mater. Res. Bull. 27, 217-221. For example: Lucovsky, G. (2003) Electronic structure of transition metal/rare earth high-K gate dielectrics: interfacial band alignments and intrinsic defects, Microelectron. Reliab. 43, 1417-1426. Zhu, W., Han, J.-P., and Ma, T.P. (2004) Mobility measurement and degradation mechanisms of MOSFETs made with ultrathin high-k dielectrics, IEEE Trans. El. Dev. 51, 98-105. Fischetti, M., Neumayer, D., and Cartier, E. (2001) Effective electron mobility in Si inversion layers in MOS systems with a high-k insulator: the role of remote phonon scattering, J. Appl. Phus. 90, 4587-4608. Maria, J.-P., Wicaksana, D., Kingon, A.I., Busch, B., Schulte, H., Garfunkel E., and Gustafsson, T., (2001) High temperature stability in lanthanum and zirconium-based gate dielectrics, J. Appl. Phys. 90, 3476-3482. Stemmer, S., Chen, Z., Keding, R., Maria, J.-P., Wicaksana, D., and Kingon, A.I., (2002) Stability of ZrO2 layers on Si (001) during high temperature anneals under reduced oxygen partial pressures, J. Appl. Phys. 92, 82-86. Data compiled by C M Osburn, NCSU, and privately communicated. Chen, P.J., Cartier, E., Carter, R.J., et al. (2002) Thermal stability and scalability of Zr-aluminate-based high-k gate stacks, 2002 Symposium on VLSI Technology Digest of Technical Papers, pp. 192-193. Misra, V., Lucovsky, G., and Parsons, G. (2002) Issues in high-k gate stack interfaces, Mater. Res. Bull. 27, 212-216. Zhong, H., Hong, S.N., Suh, Y.-S., Lazar, H., Heuss, G., and Misra, V. (2001) Properties of Ru-Ta alloys as gate electrodes for NMOS and PMOS devices, 2001 Symposium on VLSI Technology, Digest of Technical Papers, pp. 49-53. Guha, S., Cartier, E., Gribelyuk, M.A., Bojarczuk N.A., and Copel, M.C., (2000) Atomic beam deposition of lanthanum- and yttrium-based oxide thin films for gate dielectrics, Appl. Phys. Lett. 77, 2710-2712. Stemmer, S., Maria, J.-P., and Kingon, A.I. (2001) Structure and stability of La2O3/SiO2 layers on Si(001),” Appl. Phys. Lett. 79, 102-104. Copel, M., Cartier, E., Narayanan, V., Reuter, M.C., Guha, S., and Bojarczuk, N. (2002) Characterization of silicate/Si(001) interfaces, Appl. Phys. Lett. 81, 4227-4229. Park Y. and Kim, K. (2001) COB stack DRAM cell technology beyond 100nm technology node, International Electron Devices Meeting 2000, Technical Digest, pp. 391-394. Hiratani, M., Hamada, T., Iijima, S., Ohji, Y., Asano, I., Nakanishi, N. and Kimura, S. (2001) A heteroepitaxial MIM-Ta2O5 capacitor with enhanced dielectric constant for DRAMS of G-bit generation and beyond, 2001 Symposium on VLSI Technology, Digest of Technical Papers, pp. 41-42. Fukuzumi, Y., Suzuki, T., Sato, A., Ishibashi, Y., Hatada, A., Nakamura, K., Tsunoda, K., Fukuda, M., Lin, J., Nakabayashi, M., Minakata, H., Shimada, A., Kurahashi, T., Tomita, H., Matsunaga, D., Hieda, K., Hashimoto, K., Nakamura, S. and Kohyama, Y. (2000) Liner-supported cylinder (LSC) technology to realize Ru/Ta2O5/Ru capacitor for future DRAMs, International Electron Devices Meeting 2000, Technical digest, pp. 793-796.

50 29. Lin, J., Suzuki, T., Minakata, H., Shimada, A., Tsunoda, K., Fukuda, M., Kurahashi, T., Fukuzumi, Y., Hatada, A., Sato, A., Sun, P.H., Ishibashi, Y., Tomita, H., Nishikawa, N., Ito, E., Liu, W.C., Chu, C.M., Suzuki, R., Nakabayashi, M., Matsunaga, D., Hieda, K., Hashimoto, K., Nakamura, S., Kohyama, Y., and Shiah, C.M. (2001) Backend process for cylindrical Ru/Ta2O5/Ru capacitor for future DRAM, Solid-State and Integrated-Circuit Technology, Proceedings, pp. 183-188. 30. Kim, W.D., Kim, J.W., Won, S.J., Nam, S.D., Nam, B.Y., Yoo, C.Y., Park, Y.W. Lee, S.I., and Lee, M.Y. (2000) Development of CVD-Ru/Ta2O5/CVD-TiN capacitor for multigigabit-scale DRAM generation, 2000 Symposium on VLSI Technology, Digest of Technical Papers, pp. 100-101. 31. Nakamura, Y., Asano, I., Hiratani, M., Saito, T., and Goto, H. (2001) Oxidation-resistant amorphous TaN barrier for MIM-Ta2O5 capacitors in giga-bit DRAMs, 2001 Symposium on VLSI Technology, Digest of Technical Papers, pp. 39-40. 32. Takeuchi, M., Inoue, K., Sakao, M., Ssakoh, T., Kitamura, C., Arai, S., Iizuka, T., Yamamoto, T., Shirai, H., Aoki, Y., Ijamada, M., Kubota, R., and Kishi, S. (2001) A 0.151 µm logic based embedded DRAM technology featuring 0.425 µm m2 stacked cell using MIM (Metal-Insulator-Metal) capacitor, 2001 Symposium on VLSI Technology, Digest of Technical Papers, pp. 29-30. 33. Kim, Y.K., Lee, S.H., Choi, S.J., Park, H.B., Seo, Y.D., Chin, K.H., Kim, D., Lim, J.S., Kim, W.D., Nam, K.J., Cho, M.-H., Hwang, K.H., Kim, Y.S., Kim, S.S., Park, Y.W., Moon, J.T., Lee, S.I., and Lee, M.Y., (2000) Novel capacitor technology for high density stand-alone and embedded DRAMs, International Electron Devices Meeting 2000, Technical digest, pp. 369-372. 34. Park, I.-S., Lee, B.T., Choi, S.J., Im, J.S., Lee, S.H., Park, K.Y., Lee, J.W., Hyung, Y.W., Kim, Y.K., Park, H.S., Park, Y.W., Leem, S.I., and Lee, M.Y. (2000) Novel MIS Al2O3 capacitor as a prospective technology for Gbit DRAMs, 2000 Symposium on VLSI Technology, Digest of Technical Papers, pp. 4243. 35. Kim, Y.K., Lee, S.M., Park, I.S., Park, C.S., Lee, S.I., and Lee, M.Y. (1998) Novel poly-Si/Al2O3/poly-Si for high density DRAMs, 1998 Symposium on VLSI Technology, Digest of Technical Papers, pp. 52-53. 36. Lutzen, J., Birner, A., Goldbach, M., Gutsche, M., Hecht, T., Jakschik, S., Orth, A. Sanger, A., Schroeder, U., Seidl, H., Sell, B., and Schumann, D. (2002) Integration of capacitor for sub-100-nm DRAM trench technology, 2002 Symposium on VLSI Technology Digest of Technical Papers, 178-179. 37. Lee, J.-H., Kim, Y.-S., Jung, H.-S., Lee, J.-N.-I., Kang, L.-K., and Suh, K.-P. (2002) Practical next generation solution for stand-alone and embedded DRAM capacitor, 2002 Symposium on VLSI Technology Digest of Technical Papers, 114-115. 38. Bottger, U. and Summerfelt, S. (2003) Ferroelectric Random Access Memories, in R. Waser (ed.), Nanoelectronics and Information Technology, Wiley-VCH Verlag GmbH & Co., pp. 565-588. 39. See for example: Sinharoy, S., Buhay, H., Francombe, m M.H., and Lampe, D.R. (1993) BaMgF4 thin film development and processing for ferroelectric FETs, Integr. Ferroelectr. 3, 217-223. 40. Ishiwara, H. (2001) Recent progress of FET-type ferroelectric memories, Integr. Ferroelectr. 34, 11-20. 41. Fitsilis M., Kohlstad, H. Waser, R., et al (2004) A new concept for using ferroelectric transistors in nonvolatile memories, Integr. Ferroelectr. 60, 45-58.

UNSOLVED PROBLEMS IN FERROELECTRICS FOR SCANNING PROBE MICROSCOPY J.F. SCOTT Symetrix Centre for Ferroics, Earth Sciences Department Cambridge University, Cambridge CB2 3EQ, U.K.

Contents 1.

2.

3.

4.

5.

6.

Introduction and deposition techniques 1.1. Nanophase deposition techniques 1.1.1. Electron-beam direct writing 1.1.2. Focussed-ion-beam processing 1.1.3. Self-assembly Topography of new films 2.1. Hafnia and zirconia 2.1.1. Hafnia HfO2 2.1.2. Zirconia ZrO2 2.2. Zircon ZrSiO4 Ferroelectrically filled porous Si and Al2O3 3.1. Porous silicon 3.2. Misted deposition 3.3. Strontium bismuth tantalate results Coherent nucleation of nano-domains 4.1. Gruverman-Shur data on lead germanate 4.2. E-field model (Shur) 4.3. Ripple model 4.4. Gross-Pitaevski model Perimeter effect 5.1. Fringing field model (Chu et al.) 5.2. Phase transition model (Tagantsev) 5.3. Lead zirconate-titanate data 5.4. Ballistic model (Dawber, Jung, and Scott) Ultra-thin polyvinylidene-trifuoroethylene (PVDF) films 6.1. Langmuir-Blodgett film data of Bune et al. 6.2. Kay-Dunn Theory 6.3. Inhomogeneous nucleation theory (Chandra et al.) 6.4. Screening corrections 6.4.1. Ku and Ullman 6.4.2. Simmons 6.4.3. Black and Welser 6.4.4. Dawber and Scott

51 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 51-73. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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7.

8.

6.4.5. 1/C versus thickness results 6.4.6. Relation to Tilley-Zeks theory Strontium titanate/barium titanate superlattices 7.1. Bowman, Gregg, et al. 7.2. BaTiO3 thin film results 7.3. Curie temperatures 7.4. Polarization directions 7.5. X-Ray results (Rios et al.) Conclusions

1. Introduction and deposition techniques Beginning in the 1950s every large US microelectronics company (Bell Labs, IBM, Ford, RCA, etc.) was involved in ferroelectrics t research. The main driving force was the idea that the +P polarization state and the –P polarization state of a ferroelectric could be used to encode the “1” and “0” of the Boolean algebra in which modern digital computers operate. At that time, however, ferroelectrics were available only as single crystals or rather thick ceramics. Since a typical coercive field for switching a ferroelectric from +P to –P (or vice versa) is ca. 40 kV/cm, a 1-mm thick device would have an operating voltage of 4000 Volts! Moreover, the devices were expensive. Therefore as silicon DRAM (dynamic random access memories) devices developed rapidly, ferroelectric RAMs were left on the back-burner as objects of mere academic novelty. This changed rapidly through the 1980s as silicon oxide films as thin as 20 nm were fabricated in pinhole-free 6” commercial wafer form. At that point the advantages of ferroelectric memories over Si DRAMs was recognized once again: They are nonvolatile (the memory does not need refreshing, like DRAMs, and does not forget if power is interrupted); they are radiation hard, no single event upset – SEU; and they are lighter in weight than Core magnetic memories, and 1000x faster to erase and rewrite than are EEPROMs – electrically erasable programmable read-only memories). The result has been a ferroelectrics renaissance. Ferroelectric RAMs are now used in smart debit cards at the 16 kbit and 64 kbit level and FRAMs up to 4 Mbit (Figs. 1-3); in SONY Playstation 2 (Fig. 4) and telecommunications. The highest density ferroelectric (FE) chips available are 4 Mbit from Samsung (using chemical solution deposition lead zirconate titanate – PZT – ceramics ca. 40 nm in grain size) and 4 Mbit from Panasonic (using strontium bismuth tantalate – SBT). A fully commercial 8 Mbit ferroelectric RAM is scheduled for production by Infineon (Japan) and Toshiba on 1 September 2003, using sputtered PZT. At present the road map for FRAM technology is well established: by 2008 the linewidth requirements are 0.1 microns; technology node is 70 nm; feature size F is 0.13 microns; 256 Mbit is the size; and complete cycle time is 16 ns. Therefore for this technology, nano-scale is not just a trendy buzz-word; it is a very strict imperative.

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Figure 1. Samsung 64 kbit FRAM.

Figure 2. Samsung 512 kbit FRAM block.

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Figure 3. Samsung 4Mbit FRAM cross-section.

As detailed below, many of the nano-scale ferroelectric systems below have yet to be investigated by scanning probe techniques. A primary aim of this book chapter is to focus attention on specific questions relating to them.

Figure 4. Sony Playstation 2 with Toshiba EEPROM and Fujitsu FRAM.

1.1. NANOPHASE DEPOSITION TECHNIQUES The three main techniques for 20 nm – 100 nm ferroelectric cells are: electron-beam direct writing (EBDW); focussed-ion-beam (FIB) deposition; and spontaneous selfpatterning.

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1.1.1. Electron-beam direct writing EBDW consists of the use of a cannibalized SEM electron microscope gun to cut out array patterns of ferroelectric capacitors. The results, illustrated in Figs. 5 and 6, are very impressive, but it must be emphasized that a single pattern, ca. 20 x 20 µm in size, takes about 24-72 hours to fabricate. Therefore this technique is not at present suitable for commercial production.

Figure 5. PZT nanoscale array produced by e-beam direct writing. Array on 1 micron scale (bar at lower right) (M. Alexe, private communication).

Figure 6. PZT single cell at 100 nm scale (M. Alexe, private communication).

These images may be contrasted with the results of focussed ion beam patterning (FIB) in Fig. 7 (PZT), and with self-patterning, which in bismuth-excess SBT and bismuth titanate (both PLD and CVD deposition), Fig. 8.

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Figure 7. PZT focussed ion beam array.

Figure 8. Bismuth oxide nano-electrodes self-patterned on bismuth titanate.

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1.1.2. Focussed-ion-beam processing The FIB processing of nanoscale ferroelectric capacitors is analogous to that of EBDW, except that an ion beam is rastered rather than an electron beam, as illustrated in Fig. 9. Some of the best results have been obtained1,2 at the University of Maryland by Aggarwal et al. (2000) and Ganpule et al. (1999). The resulting depth profile is shown schematically in Fig. 10. Note that lateral sizes as small as 20 nm on edge can be produced in this way. Fig. 11 shows actual devices fabricated in different diameters.

Figure 9. Schematic diagram of the focussed ion beam process (R. Ramesh, private communication).

Figure 10. (a) (left) Schematic cross section of typical FIB depth profile; (b) Integration into a nano-phase FRAM device.

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Figure 11. Top to bottom, FIB PZT devices of successively smaller lateral dimensions: 1.0 micron; 260 nm; 100 nm.

1.1.3. Self-assembly Spontaneous self-patterning of ferroelectrics was invented by3,4 Alexe et al. (1998) and Scott et al. (1998). The switching of these cells was developed by5 Alexe et al. (1999). Unfortunately no scanning probe measurements have been reported on them. The theory of such self-assembly was based upon the original work6 by Andreev (1981), which showed that islands forming on solid state surfaces would be mutually repulsive. This was developed into a full theory by Shchukin et al., originally in a short letter and later in a 1999 monograph.7 The predicted ordering in Shchukin’s work is an array of pyramids with {111} faces aligned along the [100] axes of the underlying Si single-crystal substrate. This is shown theoretically in Fig. 12 and confirmed for bismuth oxide in Fig. 13. Shchukin’s theory is for zero temperature; a finite-temperature thermodynamic model was produced8 by Williams et al. (2000). Important in Williams’ work is a distribution diagram, similar to a phase diagram. Three kinds of surface island structures are formed: small pyramids, larger domes (truncated pyramids), and very large “superdomes”. The number density of each depends upon annealing temperature

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and concentration of wetting material (bismuth in our case). The direct application of the Shchukin-Williams model to bismuth oxide nano-electrodes self-patterned onto SBT or bismuth titanate is underway by Dawber (2003).9 Scanning probe techniques have not been used on these systems yet.

Figure 12. Shchukin-Williams model for pyramid self-patterning

a

b Figure 13. (a). Plan view of spontaneous self-patterning of δ-Bi2O3 nanophase electrodes on the surface of PLD-deposited bismuth titanate. These are square pyramids with edges along the [100], [010] axes of the underlying silicon substrate, atop a metallic Bi wetting layer, and satisfy in detail the theory of Shchukin et 7 al. ; (b) Cross-section.

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2. Topography of new films 2.1. HAFNIA AND ZIRCONIA Hafnia HfO2 and zirconia ZrO2 are under active study for gate oxides in dynamic random access memories and related devices.10 In thin-film form they offer chemistry compatible with Si/SiO2 surfaces and very good conformal coverage. A good recent review is by Wilk et al.;10 see also Misra, Heuss and Zhong, and Lee, Jeon, and Hwang.10 No careful studies of film topography has been made in any of these systems, either by scanning probe spectroscopy or other techniques. In the present chapter we report only preliminary studies on the Hf- and Zr-butoxides. 2.1.1. Hafnia HfO2 We have produced thin film hafnia via mist deposition, using novel precursors: Hafnium and zirconium tertiary butoxides; a) hafnium tri-isopropoxy tetramethylheptanedionate, hafnium dimethylamide, and c) hafnium 2-ethylhexanote. Here we present preliminary results on HfO2 and ZrO2 films deposited using butoxide precursors, Hf(OtBu)4 and Zr(OtBu)4. The resulting film did not have the tetragonal phase often seen in hafnia films, but instead was roughly equal mixtures of orthorhombic, monoclinic, and amorphous, as shown by the XRD data in Fig. 14, from Morrison et al. (2003).11 Other recent work on HfO2 on Si is by Lin et al. (2002).12 The monoclinic and tetragonal phases of both hafnia and zirconia have been analyzed by Quintard et al. (2002).13

Figure 14. XRD of HfO2 film.

2.1.2. Zirconia ZrO2 XRD results for a Zr-deposited film heated at 800 °C for 30 min is shown in figure 6. A weak, broad peak is observed at ca. 2θ = 29.3° and is attributed to the 111 reflection of tetragonal ZrO2. It is clear that the film is not as crystalline as the similar Hf-deposited film, figure 15. This is unusual given the similar chemistries of Zr and Hf and requires

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further investigation. Recent studies of zirconia thin films are by Wang et al. (2002) and Chang and Lin (2001).14

Figure 15. XRD of zirconia film.

2.2. ZIRCON ZrSiO4 Zircon is of current interest as an encapsulation ceramic for plutonium reactor waste15 and other high-level radioactive spent-fuel disposal. XRD of our mist-deposited zircon films is shown in Fig. 16.

Figure 16. XRD of zircon film.

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3. Ferroelectrically filled porous Si and Al2O3 3.1. POROUS SILICON Silicon and alumina photonic crystals have been successfully fabricated and characterized for several years, usually by photolithography (etching of Si is greatly enhanced under illumination). A good review is by Birner et al. (2001).The extension of photonic structures from two-dimensional [2D] to three-dimensional [3D] was initiated by Schilling et al. (2001), using modulated pore diameters. The filling of Si pores with Si3N4 via conventional gas-phase CVD was discussed by Ottow et al. (1996).16 In our studies we replace CVD with a misted deposition system discussed elsewhere.17 This has significant advantages. 3.2. MISTED DEPOSITION Misted deposition is a kind of liquid-phase CVD in which submicron droplets of stoichiometric precursor solutions are delivered by a MHz atomiser onto a substrate. It was first reported17 by McMillan et al. (1992). For pore filling in silicon or alumina this provides a great advantage over CVD that harkens back to the days of Millikan and the Millikan oil-drop experiment. As students of physics will recall, fine droplets of liquids (including rain) are usually charged, with 1e or 2e or…5e per droplet. It appears that, for reasons yet unclear, that the internal walls of pores in Si are charged positively. Therefore the mist droplets are electrostatically attracted to coat these walls. Alternative models relying purely on capillary action and minimisation of surface energy are equally viable hypotheses at this stage. This produces uniform thickness (ca. 40 nm) coatings down the walls of ca. 100-micron pores, a remarkable result for pore aspect ratios used (>25). Results are shown in Figure 17 for SBT from Morrison et al. (2002).11

Figure 17. SEM cross-section of SBT nanotubes.

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3.3. STRONTIUM BISMUTH TANTALATE RESULTS SEM photos of porous Si filled with strontium bismuth tantalate are shown in Fig. 18. Here 4.3-micron outside diameter pores have been ffilled uniformly to a depth of 80 microns, with a resulting wall thickness for the SBT of ca. 40-60 nm. No scanning probe studies of these microstructures have been done yet; nor is it known whether the polarization direction is along the tube or through the tube wall, normal to the surface.

Figure 18. SBT nanotubes, plan view.

4. Coherent nucleation of nano-domains It was discovered in Sverdlovsk two decades ago by Gruverman (1983)18 that nanodomains nucleate coherently in front of macroscopic domain walls in ferroelectrics. This was initially observed in lead germanate, where domains are easily observable optically with high contrast. The key observation is that there is a critical field E0 above which this process is observed. The process has now been observed in several ferroelectrics other than lead germanate, including gadolynium molybdate and lithium niobate. 4.1. GRUVERMAN-SHUR DATA ON LEAD GERMANATE The initial data are summarised in Fig. 19, with scale of about 50 microns per cm. In the figure the large domain is moving downwards at a velocity of ca. 10 cm/s. This is a rather high velocity for a domain wall in a ferroelectric. Note in particular that the macroscopic domain wall is NOT atomically flat, nor even optically flat. It has a distinct wavelength λ.

Figure 19. Photograph of nanodomains nucleating in front of macroscopic domain wall.

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4.2. E-FIELD MODEL (SHUR) It has been suggested by Shur (2002)19 that this nanodomain nucleation requires a critical local field but not a critical domain wall velocity, and that the field will depend strongly upon electrode configurations and geometries. 4.3. RIPPLE MODEL I have proposed elsewhere20 that the data shown in Fig. 19 are due to the velocity (downwards in the figure) exceeding the ripple velocity of capillary waves in the domain wall. Under these conditions the domain wall will be strongly damped by creation of ripples of a precise wavelength. The ripple can be seen in Fig. 19 (moving horizontally) and its wavelength can be visually measured. The abrupt decay of domain wall energies at high fields has been measured recently by Zolotoyabko et al. (2002)21 but no mechanism was suggested by them. The mechanism I hypothesize is different from normal acoustic phonon drag, used by Dawber, Jung and Scott (2002)22 for the PZT perimeter effect, discussed below. It is closely analogous to Cerenkov radiation or “bow waves” from a ship: whenever the speed of an object through a continuous medium exceeds the velocity of waves in that medium, energy can be delivered into the waves. In our case, when the domain wall velocity exceeds the speed of domain wall ripples (dependent upon wavelength, which in turn is determined by other parameters, such as distance between pinning sites), a new relaxation process sets in. Since the ripple velocity in a domain wall is given by v = (Tk/ρ), where T is the surface energy of the domain (ca. 7 ergs/cm2 in most oxide ferroelectrics), and ρ is the density (ca. 7 gm/cc), there are no adjustable parameters in the model (k = 2π/λ is measured visually). For λ = 100 microns, as shown, v (critical) in Pb5Ge3O11 will be about 10 cm/sec, as measured experimentally by Gruverman (1983-6). No scanning probe techniques have been used to test these models or to image the nano-domains forming in front of an advancing macroscopic domain. It would be of greatt interest merely to measure the dependence of domain wall curvatures (e.g., wavelengths) upon applied field and/or as a function of velocity. 4.4. GROSS-PITAEVSKI MODEL Domain wall damping into acoustic phonons can be viewed as a particular example of the Gross-Pitaevski model for decay into a boson sea; in this case, the bosons are primarily large wave-vector acoustic phonons near the Brillouin zone boundary. This line of theory is developed by Frisch et al. (1992).23

5. Perimeter effect Dawber, Jung and Scott (2002) have found11 that PZT thin-film capacitors 180 microns to 0.5 microns on edge exhibit a dielectric loss peak with frequency at maximum given by f(max) = bp, (1)

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where p is the cell perimeter and b is a constant that depends on the material.. This holds true for square cells or for rectangular ones, so it is truly dependent upon perimeter and not cell area. 5.1. FRINGING FIELD MODEL (CHU ET Al.) Durkan, Chu et al. have shown24 (2000) that fringing fields in PZT films contribute to significantly to dielectric response but only for aspect ratios (lateral width to thickness) less than 5:1. Since the smallest aspect ratio employed in the present study is 1.1 microns/0.17 microns = ca. 6.5, this is nott important here, but for the new 0.5 x 0.5 micron 32 Mbit FRAMs from Samsung (173 nm thick), it needs to be considered. 5.2. PHASE TRANSITION MODEL (TAGANTSEV) Tagantsev (2002) has presented25 a rather different model to explain “doughnut-like” shapes of back-switching in PZT thin films. In his model there are two phases encountered for thin-film PZT capacitors on substrates, and it is the traversing of the phase boundary into one monoclinic phase that is responsible. While such a phase transition might introduce a large dielectric loss peak, as we measure in our PZT experiments, we can not relate the quantitative details of this model to our data. 5.3. LEAD ZIRCONATE-TITANATE DATA Data are shown in Fig. 20 and the cell structure in Fig. 21. Note that in these cells the top electrode and PZT layer are etched separately, so that a pedestal structure results. This may be important for the mechanism.

Figure 20. Dielectric data versus frequency for PZT cells of different perimeters.

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Figure 21. Capacitor structure for data shown in Fig. 20.

5.4. BALLISTIC MODEL (DAWBER, JUNG, AND SCOTT) Dawber et al. assume that the perimeter effectt results from domains nucleating along the stress-free edges of the PZT (there is a large lattice misfit stress in the interior of the cell) and propagates either towards the centre or around the perimeter. Because the ratio of cell perimeter to average radius is nearly the same (within 10%) for all rectangles, the two mechanisms are thus far indistinguishable. The model assumes that domain walls give up energy to pairs of large-k acoustic phonons, which are known in PZT to have a one-phonon density of states peak at 100 cm-1. This number agrees with quantitative calculations for a model in which the domain wall damping is acoustic phonon drag of form proportional to velocity squared. This is the accepted model for acoustic phonon gain (or loss) in semiconductors and also the model for damping of a mechanical object moving through a viscous continuum. No scanning probe techniques have been yet used to look for the domain motion in these PZT perimeters. 6. Ultra-thin polyvinylidene-trifuoroethylene (PVDF) films 6.1. LANGMUIR-BLODGETT FILM DATA OF BUNE ET AL. In 1998 Bune et al. reported26 a remarkable set of data on polyvinylidene difluoride copolymers with trifluoroethylene (PVDF-TFE). The observed ferroelectric switching in films as thin as 0.9 nm (two molecular layers). This thickness is sufficiently thin that many models of ferroelectrics (e.g., Batra and Silverman, 1973) predict27 that depolarization fields enter the films from the surfaces, destabilize the ferroelectricity, and thus prevent any switching or hysteresis. In addition to disproving such theories, the

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data of Bune and more recently of Ducharme et al. and Fridkin et al.26 showed two other things. First, the switching was extremely slow (seconds) in the thinnest films, compared with hundreds of microseconds in thin films of the same material; this was not readily understood nor had it been predicted. Second, there were two peaks in the dielectric response versus temperature, rather than the one peak at TC expected. Subsequent work by this Moscow-Nebraska group revealed that the dependence of coercive field Ec upon film thickness d was complicated and did not satisfy the KayDunn Law. 6.2. KAY-DUNN THEORY In 1962 Kay and Dunn published28 a model for Ec(d). It predicted that Ec(d) = a d-2/3. This law works surprisingly well for a great variety of ferroelectrics and consequently it has been used widely for 40 years. However, it is derived from assumptions that do not correspond to real materials: First, it assumes 100% homogeneous nucleation, whereas real ferroelectrics generally exhibit 100% inhomogeneous nucleation [as shown by Shur et al. via synchronized electrical t switching pulses and laser flash photography, the nucleation sites are always the same, so that homogeneous nucleation and/or spinodal decomposition is not involved]. Second, Kay and Dunn made a low-field expansion in E that can not be justified for thin films where E is very large even for small applied voltages V. Third, as admitted by Kay and Dunn in their paper, the activation energies extracted from fitting real data to their model are about two orders of magnitude larger than measured independently. 6.3. INHOMOGENEOUS NUCLEATION THEORY (CHANDRA ET AL.) Chandra et al. (2002) have rederived29 the Kay-Dunn Law from scaling arguments, assuming inhomogeneous nucleation and without the small-field assumption. This is a case where the underlying physics transcends the restrictions used in the original derivation (not uncommon in physics). 6.4. SCREENING CORRECTIONS In order to compare real switching data for ferroelectrics with the Kay-Dunn Law or any other law, it is first necessary to correct these data for screening in the metal electrodes. Although the screening length in a metal electrode such as Al, Au, or Pt is only about 0.05 nm, it can play a surprisingly large role. Depending upon whether the switched polarization 2P of the ferroelectric is large or small compared with the displacement vector D (and hence to the dielectric constant e), the effect of electrode screening can be of either sign; that is, it can increase or decrease the apparent coercive field measured. In the former case it simply occurs that a significant percentage of the voltage drop is in the metal and not across the ferroelectric film. The magnitude of the voltage drop percentage in the metal increases as the film thickness decreases. For this reason decreasing the ferroelectric film thickness to extremely small values is not a useful way to maximize capacitance for DRAM capacitors using PZT. Furthermore, using an oxide electrode such as SrRuO3 with long screening lengths, or a semiconductor electrode such as p+ Si will be very disadvantageous.

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Note that it is not necessary to use the Fermi-Thomas model of the metal screening length. More complex models can be used, and in particular for liquid electrodes (such as Hg or ITO) one should revert to the original double-layer models of Helmholtz (1894) or Gouy or Stern (1946). 6.4.1. Ku and Ullman The original theory of voltage drop and screening in the metal electrodes t of a dielectric capacitor were given by Ku and Ullman (1964).30 6.4.2. Simmons Simmons (1965) extended their work and recast it into a simpler analytic form.31 6.4.3. Black and Welser Black and Welser used the models of Ku and Ullman and of Simmons to explain their capacitance data on barium strontium titanate (BST) thin films for DRAM applications.32 However, they made an error, discussed below, of assuming the dielectric constant of the strontium ruthenate electrode is ca. 100. As discussed below, the Simmons model is a free-electron t model, and as such requires ε = 1 in the metal for self-consistency. Hence while the metal screening length enters the equations in an important way, the metal dielectric constant must be unity. 6.4.4. Dawber and Scott Dawqber and Scott (2001) applied these models successfully to tunnelling currents in ferroelectrics.33 Here the electrons must tunnel nott only through a barrier in the ferroelectric (BST, PZT or SBT) near the electrode, but it must tunnel through the depletion layer in the metal electrode as well. The experimental results on samples from Gregg and Bowman fit the screening correction model of Simmons. More recently Dawber and Scott have given a simple derivation of Simmons model which shows clearly why ε = 1 in the metal electrode.9 6.4.5. 1/C versus thickness results One of the classic tests of interface capacitance models is simply to plot reciprocal capacitance 1/C versus film thickness d. For a perfect capacitor on an ideal metal this gives a straight line with intercept through the origin. Many scientists have used this to test for interfacial “dead” layers. Our fitting of the PVDF-TFE data of Ducharme et al. gives a nonzero intercept numerically compatible with the Fermi-Thomas screening length of their aluminum electrodes. It does not permit a zero intercept. Similar results have been demonstrated for other ferroelectric films, most recently by Kingon et al. (2002).34 6.4.6. Relation to Tilley-Zeks theory The theory of polarization dependence P(z) upon depth z in a ferroelectric film was developed by Tilley and Zeks,35 and recently reviewed,36 and is based upon earlier work by Mills37 and by Lubensky and Rubin.38 It is a mean field theory that begins with the addition of polarization gradient terms in the free energy (similar to those used to describe incommensurate ferroelectrics). These gradient terms are phenomenological. In a real ferroelectric they may be dominated by oxygen vacancy gradients near the surfaces or electrodes. The resulting free energy is then minimized to find the

69

equilibrium condition, resulting in two Euler-Lagrange equations. In general P(z) is not a constant throughout the ferroelectric film; as shown in Fig. 22, it either increases at the electrode interfaces (superpolarized surfaces, resulting in TC higher than in bulk) or it decreases (depolarized interfacial surfaces). These two cases are described by an extrapolation length δ, which is negative for depolarized surfaces and positive for superpolarized surfaces. Unfortunately the Tilley-Zeks theory yields neither the sign nor magnitude of δ. I have suggested elsewhere that in real systems d may be the FermiThomas (or other) screening length in the metal electrode. This would require a negative sign and give an exact numerical value 9, e.g. 0.04 nm for aluminum). This connection between the Tilley-Zeks theory and the theory of Ullman-Ku-Simmons is under study. As shown39 by Scott et al. (1988) and Duiker et al. (1990), for first-order phase transitions (as in PVDF-TFE) some values of δ result in separate phase transition temperatures for the interior off the films and for the region near their electrodes. Typically for this to occur δ must be approximately equal to or less than the polarization correlation length in the ferroelectrics, which is very small. This aspect of the extended Tilley-Zeks model seems to fit the two-peaked dielectric data of Bune et al. in PVDFTFE very well. We do NOT think thatt their data imply a domain-free “Landau” switching of the films; despite assertions to the contrary, such switching would occur at higher fields than they employ and be very fast.

Figure 22. Polarization versus depth in a depolarised thin film.

Note that none of the theories discussed in this section (Ku and Ullman; Simmons; Chandra et al.; Dawber and Scott; Tilley and Zeks) include interfacial misfit lattice strain; however, this has been incorporated analytically39 within the Tilley-Zeks model by Zhang, Yin, Zhang and Scott (2001). In real ferroelectric films the increase or decrease in Curie temperature is probably dominated by interface misfit strain and not by the effects described in the strain-free Tilley-Zeks model. This is best illustrated in BaTiO3 thin films, in which Tc can be increased by 300-500K above bulk on SrTiO3 substrates.40

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7. Strontium titanate/barium titanate superlattices 7.1. BOWMAN, GREGG, ET AL. Superlattices of many ferroelectrics have been produced, usually by molecular beam epitaxy, but also by other techniques; see especially the recent work of Bowman, Gregg et al. (Oneill 2000; Corbett 2001).41,42 The interest in these superlattices from a basic science point of view is in their Curie temperatures and their polarization directions. It would be especially useful to understand how these properties, especially TC, depend upon the superlattice periodicity. Here we denote periodicity by, for example, 4/4 or 10/10, where the first number is the number of unit cells of BaTiO3 and the second the number of SrTiO3 unit cells that make up the repeat units. Periodicities from 4/4 to ca. 100/100 have been studied in our work. See also the recent paper by Kim et al. (2002).43 7.2. BaTiO3 THIN FILM RESULTS Thin films of barium titanate exhibit very large upward shifts in Curie temperature if deposited on single-crystal strontium titanate substrates. The smaller lattice constant of strontium titanate compared with that of barium titanate shifts TC up by several hundred degrees, as shown by Li et al. (1999).40 7.3. CURIE TEMPERATURES Generally the Curie temperatures are not known as a function of periodicity. For barium titanate/strontium titanate Jiang et al. (2003)44 have found TC = 276 ºC for 10/10 superlattices and 326 ºC for 30/30 periods. 7.4. POLARIZATION DIRECTIONS Rather surprisingly, Jiang et al. find44 that not only are the strontium titanate layers ferroelectric at room temperature in barium titanate/strontium titanate superlattices of small periodicity (10/10, 30/30…) but that the polarization direction lies along [110] rather than the expected [001]. While this does not minimize the electrostatic ∇D term in the free energy, it apparently minimizes strain. 7.5. X-RAY RESULTS (RIOS ET AL.) Rios et al. (2002) have observed via XRD techniques45 a small 0.04% orthorhombic distortion from the expected tetragonal structure of strontium titanate layers in these superlattices. The lattice constants determined for the 30/30 superlattice are: a = 0.39366(1) nm; b = 0.39352(1) nm; c = 0.38566(2) nm. No scanning probe techniques have been used on these systems. 8. Conclusions In this paper I have tried to review seven different anomalies in ferroelectric thin films which have not yet been studied via scanning probe techniques. I hope that readers will be stimulated to try to elucidate some of them.

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References 1.

2. 3.

4. 5.

6. 7. 8. 9. 10.

11.

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15. 16.

17. 18.

19.

Aggarwal, S., Ganpule, C.S., Jenkins, I.G., Nagaraj, B., Stanishevsky, A., Melngailis, J., Williams, E., and Ramesh, R. (2000) High density ferroelectric memories: Materials, processing and scaling High density ferroelectric memories: Materials, processing and scaling, Integr. Ferroelec. 29, 213-225. Ganpule, C.S., Stanishevsky, A., Su, Q., Aggarwal, S., Melngailis, J., Williams, E., and Ramesh, R. (1999) Scaling of ferroelectric properties in thin films, Appl. Phys. Lett. 75, 409-411. Alexe, M., Scott, J.F., Curran, C., Zakharov, N.D., Hesse, D., and Pignolet, A. (1998) Self-patterning nano-electrodes on ferroelectric thin films for gigabit memory applications, Appl. Phys. Lett. 73, 1592 1594. Scott, J.F., Alexe, M., Zakharov, N.D., Pignolet, A., Curran, C., and Hesse, D. (1998) Nano-phase SBTfamily ferroelectric memories, Integr. Ferroelec. 21, 1-14. Alexe, M., Harnagea, C., Hesse, D., and Goesele, U. (1999) Patterning and switching of nanosize ferroelectric memory cells, Appl. Phys. Lett. 75, 1793-1795; Schilling, J., Muller, F., Matthias, S., Wehrspohn, R.B., Gosele, U., and Busch, K. (2001) Three-dimensional photonic crystals based on macroporous silicon with modulated pore diameter, Appl. Phys. Lett. 78, 1180-1182. Andreev, A.F. (1981) JETP Lett. 53, 1063. Shchukin, V.A. and Bimberg, D. (1999) Spontaneous ordering of nanostructures on crystal surfaces, Rev. Mod. Phys. 71, 1125-1171. Williams, R.S., Medeiros-Ribeiro, G., Kamins, T.I., and Ohlberg D.A.A. (2000) Thermodynamics of the size and shape of nanocrystals: Epitaxial Ge on Si(001), Ann. Rev. Phys. Chem. 51, 527. Dawber, M. and Scott, J.F. (2001) Calculation of Schottky barrier height of platinum/lead zirconate titanate interface, Integr. Ferroelec., 38, 805-813. Wilk, G.D., Wallace, R.M., and Anthony, J.M. (2001) High-k gate dielectrics: Current status and materials properties considerations, J. Appl. Phys. 89, 5243-5275; see also Wallace, R.M. and Wilk, G.D. (2002) Alternative gate dielectrics for microelectronics, MRS Bull. 27, 186-187; Wallace, R.M. and Wilk, G.D. (2002) High-k Gate Dielectric Materials, MRS Bull. 27, 192-197; Lee, H., Jeon, S., and Hwang, H. (2001) Electrical characteristics of a Dy-doped HfO2 gate dielectric, Appl. Phys. Lett. 79, 2615-2617; Misra, V., Heuss, G.P., and Zhong, H. (2001) Use of metal-oxide-semiconductor capacitors to detect interactions of Hf and Zr gate electrodes with SiO2 and ZrO2, Appl. Phys. Lett. 78, 4166-4168. Morrison, F.D., Scott, J.F., Alexe, M., Leedham, T.J., Tatsuta, T., and Tsuji, O. (2002) Use of the 'mist' (liquid-source) deposition system to produce new high-dielectric devices: ferroelectric-filled photonic crystals and Hf-oxide and related buffer layers for ferroelectric-gate FETs, Microelectron. Eng. 66, 591599. Lin, Y.-S., Puthenkovilakam, R., and Chang, J.P. (2002) Dielectric property and thermal stability of HfO2 on silicon, Appl. Phys. Lett. 81, 2041-2043. Quintard, P.E., Barberis, P., Mirgorodsky, A.P., and Merle-Mejean, T. (2002) Comparative latticedynamical study of the Raman spectra of monoclinic and tetragonal phases of zirconia and hafnia, J. Am. Ceram. Soc. 85, 1745-1749. Wang, J. C., Chiao, S.H., Lee, C.L., Lei, T.F., Lin, Y.M., Wang, M.F., Chen, S.C., Yu, C.H., and Liang, M.S. (2002) A physical model for the hysteresis phenomenon of the ultrathin ZrO2 film, J. Appl. Phys. 92, 3936-3940; Chang, J. P. and Lin, Y.-S. (2001) Highly conformal ZrO2 deposition for dynamic random access memory application, J. Appl. Phys. 90, 2964-2969. Ewing, R.C. (2001) The design and evaluation off nuclear-waste forms: Clues from mineralogy, Canadian Mineralogist 39, 697-715. Ottow, S., Lehmann, V., and Foll, H. (1996) Development of three-dimensional microstructure processing using macroporous n-type silicon, Appl. Phys. A 63, 153-159; see also Smith, R.L. and Collins, S.D. (1992) Porous Silicon Formation Mechanisms, J. Appl. Phys. 71, R1-R22; as well as Birner, A., Wehrspohn, R.B., Gosele, U., and Busch, K. (2001) Silicon-based photonic crystals, Adv. Mater. 13, 377388. MacMillan, L.D., Paz De Araujo, C.A., Roberts, T., Cuchiaro, J., Scott, M.C., and Scott, J.F. (1992) Integr. Ferroelec. 2, 351. Gruverman, A. (1986) Ph.D. thesis, Univ. Urals, Sverdlovsk (Ekaterinburg), USSR; Shur, V.Ya., Gruverman, A., Kuminov, V.P., and Tonkachyova, N.A. (1990) Dynamics of Plane Domain-walls in Lead Germanate and Gadolinium Molybdate, Ferroelecrics. 111, 197-206. Shur, V.Ya., Baturin, I.S., Shishkin, E.I., and Belousova, M.V. (2003) New approach to analysis of the switching current data in ferroelectric thin films, Ferroelectrics 291, 27-35.

72 20. Scott, J.F. Dawber M, Jiang AQ, Morrison FD (2003) Ferroelectrics 286: 945-957; Scott, J.F. (2003) Domain wall kinetics: Nano-domain nucleation in lead germanate and Tilley-Zeks theory for PVDF Ferroelectrics 291, 205-215; Scott, J.F. (2003) New ferroelectric thin-film results: Electrode effects and photonic crystals, Ferroelectrics 293, 33-41. 21. Zolotoyabko, E., Quintana, J.P., Hoerman, B.H., and Wessels, B.W. (2002) Fast time-resolved x-ray diffraction in BaTiO3 films subjected to a strong high-frequency electric field, Appl. Phys. Lett. 80, 31593161. 22. Dawber, M., Jung, D.J., and Scott, J.F. (2002) Perimeter effect in very small ferroelectrics, Appl. Phys. Lett. 82, 436-438; see also Jung, D.J., Dawber, M., Ruediger, A., Scott, J.F., Kim, H.H., and Kim, K. (2002) Dielectric loss peak due to platinum electrode porosity in lead zirconate titanate thin-film capacitors, Appl. Phys. Lett. 81, 2436-2438. 23. Frisch, T., Pomeau, Y., and Rica, S. (1992) Transition to Dissipation in a Model of Superflow, Phys. Rev. Lett. 69, 1644-1647; see also Landau, L.D. and Lifshitz, E.M. (1987) Fluid Mechanics, Pergamon, Oxford. 24. Durkan, C., Welland, M.E., Chu, D.P., and Migliorato, P. (2000) Scaling of piezoelectric properties in nanometre to micrometre scale, Electron. Lett. 36, 1538-1539. 25. Astafiev, K., Sherman, V., Tagantsev, A., Setter, N., Rivkin, T., and Ginley, D. (2002) Investigation of electrical degradation effects in ferroelectric thin film based tunable microwave components, Integ. Ferroelec. 49, 103-112. 26. Bune, A.V., Fridkin, V.M., Ducharme, S., Blinov, L.M., Palto, S.P., Sorokin, A.V., Yudin, S.G., Zlatkin, A. (1998) Two-dimensional ferroelectric films, Nature 391, 874-877. 27. Batra, I.P. and Silverman, B.D. (1972) Thermodinamic Stability of Thin Ferroelectric Films, Solid State Comm. 11, 291. 28. Kay, H.F. and Dunn, J.W. (1962) Thickness Dependence of Nucleation Field of Triglycine Sulphate, Phil. Mag. 7, 2027. 29. Chandra, P., Dawber, M., Littlewood, P., and Scott, J.F. Nature Mater., in press. 30. Ku, H.Y. and Ullman, F.G. (1964) Capacitance of Thin Dielectric Structures, J. Appl. Phys. 35, 265; see also Mead, C.A. (1961) Anomalous Capacitence of Thin Dielectric Structures, Phys. Rev. Lett. 6, 545546. 31. Simmons, J.G. (1965) An Analytic from of Ku and Ullmans Equations (Electric Field Penetration of Tunnel Junction Electrodes - T), Appl. Phys. Lett. 6, 54; Simmons, J.G. (1967) Incorporation of ElectricField Penetration of Electrodes in Theory of Electron Tunneling Through a Dielectric Layer, Brit. J. Appl. Phys. 18, 269. 32. Black, C.T. and Welser, J.J. (1999) Electric-field penetration into metals: Consequences for highdielectric-constant capacitors IEEE T. Electron. Dev. 46, 776; see also Hwang, C.S. (2002) Thicknessdependent dielectric constants of (Ba,Sr)TiO3 thin films with Pt or conducting oxide electrodes, J. Appl. Phys. 92, 432-437. 33. Dawber, M., Sinnamon, L.J., Scott, J.F., and Gregg, J.M. (2002) Electrode field penetration: A new interpretation of tunneling currents in barium strontium titanate (BST) thin films, Ferroelectrics 268, 455460. 34. Kingon, A.I. (2002) Thickness, strain, and temperature-dependent properties of barium strontium titanate thin films, Proceedings of the 13th IEEE ISAF 2002, 151-154; see also part I of this book by A.I. Kingon. 35. Tilley, D.R. and Zeks, B. (1984) Landau Theory of Phase-Transitions in Thick-Films, Solid State Comm. 49, 823-827. 36. Zhong, W.L., Wang, Y.G., and Zhang, P.L. (1998) Ferroelec. Rev. 1, 131. 37. Mills, D.L. (1971) Surface Effects in Magnetic Crystals near the Ordering Temperature, Phys. Rev. B 3, 3887-3895. 38. Lubensky, T.C. and Rubin, M.H. (1975) Critical phenomena in semi-infinite systems. II. Mean-field theory, Phys. Rev. B 12, 3885-3901. 39. Scott, J.F., Duiker, H.M., Beale, P.D., Poulighy, B., Dimmler, K., Parris, M., Butler, D., and Eaton, S. (1988) Properties of Ceramic KNO3 Thin-Film Memories, Physica B 150, 160-167; Duiker, H.M., Beale, P.D., Scott, J.F., de Araujo, C.A.P., Melnick, B.M., Cuchiaro, J.D., McMillan, L.D. (1990) Fatigue and Switching in Ferroelectric Memories - Theory and Experiment, J. Appl. Phys. 68, 5783; Zhang, J., Yin, Z., Zhang, M.-S., and Scott, J.F. (2001) Size-driven phase transition in stress-induced ferroelectric thin films, Solid State Comm. 118, 241-246. 40. Li, C., Chen, Z., Cui, D., Zhou, Y., Lu, H., Dong, C., Wu, F., and Chen H. (1999) Phase transition behavior of BaTiO3 thin films using high-temperature x-ray diffraction, J. Appl. Phys. 86, 4555-4558.

73 41. O'Neill, D., Bowman, R.M., and Gregg, J.M. (2000) Dielectric enhancement and Maxwell-Wagner effects in ferroelectric superlattice structures, Appl. Phys. Lett. 77, 1520-1522. 42. Corbett, M.H., Bowman, R.M., Gregg, J.M., and Foord, D.T. (2001) Enhancement of dielectric constant and associated coupling of polarization behavior in thin film relaxor superlattices, Appl. Phys. Lett. 79, 815-817. 43. Kim, J., Kim, Y., Kim, Y.S., Lee, J., Kim, L., and Jung, D. (2002) Large nonlinear dielectric properties of artificial BaTiO3/SrTiO3 superlattices, Appl. Phys. Lett. 80, 3581-3583. 44. Jiang, A.Q., Scott, J.F., et al. (2002) J. Appl. Phys., in press. 45. Rios, S., Ruedger, A., Scott, J.F., et al. (2002) Appl. Phys. Lett., submitted.

Part II – Fundamentals of Scanning Probe Techniques

PRINCIPLES OF BASIC AND ADVANCED SCANNING PROBE MICROSCOPY

D.A. BONNELL1, R. SHAO Department of Materials Science and Engineering, University of Pennsylvania, 3231 Walnut l ST Philadelphia, PA 19104 USA

Contents 1. 2.

Introduction Basic Concepts of Scanning Probe Microscopy 2.1. Electrostatic Force Microscopy 2.2. Magnetic Force Microscopy 3. Advanced Scanning Probe Microscopy: Exploiting Multiple Modulations 3.1. Scanning Spreading Resistance Microscopy and Scanning Capacitance Microscopy 3.2. Scanning Surface Potential Microscopy 3.3. Scanning Impedance Microscopy 3.4. Scanning Gate Microscopy 3.5. Piezoresponse Force Microscopy 3.6. Scanning Microwave and Dielectric Microscopy 4. Applications 4.1. Transport in Single Molecules and Nano Wire/Tubes 4.2. Domain Interactions in Ferroelectric Thin Films 5. Other Techniques and Future Directions

Abstract Understanding the behavior of complex materials such as organic self-assembled monolayers, molecular and nano wires, transition metal oxide thin films, is facilitated by probes of local properties. Recent extensions of scanning probe microscopy that extract electrical potential, capacitance, dielectric constant, electromechanical coupling coefficients and impedance, are described. In most cases, these complex properties are accessed by stimulations and/or response function detection with multiple frequency modulations. Several illustrative example include determination of the electronic

1

Corresponding author. Phone: (215)898-6231, Fax: (215)573-2128. Email address: [email protected] 77 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 77-101. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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structure of individual defects in a carbon nanotube, ferroelectric domain interactions in oxide thin films, and electric potential of an alkanethiol on metal. Keywords: Scanning probe microscopy, multiple modulation, spatial resolution, complex materials, molecular wires, ferroelectric domains.

1. Introduction The quest toward understanding the behavior of condensed matter has relied on measuring structure, bonding, and properties at increasingly local levels. This has driven advances in techniques that probe both soft and hard materials directly as well as indirectly. Examples include the suite of imaging and spectroscopy based on electron optics, spatially confined synchrotron radiation, optical spectroscopies such as ‘micro’ Raman and FTIR, etc. While structure and bonding based probes have accessed molecular and atomic scales for decades, local determination of properties was elusive. The emergence of scanning probes filled this gap to some extent. Technique AFM EFM MFM SSPM(KPM) SCM SCFM SSRM SGM SIM

NIM PFM NPFM SNDM NFMM

TABLE I. Properties and Modulated Operation Modes of Scanning Probes Mode Property References nc/ic, mech, [2-5] VdW interaction, topography phase/amp electrostatic force nc, mech, [2-5] phase/amp magnetic force, current flow nc, mech, [2-5] phase/amp nc, elec, 1st potential, work function, adsorbate [2-5, 27-44] harmonic enthalpy/entropy c, F, cap sensor capacitance, relative dopant density [2-12, 15-17, 1924] c, elec, 3rd dC/dV, dopant profile [26] harmonic c, F, dc current resistivity, relative dopant density [2-5, 12-14, 18] nc, elec, amp current flow, local band energy, contact [2-5,53-59] potential variation nc, elec, interface potential, capacitance time constant, [51, 53] phase/amp local band energy, potential, current flow (in comb. w/ SSPM) c,F, freq spectrum interface potential, capacitance, time constant, [55] dopant profiling c, elec, phase/amp d33 [2-5, 60-65] c, elec, 2nd switching dynamics [66] harmonic c, F, 1stt or 3rd dC/dV, dielectric constant [67, 68] harmonic c, F, phase microwave losses, d33 [69-71]

Mode: c= contact, nc = non contact, ic = intermittent contact, mech = cantilever is driven by a mechanical oscillation, elec = cantilever is driven or responds to an oscillating electrical signal, F = constant force feedback, phase = detection or feedback on phase, amp = detection of amplitude at preset frequency.

The utility of local probes is illustrated by the fact that only 10 years after commercial SPMs became available upwards of 1750 papers per year [1] are published

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that cite scanning probes as a key word. Several monographs have summarized the state of this field and various times and books are available that provide introduction to the field and general overviews [ 2 , 3 , 4 , 5 ]. Most applications utilize SPM as a straightforward qualitative mapping tool. Some researchers interested in complex behavior of solids have examined fundamental probe-surface interactions and extended SPM to probe local electronic transport, dielectric, ferroelectric, magnetic properties. It is convenient to categorize scanning probe techniques in terms of probe-sample contact, source of cantilever oscillation, and feedback function. Various scanning probe techniques are described in these terms in Table I, which also lists the properties that can be accessed with each. For example, conventional atomic force microscopy (AFM) can be a non-contact, mechanically driven oscillation with amplitude or phase based detection. Piezoelectric force microscopy (PFM) is a contact, electrically driven oscillation, with phase based detection. Combined with detection and/or feedback of not only 1stt harmonic response functions, but also 2ndd and 3rdd harmonic response functions, a wide range of properties can be accessed. An underlying theme of the newest developments is the use of multiple signal modulations or high order harmonics of modulated signals. This paper first summarizes the basic concepts that are needed in order to understand the more sophisticated uses of SPM presented in this volume. Several advanced techniques are described and a few examples are used to illustrate the applications. 2. Basic Concepts of Scanning Probe Microscopy SPM is based on the interaction of a probe tip and a surface, which for completeness are summarized here. A comprehensive treatment of surface and intermolecular forces can be found in an excellent book by J. Israelachevili [6]. The fundamental interaction at short distances is derived from van der Waals forces. At distances of a few nms, van der Waals forces are sufficiently strong to move macroscopic objects such as AFM cantilevers. Van der Waals interactions consist of three components: polarization, induction, and dispersion. Polarization refers to permanent dipole moments such as exist in water molecules or in BaTiO3. Induction refers to the contribution of induced dipoles. Dispersion is due to instantaneous fluctuations of electrons, which occur at the frequency of light causing optical dispersion, thus the name. A reasonable approximation of the interaction is obtained assuming that dispersion dominates and is isotropic, additive, and non-retarded. Under these assumptions the van der Waals potential between two planes is -A/12ρZ2, and between a sphere and a plane is -AR/6Z; in which Z is the distance between objects, b R is the radius of the sphere, and A is the Hamaker constant. The complete interaction must include the repulsive as well as attractive terms. The Hamaker constant is the term that characterizes the properties of the materials, including collective interactions and polarization if calculated from dielectric and optical properties. It is ª 3kT § ∈1 − ∈3 ·§ ∈2 − ∈3 · 3ηω « ¨ ¸ ¨ ¸ + A= 4 ¨© ∈1 + ∈3 ¸¹¨© ∈2 + ∈3 ¸¹ 8 2 «« ¬

(

)( 1

2

(

)(( ª ) «¬( 1

2

)

) +( 1

2

º » » 1 2º »¼ »¼

)

(1)

80

where k is the Boltzmann constant, T is temperature, ρ is the density, εi is the dielectric constant in medium i (i.e. sample (1), tip (2), or intervening material (3)), and ni is refractive index. The Hamaker constant is on the order of 10-19 J for most solids. In many situations long range forces act in n addition to short range forces between two surfaces. Examples of long range interactions include electrostatic attraction or repulsion, current induced or static magnetic interactions, and forces due to the surface energy of water condensed between the sample and tip. Very close to the surface these forces are much smaller than those due to van der Waals interactions and usually contribute little to the signal. Farther from the surface, the van der Waals interactions decay rapidly to the point of being negligible. In this regime long range forces are still significant. This difference in decay length provides a means to distinguish the two types of interactions. The general relations describing the force experienced by a tip above a homogeneous surface for electrostatic and magnetic interactions are described in Eqs. 3 and 4, where ∆V is the difference in potential between sample and tip, C is the tipsample capacitance as a function of separation (z), Bsample is the magnetic field emanating from the sample surface, and m is the magnetic dipole of the tip [7, 8]. Conducting and/or magnetic tips are obviously necessary to access electrical or magnetic fields: 1 (2) )2 ∂ C Felectrosta tic = − ( 2 ∂z (3) Fmagnetosta tic = ∇ m • B sample

(

)

These relations are oversimplifications but suffice to describe the operating principles of electrostatic force imaging (EFM) and magnetic force imaging (MFM). The underlying principle of AFM is that the interactions between the end of a probe tip that is mounted on a cantilever result in a response in the cantilever, notably a deflection. The mechanical resonant frequency off the cantilever is determined by the dimensions of the structure and the properties of materials from which it is made. This resonant is related to the cantilever spring constant according to Eq. 4, where the cantilever is conceptually treated as a classical, 1-dimensional, lightly loaded, “fixedfree” beam: (4) t E k ω

o

=

λ2

ρ

=

m

eff

The vibration amplitude, A, detected at a given frequency, ω, changes as a function of the force gradient as shown in Eq. 5. §ω A ( ω = ω o′ ) = a o Q ¨¨ o © ω o′

∂F § · ∂z ¸¸ ≈ a o Q ¨ 1 − ¨ 2k ¹ ©

· ¸ ¸ ¹

−1

(5)

where Q is the quality factor and k is the spring constant of the cantilever. Note that by measuring the change in amplitude both the magnitude and sign of the force gradient are determined. As shown in Fig. 1, SPM techniques detecting electrostatic or magnetic interactions are usually implemented in lift mode: typically, the topography is first obtained in contact mode after which the tip is separated a preset distance from the surface (chosen such that short range forces do not contribute to the electrostatic or magnetic forces of

81

interest). At this preset sample-tip separation long range interactions are measured, usually with an ac technique.

Surface Figure 1. Schematic of lift mode SPMs, e.g. EFM, SSPM, MFM, SIM, etc.

2.1. ELECTROSTATIC FORCE MICROSCOPY Electrostatic Force Microscopy (EFM) is commercially available and in terms of Table I is a noncontact tool with detection by y mechanical oscillation. While scanning at a constant tip/surface separation (usually 50~100nm), the cantilever is mechanically oscillated at its resonant frequency and constant amplitude. Under the attractive electrostatic interaction, the motion of the cantileverr can be approximated by a simple harmonic oscillator: (6) meff z + γz + kz = F (z ) where meff is effective mass of the tip, γ is damping coefficient. And electrostatic force F(z) can be further expanded: dF F ( z ) = F ( z 0 + dz ) = F ( z 0 ) + dz (7) dz z = z0

where z0 is the equilibrium position (approx. lift height) about which the cantilever is oscillating. The resonant frequency is shifted from that of a free cantilever (Eq. 4):

∆ω =

ω 0 dF 2k dz

(8)

The frequency shift results in phase and amplitude shifts relatively to a freely oscillating cantilever:

82

∆φ =

Q dF Q dF and ∆A A = CA0 k dz k dz

(9)

Thus imaging ∆φ or ∆A is equivalent to imaging the force gradient variation on the surface. In order to enhance the quality of imaging, a dc bias is usually applied to the cantilever. 2.2. MAGNETIC FORCE MICROSCOPY A ferromagnetic AFM tip can be used to sense the local magnetic field distribution. High resolution Magnetic Force Microscopy (MFM) has been developed to image ferromagnetic domain structures, study the quality of magnetic record media, superconducting current, as well as the integrity of microelectronic circuits. Similar to EFM, MFM is based on detecting dynamic response of a magnetized AFM cantilever to the magnetic force gradient. MFM is also realized in dual pass lift mode. In the first pass, the surface topography is acquired. In the second pass, the cantilever is mechanically driven, and during the scanning a constant separation with the surface is maintained. The magnetic force gradient results in a change in the cantilever resonant frequency, ∆ω =

Fm ω 0 dF 2k dz

, where Ȧ0 is the resonant frequency without external forces

and k the spring constant. The magnetic force Fm in the point probe model is given by Eq. 3. For more than a decade, MFM has been the primary tool of imaging magnetic domains. However, it is discovered that oftentimes with the strong magnetic field produced by the scanning probe can alter the domains in the sample. The result is a big difference between two consecutive images. Another major application of MFM is the imaging of current carrying microcircuits. Intensive study in the area has been done by R. Yongsunthon, E. D. Williams et al. [9, 10] In one of their works, current crowding effect was observed near the defects in a microfabricated connection, leading to further study of electromigration and circuit failure mechanism. One should keep in mind, however, when scanning across a biased device, electrostatic force also contributes to the cantilever oscillation (as in SSPM). To achieve quantitative magnetic interaction, MFM is further developed to include an electrostatic force correction mechanism, which is realized by combining conventional MFM with SSPM [11]. By adjusting a dc bias applied to the cantilever, the contribution from electrostatic force to the tip oscillation can be nullified.

3. Advanced Scanning Probe Microscopy: Exploiting Multiple Modulations

3.1. SCANNING SPREADING RESISTANCE MICROSCOPY AND SCANNING CAPACITANCE MICROSCOPY One of the first measurements to emerge from electrostatic force microscopy was motivated by the need to map dopant profiles in semiconductor devices with decreasing dimension. Two contact techniques, Scanning Spreading Resistance Microscopy

83

(SSRM) [12, 13, 14] and Scanning Capacitance Microscopy (SCM) [12, 15, 16, 17], were developed. Actually, both SSRM and SCM map carrier concentration from which dopant concentration is calculated. In SSRM a conducting tip is biased with respect to the sample and the current through the tip/surface contact is detected under force feedback control. The amount of current is determined by the local spreading resistance of the surface, which is related to the local conductivity ı and the contact radius a by 1 . As ı is a function of the carrier concentrations p, n and carrier mobilities µ , p R= 4 aσ µn: σ = q ( pµ p + nµ n ) , the concentration of the major carriers can be estimated. Because the native oxide layer on Si hinders tip/surface contact, carrier profiling usually requires a tip coated with a hard material (e.g. doped diamond) and a high spring constant cantilever to provide strong indentation forces (~20µN) [12, 18]. In SCM, by contrast, the carrier concentration is profiled via the detection of tipsurface capacitive coupling, and thus a good quality oxide surface layer is required. As the tip is scanned across the sample, a high frequency capacitance sensor detects the tipsample capacitance. An ac voltage applied to the tip induces the depletion and accumulation of carriers, resulting in a change in measured capacitance, ǻC. In a semiconductor, the depletion/accumulation width is inversely related to the carrier concentration, so mapping of ǻC/ǻV yields the carrier concentration profile. Difficulties in quantification arise if the dopant concentration is non-uniform, and when spatial resolution decreases due to low dopant concentrations. A recent modification incorporates an additional feedback that adjusts ǻV to maintain a constant ǻC during the scan, maintaining a constant depletion width. Analysis of SCM results are mathematically challenging, usually requiring 2-D or 3-D numerical approaches [19, 20, 21, 22]. Williams et al. were the first to map the capacitance distribution in a FET device, showing that the source, drain and gate are evident and the difference in majority carrier (n or p type) is indicated (Fig. 2). The device is 180 nm and capacitance gradients over tens of nm are clearly visible. H. Edwards et al. have imaged the capacitance variations at a pn junction in a MOS-FET with spatial resolution of ±30nm, on the same order as that of the Si depletion width [23]. The potential impact of the ability to determine local capacitance is obvious but not limited to conventional semiconductors. SCM, in conjunction with SSPM has also been used to characterize grain boundary properties in ceramics such as a CdTe film [24]. Both SSRM and SCM suffer from limitations in spatial resolution due to the need for high load tip contact and from limitations in capacitance inherent in the capacitance sensor. In addition there is an intrinsic trade off in that as contact area is reduced to increase spatial resolution, the signal to noise ratio deteriorates. Two approaches are currently being explored to solve this problem. One solution is to improve the sensitivity of the sensor [25]. An alternative proposed by Kobayashi et al. [26] is to eliminate the sensor entirely and apply an oscillating electric field to a conducting tip in contact. The third harmonic of this signal, if it can be detected, gives dC/dV directly. This technique is related to high order PFM and high order Scanning Nonlinear Dielectric Microscopy, discussed below.

84

Figure 2. Scanning Capacitance Microscopy of a 0.18 nm FET. Contrast in capacitance is related to local dopant concentration. Source, drain and gate are indicated, along with the channel. This data demonstrates that spatial resolution on the order of tens of nms is necessary to characterize variations in this device. (courtesy of C. C. Williams)

3.2. SCANNING SURFACE POTENTIAL MICROSCOPY Scanning Surface Potential Microscopy (SSPM), sometimes referred to as Kelvin Probe Microscopy (KFM), was developed to map the work function variation of a surface [27, 28]. In SSPM, the tip/surface separation is maintained constant (typically tens to 100 nanometers) with the actuator disengaged, so that the tip is no longer mechanically driven. In the meantime, a biasing voltage Vtip = Vdc + Vac cos(ωt ) is applied to the tip. Let the surface potential of the sample be Vs, the capacitive is given by

Fcap ( z ) =

1 ∂C (Vtip − Vs ) 2 2 ∂z

(10)

where C is the tip/surface. A lock-in amplifier picks the first harmonic of the tip oscillation signal under the capacitive force Fcap. By expanding equation (1), the first harmonic component (F F1Ȧ) of Fcap is obtained:

F1ω =

∂C (Vdc − Vs )Vac ∂z

(11)

By properly adjusting the dc bias component Vdc using a feedback loop, F1Ȧ can be nullified, i.e. F1Ȧ=0, which is satisfied when Vdc=V Vs. Because of its capability of resolving potential contrast down to several millivolts, this technique has found wide applications in studying electrically active interfaces, surface desorption phenomena, as well as self-assembled monolayers. The apparent ease with which potential variations

85

can be mapped belies the challenges in interpreting images on electrically [29, 30] and topographically [ 31 ] inhomogeneous surfaces. In the case of semiconductor and dielectric surfaces the electrostatic properties of a surface are not characterized solely by intrinsic potential and topography. SSPM images of these surfaces should be interpreted in terms of effective surface potential that includes capacitive interactions per se, along with contributions from surface and volume bound charges, double layers and remnant polarization [32, 33, 34, 35]. For semiconductor surfaces without Fermi level pinning, tip-bias induced band bending [ 36 ] can lead to a non-linear surface potential dependence on voltage [37], further complicating quantification of experimental results. Despite these challenges the obvious need to examine variations in local potential in electronic nano devices spurred efforts to overcome some of the obstacles with careful analytical treatments that determined limits in quantification, and allowed complex materials to be addressed. Progress was slow immediately after the introduction of the technique, but the late nineties saw SSPM applied to semiconductor [38, 39], organic [40] and ferroelectric [41, 42] surfaces, as well as to defects [43, 44], and photoinduced [45, 46] and thermal phenomena [47, 48] in these materials. An illustrative example is the attempt to determine what the surface potential of an organic self assembled monolayer (SAM) represents. Figure 3a shows how a pattern of two SAMs with dramatically different properties yields a clear difference in the surface potentials. The difference in potential for this particular comparison of dodecane thiol vs PETB is 80mV. The question arises as to what this value represents. At the limit of very short chain lengths the potential might be treated as a substrate work function modified by adsorbates. At the limit of very long chain lengths, it might be the work function presented by the polymer. In the intermediate range it must be a convolution of the substrate work function, the dipole in the molecule-metal bond, and the dipole of the molecule to which the end group (charged or uncharged) contributes. Eng at al began to address this problem with a comparison of potential as a function of chain length [49].

(a) Figure 3. (a) SSPM of a SAM pattern on Au (111) produced with two types of organic molecules. (courtesy of R. Alvarez) (b) A quantitative relationship between the alkane thiol chain length and surface potential (Courtesy of L. Eng reprinted with permission of the American Chemical Society)

Figure 3b shows a direct correlation between chain length and potential which indicates that the molecular dipole plays a large role. While most investigators assume that the metal/molecule bond is sufficiently screened so as not to contribute, Alvarez et al determined that substrates do influence the result [ 50 ]. And experimental

86

measurement and first principles calculations have not yet converged. It is fair to say that at this point that absolute values of potential can not be quantified, but variations in potential can be determined with energy resolution of 2-4 meV. It goes without saying that understanding potential variations in SAMs will be invaluable as complex molecules and patterning is explored. 3.3. SCANNING IMPEDANCE MICROSCOPY In order to explore mechanisms of transport in complex devices it is necessary to determine time and temperature dependence as well as voltage dependent response functions. In macroscopic systems this is done with frequency dependent perturbations or ultra fast probes of relaxation from excited states. The first introduction of frequency dependence in scanning probes is referred to as Scanning Impedance Microscopy (SIM) [51]. This is a noncontact, first harmonic detection in which the oscillating electrical signal is applied to the sample instead of the tip. In SIM, a lateral ac signal is applied to the sample and the tip is either grounded or dc biased. Same as in SSPM, the tip/sample separation is maintained constant with the mechanical actuator disabled. According to Eq. (1), the capacitive interaction is expressed as:

F=

1 ∂C 1 ∂C (Vdc − Vs ) 2 = [Vdc − Vs0 − Vac cos(ωt + ϕ L )]2 2 ∂z 2 ∂z

(12)

where Vs, and Vs0 are local surface potentials after and before the lateral bias is applied, Vac is the local local oscillation of the surface potential, and ijL is the phase lag relative to the power source. The amplitude and the phase lag ijc of the first harmonic cantilever response are

A1ω =

| F1ω | 1 m (ω − ω 0 ) 2 + γ 2ω 2

and

ϕ c = arctan(

γω ) ω − ω0

(13)

(14)

where | F1ω |= ∂C (Vddc − Vs0 )Vaac , is the amplitude of the first harmonic capacitive force, ∂zz m is the effective mass, Ȗ the damping coefficient. From (3), ijc is independent of the location of the tip. The total phase lag of the cantilever oscillation relative to the driving source is ϕ = ϕ L + ϕ c . During the lift scan, A1Ȧ and ij are acquired simultaneously with the lock-in technique. Assuming SIM is used to image an electrical component (e.g. an 1 active interface) with an impedance of Z = =| Z | exp( j∆ϕ ) , as illustrated in 1 + jω C R Fig.4, ∆ϕ = ϕ 2 − ϕ1 , and

| Z | ∝ | A2 exp( jϕ 2 ) − A1 exp( jϕ 1 ) | ∝ | V2 − V1 | (15) The greatly enhance spatial resolution has enabled SIM as a quick tool to check the connection reliability of a ‘nano-circuit’ built from carbon nanotubes and/or nanowires.

87

In combination with other SPM techniques, SIM has revealed dielectric constant suppression at an oxide grain boundaries [52], mapped the rectifying behavior of a Schottky junction, and characterized defect mediated transport across a nanotube [53]. The application of this technique to interfaces f is treated in detail by Bonnell and Kalinin elsewhere in this volume [54].

ϕ2, A2

ϕ1, A1 Ri

R

R

Ci

Interface

V (a)

(b)

Figure 4. Principle of SIM: an electroactive interface can be modeled as a simple RC element. The phases (ij1, ( 1, A2) of tip oscillation on both sides of the interface are related by Eqs. (13), (14) and ij2) and amplitudes (A (15).

ș

|Z| 1

2

3

3

3 µm

(a)

2

1

(b)

2.0x10

5

1.5x10

5

1.0x10

5

5.0x10

4

2V 3V 5V

Zi(Ω)

(c)

0.0 0

1x10

5

2x10 Z r( Ω )

5

3x10

5

4x10

5

Figure 5. Contrast Mechanism Maps of PFM as a function of contact radius and indentation force. SI - strong indentation regime, CSI - contact limited strong indentation, LE - linear electrostatic regime, NE - nonlinear electrostatic regime, PD - plastic deformation. Dotted line delineates the region where stress-induced switching is possible. The maps are constructed for (a) good tip-surface contact (w = 0.1 nm) and (b) bad contact (w = 1 nm).

88

The second approach to accessing frequency dependent transport expands the frequency range to many orders of magnitude. Nano Impedance Spectroscopy (NIS) [55] is a contact probe with force feedback in which the oscillating signal is applied to the tip; phase and amplitude are detected at the sample. The frequency dependent measurement can be implemented in two terminal or three terminal configurations depending on the device geometry Figure 5 (c) illustrates the local impedance between a tip and lateral electrode, across a ZnO grain boundary. The electrode contact and grain boundary properties are distinguished as two semi circles in the impedance spectra, allowing the actual properties of the grain boundary to be characterized. The local boundary potential, capacitance, charge and depletion lengths can be extracted with spatial resolution on the order of tens of nms. In a configuration with an electrode under the sample, the impedance of individual grains can be mapped, Figure 5 (a) and (b). 3.4. SCANNING GATE MICROSCOPY A clever modification of EFM has been developed recently in order to characterize lateral transport in low dimensional systems. Electrodes are connected, for example by e-beam lithography, to the feature of interest, which might be a quantum well, nanotube, or molecular wire. The tip, which is scanned over the structure while current is flowing between the electrodes, perturbs the system acting as a local gate. The signal that is mapped as an image is the relative differences in current through this three terminal device as a function of the tip position. Some samples are amenable to configurations that also contain a back gate, in which case the tip gate and back gate can both be varied. This technique, referred to as Scanning Gate Microscopy, is a non contact, electrically driven probe system [53, 56, 57, 58, 59]. Due presumably to the geometric constraints of this configuration SGM has been applied most frequently to carbon nanotubes. An example is shown in section 3.1. The next generation of local transport measurement will utilize individually addressable multiple tip scanning probes to explore alternative configurations of 2, 3, and 4 point transport measurements at nm length scales. 3.5. PIEZORESPONSE FORCE MICROSCOPY As the ability to synthesize complex functional materials increases in sophistication, characterization of nonlinear properties such as piezoelectric, ferroelectric, and ferromagnetic responses at small scales is becoming important. Piezoelectric Force Microscopy (PFM) is a scanning probe that holds the promise of determining electromechanical coupling coefficients at local scales. PFM is a contact, electrically driven V0cos(Ȧ ( t) is probe technique with feedback based on phase lag. An ac voltage Vac=V applied upon the tip and ferroelectric surface junction. If the material is piezoelectric, the field results in a local deformation of the surface that oscillates the tip, i.e. piezoresponse [60, 61]: PR = A0 cos(ωt + δ ) , where A0 is the amplitude of PR, which is proportional to the electromechanical coefficients, and į is the phase difference between the driving signal and the piezoresponse, is determined by the orientation of the domain. For downward oriented domains, į = -180o and for upward oriented domains, į = 0o.

89

The phase therefore indicates the orientation of atomic polarization, while the amplitude gives the magnitude of the coupling coefficient. In reality, electrostatic interactions between the tip and the surface, as well as between the cantilever and the surface are not negligible, as shown by a variable temperature measurement done by Luo et al. A rigorous treatment of the imaging contrast includes simultaneous electrostatic and electromechanical interactions. The total contribution to the measured amplitude A should be (16) A = Ael + Apiezo + Anloc where Aell is from the tip-surface electrostatic force, Apiezo is from the electromechanical interaction and Anloc is the nonlocal canilever effect. Obviously, reliable piezoresponse results are obtained only if electrostatic interactions are minimized. Taking an approach familiar to materials scientists, the analytical solutions of these interactions can be presented as contrast mechanism maps (Fig. 6) that relate experimental conditions to the properties of the material and delineate the conditions under which quantitative measurements can be obtained [62]. This is critical since it has been shown that under some experimental conditions the response has no connection to local properties.

Figure 6. NIS imaging of a junction of ZnO grains (a) |Z| (b) θ; impedance spectrum shows two relaxation processes (c).

The dynamic switching of a domain is studied via the dc voltage dependence of PR, which is usually accomplished by simultaneously applying a slow periodical dc voltage to the bottom electrode and acquiring the PR signal with the probe staying at one point. By defining PR mathematically as (17) PR = A cos δ an electromechanical hysteresis loop can be constructed by plotting the PR versus the dc voltage, due to the back and forth switching of the domain by the dc voltage (Fig. 7). For the general case there is also an in-plane component to polarization that can be accessed by measuring lateral response of the tip to field variation [ 63 , 64 ]. Furthermore, the piezoelectric response is a tensorial function, the complexity of which depends on the symmetry of the compound and the orientation of the grain or crystal. Harnagea et al have shown that even for BaTiO3 with relatively high symmetry it is not

90

possible to deduce domain orientation from out of plane PFM alone [65]. Either the grain orientation or the in-plane component must also be known. In spite of these challenges, PFM has become the preferred method of characterizing ferroelectric thin films and exploring the physics of polarization dynamics. The relationship between higher order harmonics, 2ndd and 3rd, of the PR function and time dependence of domain switching is being developed into a probe of switching dynamics in Nonlinear Piezo Force Microscopy (NPFM) [66].

PR(a.u.)

0.3

0.0

-0.3 -12

-8

-4

0 Bias(Volt)

4

8

12

Figure 7. Electromechanical hysteresis loop acquired with PFM on PZT 20/80 thin film.

3.6. SCANNING MICROWAVE AND DIELECTRIC MICROSCOPY A complementary strategy to accessing linear and nonlinear dielectric properties is referred to as Scanning Nonlinear Dielectric Microscopy [ 67 , 68 ] and Near Field Microwave Microscopy (NFMM) [69, 70]. The approach utilizes a coaxial probe in which a sharp, center conductor ‘tip’ protrudes. The probe is actually the end of a transmission line resonator, which is coupled to a microwave source. In the configuration used in Anlage’s group the probe tip is held fixed, while the sample is supported by a spring-loaded cantilever applying a controlled normal force on the order of 50 µN between the probe tip and the sample. The concentration of the microwave fields at the tip changes the boundary condition of the resonator, and hence, the resonant frequency and quality factor. The magnitude of the perturbation depends on the dielectric properties of the sample. A simple model describes the probe and sample interaction as a circuit of the sample sheet resistance RX connected in series with the probe/sample capacitance CX. And collected frequency shift ∆f and quality factor Q are related to RX and CX according to standard microwave theory. The spatial resolution of the microscope in this mode of operation is about 1 µm. NFMM has shown promise in mapping dielectric constant variations in a number of complex oxides; i.e. ferroelectric and superconducting compounds. Anlage et al. have demonstrated that high order

91

harmonic powers acquired by NFMM can be used to spatially resolve the local nonlinearity. In their work, the grain boundary area of superconducting YBCO thin film deposited on a SrTiO3 bicrystal was spatially resolved from the ratio of the second and the third power [71]. In the early version of NFMM, the probe is scanned in contact with the sample. The damage to the probe due to the large contact force severely impairs the spatial resolution. Recently, NFMM has been developed to include an STM feedback enabling a noncontact scanning with a separation of about 1nm. The improvement in spatial resolution was significant [72]. The sensitivity of this new NFMM technique to physical properties such as dielectric loss has been demonstrated on a boron doped silicon (Fig. 8). The difference in sheet resistance RX in doped and undoped regions were resolved by ∆f and Q images.

1 µm

350

413

475

538

600

Topography in Angstroms

364

368

371

375

378

Resonator Quality Factor

100

75

50

25

0

Frequency shift (kHz)

Figure 8. Images of a boron doped silicon sample (a) STM topography; (b) Quality factor; and (c) frequency shift images. (Courtesy of S. M. Anlage and reprinted with permission of Elsvier B. V.)

4. Applications

The expanding toolbox of local probes of complex properties can provide insight regarding local phenomena in two distinct contexts. It can be used to quantify properties of structures that are so small as to exhibit quantum mechanical or continuum based size dependent behavior. It can be used to probe local variations in larger systems in which these variations influence global behavior. An example of each follows. 4.1. TRANSPORT IN SINGLE MOLECULES AND NANO WIRE/TUBES Transport in reduced dimensions is becoming increasingly important as developments in nanostructured materials enable device elements to be made of organic and biological molecules, nanotubes, nanowires, quantum dots, etc. The ultimate goal of basing device function on molecules raises the challenge of isolating and characterizing the behavior single molecules or wires. Of course it is possible, and in fact has often been accomplished, to distribute nanostructures across a suitable substrate and examine them ‘top down’ with any of the scanning probes. In some cases this provides new insight, as for example in the STM analyses of carbon nanotubes which related local density of

92

states to atomic structure [73, 74]. It must be noted that in this configuration the tipnanostructure-substrate properties, rather than the lateral properties of the structure, are measured. The simplest conceptual approach to accessing lateral transport properties of a molecule or nanowire is to connect electrodes to molecules/wires that are distributed on a surface and measure from remote connections. One of the first wires to be characterized in this manner was a single phthalene molecule [75]. Transport in DNA has been examined, as well [76, 77]. The most extensive use of this strategy has been on carbon nanotubes. Both the Dekker group and the McEuen group have used SGM or simple transport measurements in this configuration. These results provide evidence that some carbon nanotubes act as real quantum wires. The low dimensionality is expected to result in strong electron–electron, so it has been suggested that electrons in nanotubes form a correlated ground state known as a Luttinger liquid. Some low temperature measurements provide evidence for this, while some points to coherent backscattering processes. Transport in semiconductor nanotubes is even less well understood. Electrode contact potential, chemical adsorption defects [78], and nanotubes/substrate interface contacts impact properties [79]. Johnson et al. [79] and Furher et al. [80] have used local probes to demonstrate a potential for defects at nanotubes to be memory bits. (a)

600 Spot diameter (nm)

500 400

Defect 1 Defect 2 Defect 3 Defect 4

4

(b)

3

300

2

200 100 0

(c) 0

2

4 6 Tip bias (V)

8

1

Figure 9. Scanning Impedance Microscopy of a single walled carbon nanotube acquired with increasing tip voltage from left to right (a). Contrast maxima indicate regions where the p type tube is depleted of carriers. Defects are labeled in the SGM image (b). The voltage dependence of this contrast can be used to extract the valence band energy at the defect (c).

93

Adding the frequency dependence of SIM to a configuration that isolates a single wire has been used to determine the effect of defects on transport and the local electronic structure of an individual defect. For the case of a carbon nanotube, the tip is used as a local gate to drive individual defects into carrier depletion. The consequent increase in the resistance of the tube is registered as contrast variation in SGM and SIM images as shown in Figure 9(a) and (b). A relatively straight forward calculation shows that the voltage dependence of the contrast is related to the valence band energy at each defect. The tip voltage required to cause the defect to become a scattering center is the energy difference between the local valence band and the Fermi energy. This is illustrated in Figure 9 (c), which also compares the spatial resolution of SGM and SIM on a nanotubes. This approach can be generalized to all classes of molecular wires. 4.2. DOMAIN INTERACTIONS IN FERROELECTRIC THIN FILMS Complex transition metal oxides, as exemplified by the perovskite family of compounds, offer wide ranging and intricate property combinations. Compounds in this class can be insulating or superconducting, optically active or transparent, ferroelectric or ferromagnetic. Some exhibit non intuitive combinations of properties; transparent and conducting, ferroelectric and ferromagnetic, etc. Understanding behavior in these systems requires a combination of electrostatic and dielectric probes. For the case of thin films, in which grain sizes and domain sizes can be less than 100 nm, local probes are a necessity. First observations of ferroelectric domains were based on SSPM [81, 82, 83, 84, 85]. Motivated by the concept of high density ferroelectric memory, many groups used a conductive tip of a SPM to induce local domain orientation and imaged the result with SSPM. An interesting observation was that the domains appeared not be to stable. Since the fundamental limit of domain stability was under some debate this result was accepted. However it was shortly realized that atomic polarization is compensated at a surface that is exposed to ambient. The SSPM measurements reflect the sum of surface charge due to polarization and compensation charge due to adsorbates. In most cases the sign of the surface charge is that of the compensating species, i.e. opposite that of the domain polarization. Nevertheless, SSPM was used to examine domain wall motion and ferroelectric/paraelectric phase transitions. The advent of PFM, which is not sensitive to the compensation charge, promised to eliminate ambiguity; however, first measurements were not consistent. Luo et all found that the temperature dependence of PFM contrast of triglycine sulphate near the Curie temperature was similar to that of the spontaneous polarization rather than the piezoelectric coefficient [48]. The gradual change in potential was attributed to the dominance of electrostatic interactions due to the charged surface [ 86 ], since the electromechanical response based on the piezoelectric coefficient would diverge in the vicinity of the Curie temperature. Contrasting behavior was observed in the existence of a lateral PFM signal which could not result from surface charge alone [64, 87, 88], the absence of relaxation behavior that is characteristic of compensation charge [32, 89], as well as numerous observations using both EFM/SSPM and PFM [90, 91] that clearly pointed to a significant electromechanical contribution to PFM contrast. The discrepancies can be resolved with reference to Figure 4. Experiments done under

94

conditions in which the sample/tip interaction was influenced by electrostatic forces as well as piezoelectric deformation obtained different results than those with conditions squarely in SI regime where the signal is directly related to the piezoelectric coefficient. In many cases, topological information on domain structure and orientation obtained from SPM images is sufficient and numerous observations of local domain dynamics as related to polarization switching processes [92, 93, 94], ferroelectric fatigue [95, 96, 97, 98], phase transitions [48, 99, 100, 101], mechanical stresses [102], etc. have been made. However, the detailed analysis of local ferroelectric properties including hysteresis measurements [103], stress effects in thin films [104], size dependence of ferroelectric properties [105, 106], etc. requires quantitative interpretation of the SPM interaction. Many of the articles in this volume will treat various aspects of these issues.

5. Other Techniques and Future Directions

These few examples illustrate the tremendous advances that are being made in understanding linear and non linear properties of complex systems at the nanometer scale. In some cases the complexity is due to size, in others the presence of interfaces, boundaries, and domain structures. The ability to localize the measurement to nms and to access properties more involved than linear electromagnetic responses is critical to understanding these systems. The former is facilitated by using scanning probes, the latter by utilizing multiple modulation schemes. The length constraints of a single article have precluded a comprehensive overview of all techniques. The selection presented here was driven in part by the topics treated in the associated conference and by the content of the other papers in this volume. One obvious omission is the class of measurements based on magnetic interactions. For example Scanning Squid Microscopy and local Magnetic Resonance Microscopy continue to be exciting. Similarly, near field optical probes, advanced tunneling microscopies, and the suite of local mechanical property techniques are developing simultaneously. The subset of the scanning probe field presented here is very dynamic right now, perhaps being driven by the increasing interest in a wider range of complex materials. The next few years will witness another leap in understanding as the multiple modulation probes are combined with the atomic resolution capabilities of non contact AFM.

Acknowledgements

We are grateful for continued support by the Department of Energy and the National Science Foundation for that part of the work done in our labs. Clayton Williams, Charlie Johnson, Lukas Eng and Steve Anlage have graciously provided figures. We have benefited from extensive discussions with Sergei Kalinin, Lukas Eng, and Alexei Gruverman.

95

References 1.

As determined from m COMPENDEX for 2002

2.

Bonnell D.A. (ed.) (2000) Scanning probe microscopy and spectroscopy: theory, techniques and applications, 2ndd edn, New York: Wiley VCH.

3.

Wiesendanger R. ((ed.) 1994) Scanning probe microscopy and spectroscopy-methods and applications, Cambridge University Press, Cambridge, UK.

4.

Friedbacher, G., Fuchs, H. (1999) Classification of scanning probe microscopies - (technical report), Pure and applied chemistry 71, 1337-1357.

5.

Bottomley, L. (1998) Scanning Probe Microscopy. Anal Chem 70, 425R-475R (and the references therein).

6. 7.

Israelachvili, J.N. (1992) Intermolecular and Surface Forces, Academic Press, New York. Hartmann, U. (1989) The point dipole approximation in magnetic force microscopy. Phys. Lett. A 137, 475-478.

8.

Hartmann, U. (1999) Magnetic force microscopy, Annu. Rev. Mater. Sci. 29, 53-87.

9.

Yongsunthon, R., Stanishevsky, A., Williams, E.D., and Rous, P.J. (2003) Mapping electron flow using magnetic force microscopy, Appl. Phys. Lett. 82, 3287-3289.

10. Yongsunthon, R., Stanishevsky, A., Williams, E.D., et al. (2002) Test of response linearity for magnetic force microscopy data, J. Appl. Phys. 92, 1256-1261. 11. Alvarez, T., Kalinin, S.V., Bonnell, D.A. (2001) Magnetic-field measurements of current-carrying devices by force-sensitive magnetic-force microscopy with potential correction, Appl. Phys. Lett. 78, 1005-1007. 12. De Wolf, P., Stephenson, R., Trenkler, T., Clarysse, T., Hantschel, T., and Vandervorst, W. (2000) Status and review of two-dimensional carrier and dopant profiling using scanning probe microscopy, J. Vac. Sci. Technol. B 18, 361-368. 13. De Wolf, P., Snauwaert, J., Hellemans, L., Clarysse, T., Vandervorst, W., D’Olieslaeger, M., and Quaeyhaegens D. (1995) Lateral and vertical dopant profiling in semiconductors by atomic force microscopy using conducting tips, J. Vac. Sci. Technol. A 13, 1699-1704. 14. De Wolf, P., Clarysse, T., and Vandervorst, W. (1998) Low weight spreading resistance profiling of ultrashallow dopant profiles, J. Vac. Sci. Technol. B 16, 401-405. 15. Matey, J.R. and Blanc, J. (1985) Scanning capacitance microscopy, J. Appl. Phys. 57, 1437-1444. 16. Barrett, R.C. and Quate, C.F. (1991) Charge storage in a nitride-oxide-silicon medium by scanning capacitance microscopy, J. Appl. Phys. 70, 2725-2733. 17. Huang, Y., Williams, C.C., and Wendman, M.A. (1996) Quantitative two-dimensional dopant profiling of abrupt dopant profiles by cross-sectional scanning capacitance microscopy, J. Vac. Sci. Technol. A 14, 1168-1171. 18. Hantschel, T., Niedermann, P., Trenkler, T., and Vandervorst, W. (2000) Highly conductive diamond probes for scanning spreading resistance microscopy, Appl. Phys. Lett. 76, 1603-1605. 19. Marchiando, J.T. and Kopanski, J.J. (2002) Regression procedure for determining the dopant profile in semiconductors from scanning capacitance microscopy data, J. Appl. Phys. 92, 5798-5809.

96 20. Yang, J. and Kong, F.C.J. (2002) Simulation of interface states effect on the scanning capacitance microscopy measurement of p–n junctions, Appl. Phys. Lett. 81, 4973-4975. 21. Lányi, S., Török, J., and Rehurek, P. (1996) Imaging conducting surfaces and dielectric films by a scanning capacitance microscope, J. Vac. Sci. Technol. B 14, 892-896. 22. Belaidi, S., Girard, P., and Leveque, G. (1997) Electrostatic forces acting on the tip in atomic force microscopy: modelization and comparison with analytic expressions, J. Appl. Phys. 81, 1023-1030. 23. Edwards, H., McGlothlin, R., San Martin, R., U, E., Gribelyuk M., et al. (1998) Scanning capacitance spectroscopy: an analytical technique for pn-junction delineation in Si devices, Appl. Phys. Lett. 72, 698700. 24. Viscoly-Fisher, I., Cohen, S.R., and Cahen, A. (2003) Direct evidence for grain-boundary depletion in polycrystalline CdTe from nanoscale-resolved measurements, Appl. Phys. Lett. 82, 556-558. 25. Tran, T., Oliver, D.R., Thompson, D.J., and Bridges, G.E. (2002) Capacitance sensor with sub-zeptofarad (1018 cm-3 in order for the band bending to be larger then a few mV.

142

Figure 18. Position of

EF

relative to the intrinsic Fermi level (local band bending) at the surface of a silicon

sample due to a charged defect located at varying distances from the surface.

4.3. NUMERICAL ANALYSIS OF THE TIP-SEMICONDUCTOR ELECTROSTATIC FORCE As mentioned in section 1 above, the KPFM feedback loop ideally nullifies the force component at the frequency ω , therefore a calculation of the KPFM signal amounts to find the tip voltage, Vtip , that minimizes the total electrostatic force, V DC in Eq. (3). The calculation of the electrostatic force, F , can be accomplished in two different ways: the first by calculating the total electrostatic energy for two closely separated tip heights, and then calculate the derivative with respect to the tip sample distance as:

F =−

∂U ∂ 1 2 = − ³ E dV ∂z ∂z V 8π

(16)

where E is the electric field and the integral is taken over the entire volume of the tip and sample. The same result can be obtained by calculating the Maxwell stress that acts on the tip surface. Because the tip is assumed to be a perfect conductor, the Maxwell stress is a vector (rather then a tensor) that is perpendicular to the tip surface. The electrostatic force can then be calculated by integrating the Maxwell stress over the entire tip surface:

143

F =−

∂U 1 =− ∂z 2ε s

³E

2

dS Sˆ ⋅ zˆ

(17)

S

where dSˆ is a tip surface element, and zˆ is a unit vector in the z direction. In both cases the electric field is calculated from the solution of the Poisson equation of the tipsample system; both methods gave identical results. The KPFM method uses AC C modulated tip bias and the nullifying procedure is done on the ω component of the force. However, the force and as a consequence, the Poisson equation, are solved for steady state conditions. This is justified based on the following argument. It is assumed (and this was also verified experimentally) that the frequency ω is large enough that the semiconductor band edges cannot follow the AC voltage changes. Therefore, Eq. (3) can be solved for a given DC C bias. Following Hudlet et. al.18 the force at a frequency ω can be expanded in a Taylor series see from the DC C force to give:

F(

p

(ω ))

( ) + ∂FF

DC

∂Vtip

p

VAC sin(ω ) +

∂ 2 FDC 2 ∂Vtip

(

AC

(ω ))2

(18)

where FDC is the calculated DC C force and Vttip is the DC C bias applied to the tip or to

the sample by the external voltage source. The force at a frequency ω is the second term on the r.h.s. of Eq.

F(

p

(ω ))

( ) + ∂FF

DC

p

∂Vtip

VAC sin(ω ) +

∂ 2 FDC 2 ∂Vtip

(

AC

and can be neglected. The measured value VCCPD is the value Vttip electrostatic force (the analogous to V DDC , in Eq.

(ω ))2 (18) that minimizes the

(3)).

The same approach was applied in the past for a metallic sample where the sample surface potential was assumed constant.19 Here the force is calculated as the derivative of the total energy including the energy required to bend the semiconductor bands; this energy can be due to the presence of surface states and/or due to tip induced band bending. Figure 19 shows the potential distribution calculated according to the above formalism for a defect (with a total charge q) located 2 nm below the sample surface. The tip bias is 0.3 V below the potential of the intrinsic Fermi level at the semiconductor surface, i.e., Vtip =

φs q

− 0.3 .

One of the most important issues in KPFM is whether the presence of the tip has any significant effect on the semiconductor SCR, i.e. is there any tip-induced band bending at the semiconductor surface. Hudlet et. al.18 showed that at least in the one-dimensional case, the tip sample system can be modeled as two capacitors in series. They have also shown, that when the distance between the sample and the tip is reduced, the sample capacitance cannot be neglected and it changes the force acting on the tip. In order to check this effect in the three-dimensional case, we have calculated the surface band

144

bending (E F

Eis ) for a tip located 5 nm above a semiconductor surface and having a

potential of 0.6 Volt higher then the sample. The result is shown in Figure 20. The figure clearly shows that the local changes in the sample band bending induced by the biased tip are very small (less than 40 mV). This means that the change in the semiconductor SCR capacitance due to the biased tip is negligible, and that the tipsample electrostatic force is governed by the 'air capacitor'. A similar calculation (including the calculation of the electrostatic energy and the electrostatic force) was conducted for a silicon sample having surface states density of n SS = 3 ⋅ 10 11 cm -2 but with varying doping concentration, and a tip-sample bias of 0.5 V.

Figure 19. Potential distribution for the near surface defect with tip-sample bias of 0.3 V; the defect is indicated by the oval under the tip.

145

Figure 20. Local band bending ( E F

− Eis ) for a semiconductor that has no surface states and for tip-

sample distance of 5 nm and an applied bias of Vttip = 0.6 V

The results shown in Figure 21 demonstrate that the electrostatic energy increases with decreasing doping density because the energy required for the formation of the depletion region increases; this can be explained as follows. The electrostatic energy is due to the tip-air-semiconductor surface capacitor, which is in series with the SCR capacitor. In the one-dimensional case the sample SCR capacitance decreases with increasing doping level, therefore the total energy increases with the increase in doping. However, in the 3D case the energy of the air capacitor can be neglected due to the tip shape. The energy of the semiconductor SCR is proportional to the voltage square over the depletion region length, and the latter increases as the square root of the voltage. Therefore, the total energy increases with decreasing doping level as shown in Figure 21.

146

Figure 21. Electrostatic energy as function of tip-sample m distance for varying doping concentration.

All the curves in Figure 21 have the same dependence on the tip-sample distance and they differ only by a constant. This means that the electrostatic force, which is the derivative of the energy, should be the same for all the doping concentrations. This is shown in Figure 22 which shows the electrostatic force (as a function of tip-sample distance) calculated from the electrostatic energy using: F =

∂U . ∂Z Q

It must be reemphasized that in KPFM measurements the potential difference between the tip and the sample m is nullified, so typically the voltage difference between any point on the tip and the sample surface is much lower then the 0.5 Volt used in the calculations above. In summary, we have found that the dominant contribution to the tip-sample electrostatic force is from the tip-air-sample surface capacitor and the SCR capacitance can be neglected. Thus, the assumption used by Hochwitz et. al.15, Hudlet et. al.9, and Jacobs et. al.19 is correct. Furthermore, the above conclusion simplifies tremendously the simulation of semiconductors KPFM images. This is because it requires solving only the Laplace equation using fixed (potential independent) boundary conditions at the sample surface.

147

Figure 22. The absolute value of the electrostatic force as a function of tip-sample distance for all the doping concentrations shown in Figure 21.

4.4. COMPARISON WITH EXPERIMENTAL RESULTS Considerable effort, both experimental and theoretical, has been devoted to the understanding of the properties of atomic steps formed during epitaxial growth of semiconductors.20 In particular, reconstructions of step edges can locally change the electronic properties and lead to charged states in the gap; such charges have a pronounced influence on the physics of the microscopic mechanisms of diffusion and sticking of charged adatoms21 and vacancies at the step edges. Therefore, knowledge of the existence of charges along steps is important m for the optimization of growth parameters and interface properties. Measurements conducted by Sommerhalter et al.22 have shown that the atomic steps present at the GaAs(110) surface show local changes in band bending due to the presence of charged surface states. KPFM measurements across atomic steps are an ideal system to compare to our simulation since both the width of the defect and its location (right at the surface) are very well defined. Figure 23 shows a KPFM measurement (dots) conducted under ultra high vacuum (UHV) conditions of an atomic step on a cleaved GaP (110) surface.23 The CPD calculation (solid line) was carried out in the following way. First we have calculated the CPD resulting from a Gaussian potential distribution 0.1 V high, and 50 nm (standard deviation) wide.

148

Figure 23. KPFM measurement on UHV cleaved p-GaP (110) (dotted line) together with the calculated CPD (solid line) for a tip height of 6 nm.

The results are shown in Figure 24 for seven different tip-sample distances. The figure shows that for a tip height of 5 nm (the experimental scanning height in the UHV KPFM measurements) the CPD signal is reduced by a factor of around 2.5, and its width is increased by a factor of 2 relative to the theoretical surface potential represented by the top curve. The surface charge density at the step was extracted by fitting the measured CPD to the calculated surface potential, based on two main assumptions: 1) The surface states induced by the atomic steps are positioned only at the sample surface; 2) Average surface states densities are assumed (and not quantum mechanical distributions(, therefore, we assumed that the surface states are distributed in a width of 2 nm (the lower limit for the validity of the Poisson equation) The above assumptions reduce the number of unknown (‘fitting’) parameters to two: the surface charge density at the atomic step, and the background surface charge density. The calculation is conducted by assuming two initial values for these two surface charge densities, and then the 3D semiconductor surface potential distribution is calculated using Eq. (6). We then calculate the measured CPD in the presence of the tip and compare the result with the measurement; this procedure is repeated, by changing both the background surface charge, and the charge at a step until a good visual fit like the one shown in Figure 23 is obtained. This fit was obtained using a background and defect surface charge densities, of and nssbackground = (6 ± 2) ⋅1011 cm -2

nssstep = (6 ± 1) ⋅1012 cm-2 respectively. Such a reasonable fit could not be achieved using other combinations of charge densities. This is due to the fact that the width of the Gaussian CPD profile at the step is much more sensitive to the background charge

149

density (screening effect), while its magnitude is more sensitive to the charge density at the step itself. In addition these results are in a good agreement with measurements of s

the surface built-in potential, Vbi (see 2.1 above)~ 1.2 eV, V of a GaP pn junction measured

n

background ss

on

the

same sample. A surface charge density of -2 corresponds to band bending of eV on both the p≅ 0.4 = (6 ± 2) ⋅10 cm 11

s

and n-sides of the junction thus giving a Vbi of around 1.3 eV.

Figure 24. Calculated CPD for a Gaussian surface potential distribution for seven different tip heights.

In summary, we have presented a method for obtaining the real potential distribution from KPFM measurement. It was demonstrated that this method could be used to extract the charge density at a local defect (atomic step) and under certain conditions, the background surface charge density. The latter is critical information necessary for the development of two dimensional doping measurement methods.

5. Conclusions This review described the use of the KPFM technique to image potential distribution across semiconductor surfaces and devices. The important contribution of this work can be divided into three main parts: a) Measurement of operating semiconductor devices. b) Development of the KPFM technique to measure minority carrier diffusion length in semiconductors.

150

c) Development of a general methodology for calculating tip-semiconductor electrostatic interaction, which can be used for extracting real CPD images from KPFM data. This work serves just as a starting point for the reconstruction of the real potential distribution on semiconductor surfaces. Faster calculation schemes should be developed, since the present computation time is a major obstacle for a practical usage of this methodology. The relation between the sample physical parameters and the surface potential distribution should be improved. For example, the charge term (the right hand-side) of the Poisson equation (Eq. (6)) is given using average charge distribution. This is not valid for very small scales (< 2 nm) and have to be replaced with charge density calculated using the Schrödinger equation. In addition, for the simulation of measurements under ambient conditions, the model must incorporate a water layer on top of the semiconductor surface.

6. Acknowledgement This research was supported by the Israel Science foundation administered by the Israel Academy of Sciences and Humanities-Recanati and IDB group foundation, and by grant 9701 of the Israel Ministry of Sciences. R. S. was supported by Eshkol special scholarship of Israel Ministry of Sciences. The authors acknoweledge the collaboration with Th. Glatzel, and S. Sadewasser from the Hahn Mietner Institute (Berlin) in the UHV KPFM measurements.

References 1. Nonenmacher, M., O’Boyle, M.P,. and Wickramasing H.K. (1991) Kelvin probe force microscopy, Appl. Phys. Lett. 58, 2921-2923. 2. Leng, Y., Williams, C.C., Su, L.C., and Stringfellow G.B. (1995) Atomic ordering of GaInP studied by Kelvin probe force microscopy, Appl. Phys. Lett. 66, 1264-1267. 3. Kikukawa, A., Hosaka, S., and Imura R. (1995) Silicon pn junction imaging and characterizations using sensitivity enhanced Kelvin probe force microscopy, Appl. Phys. Lett. 66, 3510-3512. 4. Vatel, O.and Tanimoto M. (1995) Kelvin probe force microscopy for potential distribution measurement of semiconductor devices J. Appl. Phys. 77, 2358 -2362. 5. Chavez-Pirson, A., Vatel, O., Tanimoto, M., Ando, H., Iwamura, H., and Kanbe H. (1995) Nanometer-scale imaging of potential profiles in optically excited n-i-p-i heterostructure using Kelvin probe force microscopy, Appl. Phys. Lett. 67, 3069-3071. 6. Mizutani, T., Arakawa, M., and Kishimoto, S. (1997) Two-dimensional potential profile measurement of GaAs HEMT’s by Kelvin probe force microscopy, IEEE Elec. Dev. Lett. 18, 423-425; Arakawa, M., Kishimoto, S., and Mizutani, T. (1997) Kelvin probe force microscopy for potential distribution measurements of cleaved surface of GaAs devices, Jpn. J. Appl. Phys. 36, 1826-1829. 7. Shikler, R., Fried, N., Meoded, T., and Rosenwaks, Y. (1999) Potential Imaging of Operating Light Emitting Devices using Kelvin Force Microscopy, Appl. Phys. Lett. 74, 2972-2974; Shikler, R., Meoded, T., Fried, N., Mishori, B., and Rosenwaks, Y. (1999) Two Dimensional Surface Band Structure of Operating Semiconductor Devices, J. Appl. Phys. 86, 107-113. 8. Henning, A.K., Hochwitz, T., Slinkman, J., Never, J., Hoffman, S., Kaszuba, P., and Daghlian, C. (1995) Two-dimensional surface dopant profiling in silicon using scanning Kelvin probe microscopy, J. Appl. Phys. 77, 1888-1896.

151 9. Hudlet, S., Jean, M.S., Roulet, B., Berger, J., and Guthmann, C. (1995) Electrostatic forces between metallic tip and semiconductor surfaces, J. Appl. Phys. 59, 3308-3314. 10. Sandroff, C.J., Nottenburg, R.N., Bischoff, J.C., and Bhat, R. (1987) Dramatic enhancement in the gain of a GaAs/AlGaAs heterostructure bipolar transistor by surface chemical passivation, Appl. Phys. Lett. 51, 3335. 11. Mayergoyz, I.D. (1986) Solution of the nonlinear Poisson equation of semiconductor device theory, J. Appl. Phys. 59, 195-199. 12. Gustafsson, A., Pistol, M.E., Montelius, L., and Samuelson, L. (1998) Local probe techniques for luminescence studies of low-dimensional semiconductor structures, J. Appl. Phys. 84, 1715-1775; Vertikov, A., Ozden, I., and Nurmiko A.V. (1999) Investigation of excess carrier diffusion in nitride semiconductors with near-field optical microscopy, Appl. Phys. Lett. 74, 850-852. 13. Goodman, M. (1961) A method for the measurement of short minority carrier diffusion lengths in semicondutors, J. Appl. Phys. 32, 2550-2552. 14. Markiewicz, P., and Goh, M.C. (1994) Atomic force microscopy probe tip visualization and improvement of images using a simple deconvolution procedure, Langmuir 10, 5-7. 15. Hochowitz, T., Henning, A.K., Levey, C., Daghlian, C., and Slinkman, J. (1996) Capacitive effects on quantitative dopant profiling with scanned electrostatic force microscopes, J. Vac. Sci. Technol. B 14, 457464. 16. Mayergoyz, I.D. (1986) Solution of the nonlinear Poisson equation of semiconductor device theory, J. Appl. Phys. 59, 195-199. 17. Korman, C.E., and Mayergoyz, I.D. (1990) A globally convergent algorithm for the solution of the steady-state semiconductor device equations, J. Appl. Phys., 68, 1324-1334. 18. Hudlet, S., Saint Jean, M., Roulet, B., Berger, J., and Guthmann, C. (1995) Electrostatic forces between metallic tip and semiconductor surfaces, J. Appl. Phys. 77, 3308 -3314. 19. Jacobs, H.O., Knapps, H.F., Muller, S., and Stemmer, A. (1997) Surface potential mapping: A qualitative material contrast in SPM, Ultramicroscopy 69, 39-49; Belaidi, S., Lebon, F., Girard, P., Leveque, G., and Pagano, S. (1998) Finite element simulations of the resolution in electrostatic force microscopy, Appl. Phys. A 66, S239-S243. 20. Heinrich, M., Domke, C., Ebert, Ph., and Urban, K. (1996) Phys. Rev. B 53, 10894. 21. Doi T., and Ichikawa, M. (1995) Direct Observation of Electron Charge of Si Atoms on a Si(001) Surface, Jpn. J. Appl. Phys. 34, 25-29. 22. Sommerhalter, Ch., Matthes, Th.W., Glatzel, Th., Jäger-Waldau, A., and Lux-Steiner, M.Ch. (1999) High-sensitivity quantitative Kelvin probe microscopy by noncontact ultra-high-vacuum atomic force microscopy, Appl. Phys. Lett. 75, 286-288. 23. Glatzel, Th., Sadewasser, S., Shikler, R., Rosenwaks, Y., and Lux-Steiner, M.Ch. (2003) Kelvin Probe Force Microscopy on III-V Semiconductors: The Effect of Surface Defects on the Local Work Function, Materials Sci. and Eng. B 102, 138-142.

EXPANDING THE CAPABILITIES OF THE SCANNING TUNNELING MICROSCOPE

K.F. KELLY,* Z.J. DONHAUSER, B.A. MANTOOTH, and P.S. WEISS Departments of Chemistry and Physics, The Pennsylvania State University 152 Davey Laboratory, University Park, PA 16802-6300, USA

Contents 1. 2.

Introduction Alternating Current Coupled to the STM 2.1. Introduction 2.2. Experimental setup 3. Dopant-profiling with ACSTM 3.1. Theory 3.2. Characterization of bulk silicon 3.3. Atomic scale imaging of pn junctions 4. Feature Tracking in STM 4.1. Cross-correlation 4.2. Pixel roundoff correction 5. Application of the Feature Tracking Algorithm 5.1. Single molecule switching 5.2. Diffusion of benzene on Ag{110} 6. Conclusions

Abstract Scanning probe microscopes allow unprecedented views of surfaces and the site-specific interactions and dynamics of adsorbates. Our efforts to identify and to characterize atoms and molecules on surfaces and how it is that the scanning tunneling microscope images these surfaces and adsorbates will be discussed. We have extended the capabilities of scanning probe microscopes in several ways; two in particular will be highlighted. In the first section, recent advances in tunable microwave frequency scanning tunneling microscopy (STM) allow dopant profiling at unprecedented resolution will be presented. We apply nonlinear tunable microwave frequency scanning tunneling microscopy and spectroscopy to profiling dopants at ultrahigh resolution in *

Present Address: Department of Electrical Engineering, Rice University, Houston, TX 77251, USA

153 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 153-171. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

154

semiconductors that is sensitive to both dopantt type and density. We are then able to use a spectroscopic imaging mode to map the dopant density at the atomic scale. In the second part of this chapter, advanced image processing techniques that extend the scientific capabilities of STM will be presented. A digital image tracking algorithm based on Fourier-transform crosscorrelation has been developed to correct for instrumental drift in scanning tunneling microscope images. This tracking algorithm was used to monitor conductance changes associated with different conformations in conjugated switching molecules and to trace the diffusion of individual benzene molecules on silver.

1. Introduction With its ability to image and to manipulate single molecules and atoms, the scanning tunneling microscope has produced many new discoveries in physics, chemistry, and engineering. However, two drawbacks to this instrument are its restriction to a serial, point-by-point mode of data acquisition and the inevitable convolution of the electronic information that it acquires with surface topography. In this chapter, we discuss in detail two separate modifications to the operation of the scanning tunneling microscope to overcome both these challenges to obtain capacitive and dynamic information.

2. Alternating Current Coupled to the STM 2.1. INTRODUCTION The alternating current scanning tunneling microscope (ACSTM) extends the capabilities of the STM from a simple atomic-scale topographic and local density of states (LDOS) probe to a spectroscopic probe capable of measuring local variations in chemical, dielectric, and magnetic properties [1-10]. Spectroscopic contrast has already been observed on several surfaces [5, 6], but its interpretation is not resolved. In addition, the ACSTM extends the capabilities of the tunneling microscope beyond conducting substrates to enable imaging and local spectroscopy on insulators [1-10]. Kochanski first demonstrated ACSTM operation on insulators and semiconductors by detecting the third harmonic of the applied bias frequency in a cavity resonant at the detection frequency [1]. This scheme has since been applied by others to study substrates ranging from metals to insulatorr [2, 3] and to measure doping profiles in silicon [6]. In our previous work, we have avoided the use of a resonant cavity in favor of retaining broadband frequency response [4, 7-10]. The ACSTM is conceptually very similar to conventional STM. The sample is scanned by a piezoelectrically driven probe tip controlled by a feedback signal sensitive to the probe tip-sample gap distance. In ACSTM, high (typically microwave) frequency bias is applied to the probe tip. As described below, we can apply a bias at any frequency from dc to 20 GHz, an additional variable with which to interrogate the surface under study. The response that can be expected varies from quasi-static at low frequencies to dispersive at high frequencies. In the quasi-static limit the ACSTM can be treated as a conventional STM experiment with a time-varying dc bias. In this limit

155

there is no extra information in the frequency domain. At sufficiently high frequencies, the tunnel junction response becomes dispersive. This deviation from quasi-static behavior is due to the inability of the system to follow the rapidly oscillating ac field. By mapping the frequency and amplitude dependence of this response, we can determine the origin(s) of these effects and use this information to elucidate the local chemical, electronic, and/or structural properties of the substrate surface [8, 10]. Only a small fraction of the power incident on the ACSTM probe tip is expected to interact with the tunneling junction, therefore the response signal due to the tunneling junction will be small compared to the total incident power at the fundamental frequency. Further complications arise from non-local interactions of the probe with its surroundings, which can dominate the ACSTM response at the fundamental frequency; these non-local interactions can affect the microwave energy in the tunneling junction, thus complicating the interpretation of the local nonlinear signal as well. Minimizing the non-local interactions is desirable for any ACSTM detection scheme and is the focus of this report. 2.2. EXPERIMENTAL SETUP The STM is a well-established surface probe that is routinely capable of spatial resolution on the atomic scale. Our goal has been to incorporate the unparalleled imaging capability of the STM with an accurate dopant profiling system. The instrument is based on a custom-built alternating current scanning tunneling microscope (ACSTM) [4, 5, 7]. The instrument we have developed is tunable over a wide frequency range and has the ability to introduce up to three frequencies to the tunneling junction simultaneously. Our instrument detection system m is capable of detecting transmitted and reflected, fundamental and nonlinear alternating current signals, and can determine phase as well as magnitude information. This allows us local access to the frequency dependence, energy, and position of electronic t responses of doped semiconductors. Introducing multiple frequencies to the STM probe tip allows us to generate a difference frequency signal that is produced due to the nonlinear nature of the STM tunnel junction. Data acquisition was accomplished using a difference frequency mixing strategy. A schematic of the instrument configuration is shown in Figure 1. In addition to the conventional DC bias applied to the tunnel junction, two frequencies are introduced from tunable waveform generators. The frequencies are offset by a small amount (typically 5 kHz) that becomes the detected frequency (ω1 − ω2 = ∆ω). The mixing product (∆ω) occurs at low frequency, so it is conveniently extracted and detected using a lock-in amplifier. A reference signal for the lock-in is created by splitting off a portion of the applied AC signals prior to the tunnel junction in a directional coupler, and sending them through a diode to generate the desired difference frequency.

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Figure 1. Schematic of the difference frequency detection scheme. Two frequencies, ω1 and ω2 are generated and combined. A portion of the mixed signal is sent through at diode, which creates the nonlinear difference frequency reference signal, ∆ωreff for the lock-in amplifier. The remainder of the mixed signal is combined with the DC bias voltage and sent to the STM tip. The nonlinear nature of the STM tunnel junction and of the sample creates the difference frequency signal, ∆ωsignal. This is extracted from the tunnel current and sent to the lock-in amplifier for comparison with ∆ωref.

3. Dopant-profiling with ACSTM 3.1. THEORY Scanning probe microscopes are extremely important for characterizing semiconductors with very high spatial resolution. According to the 2001 International Technology Roadmap for Semiconductors [11], there was already an unmet critical need to be able to determine 2-D dopant profiles with 3 nm resolution in 1999, and 1 nm spatial resolution will be required by 2008. Much of the recent work in this field has focused on the development of the scanning capacitance microscope (SCM) [12-18]. These instruments have shown high sensitivity towards dopant density and type, and have accurately imaged devices on semiconductor surfaces with resolution as high as 10 nm [12, 13]. However, the lateral resolution when using capacitance detection is limited by the probe tip geometry and dopant level [12, 14, 15]. Improving spatial resolution may require the development of new scanning probe techniques. To address these and related issues, we have developed a novel dopant profiling tool based on the alternating current scanning tunneling microscope (ACSTM). We have already used this method to probe the electronic properties of self-assembled monolayers and single molecules [19-21]. The present work extends the technique to semiconductor dopant profiling. As shown below, the nonlinear AC signal is sensitive to dopant type and density, giving us a convenient means to integrate a dopant profiling system with the high-resolution STM. In this work, we first characterize the frequency and voltage response of our technique, and then use it to image patterns on doped semiconductor substrates to determine both dopant density and dopant type.

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The tip-sample distance is precisely controlled using the DC tunneling current for feedback; this prevents the metal tip from contacting the silicon substrate. This tip-airsilicon arrangement resembles a metal-insulator-semiconductor (MIS) structure. Signals resulting from an MIS structure will consist of both capacitive C(V) and conductive G(V) terms. C(V) originates from the capacitances of the air gap and silicon depletion layer, while G(V) describes losses in the AC signal from effects such as the STM tunneling current, series resistance in the silicon substrate, and sample and tip local density of states. The difference frequency detection strategy makes use of the capacitive characteristics of the doped silicon, which vary according to dopant density and type. Additionally, it allows us to tune the instrument over a range of fundamental frequencies, from 0-20 GHz, while the output signal detected remains at a constant frequency. A similar two-frequency mixing strategy designed to image p-n junctions using a microwave frequency compatible AFM has been previously reported by Schmidt et al. [22]. These particular AFM experiments used the sum and third harmonic frequencies as nonlinear mixing product signals. It was found that the sum frequency V2, signal and the third harmonic signal are proportional to dC/dV and d2C/dV respectively. We expect the difference frequency signal to be analogous to the sum frequency signal, and be proportional to dC/dV. V A complete discussion of the nature of the mixing products can be found in references 22 and 23.

Figure 2. (A) Model capacitance curves of a metal-insulator-semiconductor structure for both a p- and n-type semiconductor. (B) Corresponding dC/dV V curves, determined numerically.

Typical capacitance versus voltage curves for a MIS structure are shown in Figure 2 (A). The numerical differentials are shown in Figure 2 (B). If we consider the ACSTM tip-gap-semiconductor a MIS structure, the precise shape and magnitude of these curves would be determined by a variety of factors, including the dopant concentration, the distance of the tip from the sample (the insulator thickness), the magnitude of the tunneling current, and the probe tip geometry. Although they only model our system, we can use the curves in Figure 2 to understand qualitatively the difference frequency signal observed in images at different biases. Because we expect the magnitude and phase of the signal to be related to the differential capacitance [22, 23], we expect that the largest signals would occur at bias voltages near 0 volts, and that the signal would decrease as the magnitude of the bias increases. Additionally, we expect an 180o relative phase shift between n- and p-type regions on the semiconductor surface. While our system is significantly more complicated than a simple fixed-geometry MIS capacitor,

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the dC/dV V model is a first-order approximation for the contrast observed in difference frequency images obtained at different bias voltages. 3.2. CHARACTERIZATION OF BULK SILICON For the initial characterization experiments on uniformly doped Si, we purchased p- and n-type silicon samples from Virginia Semiconductors, Inc, Fredricksburg, VA 22401, USA. Samples were prepared by annealing at 950 oC for one hour, and then the surfaces were cleaned with a 1:1 H2O2:HCl solution. Immediately prior to all measurements with the ACSTM, we dipped the test samples in a 48% HF solution for ~2 minutes to remove the surface oxide. The experiments were performed in a custom-built tunable ACSTM, described above. All measurements were carried outt at ambient temperature and pressure. The cleaned, doped Si substrates were used to map out the frequency and voltage response as a function of dopant type and concentration. t In Figure 3, the magnitude of the difference frequency signal is plotted as a function of applied frequency and voltage, for both p- and n-type Si. The difference frequency signal is strongly dependent on the fundamental frequency.

Figure 3. Difference frequency signal magnitude for Si(100) as a function of fundamental modulation frequency and bias for (A) 0.001 ohm-cm boron-doped and (B) 1-3 ohm-cm phosphorus-doped silicon.

Figure 3 (A) shows data for highly doped p-type Si (0.001 Ω-cm boron-doped). A wide peak is centered at -0.7 V sample bias. Figure 3 (B) shows data for lightly doped n-type Si (1-3 Ω-cm phosphorous-doped). The peak for phosphorous occurs close to 0 V bias voltage. For both n- and p- type Si, the lower applied frequencies provide the largest signals. This can be attributed to attenuation of high frequency signals through the transmission lines leading into the ACSTM. Reflections of the high frequency signals can occur at transmission line connectors and in coupling to the STM tip. This results in more loss at high frequency, and less signal generated in the tunnel junction for frequencies greater than several hundred MHz. Fortunately, a large nonlinear effect is seen at low frequencies providing the difference frequency signal necessary for semiconductor characterization. It is important to note that the data displayed is the magnitude of the difference frequency signal; all phase information is neglected. Because changes in tip size and shape as well as the precise sample orientation can

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affect the phase of the difference frequency signal, between the n- and p-type Si it is difficult to compare phase information. As described below, we expect a phase difference when n- and p-type Si are compared on the same substrate. 3.3. ATOMIC-SCALE IMAGING OF pn JUNCTIONS To fabricate patterned substrates, we photolithographically prepared a stripe pattern with a 2 µm pitch. Both boron doped p-type and phosphorus doped n-type bulk silicon with concentrations of 1 x 1015 cm-3 were used as the base substrates. We implanted both n- and p-type bulk substrates with boron doses ranging from 1x1011 cm-2 to 2x1014 cm-2 or phosphorus doses ranging from 1x1011 cm-2 to 3x1014 cm-2. All boron implants were done at 35 keV and all phosphorus implants were done at 50 keV. The samples discussed below are bulk n-type Si doped with a 1x1011 cm-2 dose of boron, and bulk ptype Si doped with a 2x1013 cm-2 dose of boron. The final samples consisted of implanted 0.5 µm stripes spaced by 1.5 µm of unimplanted substrate. We then activated all implants by rapid thermal annealing at 1040 oC for 40 seconds.

Figure 4. 1.2 µm x 1.2 µm difference frequency images of a 1x10 cm boron implant stripe in a 1x10 cm phosphorus doped Si(100) substrate at varying biases. The implant stripes were defined photolithographically to have a width of 0.5 µm. Corresponding topographic images are inset. Colorbars indicate signal magnitude. 15

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A series of bias dependent images demonstrate the trends described above. Figures 4 (A-D) are difference frequency images of an n-type substrate doped with boron. The

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sample was prepared to have 0.5 µm p-type (with a nominal concentration of 1015 cm-3) stripes, with a background of 1.5 µm n-type stripes (1015 cm-3). After locating the pattern, difference frequency images were acquired at voltages ranging from -1.5 V to +1.5 V. There is a strong bias dependence observed in the difference frequency images. Between all of the images, the largest signals are seen at +0.5 V and -0.5 V. This is expected for values near zero, based on the model differential capacitance curves. The stripe feature is still evident at higher bias voltages, but the magnitude of the difference frequency signal has greatly diminished. The expected 180o phase shift between n- and p-type silicon is observed in the images acquired at high voltage, confirming our assignment of the different regions. However, the images at lower applied bias do not exhibit this phase shift. This may be due to convolution of the dC/dV V signal with contributions from the STM tunneling current or contributions of higher-order V2). The corresponding topographic images are inset capacitive products (such as d2C/dV in Figures 4 (A-D). The stripe feature is evident in all of these images. This may be an artifact of STM imaging that is due to differences in conductivity of the different regions, or it may be a physical artifact of the substrate processing steps, such as implant swelling.

Figure 5. A 1.2 µm x 1.2 µm difference frequency image of a 1x10 cm boron implant stripe in a 1x10 -3 cm boron doped Si(100) substrate. The implant stripe was defined photolithographically to have a width of 0.5 µm. The corresponding topographic image is inset. 18

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We have also demonstrated the ability of the ACSTM to distinguish between areas of differing dopant concentration for the same dopant type. Figure 5 is an image of a 1015 cm-3 p-type Si substrate nominally doped with stripes of 1018 cm-3 boron. Again, the area of higher concentration was implanted as 0.5 µm stripes with a pitch of 2 µm. In this case, the stripe feature is completely absent from the inset topographic image, yet shows up clearly in the corresponding difference frequency image. From these and related results, we ascertain that the magnitude of the difference frequency signal depends strongly on the local dopant density, and not solely on dopant type. The results

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also demonstrate the ability of the ACSTM to differentiate between areas of high and low concentration using difference frequency detection. This capability is important for imaging patterned substrates. The results and discussion presented above are a preliminary demonstration of the ability of the ACSTM to characterize semiconductor dopants using difference frequency detection. We have successfully shown the ability to distinguish between n- and p-type silicon on the nanometer scale. In addition, we have demonstrated the sensitivity of the ACSTM to dopant concentration. Future experiments with the ACSTM are underway that will determine the concentration detection limits, the maximum spatial-resolution, and will image functional semiconductor devices. 4. Feature Tracking in STM A common problem with many scanning probe microscope (SPM) techniques is drift caused by the relative motion of the sample and the probe. In applications using piezoelectric crystals for raster scanning such as most SPMs, nonlinear responses such as hysteresis and creep [24, 25] can cause the probe to drift across the surface [26]. When attempting long-term temporal analysis (i.e., many hours), the drift can be severe enough to cause the scan window to drift beyond the original region of interest. In an effort to compensate for this drift, we have developed a fast image cross-correlation (XC) technique that is used for post acquisition drift elimination to enable several types of temporal data analyses, including temporal topographic analysis and diffusion of individual molecules. This method of drift elimination is a computational solution; others have implemented instrumentation-based solutions such as tracking tunneling microscopy [27] where the tip is moved in a circle laterally (x,y ( ) and the x and y topographic gradients are measured to keep the tip located over one molecule. This method has been used to track diffusion [28, 29] and rotation [30] of individual molecules on surfaces. While instrumentation-based tracking has the capability of measuring events on the millisecond time scale, it is limited in that it can track only one molecule at a time. In contrast, tracking diffusion or other t events in an image allows monitoring of many molecules with the disadvantage of a slower time scale limited by the image acquisition rate. However, the use of fast scanning instruments (up to 20 frames/s) [31] can bring image tracking close to the time scale of instrumentation-based tracking. One of our interests is in monitoring topographic changes, such as the change in apparent height or even disappearance of a protrusion. In this latter case, the instrumentation-based tracking method, which would fail as a protrusion (or depression) is required for tracking. With XC tracking, the entire image is tracked such that the behavior of individual molecules (topographic or spatial changes) does not play a major role in the overall tracking algorithm; these features are extracted and analyzed after tracking. Image tracking is a vital factor in video compression [32] and machine vision, and as such, is currently an area of intense research [33-36]. Several algorithms have recently been developed to track three-dimensional (3D), two-dimensional (2D), and rotational motions for various microscopies [37-40]. In this work, we assume that SPM techniques will produce image sequences of constant image size and resolution. Because of the orthogonal geometry of mostt scanning probe actuators, such as the beetle-style STM

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[41], only lateral, non-rotational drift should occur. Therefore, a simple XC approach can be used to find the best alignment, and thus the drift between successive images. The maximum of the 2D XC represents the best alignment, also referred to as registration, of two images. The deviation of this maximum from the center of the image represents the drift (in pixels) between the two images. By determining the drift from a sequence of images, we have been able to select and to extract specific regions from these images for further analysis. In this article, we discuss methods to resolve temporally the conductance behavior of conjugated molecular switches inserted into alkanethiol self-assembled monolayers [42] and to analyze the diffusion of individual benzene molecules adsorbed on Ag{110} [43]. Through these analyses, statistical information on conformational switching and hopping distances, respectively, are obtained. 4.1. CROSS-CORRELATION The real-space cross-correlation C(x ( ,y , ) for two digital images, a1(I (I,JJ) and a2(I (I,JJ), can be calculated from a comparison of the subsets of each image [a1(i,j , ) and a2(i,j , )] at multiple offsets (x ( and y) specified by [44, 45]

Variations on this technique have been used previously for several applications including molecular tagging velocimetry [45], stability measurements of STMs [46], and atomic force microscopy (AFM) image alignment [39]. While this approach yields the desired XC, large images can present an exceptionally demanding computational load. In addition, this technique is limited in the available search window (limits on x and y), determined by the size of the subset images such that i + x 2. As a future improvement to this algorithm, an average velocity could be monitored to adjust k as the tracking progresses. Correlating all images to the first image was not done because subsequent images may have enough drift or change over time to the extent that they do not resemble the initial image.

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5. Application of the Feature Tracking Algorithm 5.1. SINGLE MOLECULE SWITCHING It has been observed that a family of conjugated oligomers, based on 4,4’di(ethynylphenyl)-1-benzenethiol, referred to henceforth as ‘‘switch’’ molecules, exhibit reversible conductance switching. Previously, these molecules have been used in nanopore experiments where they have been shown to exhibit negative differential resistance and bi-stable conductance states that can be reversibly switched from high conductance (ON) to low conductance (OFF) states under applied electric t fields [51-53]. To understand the physical processes governing this behavior on an individual-molecule basis, we studied these molecules using the STM as a local probe [42]. Using an insertion process described in detail in the following chapter, single switch molecules are inserted at defect sites in self-assembled monolayers (SAMs) of ndodecanethiolate, such as step edges and substrate vacancies, resulting in a low density of isolated switch molecules adsorbed on the surface (Figs. 7 and 9). When imaged, switches would spontaneously switch from their on state, appearing in STM images as a 4–5 Å protrusion, to their off state appearing as a slight protrusion out of the SAM [26, 42]. Persistence times of both states range from fractions of seconds to many hours.

Figure 9. Topographic analysis of switch molecules. (A) and (B) are frames 1 and 60, respectively, from a movie. The white protrusions are individual switch molecules, the darker large spots are substrate vacancies or SAM domain boundaries. The white boxes represent the extracted region for one switch. Imaging conditions: 2 Vsample = -1.0 V, I = 0.8 pA, 1000x1000 Å , 256x256 pixels, frame acquisition = 2.07 min/frame, delay between frames 556 s. (C) Using the real track from Fig. 3, a region was extracted from 60 frames of a movie, bicubicly interpolated (for improved display, height calculations do not use interpolated data) to twice the original pixel density and median filtered. The 60 frames displayed were extracted from a 200 frame movie acquired over 10 h; further switching was not observed after the displayed frames by this particular switch. Note that each frame has been rotated by 90° relative to (A) in the frame view. (D) The calculated height of the molecule above the SAM.

Constant-current STM images representt the convolution of topography and electronic state (conductivity) of the surface. Therefore, a change in either or both will result in an apparent change in height. To measure this change in height, a movie spanning several hours of data acquisition would be tracked. Using the drift track, individual switch molecules are extracted and processed to calculate the apparent height

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of a switch molecule above the SAM surface throughout the course of the movie (Fig. 4). Large scan areas on the order of 1500 x 1500 to 2500 x 2500 Å2 were acquired over a period of up to 25 h (400 images) to observe as many molecules and switching events as possible. While this tracking algorithm may be implemented in the acquisition software, we have currently used it only as a post acquisition tool to extract and to monitor the topography of each single switch. The topographic height of a switch is obtained by extracting a 16 x 16 pixel region from each image, and calculating the difference of the median of the 9 highest pixels in the extracted region and the median of the lowest fourth of the points from the extracted region. This method was used because of the different adsorption sites in which each switch could reside, such as in or beside a substrate vacancy, or on or next to a step edge. We found that this method gave reproducible topographic heights for most of the molecules analyzed, independent of the insertion site. An example of the time-resolved switching can be seen in Fig. 9(C) and the corresponding calculatedd topographic height difference in Fig. 9(D). A consequence of this method is that an apparent height of zero will occur only if the image were flat and featureless. Because our images have corrugation from the superlattice molecules and noise, the minimum calculated height of a switch molecule is typically 2 Å. In addition, this method works best for well-isolated molecules; if adjacent switch molecules are inside of the extracted region, their topographic height can influence the calculated apparent height of the switch molecule of interest, especially in the case of measuring a switch in an off state. Another complication for the height calculation is that the difference in height between the SAM and the switch can be sufficient to drive the STM feedback mechanism to overcompensate and ‘‘overshoot’’ the real topographic height of the molecule and possibly to oscillate while scanning over the molecule, yielding falsely large apparent heights. To minimize these problems, the extracted region is median filtered before the height calculation to eliminate any noise spikes or gross overcompensation of the feedback mechanism. To characterize the reversible conductance switching, a statistical analysis of the calculated height distributions was performed, as seen in Fig. 3 in Ref. 19. The distribution shows the bimodal characteristic of the switches and the on/off ratios of the molecules. From these data we were able to conclude that the switching observed by the STM was induced by a change in molecular conformation [42]. 5.2. DIFFUSION OF BENZENE ON Ag{110} Another application for this tracking algorithm is to record molecular adsorbate motion across surfaces. The adsorbed molecule’s motion in two dimensions over the anisotropic surface potential is of great importance in understanding chemical reactions on surfaces. Of particular interest is the study of molecular motion and binding with respect to lattice orientation, step edges, and surface defects. A model system for the investigation of substrate-adsorbate interactions is the adsorption of benzene on metal surfaces, due to benzene’s relative structural simplicity, high symmetry, and the high contrast with which it appears in STM images. Jackiw et al. have observed that benzene molecules tend to diffuse along the closepacked rows of the [101] direction for Cu{110}, and along the [001] direction for

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Ag{110} [43, 54]. In these experiments, benzene was deposited on the Cu or Ag surface at 4 K, where it adsorbed randomly on the terraces. Repeated cycles of Ar+ sputtering and annealing were used to clean the surface before exposure to the benzene molecules. STM measurements were carried out in a low temperature, ultrahigh-vacuum STM [55] with the capability of raising the temperature in a controlled manner up to 80 K [56]. Instead of calculating the height of a molecule as in the previous case, here the location of the molecule in each extracted image is calculated. This is accomplished by selecting the area in which a molecule diffuses during the STM movie [white box in Figs. 10(A) and 10(B), enlarged in Fig. 10(D)]. The object to be tracked is then selected [white box in Fig. 10(D), also shown in Fig. 10(E)]. The extracted images [Figs. 10(D) and 10(E)] are then cross-correlated to generate the XC image, Fig. 10(F). In this case, the extracted images are relatively small and the search window the full image size of Fig. 10(D) is larger than the adsorbate image Fig. 10(E), therefore it is more advantageous to calculate the diffusion track with the real space XC similar to Eq. (1). For the diffusion track XC we used the built-in Matlab XC function (xcorr2) that returns the normalized cross-correlation with an extended search window as described in Ref. 44. The maximum of Fig. 10(F) corresponds to the adsorbate location and can be tracked to monitor the diffusion of the molecule Fig. 10(G), henceforth referred to as a diffusion track. Because the adsorbate is user selected, this method can be used to track protrusions or depressions of any shape. However, the diffusion tracking, just like the drift tracking algorithm, uses the maximum of the XC image for detecting motion; in the event that another molecule diffuses into the extracted area the maxima may correspond to the new molecule rather than the desired molecule, leading to a false diffusion track.

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Figure 10. Diffusion analysis of benzene on Ag{110}. (A) and (B) are frames 1 and 19 (280 x 280 Å ), respectively, from a 64 frame movie. The white protrusions are individual benzene molecules adsorbed on Ag{110}. The white boxes represent the user selected region for a molecule. Imaging conditions: Vsample = 0.100 V, I = 0.1 nA. (C) The extracted images of the drift area, 64 frames acquired over 52 min. (D) Frame 1 of the extracted images; the white box illustrates the region extracted for (E). (F) The real space XC of (D) and (E). (G) The location track of the benzene molecule calculated from the maxima of the real space XC (F) of each extracted image.

Because this analysis involves the detection of small motions and precise positions, all drift must be completely eliminated. While this algorithm can track general diffusion as seen in Fig. 8, it does not provide the precision to eliminate all of the drift completely. When the drift track is used on a stationary object such as a step edge or an

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immobile adsorbate, the diffusion track of the object should exhibit no movement, however, deviations of 5µm) but also close to it ( ȟ) the value of the function g(R) no longer scales with R and saturates at 2ı2. The value of the parameter ȟ may be attributed to the linear dimension of islands. The roughness dependence on the length scale was investigated on the basis of the AFM topography measurements. The microscopic images, recorded with the resolution 512x512 pixels, were divided consecutively by 4, 16, 64, 256, 1024 and 4096 nonoverlapping square tiles of decreasing linear dimension R, covering the whole scanned area and σ was computed for each tile. The mean of sigma over all tiles with the same dimension R gives a final value for s(R) for a given R. This procedure ensures that the smaller images are rightly slope corrected using the average plane found for the large original image. As mentioned earlier, the growth mode was determined by the type of substrate. Au layers deposited on a glass substrate are polycrystalline. Typical 3D growth occurs and as a consequence the surface is very rough –parameter σ ranging from 4 to 6.5 nm for RT and annealing at 300 oC, respectively and ξ– from 87 to 101 nm for the same temperatures. Consecutive annealing at temperatures t in the range from 170 ºC to 350 ºC shows that the roughness has a tendency to increase, similarly to the trend shown by the correlation length, leading to the wider and higher islands of the deposited material. The quality of Au layer morphology is substantially improved by using the sapphire substrate buffered with Mo. The deposition att RT results in 2D-3D growth (StranskiKrastanov mode), typical of (111) plane with the roughness of 0.46 nm. Atomically flat islands, ca 80 nm in diameter, with six-fold symmetry are observed (see Fig. 1). The fluctuation of the surface height is of the order of double (111) plane spacing. Contrary to Au growth on glass, the vacuum annealing of an Au layer deposited on a Mo buffer improves significantly its morphology. The temperature of 170 ºC is high enough to reduce the roughness more than twice (fluctuation of the surface height is suppressed to

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single monolayer step height, sigma = 0.17 nm) and increase the diameter of atomically flat islands by the same factor from 64 to 125 nm. The annealing at higher temperatures up to 600 ºC does not affect the topography markedly. It is worth to mention that even small miscut of a sapphire wafer might change the growth of Au bottom layer from the island-like to terrace-like mode.

Figure1. The AFM image of Au 200 Å thick layer as deposited on Mo buffer (left) and after annealing at 200 ºC (right). The range of the grey scale is 2 nm.

Figure2. The relation between the correlation function g(R) and the scan size R.. Squares denote the data for Au 200 Å thick layer as-deposited on Mo buffer and triangles represent data for the same Au layer but after 30 minutes annealing at 200oC.

In Fig. 2 the relation between the correlation function g(R) and a scan size R is presented. For both Au samples (as deposited on Mo buffer and after annealing at

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200oC) g(R) first increases with the R dimension and then saturates for larger R. The initial slope is related to the Hurst dimension H, the position of the slope change gives the correlation length (islands dimension) and the curve saturation defines rms roughness σ = (g(R)/2)1/2. The evaporation of Co layers was performed at RT on the Au buffer layer, previously annealed at 200 ºC. The surface of as-deposited Co layer has the island-like structure, similar to that for annealed Au layer (see Fig. 3). Additionally, a weak structure of the islands surface is visible in the AFM image. The Co surface roughness (sigma = 0.18 nm) is comparable to that of the annealed Au layers, but the correlation length takes different values as Co layer thickness changes.

Figure 3. The AFM image of Co 15 Å thick layer as deposited on Au layer (left) and after annealing at 250 ºC(right). The range of the grey scale is 4 nm.

For 8 Å thick Co film the parameter ȟ oscillates around 90 nm, whereas for a thicker layer (15 Å) it increases up to 170 nm. Annealing at 250 ºC for 45 min causes smearing of the surface island structure. Their contours, although still visible in AFM image, are much less pronounced. The annealing process does not affect markedly the roughness of the thin Co layer whereas for the thicker one results in the surface smoothening as reported in the Ref [4]. The RHEED pattern observation allows to investigate the existence of strains due to the lattice mismatch of both components of the sandwich. The surface lattice parameter of as-deposited 8 Å Co layer is substantially enhanced up to the value of 2.72 Å, in comparison to 2.51 Å for the bulk. In 15 Å thick Co layer the strain relaxation is stronger, revealing the value of 2.59 Å. Annealing of the Co layer performed at 250 ºC evidently affects its crystalline structure. For 8 Å thick Co film the splitting of the RHEED streaks occurs (see Fig. 4), being a proof of the lattice constant relaxation to the values of 2.54 Å and 2.84 Å for Co and Au, respectively. Surprisingly, such splitting is not observed for 15 Å Co layers. For this sample the lattice parameter after annealing, measured by the distance between the RHEED streaks, is equal to 2.83 Å - very close to the Au bulk value (2.88 Å).

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Figure 4. The 12 kV RHEED streaks observed on as deposited (left) and after annealing at 250 ºC (right) 8 Å thick Co layer. Before annealing streaks are blurred, but single and corresponds to the lattice constant 2.72 Å. After annealing they are split revealing the lattice parameter 2.59 Å and 2.83 Å for outer and inner, respectively.

Figure 5. Auger spectra (energy 3 keV) for the sample: sapphire/Mo/Au/Co15 Å in the as-deposited state (top) and after annealing at 250 oC (bottom).

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On the basis of AES spectra comparison of both as deposited and annealed 15 Å Co layer it is evident that Au behaves as a surfactant. As can be clearly seen in Figure 5 the intensity of Auger signal of Au (around 220 eV) after annealing is more pronounced and at the same time a suppression of the respective intensity for Co is observed. This fact suggests that gold atoms are diffusing through the Co layer towards the surface during annealing and some of them are present at the cobalt/vacuum interface.Since Au and Co are mutually insoluble, the reverse diffusion may take place after the annealing. Most probably, due to the diffusion activated by annealing, a non-continuous Au film appears on the surface of thin Co layer, whereas in the case of the thick Co film, it is fully covered with Au. Moreover the coherence between Au and Co layers is lost, giving rise to the relaxation of strains. This is compatible with the evolution of Co surface morphology, monitored changes in the RHEED pattern and the Auger electron spectrum after the annealing of the samples. Thus it is reasonable to expect that the annealing lowers substantially magnetoelastic contribution to the magnetic properties of Co layers.

4. Conclusions The sapphire substrate buffered with Mo is ideal for growth of smooth Au/Co/Au sandwiches. The annealing above 170 ºC improves the flatness of Au surface atomically smooth areas of a few hundred nanometers in diameter occur. In the as deposited Co layers the expanding strains should result in the significant magnetoelastic contribution to magnetic properties. The thermal treatment releases strains and lattice parameter relaxes. The Au acts as a surfactant appearing on the top of Co layer after annealing at 250 ºC.

Acknowledgements This work was supported in part within European Community program ICA1-CT-200070018 (Centre of Excellence CELDIS).

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11.

12. 13. 14.

Murayama, A., Hyomi, K., Eickmann, J., and Falco, C.F., (1999) Magnetoresistance, micromagnetism, and domain-wall scattering in epitaxial hcp Co films, Phys. Rev. B 60, 15245-15250. Train, C., Mégy, R., and Chappert, C. (1999) Magnetic anisotropy and magneto-optical Kerr effect of a Pt/Co/Au(111) sandwich at low Pt thickness, J. Magn. Magn. Mat. 202, 321-326. Schneider, C.M., Bressler, P., Schuster, P., Kirschner, J., de Miguel, J.J., and Miranda, R. (1990) Curie temperature of ultrathin films of fcc-cobalt epitaxially grown on atomically flat Cu(100) surfaces, Phys. Rev. Lett. 64, 1059-1062. Huang, F., Kief, M.T., Mankey, G.J., and Willis, R.F. (1994) Magnetism in the few-monolayers limit: A surface magneto-optic Kerr-effect study of the magnetic behavior of ultrathin films of Co, Ni, and Co-Ni alloys on Cu(100) and Cu(111), Phys. Rev. B 49, 3962-3971. Zhang, R.and Willis, R.F. (2001) Thickness-Dependent Curie Temperatures of Ultrathin Magnetic Films: Effect of the Range of Spin-Spin Interactions, Phys. Rev. Lett. 86, 2665-2668. Palasantzas, G. and Krim, J. (1993) Effect of the form of the height-height correlation function on diffuse x-ray scattering from a self-affine surface, Phys. Rev. B 48, 2873-2877. Palasantzas, G. (1993) Roughness spectrum and surface width of self-affine fractal surfaces via the Kcorrelation model, Phys. Rev. B 48, 14472-14478.

CORRELATION BETWEEN THE SURFACE RELIEF AND INTERFACES STRUCTURE OF Fe/Cr SUPERLATTICES AND ELECTROMAGNETIC WAVES PENETRATION

A. RINKEVICH, L. ROMASHEV, V. USTINOV Institute of Metal Physics Ural Division of RAS 18 S.Kovalevskaya St, Ekaterinburg 620219 Russia

1. Introduction Investigation of the physical properties of magnetic metallic multilayers is one of the subject matters in modern nanophysics. Essential attention to this topic is supported by the GMR effect. Very essential information about magnetic and magnetoresistive properties of metallic nanostructures can be extracted using microwaves. Two phenomena that influence the penetration of electromagnetic waves through a multilayer are of especial interest. They are the giant magnetoresistive effect and the ferromagnetic resonance (FMR) [1-3]. Microwaves of millimeter waveband provide a way to observe simultaneously the result of two above phenomena and to estimate the constants of exchange interaction between neighboring ferromagnetic layers. Currently the effect of the structure of interfaces between the layers on the electromagnetic penetration is carefully investigated [4, 5]. Most interest is paid to the damping of the magnetic moment oscillations and to the measuring of the magnetic moments of the interfaces. Direct penetration of electromagnetic waves through the multilayers is used here in comparison to the tunneling microscopy data. The Fe/Cr superlattices, sandwiches and ultrathin Fe films are the objects of the present investigation.

2. Experimental The penetration coefficient of electromagnetic waves from the frequency interval of 25 to 38 GHz was studied experimentally. The scheme of measurements is shown in Fig.1. The sample is positioned into the cross section of a rectangular waveguide and the module of transmission coefficient is measured. The objects of investigation were the Fe/Cr superlattices with different thickness of layers, sandwiches and Fe films covered by thin chromium layer. All the samples were grown by the molecular beam epitaxy method on the MgO and Al2O3 substrates. The superlattices had the thickness of Fe layers of 2 to 28 Å and Cr layers of 9 to 18 Å. The number of bilayers varied from 6 to 50. The external magnetic field up to 16 kOe was applied in the plane of the sample always perpendicularly to the k wavevector. Two variants of the external field direction 443 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 443-448. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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were employed. The first one is shown in Fig.1. Direction of the dc magnetic field H vector is perpendicular to the microwave electric E~ vector. So, the H vector lies in the plane of microwave magnetic field H~. In the second variant the orientation H// H /E~ is realized.

Figure1. Position of a sample in the waveguide

3. Results and discussion The peculiarities in electromagnetic penetration were compared to the surface relief of multilayers and superlattices. The surface relief was obtained with the STM microscope SMM2000. The scans of thin Fe films and Fe/Cr superlattices were measured at the scanning fields from 15 x 15µm to 320 x 320 nm. The minimal Rq values (that is the square-root roughness) of the Fe/Cr superlattices are about 5 ǖ. The roughness of the hybrid-cluster nanostructures is barely higher than in superlattice samples. The typical tunnel scan is presented in Fig.2. The scan is relevant to the sandwich Al2O3/[Fe (30ǖ)/Cr (100ǖ) sample. The scanning field is 1.133 x 1.133 µm and the peak-to-peak difference in heights is 83 ǖ. In order to characterize the surface relief quantitatively the Hurst parameter is introduced [6, 7]. The square-root roughness Rq is measured at the profiles of different length L. The peak-to-peak difference in heights Rmax is measured in addition and Rmax / Rq ratio is calculated. The results are plotted in double logarithmic scale. For the several classes of surfaces the graph is a straight line. The Hurst parameter H is calculated from the slope of this line. The value of the Hurst parameter H = -0.5 corresponds to a

445

random process with the spectrum of white noise. It is generally agreed that the Hurst parameter is a measure of steadiness of the process. This parameter is used in our paper in order to characterize the surface relief of nanostructures. The Hurst plot for the socalled hybrid-cluster Al2O3/Cr(70A)/[Fe(2A)/ɋr(10Ⱥ)]50 structure is shown in Fig.3 for two tunnel scans.

Figure 2. Tunnel scan for the Al2O3/[Fe (30ǖ)/Cr (100ǖ) sample

0,8

lo g (R max / R q )

0,7

0,6

0,5

0,4

0,3 1,0

1,5

2,0

2,5

3,0

3,5

4,0

log L (n m ) Figure 3. The Hurst plot for the Al2O3/Cr(70A)/[Fe(2A)/ɋr(10Ⱥ)]50 sample

446

Microwave measurements gave the following results. It was shown that the microwave GMR is very similar to that measured on dc. The relative variations of penetration coefficient rm=[D(H)-D(0)]/D(0) versus magnetic field are usually plotted, where D(H) is the penetration coefficientt module in the field H. The magnetic field dependence of the rm was measured at several frequencies of millimeter-waveband for the [Cr(19Å)/Fe(23Å)]12/Cr(80Å)/MgO superlattice, see Fig.4. Two origins of microwave variations are evident. The first origin is due to the microwave GMR effect. Variations caused by this mechanism are negative, monotonic and non-resonant [1, 3]. The absolute values of the microwave variation due to this mechanism are almost equal to the dc GMR. So, the structure of Fe-Cr interfaces acts on the microwave penetration just in the same manner as on the dc magnetoresistance. Resonant-type variations of the penetration t coefficient are caused by the ferromagnetic resonance. They are observed only if the dc external magnetic field is parallel to the E~ vector, that is perpendicular to the plane of H~ vectors. So, the acoustic or uniform resonance mode manifests itself in the penetration coefficient. The structure of interfaces influences essentially the damping of the magnetic moments rotation.

2 ,5

rm , %

f, GHz 32 34 36

0 ,0 -2 ,5 -5 ,0 -7 ,5

- 10 ,0 - 12 ,5 - 15 ,0 - 17 ,5

0

2

4

6

8

10

12

14

Magnetic field H, kOe Figure 4. Magnetic field dependence of the relative variations of penetration coefficient

The frequency dependence of the FMR line width is shown in Fig.5 for the Fe/Cr superlattice [Cr(19Å)/Fe(23Å)]12/Cr(80Å)/MgO, the thin iron film

447

Cr(10Å)/Fe(573Å)/Al2O3, and the hybrid-cluster structure [Cr(11Å)/Fe(9Å)]40/Cr(85Å)/MgO. The definite correlation is observed between the width of the FMR line and the surface relief characteristics. In the hybrid-cluster nanostructures with very thin discontinuous Fe layers the FMR width is relatively large up to 0.7 kOe. It is safe to assume that the large width is connected with inhomogeneous demagnetizing field near the discontinuities of Fe layers. The Gilbert constant G of magnetic damping was calculated in different multilayers. In the [Cr(19Å)/Fe(23Å)]12/Cr(80Å)/MgO superlattice G = (1,1 ÷ 1,2)⋅108 s-1 and in the Cr(10Å)/Fe(573Å)/Al2O3 iron film G = (1,5 ÷ 1,7)⋅108 s-1.

0,7 [Cr(19A)/Fe(23A)]12 film Fe (573A) hybride-cluster structure [Cr(11A)/Fe(9A)]40

∆ H , kOe

0,6

0,5

0,4

0,3 26

28

30

32

34

36

38

Frequency f, GHz Figure 5. Frequency dependence of the FMR line width in multilayers

4. Conclusion The relative variation of the transmission coefficient through the Fe/Cr superlattices and thin Fe films has been studied. For the frequencies higher then 30 GHz a clear narrow line of the ferromagnetic resonance (FMR) is seen besides the non-resonant variations caused by the GMR effect. The roughness of the surface relief in the best superlattice samples and in the films is as low as 5 Å. The roughness of the hybrid-cluster structures with very thin Fe layers is a bit higher. The roughness of the interfaces influences the

448

monotonic part of microwave variation in the same manner as the dc magnetoresistance. A distinct correlation is observed between the width of the FMR line observed in the penetration coefficient and the thickness of Fe layers. In the hybride-cluster structures the Fe layers are discontinuous. So, the FMR line width is much larger because of inhomogeneous demagnetizing fields. The non-resonantt variation of the transmission coefficient is practically absent in the thin Fe film. The FMR line, however, in this sample is narrower and has larger amplitude then in superlattices. The work was partially supported by the RFBR and INTAS grants.

References 1.

2. 3. 4. 5. 6. 7.

Krebs, J.J., Lubitz, P., Chaiken, A., and Prinz, G.A. (1991) Magnetoresistance origin for nonresonant microwave absorption in antiferromagnetically coupled epitaxial Fe/Cr/Fe(001) sandwiches, J. Appl. Phys. 69, Pt. II, 4795-4797. Kuanr, B.K., Kuanr, A.V., Grunberg, P., and Nimtz, G. (1996) Swept-Frequency FMR on Fe/Cr Trilayer Ultrathin Films – Microwave Giant Magnetoresistance, Physics Letters, A 221, 245-252. Rinkevich, A.B., Romashev, L.N., and Ustinov, V.V. (2000) Radiofrequency Magnetoresistance of Fe/Cr Superlattices, JETP, 90, 834-841. Urban, R., Woltersdorf, G., and Heinrich, B. (2001) Gilbert Damping in Single and Multilayer Ultrathin Films: Role of Interfaces in Nonlocal Spin Dynamics, Phys. Rev. Letters, 87, 217204-1-4. Celinski, Z., Urquhart, K.B., Heinrich, B. (1997) Using ferromagnetic resonance to measure the magnetic moments of ultrathin films, JMMM, M 166, 6-26. Hurst, H.E., Black, R.P., and Simaika, Y.M. (1965) Long Term Storage: An Experimental Study. L.: Constable. Schroeder, M. (1991) Fractals, Chaos, Power Laws. Miniatures from an Infinite Paradise. NY: W.H.Freeman & Company.

MAGNETORESISTANCE AND MICROSTRUCTURE OF MAGNETIC THIN FILM MULTILAYERS

J. NEAMTU∗, M. VOLMER∗∗ ∗Advanced Research &Development Institute for Electrical Engineering, SplaiulUnirii 313, Bucharest 030138,Romania [email protected] ∗∗∗Transilvania University Brasov, B-dul Eroilor 29, Brasov Romania

Abstract The magnetic properties and the magnetoresistance in correlation with microstructural properties of [NiFe(t)/Cu(s)/NiFe(t)]n and [NiFe(t)/Mo(s)/NiFe(t)] multilayers have been investigated. The thickness (t) of permalloy (Ni 80%Fe 20%) layers was ranged from 4 to 12 nm, while the copper and molybdenum layers (s) was ranged from 3 to 8 nm. The multilayers exhibit magnetoresistive properties correlated with microstructure and roughness at the interface of permalloy film and cooper or molybdenum layer. By decreasing of the NiFe layer thickness and by increasing of the non-magnetic interlayer thickness, the influence of interfacial intermixing effects on magnetic properties become more important. Although the thickness of layers has the leading part for magnitude of Giant Magnetoresistance effect, the microstructural properties of interfaces and the grain boundaries scattering must not be neglected. 1. Introduction The magneto-transport properties of ferromagnetic/nonmagnetic/ferromagnetic multilayers are dependent of the thickness of thin films, the roughness and the nature of thin film. Giant Magnetoresistance (GMR) effect results from two factors: (1) spin dependence of the electronic band structure of a defect-free system, and (2) spin dependence of scattering potential [1]. The aim of this paper is to investigate the influence of nature and microstructure of spacer-layer and influence of interface roughness on the magnetoresistance properties of the multilayers [NiFe(t)/Cu(s)/NiFe(t)]n and [NiFe(t)/Mo(s)/NiFe(t)]. 2

2. Experimental Four types of samples are considered in this paper: 1) Si/SiO2/ (Permalloy Ni 80%Fe20%) monolayer films 2) Si/SiO2/Py (10 nm)/Cu (4 nm)/Py(10 nm) in which Py is Ni80Fe20. 3) Si/SiO /Py(10 nm)/Mo(6 nm)/Py(10 nm) 449 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 449-456. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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4) Si/SiO2/[Py(10 nm)/Cu (4 nm)]9/Py(10 nm) Permalloy Ni80%Fe20% monolayer films were deposited, with thickness in range 4100 nm, using high vacuum evaporation with a base pressure of 10-7Torr, on Si/SiO2 substrates. During deposition the magnetic field of 15 kA/m was applied in the plane of the substrates in order to induce an easy magnetization axis in the films. The multilayers were prepared by R.F. sputtering at a base pressure of 10-7 Torr and an argon pressure of 1.5 mTorr (target to substrate distance is 100 mm). As substrates we used Si (100) single-crystal wafers, cut to a size of 5x10 mm2, with a thickness of 0.5 mm. Prior to insertion into the sputtering machine, the substrates were chemically etched using a 2% HF solution to have a flat surface and then oxidized. The permalloy layer thickness (t) was changed from 4 to 12 nm, while the copper and molybdenum layers thickness (s) was changed from 3 to 8 nm. The number of layers n for multilayer [Ni80Fe20 (10nm)/Cu (4nm)/Ni80Fe20 (10nm)]n was up to 10. 3. Results The magnetization measurements were performed att room temperature using a vibrating sample magnetometer (VSM). The magnetoresistance effect measurements were performed at room temperature in four-point contactt geometry with the contacts in line, using a DC current of 10 mA. The magnetoresistance measurements of permalloy (Ni80Fe20) films were made for two configurations: a) magnetic field applied parallel to the current direction and b) magnetic field perpendicular to the current direction. The magnetoresistance (MR) is defined as the variation ∆R=(R0-RH) of the resistance due to magnetic field normalized by the resistance R0 at zero magnetic field: MR=∆R/R0. The magnetic properties of Si/SiO2/Py(10 nm)/Cu (4 nm)/Py (10 nm) trilayer has presented in figure 1.The magnetic field is applied in the film plane, directed along the easy axis. 0.008

Py (10 nm)/Cu (4 nm)/Py (10 nm)

0.006

Moment (e.m.u.)

0.004 HC=250 Oe M r=0.0022 u.e.m.

0.002 0.000 -0.002 -0.004 -0.006

H in plan

-0.008 -4000

-2000

0

2000

4000

H (Oe) Figure 1. Magnetization curve measured at room temperature for NiFe (10 nm)/Cu (4 nm)/ NiFe (10 nm) trilayer with magnetic field applied in the film plane directed along the easy axis.

451

S /Py (10 nm)/Mo (6 nm)/Py (10 nm) 2 0.015 Si/SiO

Moment (e.m.u.)

0.010

HC=200 Oe Mr=0.003 u.e.m.

0.005 0.000 -0.005

0 .0 1 5 0 .0 1 0 0 .0 0 5

-0.010

0 .0 0 0 -0 .0 0 5 -0 .0 1 0

-0.015

-0 .0 1 5 - 1 5 0 0 0- 1 0 0 0 0- 5 0 0 0

-4000

-2000

0 H (Oe)

0

5 0 00 1 0 00 01 5 00 0

2000

4000

Figure 2. Magnetization curves of NiFe(10 nm)/Mo(6 nm)/NiFe(10 nm) trilayer, at medium and high magnetic fields.

Figure 2 shows the magnetization curve for Py (10 nm)/Mo (6 nm)/Py (10 nm) multilayer with the magnetic field directed along the easy axis. The insert figure shows the behavior of Py (10 nm)/Mo (6 nm)/Py (10 nm) multilayer at high magnetic fields. One can see descending slope of the saturation magnetization for high magnetic fields, due to the diamagnetic contribution of Si/SiO2 substrate. 0.06

Si/SiO2/[Py(10 nm)/Cu(4 nm)]*9 - Py (10 nm)

Moment (e.m.u.)

0.04 0.02

Mr=0.021 u.e.m. HC=202 Oe

0.00 -0.02 -0.04 -0.06 -3000

-2000

-1000

0

1000

2000

3000

H (Oe)

Figure 3 Magnetization curve of [NiFe(10 nm)/Cu(4 nm)]9/NiFe(10 nm) multilayer, measured at room temperature.

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Figure 3 shows magnetization curve of [Py (10 nm)/Cu (4 nm)]9/Py (10 nm) multilayer, with magnetic field directed along the easy axis. We observe the increase of saturation magnetization, comparison with Py(10 nm)/Cu (4 nm)/ Py (10 nm) trilayer.

Flux density(T) Figure 4. Longitudinal and transversal resistance, (RB-R0)/R0, versus flux density, for NiFe film of 100 nm.

Figure 4 shows the longitudinal and transversal magnetoresistance for Ni80Fe20 monolayer of 100 nm thickness. Transversal magnetoresistance has relative high value for small field (B111> substrate, while a folding of the film is found on SAMs of 4-methylbenzylthiol on gold substrate.

5. Acknowledgments We thank Prof. Flavio Gatti of the Physics Department of University of Genoa, for support in substrate preparation and Prof. Sergio Thea, Dr. Chiara Natale of the Chemistry Department of University of Genoa, for the synthesis of gold particles.

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Work supported by the Italian Ministry t of University and Scientific and Technological Research (National Research Program m Physical properties and interfacing of single-electron devices for quantum computing”) and by the national Research Council (Grant “Nanoarchitetture organiche per l’elettronica e paradigmi computazionali”).

References 1.

Roychowdhury, V.P. and Bandyopadhyay, S. (1996) IEEE Trans.Electr.Devices, 43, 1688-1699.

2.

Osifchin, R.G., Andres, R.P., Henderson, J.I., K Kubiak, C.P., and Dominey, R.N. (1996) Synthesis of nanoscale arrays of coupled metal dots, Nanotechnology 7, 412-416.

3.

Brust, M., Walker, M., Bethell, D., Schiffrin, D.J., and Whyman, R. (1994) Synthesis of thiol-derivatised gold nanoparticles in a two-phase liquid-liquid system, J. Chem. Soc., Chem. Commun. 7, 801-802.

4.

Whetten, R.L., Khoury, J.T., Alvarez, M.M., Murthy, S., Vezmar, I., Wang, Z.L., Stepphens, P.W., r gold molecules, Adv. Mater. 8, Cleveland, C.L., Luedtke, W.D., and Landman, U. (1996) Nanocrystal 428-433.

5.

Schaaff, T.G., Shafigullin, M.N., Khoury, J.T., Vezmar, I., Whetten, R.L., Cullen, W.G., First, P.N., Gutiérrez-Wing, C., Ascensio, J., and Jose-Yacamán, M. J., (1997) Isolation of smaller nanocrystal Au molecules: robust quantum effects in optical spectra, J. Phys. Chem. B 101, 7885-7891.

6.

Heath, J.R., Knobler, C.M. and Leff, D.V. (1997) Pressure /temperature phase diagrams and superlattices of organically functionalized metal nanocrystal monolayers: the influence of particle size, size distribution and surface passivant, J. Phys. Chem. B 101, 189-197.

7.

Ingram, R.S., Hostetler, M.J., Murray, R.W., Schaaff, T.G., Khoury, J.T., Whetten, R.L., Bigioni, T.P., Guthrie, D.K., and First, P.N., (1997) 28kDa alkanethiolate-protected Au clusters give analogous solution electrochemistry and STM Coulomb staircases, J. Am. Chem. Soc. 119, 9279-9280.

8.

Ohara, P.C., Leff, D.V., Heath, J.R. and Gelbart, W.M. (1995) Crystallization of opals from polydisperse nanoparticles, Phys. Rev. Lett. 75, 3466-3469.

9.

Burghard, M., Philipp, G., Roth, S., Von Klitzing, K., Pugin, R., and Schmid, G. (1998) Multilayered Langmuir –Blodgett films of thiol substituted ultrasmall gold clusters, Adv. Mater. 10, 842-845.

10. Chen, X.Y., Li, J.R., and Jiang, L. (2000) Two-dimensional arrangement of octadecylaminefunctionalized gold nanoparticles using the LB technique, Nanotechnology 11, 108-111. 11. Hostetler, M.J., Wingate, J.E., Zhong, C-J., Harris, J.E., Vachet, R.W., Clark, M.R., Londono, J.D., Green, S J., Stokes, J.J., Wignall, G.D., Glish, G.L., Porter, M.D., Evans, N.D., and Murray, R.W. (1998) Alkanethiolate gold clusters molecukles with core diameters from m 1.5 to 5.2nm: core and monolayer properties as a function of core size, Langmuir 14, 17-30. 12. Langmuir, I. and Schaefer V.J. (1938) J. Am. Chem. Soc. 60, 1351. 13. Harrell, L.E., Bigioni, T.P., Cullen, W.G., Whetten, R.L., First, P.N. (1999) Scanning tunneling microscopy of passivated Au nanocrystals immobilized on Au(111) surfaces, J. Vac. Sci. Technol B 17, 2411-2416.

AFM OF GUANINE ADSORBED ON HOPG UNDER ELECTROCHEMICAL CONTROL A.-M. CHIORCEA, A.M. OLIVEIRA BRETT l de Ciências e Tecnologia, Departamento de Química, Faculdade Universidade de Coimbra, 3004-535 Coimbra, Portugal

Abstract The characterisation of the adsorption mechanism of guanine on a highly oriented pyrolytic graphite (HOPG) electrode surface was carried out using in situ MAC Mode Atomic Force Microscopy (MAC Mode AFM) and the electrochemical t behaviour of the guanine layer was investigated with Electrochemical AFM. Guanine adsorbs spontaneously on the HOPG substrate as a stable molecular layer, covering the surface uniformly and almost completely. The adsorption of DNA at the HOPG surface was controlled by adjusting the potential of the HOPG electrode and electrochemical adsorption provides better attachment of the guanine at the electrode surface compared to natural adsorption. The characteristics of the adsorbed guanine films were dependent on the deposition time and on the electrochemical adsorption procedure. The film was dissolved by carrying out cyclic voltammetry between 0 and + 1.3 V, after which guanine started to readsorb freely on the clean HOPG surface. The guanine molecules were held together on the substrate mainly by non-covalent interactions such as hydrogen bonding, van der Waals and hydrophobic interactions.

1. Introduction Guanine is one of the constituent purine bases of nucleic acids. Together with the other purine, adenine, and the pyrimidine bases, thymine and cytosine, it forms the units of the DNA which transport genetic information. The structure of guanine, Fig. 1, assembled at solid surfaces has been examined using a variety of approaches [1-3]. Electrochemical studies showed that guanine oxidises irreversibly at the C8-H position by a two-step mechanism involving the loss of 4H+ and 4e- and leading to 8-oxoguanine, one of the oxidation products, which is also electroactive [4]. Guanine is more easily oxidised than the other DNA bases adenine, thymine and cytosine, as it has a lower oxidation potential for the same experimental conditions [5]. Electrochemical studies with guanine at concentrations near saturation led to the formation of dimers and trimers within the oxidation products, but the oligomers are difficult to detect due to the low solubility of guanine [6, 7]. Guanine adsorbates on solid surfaces have been studied at the molecular level using AFM and STM [8-10], but problems were encountered due to guanine’s weak interaction with the substrate. However, Magnetic AC mode AFM (MAC Mode AFM) is a technique that permits the visualisation of the molecules that are weakly bonded to the substrate material [11]. 467 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 467-473. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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O HN 1

6

7

4

9

N 8

2 3

H2N

5

N

N H

Figure 1. The chemical structure of guanine

MAC mode uses a solenoid placed under the sample to make a magnetically coated AFM cantilever oscillate near its resonant frequency. As it scans the sample, the AFM tip oscillates and touches the sample surface only at the bottom of this oscillation. Because there is no need to drive the cantilever holder, cantilever chip and solution as in tapping mode, control of cantilever movement increases considerably, which enables operation at smaller oscillation amplitudes, the lateral forces being better eliminated. This paper presents results from in situ MAC Mode AFM imaging, under electrochemically controlled conditions, of adsorbed guanine molecules at the highly oriented pyrolytic graphite (HOPG) electrode.

2. Experimental Materials. Guanine was purchased from Sigma Chemical Co. and was used without further purification. The supporting electrolyte used was pH 4.5 0.2 M acetate buffer solution and was prepared using analytical grade reagents and purified water from a Millipore Milli-Q system (conductivity < 0.1 µS cm-1). Guanine was dissolved directly in buffer solution, and a concentration of 10-3 M in the supernatant was obtained in saturated solutions [12]. Highly oriented pyrolytic graphite (HOPG), grade ZYH, of rectangular shape with 15 x 15 x 2 mm dimensions, from Advanced Ceramics Co., UK, was used as substrate. The HOPG was freshly cleaved with adhesive tape prior to each experiment and was imaged by Contact Mode AFM in order to establish its cleanliness. Apparatus. AFM was performed with a Pico SPM controlled by a MAC Mode module and interfaced with a PicoScan controller from Molecular Imaging Co., USA. All the experiments were performed with a CS AFM S scanner with the scan range 6 µm in x-y and 2 µm in z directions. Electrochemical control was carried out with a potentiostat/galvanostat PicoStatTM. Silicon type II MAClevers 225 µm length, 2.8 N m-1 spring constant and 27-30 kHz resonant frequencies in liquid (Molecular Imaging Co.) were used. All images (256 samples/line x 256 lines) were taken at room temperature, at a scan rate of 1.95 lines s-1. The images were processed by flattening in order to remove the background slope and the contrast and brightness were adjusted. Both Atomic Force Microscopy and voltammetric experiments were carried out in a one-compartment Teflon cell of approximately 12.5 mm internal diameter holding the HOPG sample – the working electrode – on the base. A Pt wire counter electrode and a Ag wire as quasi-reference electrode were placed in the cell, dipping approximately 5 mm into the solution. All images were visualised three-dimensionally using the Scanning Probe Image Processor SPIP, version 2.3011, Image Metrology ApS. Origin (version 6.0) from Microcal Software was used for the presentation of the experimental voltammogram reported in this work.

469

Sample preparation. For the guanine samples prepared by free adsorption 500 µl solution of guanine were placed in the AFM cell and the guanine was left to adsorb at the HOPG surface over periods from 5 min to 1h. For the guanine samples prepared under electrochemically controlled conditions the following procedure was carried out. The freshly cleaved HOPG was first examined by cyclic voltammetry in pH 4.5 0.2 M acetate buffer solution, with potential window from 0 to + 1.3 V (vs. Ag wire), in order to establish its cleanness. Then 5 successive cyclic voltammograms were registered in saturated guanine solution between 0 and + 1.3 V (vs. Ag wire), scan rate 0.1 V s-1, followed by the application of a positive potential of +0.75 V (vs. Ag wire) to the HOPG electrode during 5 min. The HOPG with guanine adsorbed by both immobilisation procedures was immediately imaged by in situ MAC Mode AFM.

3. Results and Discussion 3.1 FREE ADSORPTION OF GUANINE WITH THE ELECTROCHEMICAL CELL AT OPEN CIRCUIT The HOPG electrode was modified by a ffilm of guanine obtained by free adsorption from a saturated guanine solution at different adsorption times using the method described in the experimental section. The results, in 10-3 M saturated guanine in pH 4.5 0.2 M acetate buffer solution, indicate that the guanine molecules adsorb spontaneously on the HOPG surface and the AFM images obtained in situ reveal a good surface coverage. The adsorption of guanine molecules on the HOPG surface occurs very fast. After introducing the guanine solution into the electrochemical cell, the molecules immediately start to condense in small nuclei that appear as bright spots in the images, without forming a well packed structure. Afterr 5 min of free adsorption the molecules appeared condensed in small nuclei of 2–3 nm height and 20 to 40 nm diameter, covering the HOPG surface uniformly, Fig. 2A. Considering the dimensions of a guanine molecule these aggregates correspond to several hundred guanine molecules lying flat on the surface [9]. The effect of altering the time of exposure of the HOPG surface to guanine solution was investigated using periods from 5 min to 1 h and the guanine layer appears to reorganise over time. After an exposure of 1 h the MAC Mode AFM shows that the electrode surface is completely covered by a very thick film of guanine, Fig. 2B. The topography of the film shows nuclei of different sizes from 3 to 6 nm height and 20 to 40 nm diameter that are organised in clusters of approximately 100 nm diameter. The molecules are stabilised on the HOPG surface by hydrophobic interactions between the hydrophobic aromatic rings of the guanine molecules and the hydrophobic carbon surface. Performing one cyclic voltammogram from 0 to + 1.3 V (vs. Ag wire), scan rate 0.05 V s-1, the guanine layer starts to dissolve byy forming small pits inside the guanine layer with dimensions of approximately 30 nm diameter and 1−3 nm deep, Fig. 2C. After 5 min of successive cyclic voltammograms the film dissolves completely, Fig. 2D, showing an almost clean HOPG surface.

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Figure 2. In situ MAC Mode AFM topographical images of the guanine layer, in a solution of 10-3 M saturated guanine in pH 4.5 0.2 M acetate buffer, prepared by free adsorption onto HOPG. (A) 5 min free adsorption and (B) 1 h free adsorption. (C) After one cyclic voltammogram from 0 to + 1.3 V (vs. Ag wire), scan rate 0.05 V s-1, the guanine layer starts to dissolve. (D) The film is completely removed after 5 min of successive cyclic voltammograms.

3.2 ELECTROCHEMICAL DEPOSITION - THE INFLUENCE OF THE HOPG POTENTIAL IN THE PROCESS OF GUANINE ADSORPTION The electrochemical potential applied to the HOPG substrate influences the processes of nucleation and growth of the guanine adsorbates on the HOPG surface. The surface characteristics of the electrode modified by guanine according to the electrochemical procedure described in the experimental section was investigated. Guanine adsorbs freely on the surface at the moment of injection of the solution in the AFM cell. In order to induce desorption of the adsorbed guanine and clean the surface electrochemically, 5 cyclic voltammograms were performed in the saturated guanine solution between 0 and + 1.3 V (vs. Ag wire), scan rate 0.1 V s-1. Subsequently, the HOPG was held for 5 min at the positive potential of +0.75 V (vs. Ag wire), which corresponds to the oxidation potential of guanine. The guanine molecules condensed into larger nuclei of 90–150 nm diameter and 10–30 nm height, Fig. 3. The nuclei appear grouped together in intercalated polymer-like chains of many different lengths, some longer than 1 µm, uniformly distributed over the HOPG surface, Fig. 4A.

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Figure 3. In situ MAC Mode AFM topographical images of the guanine layer, in the solution of 10-3 M saturated guanine in pH 4.5 0.2 M acetate buffer, on HOPG.

Figure 4. Sequence of 3 consecutive MAC Mode AFM topographical images obtained in situ in 10-3 M saturated guanine in pH 4.50 0.2 M buffer acetate solution and showing the dissolution of the guanine film. The images were taken (A) before, (B) during and (F) after performing one cyclic voltammogram that induced desorption of the guanine layer. The cyclic voltammogram from 0 to + 1.3 V (vs. Ag wire), scan rate 0.1 V s-1, was performed inside the AFM cell during scanning (B).

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Guanine oxidation can induce the formation of dimers and trimers at the HOPG electrode [6, 7]. The guanine molecules are stacked at the surface together with guanine oligomers, forming a complex polymer chain. All the components interact between themselves and with the HOPG surface by hydrogen bonding, van der Waals forces, and hydrophobic interactions. The stability of the guanine film obtained during controlled oxidation was very much enhanced when compared with spontaneous condensation, due to the electrostatic interaction of guanine rings with the positively charged HOPG surface. The cyclic voltammogram in the saturated guanine solution, Fig. 4B, shows a peak at + 0.76 V (vs. Ag wire) which corresponds to the oxidation reaction of guanine. During the cyclic voltammetric scan it was observed that the guanine layer was stable between 0 and + 0.75 V (vs. Ag wire). Increasing the potential led to abrupt changes in the guanine film. The products of oxidation of guanine are 8-oxoguanine and guaninedimers [6]. The 8-oxoguanine is oxidised at lower potentials and the oxidation products are hydrolysed and go to the solution. At approximately + 0.9 V all polymer-like chains completely disappeared, only the strongly adsorbed guanine-dimers remaining adsorbed at the electrode surface, Fig. 4C. After the film was dissolved by carrying out cyclic voltammetry the guanine layer began to grow immediately again by spontaneously adsorption, the undissolved guanine-dimers serving as nucleation centres. After approximately 30 min the guanine layer was completely reformed, Fig. 5.

Figure 5. In situ MAC Mode AFM topographical images obtained in the solution of 10-3 M saturated guanine in pH 4.50 0.2 M acetate buffer showing reconstruction of the guanine layer after approximately 30 min free adsorption.

4. Conclusion The process of adsorption of guanine at the HOPG surface can be controlled by the applied potential and the electrochemical deposition method provides better attachment of the molecules at the HOPG surface compared to passive adsorption. The characteristics of the guanine layer and the apparent height of the film depend on the applied potential. The guanine is adsorbed onto the HOPG surface only by non-covalent interactions such as hydrogen bonding, electrostatic, van der Waals and hydrophobic interactions. MAC Mode AFM in an electrochemically controlled environment is capable of showing in situ the surface morphological structure of guanine adsorbates and may

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contribute to the understanding of the mechanism of adsorption and the nature of the guanine-electrode surface interaction.

Acknowledgements Financial support from Fundação para a Ciência e Tecnologia (FCT), Ph.D. grant PRAXIS XXI/ BD/ 19728/99 (A.-M. C.), POCTI (co-financed by the European Community Fund FEDER) ICEMS (Research Unit 103) and European Projects ERBICT15-CT960804 and QLK3-2000-01311 are gratefully acknowledged.

References 1.

Sowerby, S.J., Edelwirth, M. and Heckl, W.M. (1998) Self-assembly at the prebiotic solid-liquid interface: Structures of self-assembled monolayers of adenine and guanine bases formed on inorganic surfaces, J. Phys. Chem. B 102, 5914-5922. 2. Tao, N.J., Shi, Z. (1994) Monolayerr guanine and adenine on graphite in NaCl solution: A comparative STM and AFM study, J. Phys. Chem. 98, 1464-1471. 3. Oliveira-Brett, A.M., Brett, C.M., Silva, L.A. (2002) An impedance study of the adsorption of nucleic acid bases at glassy carbon electrodes, Bioelectrochemistry 56 (1-2), 33-35. 4. Brett, C.M.A., Oliveira Brett, A.M., Serrano, S.H.P. (1994) On the adsorption and electrochemical oxidation of DNA at glassy carbon electrodes, J. Electroanal. Chem. 366, 225-231. 5. Oliveira Brett, A.M., Serrano, S.H.P., Piedade, J.A.P. (1999) Electrochemistry of DNA, in: R.G. Compton, G. Hancock (eds.), Comprehensive Chemical Kinetics, vol. 37, Elsevier, Amsterdam, cap. 3, pp. 91-199. 6. Oliveira-Brett, A.M., Diculescu, V., Piedade, J.A. (2002) Electrochemical oxidation mechanism of guanine and adenine using a glassy carbon microelectrode, Bioelectrochemistry 55 (1-2), 61-62. 7. Subramanian, P., Dryhurst, G. (1987) Electrochemical oxidation of guanosine - formation of some novel guanine oligonucleosides, J. Electroanal. Chem. 224, 137-162. 8. Heckl, W.M., Smith, D.P.E., Binning, G., Klagges, H., Hansch, T.W. and Maddocks, J. (1991) Twodimensional ordering of the DNA base guanine observed by scanning tunneling microscopy, Proc. Natl. Acad. Sci. USA 88, 8003-8005. 9. Tao, N.J., DeRose, J.A., Lindsay, S.M. (1993) Self-assembly of molecular superstructures studied by in situ scanning tunneling microscopy: DNA bases on Au (111), J. Phys. Chem. 97, 910– 919. 10. Tao, N.J., Shi, Z. (1994) Potential induced changes in the electronic states of monolayer guanine on graphite in NaCl solution, Surf. Sci. Lett. 301, 217-223. 11. Han, W., Lindsay, S.M., Jing, T.W., (1996) A magnetically driven oscillating probe microscope for operation in liquids, Appl. Phys. Lett. 69, 4111-4113. 12. Oliveira Brett, A.M., Matysik, F.-M. (1997) Sonoelectrochemical studies of guanine and guanosine, Bioelectrochemistry and Bioenergetics 42, 111-116.

DYNAMICS IN MODEL MEMBRANES AND DNA-MEMBRANE COMPLEXES USING TEMPERATURE CONTROLLED ATOMIC FORCE MICROSCOPY

Z.V. LEONENKO, D.T. CRAMB Department of Chemistry, University of Calgary, 2500 University Drive NW, Calgary, AB T2N 1N4, Canada [email protected]

Abstract In this paper we review our recent work studying biomolecular self-assembly using temperature controlled atomic force microscopy. In particular, we examined supported planar bilayers (SPBs), DNA-SPBs complexes, and their transitions during heating the system above the melting transition temperature. 1. Introduction The invention of atomic force microscopy has proven to be invaluable to the study of so-called “soft” biological samples1. Single macromolecules, like DNA and proteins as well self – assembled structures, like lipid membrane can be directly visualized by AFM. Based on Van der Waal’s interactions between the sample and scanning probe, AFM has overcome the restrictions of STM for studying nonconductive samples, allowing the study biomolecules in their native liquid environment. In particular, the imaging of such “soft” samples has benefited from Tapping Mode or oscillating mode 24 , where the probe oscillates, driven acoustically or magnetically. The oscillating probe makes only intermittent contact with the sample, m and minimizes shear forces, which in contact mode AFM can cause sample damage or displacement. In this mode, changes in amplitude and phase shift of the oscillating probe can be measured. Monitoring of the changes in oscillation amplitude produces topographic information. The differences in phase between the input and output oscillations upon interaction with the sample surface can give information about differences in chemical (hydrophobic or electrostatic) and mechanical (elastic) properties of the sample surface at the nanometer scale.

2. Materials and Methods We investigated several different kinds of phospholipids: 1,2-Dioleoyl-sn-Glycero-3Phosphocholine (DOPC), 1,2-Dipalmitoyl-sn-Glycero-3-Phosphocholine (DPPC), 1,2Dioleoyl-sn-Glycero-3-Phosphoethanolamine (DOPE), 1,2-Dipalmitoyl-sn-Glycero-3475 P.M. Vilarinho et al. (eds. ), Scanning probe Microscopy: Characterization, Nanofabrication, and Device Application of Functional Materials, 475-483. © 2005 Kluwer Academic Publishers. Printed in the Netherlands.

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Phosphoethanolamine (DPPE), 1,2Dioleoyl-3-Trimethylammonium-Propane (DOTAP), 1,2- Dipalmitoyl-3-Trimethylammonium-Propane (DPTAP). DOTAP and DPTAP are positively charged lipids, DOPE, DPPE, DOPC are neutral zwitterionic lipids. Phospholipids DOPC, DPPC (lyophilized or in chloroform solution), DOPE, DPPE, DOTAP (chloroform solution), and DPTAP were purchased from Avanti Polar Lipids Inc., Alabaster, AL and were used without further purification. Tris-EDTA (TE), phosphate buffer and distilled, Nanopure water were used in the preparation of all vesicles and DNA solutions. Freshly cleaved ASTMV-2 quality, scratch-free ruby mica (Asheville-Schoonmaker Mica Co., Newport News, VA) was used as the solid support. The 14 base pair oligonucleotide (ODN) (ATATAAATTTATAT) was obtained, desalted, from UCDNA services (University off Calgary, AB). The longer 1000 base pair DNA was calf thymus (Sigma), type II, fragmented by ultra-sonication. In our work, we employed MAC (magnetic A/C) mode, where the magnetically coated probe oscillates near its resonant frequency driven by an alternating magnetic field. All images were taken using a Pico SPM microscope with an AFMS-165 scanner (Molecular Imaging Inc., Phoenix, Az.). Au-Cr coated Maclevers® (Molecular Imaging Inc., Phoenix, Az.) were used for MAC mode imaging. Their specifications are: length 85 µm, force constant 0.5 N/m and resonant frequency 38-40 kHz in water. The standard MAC mode fluid cell (Molecular Imaging) was used throughout. The scanning speed was 2-3 lines per second. The height scale was calibrated using colloidal gold spheres of well-defined size5. For elevated temperature experiments the AFM Temperature Controller and Hot Mac Mode Stage from Molecular Imaging Inc., were used. Temperature was varied from 22 to 70 °C, with a 1 °C per minute ramp. For our AFM experiments we formed supported planar bilayers on mica by method of vesicle fusion. When sonicated, lipid bilayers spontaneously form closed spherical structures, called vesicles or liposomes. We apply vesicle solution to the mica support and when vesicles adsorb to the mica they fuse and form a planar supported bilayer. In some cases mica was modified with APTES (2-aminopropyltriethoxysilane) and PLL (poly-L-lysine). We also used Mg2+ ions as binding bridges between DNA and regular mica. Supported planar bilayers were prepared for AFM imaging by method of vesicle fusion6-7. Aliquots of liposome solution were deposited on modified or unmodified freshly cleaved mica. After a controlled period of time the mica was gently rinsed with ultrapure water. Solutions of DNA (100-200 µL, 2 – 1000 µg/mL) were pipetted onto the wet phospholipid bilayer surface and were left to incubate at 4oC for 1- 24 hours. The excess DNA was gently rinsed away and the surface was imaged under water in the liquid cell, at room temperature t and at various higher temperatures.

3. Results and Discussions Supported planar phospholipid bilayers are widely used as a model for studying biological membranes. Supported planar bilayers (SPBs) are composed of phospholipids adsorbed to a planar hydrophilic solid support. In water environment most phospholipids spontaneously form bilayer structures t hiding hydrophobic tails inside and pointing hydrophilic heads outside the bilayer - water interface. Understanding the physical and chemical properties of SPBs is critical to our understanding of membrane structure-function relationships. Model membranes are also good candidates for

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nanotechnology applications in nanosensor development. They can serve as a template for the incorporation of proteins and receptors and reduce non-specific interactions. In the following sections, we present our results on supported phospholipid bilayers (SPB) and structural changes in SPB and SPB-DNA complexes during transition above melting transition temperature (Tm)8-10. 3.1. SUPPORTED PLANAR BILAYER AND VESICLES ON MICA. Using MAC-mode AFM, we examined the process of supported planar bilayer preparation on mica via vesicle fusion7, Figure 1 a) shows disk-like surface features associated with single DOPC vesicle deposition. In the Figure 1 b) a complete bilayer was formed. Interestingly, this bilayer developed defects that were found to be quite dynamic in nature and allowed us to measure the bilayer thickness.

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b

Figure 1. A three-dimensional rendering of the AFM topography image of pure DOPC vesicles on APTES modified mica surface in aqueous solution, a), DOPC supported planar bilayer, b).

3.2. INTRODUCTION TO DNA-CATIONIC LIPID INTERACTIONS. Positively charged or Cationic Lipids (CL) are promising as gene delivery vehicles. They can deliver gene material to the cell and can be targeted to a specific tissue. There are several advantages to using liposomes over other drug delivery systems, the major being nontoxicity11. The process of gene delivery includes lipoplex formation between DNA and cationic lipids. The mechanism by which genetic material is delivered to the nucleus is not entirely clear. Moreover, despite their promise as gene transfer reagents, the phase dependence of DNA-cationic lipid interactions has not been extensively studied at the single molecule level. The formation and stability of the lipoplexes depend on many factors; examples of which are the type of helper lipids12, and the solvent environment. Several recent studies have investigated phospholipid - DNA interactions in solution13-19. However, AFM can be a powerful tool to study DNA interaction and complex formation at the surface of a supported phospholipid bilayer. There have been only a few studies using AFM to examine DNA on supported cationic

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bilayers, using gel phase bilayers20-21. Using fluid phase bilayers, we hope to elucidate stages in the process of complexation at the single DNA level in a real time. Clearly, there is a considerable challenge in using AFM to study DNA on supported planar bilayers in the liquid phase because of fluidity of bilayer. Such bilayer systems are challenging for current AFM technology, since they are soft and dynamic. However, because of their ease of deformation, fluid bilayers must be very sensitive to the presence of DNA and may easily respond to perturbations caused by adsorbed DNA and serve as a good model system to study DNA-CL lipoplex formation. 3.3. DNA ADSORPTION ON MICA. Since DNA in TE buffer solution does not bind to regular mica (negative surface charge), the surface was modified with APTES or PLL (positive surface charge). We also used Mg2+ ions as binding bridges between DNA and regular mica22,23. A height of 2.0 ± 0.2 nm was measured for 14 base pair ODN on mica. Small DNA appear as flat ovoids, due to the convolution of the tip radius (~10–40 nm) onto the surface feature. This results in the width of the DNA being overestimated. The height measurements, however, are not limited by the tip radius and give a better representation of the true DNA diameter on a solid surface. Our data for large DNA adsorbed on solid mica surface give a height of 2.0 ± 0.1 nm. Without surface modification or use of Mg2+, no DNA was observed to bind. Figure 2 shows topography images of DNA adsorbed on mica surface imaged by MAC mode AFM in a liquid cell.

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b

Figure 2. 14 basepair oligodeoxynucleotides (double stranded) on APTES-modified mica, a). 3000 base pair 2+ linearized plasmid DNA bridged onto mica by Mg , b).

3.4. DNA ADSORPTION ON SUPPORTED PLANAR BILAYER. The same DNA were applied to the bilayer surface8 – Figure 3. Figure 3. (a) Solution of (2 mg/ml) 14 base pair DNA was exposed to a DOTAP/DOPE SPB for 30 minutes and then gently rinsed, the oligodeoxynucleotide was found to associate with the surface. The height of 14 base pair DNA on the supported planar bilayer is 1.3 ± 0.2 nm.

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a

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c

Figure 3. 14 basepair oligodeoxynuleotides adsorbed onto a DOTAP/DOPE supported bilayer, a). 1000 base pair linearized plasmid DNA adsorbed on a DOTAP/DOPE supported bilayer, b). 1000 base pair linearized plasmid DNA adsorbed onto a DOTAP/DOPC supported bilayer, c). Image scale 1500 nm.

Adsorption of 14 base pair DNA does not appear to disturb the bilayer. Only a few defects were observed after removal of the excess of DNA from the cell. This effect is comparable to rinsing the sample cell with buffer solution containing no DNA. (b) 1000 base pair DNA adsorption onto DOTAP/DOPE changed the bilayer dramatically. Large DNA strips bilayer from the surface changing the structure of bilayer, we observed aggregate formation, which were a double bilayer thick. When negatively charged DNA adsorbs to the bilayer it interacts preferentially with positively charged DOTAP lipids, possibly causing demixing of lipids in a fluid DOTAP/DOPE bilayer. As a result of this demixing process regions of pure DOTAP and DOPE are formed. Pure DOPE readily transforms to a hexagonal phase and causes wrapping of a bilayer around DNA forming big aggregates, which can be easily detached from the surface. In a control experiment, when DNA-free (TE) buffer solution was applied to supported bilayer, no bilayer removal was observed. (c) In order to assess the role off the helper lipid, DOPE on the above-mentioned behaviour, we examined the effectt of DNA adsorption to a binary system with a different zwitterionic helper lipid, DOPC. The addition of DNA to the supported DOTAP/DOPC bilayer shows adsorption of DNA at the bilayer surface and partially removes a bilayer from mica, but not as severely as it does for DOTAP/DOPE. No changes of bilayer thickness were observed. Statistical analysis of our data on the height of DNA adsorbed on a fluid phase SPB gives 0.8 - 1.5 nm8. Resolution and the height of DNA were also low, likely due to the partial penetration of DNA into the fluid phase bilayer. To elucidate the reason for low contrast and height of DNA we formed the complex of DNA on a gel phase mixed DPTAP/DPPE bilayer9. Upon heating the SPB above Tm, it will become analogous to DOTAP (or DOTAP/DOPE). In this case DNA can be clearly seen and the height of DNA is similar to that observed on solid support, Figure 4 (a).

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a

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c

Figure 4. Heating supported mixed DPTAP/DPPE bilayer with adsorbed DNA: a) room temperature, 22°C; b) heating to 50°C; c) cooling back to room temperature, 23°C. Image scale 850 nm.

Heating the complex of DNA on the DPTAP/DPPE bilayer shows that the contrast and resolution of DNA adsorbed on the bilayer were lost at elevated temperature (data not shown). After slow cooling, the contrast t and resolution of DNA were restored. In the fluid phase bilayer DNA can easily penetrate into the head group area of bilayer, which leads to low resolution and height. In a mixed DPTAP/DPPE bilayer, additional to the changes in height and resolution of DNA we observed structural changes in the bilayer9, Figure 4. DNA, initially adsorbed on the higher domains after bilayer melting DNA was found on the lower domains, while higher domains did not contain DNA. We assume that during bilayer melting the demixing of lipids occurs and DNA appear at the lower (DPTAP enriched domains), and (DPPE enriched) higher domains do not contain DNA. 3.5. PHASE TRANSITIONS IN SPB AND DNA-SPB COMPLEXES. An understanding of structure-function relations of biomembranes ultimately relies on knowledge of the properties of lipid bilayers. The complex structural dynamics of membranes is related to specific membrane functions. Such structural parameters as average interfacial area per lipid, bilayer thickness and disorder of hydrophobic tails are very important for understanding intermolecular interactions in membranes. All these membrane parameters change during phase transition. The phase behaviour is dominated by the main L - L (gel-fluid) phase transition. In gel phase, hydrophobic tails are more ordered and the mobility of lipids is reduced. In fluid phase hydrophobic tails are very flexible and disordered, which decreases the bilayer thickness. A gel phase bilayer can be easily transformed to the fluid phase by melting a bilayer above melting transition temperature. Below the main transition, a basic equilibrium structure is the gel - subgel (crystalline) Lc phase24. A large number of intermediate stable, metastable, and transient lamellar gel structures are adopted by different lipids24 – with perpendicular or tilted chains, with interdigitated or partially interdigitated chains. We observed several phase transitions in the DPPC gel phase bilayer upon heating and cooling back to room temperature, Figure 5. During heating of the DPPC bilayer, a broad main transition was observed at 42-50°C. Dynamic coexistence of 2 domains was observed during this temperature interval. Further increase in temperature leads to formation of lower domains with this additional transition complete at 60°C. The second transition is likely attributed to the formation of fluid disordered phase formation, possibly with

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interdigitated lipid chains. Slow cooling the system back restores bilayer reversibly to the initial thickness - Figure 5.

a

b

c

d

e

f

Figure 5. The AFM topography images showing phase transition in DPPC bilayer in liquid cell upon heating and cooling back. A heating DPPC bilayer: a) room temperature, 22°C; b) heating to 50°C; c) heating to 52°C, d) heating to 54°C, e) heating to 60°C, f) cooling back to 54°C. Image scale 1500 nm.

Phase transitions in a mixed DPTAP/DPPE bilayers are interesting in that we observed irreversible changes9. The supported bilayers formed from water solutions were planar, contained defects and elevated regions (domains) – Figure 6. The difference in the thickness of the two domains was 1.7 nm. The difference here is likely due to domains rich in DPPE versus domains rich in DPTAP. DPPE is known to have a thickness of 6 nm. Whereas an annealed DPTAP supported bilayer has been observed to have a thickness between 4 and 5 nm. The two lowest domains in Fig 6 could be domains in the enriched DPTAP bilayer similar to those observed by Longo and coworkers with difference between two domains being 1.4 nm10. We observe the difference in two lower domains being 1.2-1.4 nm and assume that they could appear during non-equilibrium bilayer formation by vesicle fusion, because the vesicle solution was preheated during sonication. After heating bilayer to 70°C we observe that domains and defects disappear and the surface was covered with a flat bilayer with only few small defects.

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c

Figure 6. Heating supported mixed DPTAP/DPPE bilayer, formed from water vesicle solution: a) room temperature, 22°C, b) heating to 70°C, c) cooling back to room temperature, 23°C. Image scale 800 nm.

After cooling the system back to room temperature, we observed a redistribution of the domains. We observed the formation of only two domains (5.5 nm and 4.0 nm) in addition to defect formation. This is suggestive again of the formation of DPPE rich and DPTAP rich bilayer areas, respectively. Various experimental and theoretical studies have suggested the existence of domains in both biological and model membranes25. Nonequilibrium domain formation, which offers local geometrical environment for membrane processes, may be of considerable importance for the functionality of biological membranes.

4. Conclusions Here we demonstrated that atomic force microscopy is a powerful method to investigate phase transitions in phospholipid bilayers and structural transformations in macromolecular membrane complexes. During dynamic processes, temperature control AFM can provide additional information and insight into membrane behaviour. For example, we believe that melting a complex of gel phase bilayer with DNA provides a reproducible way to form and image by AFM complex of DNA adsorbed on a fluid phase bilayer, which is very challenging for current AFM techniques. By choosing appropriate phospholipids, which may interact with DNA in the fluid phase, it can be possible to observe and control the process of complexation.

Acknowledgement This research has been financially supported by NSERC, ACB and the University of Calgary. We thank the staff at Molecular Imaging for their continued support.

References 1. Morris, V.J., Kirby A.R., and Gunning A.P. (2001) Atomic force microscopy for Biologists, Imperial College Press, London. 2. Han, W., Lindsay, S.M., and Jing, T.A (1996) Magnetically driven oscillating probe microscope for operation in liquids, Appl. Phys. Lett. 69, 1-3.

483 3. Han, W. and Lindsay, S.M. (1998) Probing molecular ordering at a liquid-solid interface with a magnetically oscillated atomic force microscope, Appl. Phys. Lett. 72, 1656-1658. 4. Fasolka, M.J., Mayes, A.M. and Magonov, S.N. (2001) Thermal enhancement of AFM phase contrast for imaging diblock copolymer thin film morphology, Ultramicroscopy 90, 21-31. 5. Vesenka, J., Manne, S., Giberson, R., Marsh, T., and Henderson, E. (1993) Colloid gold particles as an incompressible atomic force microscope imaging standard for assessing the compressibility of biomolecules, Biophys. J. 65, 992-997. 6. Brian, A.A. and McConnell, H.M. (1984) Allogenic stimulation of cytoxic T cells by supported planar membranes, Proc. Natl. Acad. Sci. U.S.A. 81, 6159-6163. 7. Leonenko, Z.V., Carnini, A., and Cramb, D.T. (2000) Supported planar bilayer formation by vesicle fusion: the interaction of phospholipid vesicles with surfaces and the effect of gramicidin on bilayer properties using atomic force microscopy, Biochim. Biophys. Acta. 1509, 134-147. 8. Leonenko, Z., Merkle, D., Lees-Miller, S.P. and Cramb, D. (2002) Lipid Phase Dependence of DNA Cationic Phospholipid Bilayer Interactions examined using Atomic Force Microscopy, Langmuir, 18, 48734884. 9. Leonenko, Z. and Cramb, D. (2002) Effect of DNA adsorption on the phase cycling of a supported phospholipid bilayer, Nanoletters, 2, 305-309. 10. Leonenko, Z.V., Ma, H., Dahms, T.E.S., and Cramb, m D.T., (2004) Investigation Of Temperature Induced Phase Transitions In DOPC And DPPC Phospholipid Bilayers Using Temperature-Controlled Scanning Force Microscopy, Biophys. J. (in press). 11. Sorgi F.L. and Huang, L. (1997) Drug delivery applications, In Lipid polymorphism and membrane properties. Current Topics in Membranes. 44, 449-475. 12. May, S., Harris, D., and Ben-Shaul, A. (2000) The Phase Behavior of Cationic Lipid-DNA Complexes, Biophys. J. 78, 1681-1697. 13. Koltover, I., Salditt, T., Rädler, J.O., and Safinya, C.R. (1998) An inverted hexagonal phase of cationicDNA complexes related to DNA release and delivery”, Science. 281, 78-81. 14. Rädler, J.O., Koltover, I., Jamieson, A., Salditt, T., and Safinya, C.R. (1998) Structure and interfacial aspects of self-assembled cationic lipid-DNA gene carrier complexes, Langmuir. 14, 4272-4283. 15. Wong, F.M.P., Reimer, D.L., and Bally, M.B. (1996) Cationic lipid binding to DNA: characterization of complex formation, Biochem. 35, 5756-5763. 16. May, S. and Ben-Shaul, A. (1997) DNA-lipid complexes: stability of honeycomb-like and spaghetti-like structures, Biophys. J. 73, 2427-2440. 17. Dan, N. (1998) The structure of DNA complexes with cationic liposomes-cylindrical or lamellar?, Biochim. Biophys. Acta. 1369, 34-38. 18. Wagner, K., Harries, D., May, S., Kahl, V., Radler, J. O., and Ben-Shaul, A. ( 2000 ) Counterion release upon cationic lipid-DNA complexation, Langmuir.16, 303-306. 19. Dan, N. (1997) Multilamellar structures of DNA complexes with cationic, liposomes Biophys. J. 73, 1842-1846. 20. Mou, J., Czajkowsky, D.M., Zhang, Y., and Shao, Z. (1995) High Resolution Atomic Force Microscopy of DNA: the pitch of the double helix, FEBS Lett. 371, 279-282. 21. Fang, Y. and Yang, J. (1997) Effect of Cationic Strength and Species on 2-D Condensation of DNA, J. Phys. Chem. B 101, 3453-3456. 22. Hansma H. G. and Laney, D. E. (1996) DNA binding to mica correlates with cationic radius: assay by atomic force microscopy, Biophys. J. 70, 1933-1939. 23. Han, W., Dlakic, M., Zhu, Y.J., Lindsay, S.M., and Harrington, R.E. (1997) Strained DNA is kinked by 2+ low concentrations of Zn , Proc. Natl. Acad. Sci. U.S.A. 94, 10565-10570. 24. Tenchov, B., Koynova, R., and G.Rapp, (2001) New ordered metastable phases between the gel and subgel phases in hydrated phospholipids, Biophys. J. 80, 1873-1890. 25. Kinnunen, P.K.J. and Mouritsen, O.G., (1994) Special Issue of Functional Dynamics of Lipids in Biomembranes, Chem. Phys. Lipids, 73, 1-2.

INDEX OF KEYWORDS

Part I – Fundamentals of Functional Materials FUNCTIONAL MATERIALS, P. VILARINHO Keywords: Dielectrics, piezoelectrics, pyroelectrics, ferroelectrics, ferroelectrics, relaxor ferroelectrics, preparation, properties, applications.

incipient

SCALING OF SILICON-BASED DEVICES TO SUBMICRON DIMENSIONS, A.I. KINGON Keywords: Scaling, silicon-based microelectronics, nanoelectronics, functional electronics, new materials. UNSOLVED PROBLEMS IN FERROELECTRICS FOR SCANNING PROBE MICRSOCOPY, J.F. SCOTT Keywords: Ferroelectric thin films, lead zirconate titanate (PZT), hafnia, zirconia, strontium bismuth tantalite, strontium titanate/barium titanate superlattices, Fe-RAM, DRAM Part II – Fundamentals of Scanning Probe Techniques PRINCIPLES OF BASIC AND ADVANCED SCANNING PROBE MICROSCOPY, D.A. BONNELL, R. SHAO Keywords: Scanning probe microscopy, multiple modulation, spatial resolution, complex materials, molecular wires, ferroelectric domains. NANOSCALE PROBING OF PHYSICAL AND CHEMICAL FUNCTIONALITY WITH NEAR-FIELD OPTICAL MICROSCOPY, L.M. ENG Keywords: Near-field optical microscopy, organic and inorganic materials, microscopical and sensoric applications NANOSCALE ELECTRONIC MEASUREMENTS OF SEMICONDUCTORS USING KELVIN PROBE FORCE MICROSCOPY, Y. ROSENWAKS AND R. SHIKLER Keywords: Kelvin probe force microscopy, semiconductor GaP, minority-carrier diffusion length, potential distribution EXPANDING THE CAPABILITIES OF THE SCANNING TUNNELING MICROSCOPE, K.F. KELLY, Z.J. DONHAUSER, B.A. MANTOOTH, AND P.S. WEISS Keywords: Scanning tunneling microscopy, semiconductors, dopant profiling, advanced image processing techniques

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FUNCTIONS OF NC – AFM ON ATOMIC SCALE, S. MORITA, N. OYABU, T. NISHIMOTO, R. NISHI, O. CUSTANCE, I. YI, Y. SUGAWARA Keywords: Noncontact atomic force microscope, atomic force mechanisms, threedimensional mapping tool, atom manipulation tool

Part III – Application of Scanning Tecnhiques to Functional Materials SCANNING PROBE MICROSCOPY OF PIEZOELECTRIC AND TRANSPORT PHENOMENA IN ELECTROCERAMIC MATERIALS, S.V. KALININ AND D.A. BONNELL Keywords: Scanning Probe Microscopy, electronic ceramic materials, ZnO, Nb-doped SrTiO3, BaTiO3, BiFeO3, transport properties, electric phenomena, ferroelectric domain structure SFM-BASED METHODS FOR FERROELECTRIC STUDIES, A. GRUVERMAN Keywords: Scanning force microscopy, piezoresponse force microscopy, ferroelectrics, films, domain imaging, polarization mechanism SCANNING TUNNELING SPECTROSCOPY. Local density of states and spin distribution of interacting electron systems, M. MORGENSTERN Keywords: Scanning tunneling spectroscopy, electron t systems in InAs, domains in ferro-magnetic particles. NANOINSPECTION OF DIELECTRIC AND POLARIZATION PROPERTIES AT INNER AND OUTER INTERFACES IN FUNCTIONAL FERROELECTRIC PZT THIN FILMS, L.M. ENG Keywords: piezoresponse force microscopy, Kelvin probe force microscopy, pull-off force spectroscopy, PZT thin films, polarization profile, local dielectric properties MICROSCALE CONTACT CHARGING ON A SILICON OXIDE, S. MORITA, T. UCHIHASHI, K. OKAMOTO, M. ABE, Y. SUGAWARA Keywords: Electrostatic force microscopy, noncontactt atomic force microscopy, Kelvin probe force microscopy, silicon oxide, GaAs, contact charging, charge distribution CONSTRUCTIVE NANOLITHOGRAPHY, S.R. COHEN, R. MAOZ, AND J. SAGIV Keywords: Scanning Probe Microscopy, constructing nanolithography, self-assembled monolayers (SAMs), silane-based SAMs NANOMETER-SCALE ELECTRONICS AND STORAGE, K.F. KELLY, Z.J. DONHAUSER, P.A. LEWIS, R.K. SMITH, AND P.S. WEISS Keywords: Scanning tunneling microscopy (STM), self-assembly techniques, alkanethiolate SAMs, nanoscale electronic devices

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Part IV – Contributed papers STM TIPS FABRICATION FOR CRITICAL DIMENSION MEASUREMENT, A.PASQUINI, G.B.PICOTTO, M. PISANI Keywords: Fabrication, STM tips, Tungsten (W), electrochemical reaction SCANNING PROBE MICROSCOPY CHARACTERIZATION OF FERROELECTRICS DOMAINS AND DOMAINS WALLS IN KTiOPO4, C. CANALIAS, R. CLEMENS, J. HELLSTRÖM, F. LAURELL, J. WITTBORN, H. KARLSSON Keywords: Scanning probe microscopy, inverse piezoelectric effect, KTiOPO4, ferroelectric domains, domain imaging, single crystals and waveguides. IMAGING LOCAL DIELECTRIC AND MECHANICAL RESPONSES WITH DYNAMIC HETERODYNED ELECTROSTATIC FORCE MICROSCOPY, D. R. OLIVER, K.M. CHENG, A. PU, D.J. THOMSON AND G.E. BRIDGES Keywords: scanning probe microscopy, electrostatic force microscopy, polarization dynamics, nanotechnology, micro-electro-mechanical systems. AFM PATTERNING OF SrTiO3-δ THIN FILMS AND DEVICE APPLICATIONS, L. PELLEGRINO. Keywords: Atomic Force Microscope, conducting SrTiO3-δ thin films, functional electronics conducting element. NANOSCALE INVESTIGATION OF A RAYLEIGH WAVE ON LiNbO3, J. YANG AND R. KOCH Keywords: Surface acoustic waves, elastic properties of solid, scanning tunneling microscopy. SCANNING CAPACITANCE FORCE MICROSCOPY AND KELVIN PROBE FORCE MICROSCOPY OF NANOSTRUCTURES EMBEDDED IN SiO2, G. TALLARIDA, S. SPIGA, M. FANCIULLI Keywords: Scanning capacitance force microscopy, Kelvin probe force microscopy, Sn nanostructures, silicon oxide thin films ELECTRICAL CHARACTERISATION OF III-V BURIED HETEROSTRUCTURE LASERS BY SCANNING CAPACITANCE MICROSCOPY, O. DOUHÉRET, K. MAKNYS AND S. ANAND Keywords: Scanning capacitance microscopy, GaAs/AlGaAs heterostructures, GaInP:Fe, optoelectronic devices PROBING THE DENSITY OF STATES OF HIGH TEMPERATURE SUPERCONDUCTORS WITH POINT CONTACT TUNNELING SPECTROSCOPY, L. OZYUZER, J.F. ZASADZINSKI, N. MIYAKAWA, K.E. GRAY Keywords: Point contact tunneling spectroscopy, I-V and dI/dV-V characteristics, double CuO2 layer Bi2Sr2CaCu2O8+į, high-Tc superconductivity

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ANNEALING INFLUENCE ON CO ULTRATHIN FILM MORPHOLOGY IN MBE GROWN Co/Au BILAYERS, A. WAWRO, L.T. BACZEWSKI, P. PANKOWSKI, P. ALESZKIEWICZ, M. KISIELEWSKI, I. SVEKLO, A. MAZIEWSKI Keywords: Atomic force microscopy, Auger electron spectroscopy, reflection high energy electron diffraction, Au/Co/Au, magnetic applications. CORRELATION BETWEEN THE SURFACE RELIEF AND INTERFACES STRUCTURE OF FE/CR SUPERLATTICES AND ELECTROMAGNETIC WAVES PENETRATION, A.RINKEVICH, L.ROMASHEV, V.USTINOV Keywords: Penetration of electromagnetic waves, tunneling microscopy, ferromagnetic resonance, Fe/Cr superlattices, thin films, magnetic properties. MAGNETORESISTANCE AND MICROSTRUCTURE OF MAGNETIC THIN FILM MULTILAYERS, JENICA NEAMTU AND M. VOLMER Keywords: magnetoresistance, nickel iron permalloy, multilayers, thickness, microstructure SPM INVESTIGATION OF THIOLATED GOLD NANOPARTICLE PATTERNS DEPOSITED ON DIFFERENT SELF-ASSEMBLED SUBSTRATES, F. SBRANA, M. T. PARODI, D. RICCI, E. DI ZITTI Keywords: Scanning Probe Microscopy, Transmission Electron Microscopy, thiolated gold nanoparticles, self-assembled monolayers. AFM OF GUANINE ADSORBED ON HOPG UNDER ELECTROCHEMICAL CONTROL, A.-M. CHIORCEA AND A.M. OLIVEIRA BRETT Keywords: Atomic Force Microscopy, Magnetic AC Mode, electrochemical deposition, guanine, adsorption mechanism. DYNAMICS IN MODEL MEMBRANES AND DNA-MEMBRANE COMPLEXES USING TEMPERATURE CONTROLLED ATOMIC FORCE MICROSCOPY, Z.V. LEONENKO AND D.T. CRAMB Keywords: Atomic force microscopy, phospholipid bilayers, macromolecular membrane complexes, phase transition.