Gallium Oxide: Technology, Devices and Applications (Metal Oxides) 0128145218, 9780128145210

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Gallium Oxide

Metal Oxides Series The Metal Oxides Book Series Edited by Ghenadii Korotcenkov Forthcoming Titles l

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Metal Oxide Powder Technologies, Yarub Al-Douri, 9780128175057 Metal Oxide Glass Nanocomposites, Sanjib Bhattacharya, 9780128174586 Tin Oxide Materials, Marcelo Ornaghi Orlandi, 9780128159248 Cerium Oxide, Salvatore Scire, Leonardo Palmisano, 9780128156612 Solution Processed Metal Oxide Thin Films for Electronic Applications, Zheng Cui, 9780128149300 Metal Oxides for Non-volatile Memory, Panagiotis Dimitrakis, Ilia Valov, 9780128146293 Colloidal Metal Oxide Nanoparticles, Sabu Thomas, Anu Tresa Sunny, Prajitha V, 9780128133576 Gallium Oxide, Stephen Pearton, Fan Ren, Michael Mastro, 9780128145210 Metal Oxide Nanostructures, Daniela Nunes, Lidia Santos, Ana Pimentel, Pedro Barquinha, Luis Pereira, Elvira Fortunato, Rodrigo Martins, 9780128115121 Metal Oxides in Energy Technologies, Yuping Wu, 9780128111673 Gas Sensors Based on Conducting Metal Oxides, Nicolae Barsan, Klaus Schierbaum, 9780128112243 Metal Oxide Nanostructured Phosphors, H. Nagabhushana, Daruka Prasad, S.C. Sharma, 9780128118528 Metal Oxide-Based Nanostructured Electrocatalysts for Fuel Cells, Electrolyzers, and Metal-Air Batteries, Teko W. Napporn, Yaovi Holade

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Metal Oxide-Based Photocatalysis, Adriana Zaleska-Medynska, 9780128116340 Metal Oxides in Heterogeneous Catalysis, Jacques C. Vedrine, 9780128116319 Magnetic, Ferroelectric, and Multiferroic Metal Oxides, Biljana Stojanovic, 9780128111802 Iron Oxide Nanoparticles for Biomedical Applications, Sophie Laurent, Morteza Mahmoudi, 9780081019252 The Future of Semiconductor Oxides in Next-Generation Solar Cells, Monica Lira-Cantu, 9780128111659 Metal Oxide-Based Thin Film Structures, Nini Pryds, Vincenzo Esposito, 9780128111666 Metal Oxides in Supercapacitors, Deepak Dubal, Pedro Gomez-Romero, 9780128111697 Transition Metal Oxide Thin Film-Based Chromogenics and Devices, Pandurang Ashrit, 9780081018996

Metal Oxides Series

Gallium Oxide Technology, Devices and Applications

Edited by

Stephen Pearton Fan Ren Michael Mastro Series Editor

Ghenadii Korotcenkov

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States © 2019 Elsevier Inc. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/ permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein).

Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-12-814521-0 For information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Mathew Deans Acquisition Editor: Kayla Dos Santos Editorial Project Manager: Gabriela Capille Production Project Manager: Swapna Srinivasan Cover Designer: Miles Hitchen Typeset by SPi Global, India

List of contributors

Kwang Hyeon Baik Department of Materials Science and Engineering, Hongik University, Jochiwon, South Korea Kristen Bevlin Plasma-Therm, St. Petersburg, FL, United States Patrick H. Carey IV Department of Chemical Engineering; Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States Kalyan K. Das Materials Science & Engineering, North Carolina State University, Raleigh, NC, United States Roberto Fornari Department of Mathematical, Physical and Computer Sciences, University of Parma; IMEM-CNR Institute, Parma, Italy W. Beall Fowler Department of Physics, Lehigh University, Bethlehem, PA, United States Dwarakanath Geerpuram Plasma-Therm, St. Petersburg, FL, United States Krishnendu Ghosh Electrical Engineering Department, University at Buffalo, Buffalo, NY, United States B.P. Gila Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States N.C. Giles Department of Engineering Physics, Air Force Institute of Technology, Wright-Patterson Air Force Base, Dayton, OH, United States L.E. Halliburton Department of Physics and Astronomy, West Virginia University, Morgantown, WV, United States D. Hays Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States Hisao Hidaka Department of Chemistry, Meisei University, Hodokubo, Hino-shi, Tokyo, Japan

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List of contributors

Ching-Hwa Ho Graduate Institite of Applied Science and Technology, National Taiwan University of Science and Technology, Taipei, Taiwan Soohwan Jang Department of Chemical Engineering, Dankook University, Yongin, South Korea Lai Jiang Materials Science & Engineering, Carnegie Mellon University, Pittsburgh, PA, United States Sunwoo Jung Department of Chemical Engineering, Dankook University, Yongin, South Korea Rohit Khanna Plasma-Therm, St. Petersburg, FL, United States Jihyun Kim Department of Chemical and Biological Engineering, Korea University, Seoul, South Korea Janghyuk Kim Department of Chemical and Biological Engineering, Korea University, Seoul, South Korea Suhyun Kim Department of Chemical and Biological Engineering, Korea University, Seoul, South Korea Sriram Krishnamoorthy Electrical and Computer Engineering, The University of Utah, Salt Lake City, UT, United States Mukesh Kumar Department of Physics; Nanoscale Research Facility, Indian Institute of Technology Delhi, New Delhi, India Sudheer Kumar Department of Physics, Indian Institute of Technology Delhi, New Delhi, India Vikram Kumar Nanoscale Research Facility, Indian Institute of Technology Delhi; Solid State Physical Laboratory, New Delhi, India Luke A.M. Lyle Materials Science & Engineering, Carnegie Mellon University, Pittsburgh, PA, United States M.A. Mastro U.S. Naval Research Laboratory, Washington, DC, United States David J. Meyer U.S. Naval Research Laboratory, Washington, DC, United States Rangarajan Muralidharan Centre for Nano Science and Engineering (CeNSE), Indian Institute of Science (IISc), Bangalore, India

List of contributors

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Digbijoy N. Nath Centre for Nano Science and Engineering (CeNSE), Indian Institute of Science (IISc), Bangalore, India Neeraj Nepal U.S. Naval Research Laboratory, Washington, DC, United States Sooyeoun Oh Department of Chemical and Biological Engineering, Korea University, Seoul, South Korea Stephen Pearton Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States Lisa M. Porter Materials Science & Engineering, Carnegie Mellon University, Pittsburgh, PA, United States Anamika Singh Pratiyush Centre for Nano Science and Engineering (CeNSE), Indian Institute of Science (IISc), Bangalore, India Ying Qin Department of Physics, Lehigh University, Bethlehem, PA, United States Siddharth Rajan Department of Electrical and Computer Engineering, The Ohio State University, Columbus, OH, United States C.V. Ramana Department of Mechanical Engineering, University of Texas at El Paso, El Paso, Texas, United States F. Ren Department of Chemical Engineering; Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States D. Scott Katzer U.S. Naval Research Laboratory, Washington, DC, United States R. Singh Department of Physics; Nanoscale Research Facility, Indian Institute of Technology Delhi, New Delhi, India Uttam Singisetti Electrical Engineering Department, University at Buffalo, Buffalo, NY, United States Michael Stavola Department of Physics, Lehigh University, Bethlehem, PA, United States Marko J. Tadjer U.S. Naval Research Laboratory, Washington, DC, United States Li-Chia Tien Department of Materials Science and Engineering, National Dong Hua University, Hualien, Taiwan

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List of contributors

Tohru Tsukamoto Department of Chemistry, Meisei University, Hodokubo, Hino-shi, Tokyo, Japan Holger von Wenckstern Universit€at Leipzig, Felix-Bloch-Institut Festk€ orperphysik, Halbleiterphysik, Linnestraße 5, Leipzig, Germany

f€ur

Philip Weiser Department of Physics, Lehigh University, Bethlehem, PA, United States Jiancheng Yang Department of Chemical Engineering; Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States

Series editor’s biography

Ghenadii Korotcenkov earned his PhD in material sciences from the Technical University of Moldova, Chisinau, Moldova, in 1976 and his Doctor of Science degree in physics from the Academy of Science of Moldova in 1990 (Highest Qualification Committee of the USSR, Moscow). He has more than 40 years of experience as a scientific researcher. For a long time, he was the leader of the gas sensor group and manager of various national and international scientific and engineering projects carried out in the Laboratory of Microand Optoelectronics, Technical University of Moldova. His research had financial support from international foundations and programs such as the CRDF, the MRDA, the ICTP, the INTAS, the INCO-COPERNICUS, the COST, and the NATO. From 2007 to 2008, he was an invited scientist in the Korea Institute of Energy Research, Daejeon, South Korea. After which, until the end of 2017 G. Korotcenkov was a research professor at the School of Materials Science and Engineering at the Gwangju Institute of Science and Technology, Gwangju, South Korea. Currently G. Korotcenkov is the chief scientific researcher at the Department of Physics and Engineering at the Moldova State University, Chisinau, Rep. of Moldova. Specialists from the former Soviet Union know G. Korotcenkov’s research results in the field of study of Schottky barriers, MOS structures, native oxides, and photoreceivers on the basis of III–V compounds such as InP, GaP, AlGaAs, and InGaAs. His present scientific interests starting from 1995 include material sciences, focusing on the metal oxide film deposition and characterization, surface science, and the design of thin film gas sensors and thermoelectric convertors. These studies were carried out in cooperation with scientific teams from Ioffe Institute (St. Petersburg, Russia), University of Michigan (Ann Arbor, United States), Kiev State University (Kiev, Ukraine), Charles University (Prague, Czech Republic), St. Petersburg State University (St. Petersburg, Russia), Illinois Institute of Technology (Chicago, United States), University of Barcelona (Barcelona, Spain), Moscow State University (Moscow, Russia), University of Brescia (Brescia, Italy), Belarus State University (Minsk, Belarus), and South-Ukrainian University (Odessa, Ukraine). G. Korotcenkov is the author or editor of 38 books, including the 11-volume Chemical Sensors series published by the Momentum Press (United States), 15-volume

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Series editor’s biography

Chemical Sensors series published by Harbin Institute of Technology Press (China), 3-volume Porous Silicon: From Formation to Application published by CRC Press (United States), 2-volume Handbook of Gas Sensor Materials published by Springer (United States), and 3-volume Handbook of Humidity Measurement, which is being published by CRC Press (United States). In addition, at present, G. Korotcenkov is a series’ editor of Metal Oxides series, which is published by Elsevier. G. Korotcenkov is author and coauthor of more than 600 scientific publications, including 30 review papers, 38 book chapters, and more than 250 articles published in peer-reviewed scientific journals (h-factor ¼ 40 [Scopus] and h-factor ¼ 47 [Google Scholar citation]). He is a holder of 17 patents. He has presented more than 200 reports at national and international conferences, including 17 invited talks. G. Korotcenkov was co-organizer of several international conferences. His name and activities have been listed by many biographical publications, including Who’s Who. His research activities are honored by an Award of the Supreme Council of Science and Advanced Technology of the Republic of Moldova (2004); Prize of the Presidents of the Ukrainian, Belarus, and Moldovan Academies of Sciences (2003); and National Youth Prize of the Republic of Moldova in the field of science and technology (1980), among others. G. Korotcenkov also received a fellowship from the International Research Exchange Board (IREX, United States, 1998), Brain Korea 21 Program (2008–2012), and Brain Pool Program (Korea, 2007–2008 and 2015–2017).

Editors biography

Stephen Pearton is a distinguished professor and alumni chair of Materials Science and Engineering at the University of Florida, Gainesville, FL, United States. He has a PhD in Physics from the University of Tasmania and was a postdoc at UC Berkeley prior to working at AT&T Bell Laboratories in 1994–2004. His interests are in the electronic and optical properties of semiconductors. He is a Fellow of the IEEE, AVS, ECS, TMS, MRS, SPIE, and APS. He was the recipient of the J.J. Ebers Award from IEEE, Gordon Moore Award from ECS, John Thornton Award from AVS, Adler Award from APS, and the Bardeen Award from TMS. Michael Mastro is a scientist at the US Naval Research Laboratory in Washington DC, United States. His research interests are in fabrication and modeling physics of high-power transistors as well as light emitting, photonic crystal, plasmonic, metamaterial, and photovoltaic devices, fabrication of nanostructured devices, development of numerical models, and neural network algorithms to describe nanoscale semiconductor physics. He has a PhD in Chemical Engineering from University of Florida and BS from John’s Hopkins University. He is the co-editor of the very popular book “III–V Compound Semiconductors: Integration with Silicon-Based Microelectronics” and won an outstanding paper award from Japan Society of Applied Physics in 2014. Fan Ren is a distinguished professor of Chemical Engineering at the University of Florida, Gainesville, FL, United States. He joined UF in 1997 after 12 years as a member of Technical Staff at AT&T Bell Laboratories, where he was responsible for highspeed compound semiconductor device development. He is a fellow of AIChE, ECS, IEEE, APS, MRS, and AVS. He was the recipient of the Gordon Moore Award from ECS and the Albert Nerken Award from AVS.

Preface to the series

The field of synthesis, study, and application of metal oxides is one of the most rapidly progressing areas of science and technology. Metal oxides are one of the most ubiquitous compound groups on earth, which has large variety of chemical compositions, atomic structures, and crystalline shapes. In addition, metal oxides are known to possess unique functionalities that are absent or inferior in other solid materials. In particular, metal oxides represent an assorted and appealing class of materials, properties of which exhibit a full spectrum of electronic properties—from insulating to semiconducting, metallic, and superconducting. Moreover, almost all the known effects including superconductivity, thermoelectric effects, photoelectrical effects, luminescence, and magnetism can be observed in metal oxides. Therefore, metal oxides have emerged as an important class of multifunctional materials with a rich collection of properties, which have great potential for numerous device applications. Specific properties of the metal oxides, such as the wide variety of materials with different electrophysical, optical, and chemical characteristics, their high thermal and temporal stability, and their ability to function in harsh environments, make metal oxides very suitable materials for designing transparent electrodes, high-mobility transistors, gas sensors, actuators, acoustical transducers, photovoltaic and photonic devices, photoand heterogeneous catalysts, solid-state coolers, high-frequency and micromechanical devices, energy harvesting and storage devices, nonvolatile memories, and many others in the electronics, energy, and health sectors. In these devices metal oxides can be successfully used as sensing or active layers, substrates, electrodes, promoters, structure modifiers, membranes, and fibers, that is, can be used as active and passive components. Among other advantages of metal oxides are the low fabrication cost and robustness in practical applications. Furthermore, the metal oxides can be prepared in various forms such as ceramics, thick films and thin film. In addition thin-film deposition can be used in deposition techniques that are compatible with standard microelectronic technology. Last factor is very important for large-scale production, because the microelectronic approach promotes low cost for mass production, offers the possibility of manufacturing devices on a chip, and guarantees good reproducibility. Various metal oxides nanostructures, including nanowires, nanotubes, nanofibers, core-shell structures, and hollow nanostructures can also be synthesized. As it is known, the field of metal-oxide nanostructured morphologies (e.g., nanowires, nanorods, nanotubes, etc.) has become one of the most active research areas within the nanoscience community. The ability to create a variety of metal oxide-based composites and the ability to synthesize various multicomponent compounds significantly expand the range of properties that metal oxide-based materials can have, making metal oxides a truly

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versatile multifunctional material for widespread use. As it is known small changes in their chemical composition and atomic structure can be accompanied by the spectacular variation in properties and behavior of metal oxides. Even now, advances in synthesizing and characterizing techniques are revealing numerous new functions of metal oxides. Taking into account the importance of metal oxides for progress in microelectronics, optoelectronics, photonics, energy conversion, sensor, and catalysis, a large number of various books devoted to this class of materials have been published. However, one should note that some books from this list are too general, some are collections of various original works without any generalizations, and other ones were published many years ago. But, during past decade a great progress has been made on the synthesis as well as on the structural, physical, chemical characterization and application of metal oxides in various devices, and a large number of papers have been published on metal oxides. In addition, till now many important topics related to metal oxides study and application have not been discussed. To remedy the situation in this area, we decided to generalize and systematize the results of research in this direction and to publish a series of books devoted to metal oxides. One should note that proposed book series “Metal Oxides” is the first one, devoted to consideration of metal oxides only. We believe that combining books on metal oxides in a series could help readers in searching required information on the subject. In particular, we believe that the books from our series, which have a clear specialization by its content, will provide interdisciplinary discussion for various oxide materials with a wide range of topics, from material synthesis and deposition to characterizations, processing, and then to device fabrications and applications. This book series is prepared by a team of highly qualified experts, which guarantees it a high quality. I hope that our books will be useful and comfortable in use. I also hope that readers will consider the “Metal Oxides” book series like an encyclopedia of metal oxides that enables to understand the present status of metal oxides, to estimate the role of multifunctional metal oxides in design of advanced devices, and then based on observed knowledge formulate new goals for the further research. The intended audience of present book series is scientists and researchers, working or planning to work in the field of materials related to metal oxides, that is, scientists and researchers whose activities are related to electronics, optoelectronics, energy, catalysis, sensors, electrical engineering, ceramics, biomedical designs, etc. I believe that the “Metal Oxides” book series will also be interesting for practicing engineers or project managers in industries and national laboratories, who would like to design metal oxide-based devices, but don’t know how to do it, and how to select optimal metal oxide for specific applications. With many references to the vast resource of recently published literature on the subject, this book series will be serving as a significant and insightful source of valuable information, providing scientists and engineers with new insights for understanding and improving existing metal oxidebased devices and for designing new metal oxide-based materials with new and unexpected properties.

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I believe that the “Metal Oxides” book series would be very helpful for university students, postdocs, and professors. The structure of these books offers a basis for courses in the field of material sciences, chemical engineering, electronics, electrical engineering, optoelectronics, energy technologies, environmental control, and many others. Graduate students could also find the book series to be very useful in their research and understanding features of metal oxides synthesis, study and application of this multifunctional material in various devices. We are sure that all of them will find information useful for their activity. Finally, I thank all contributing authors and book editors who have been involved in the creation of these books. I am thankful that they agreed to participate in this project and for their efforts in the preparation of these books. Without their participation, this project would have not been possible. I also express my gratitude to Elsevier for giving us the opportunity to publish this series. I especially thank all team of editorial office at Elsevier for their patience during the development of this project and for encouraging us during the various stages of preparation. Ghenadii Korotcenkov

Preface

The properties of Ga2O3 have been investigated over a long period. For example, the phase equilibria of the Al2O3-Ga2O3-H2O system was first reported in 1952, just 5 years after the first demonstration of transistors in the Ge system. That initial work also identified the different polymorphs (i.e., different forms or crystal structures) of Ga2O3 and established their regions of stability. There are five commonly identified polymorphs of Ga2O3, labeled as α, β, γ, δ, and ε. These are known, respectively, as corundum (α), monoclinic (β), defective spinel (γ), and orthorhombic (ε), with the δ-phase commonly accepted as being a form of the orthorhombic phase. Among these different phases of Ga2O3, the orthorhombic β-gallia structure (β-phase or β-Ga2O3) is the most stable crystal structure and has attracted most of the recent attention. It is also the subject of this book. The melting point of Ga2O3 is about 1800°C and at such temperature only β-phase is stable among the five polymorphs. The different polymorphs can be either insulators or conductors, depending on the growth conditions. The resulting crystals are layered materials, similar to the behavior of GaSe and GaTe. The original studies in the 1950s and 1960s indicated that the other polymorphs of Ga2O3 convert to the stable β-form with heat treatment. β-Gallium oxide (Ga2O3) is emerging as a viable candidate for certain classes of power electronics, solar blind UV photodetectors, solar cells, and sensors with capabilities beyond existing technologies due to its large bandgap. Research on β-Ga2O3 and to a lesser extent some of the other polymorphs, has been accelerating worldwide, especially since 2010. This is because its ultrawide-band-gap (5 eV) has attracted considerable attention from the field of power electronics and ultraviolet optical device engineering, and 6-in. high-quality single crystals are now available owing to the development of liquid-solution technologies. Fig. 1 shows the exponential growth in the number of publications on Ga2O3 as a function of time on logarithmic scale. The observed exponential growth is evidence of constantly increasing interest in Fig. 1 Number of publications Ga2O3 as a function of time. Data from the ISI Web of Knowledge search engine and was collated by Dr. Marko Tadjer of Naval Research Laboratory. The data were then plotted and supplied by courtesy of Dr. Lucas Lacasa, EPSRC Early Career Fellow and Senior Lecturer in Applied Mathematics, Queen Mary University of London (http://www.maths.qmul.ac. uk/lacasa/main.html).

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the topic. It is expected this trend will continue for the foreseeable future as research funding for power device development increases, which in turn will be driven by funding for development of bulk and epitaxial growth. Since melt growth techniques can be used to grow bulk crystals of β-GaO3, the cost of producing larger area, uniform substrates is potentially lower compared with the vapor growth techniques used to manufacture bulk crystals of GaN and SiC, or more exotic wide-band-gap semiconductors like diamond. The main melt growth methods reported to date have included Czochralski (CZ), float zone (FZ), vertical Bridgman (VB)/vertical gradient freeze (VGF), and edge-defined film-fed growth (EFG) methods. An illustration of the relative cost of β-Ga2O3 grown by the EFG technique to a large diamond crystal of comparable value is shown in Fig. 2, where the Ga2O3 is already a thousand times less expensive than diamond. There are aggressive moves worldwide to lower the cost of bulk Ga2O3 substrates in order to accelerate the competitiveness of power electronics based on this material with that of SiC and GaN. The performance of technologically important high-voltage rectifiers and enhancement-mode metal-oxide field effect transistors benefit from the larger critical electric field of β-Ga2O3 relative to either SiC or GaN. The estimated electric field breakdown field strength (EB) is in the range of 5–9 MV.cm and is commonly quoted as 8 MV cm 1 for β-Ga2O3. This value is roughly a factor of two larger than the theoretical limits of SiC and GaN. In terms of power switching device figures-of-merit, the large EB for Ga2O3 leads to a Baliga figure of merit almost three times larger than

Fig. 2 Approximate cost of two representative ultrawide large bulk crystals of diamond (left) or Ga2O3 (right). Courtesy of Dr. Marko Tadjer and Michael Mastro (Naval Research Laboratory, Washington DC, United States).

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for GaN and SiC. Of course, this is a rule-of-thumb comparison only and it does not take into account limiting factors such as the low thermal conductivity of Ga2O3 and the need to develop the epitaxial growth, contacting, and gate dielectrics needed for a complete device technology. Similarly, the absence of facile p-type doping in Ga2O3, which may be a fundamental issue resulting from the band structure, makes it very difficult to simultaneously achieve low turn-on voltages on power switching devices. For simple switching devices like Schottky rectifiers, the ability to reduce DC conduction losses in by minimizing on-resistance (Ron) and maximizing breakdown voltage (VB) is attractive for high-power and high-voltage applications and, when combined with advanced scaling techniques, high-speed switches for low/medium-power applications. The purpose of this book is to provide an update on recent advances in all areas of Ga2O3 materials, processing, and device development. These include epitaxial growth using advanced methods like metal organic chemical vapor deposition and molecular beam epitaxy, which are the workhorse thin film growth techniques for other compound semiconductors and because of their large installed base, will need to perform the same function for Ga2O3 transistor technologies. Recent work has shown the ability to control the electron densities in epitaxial layers from 1015 to 1020 cm 3 using Sn, Si, or Ge as the donor dopants. In terms of p-type doping, theory indicates that that all the expected acceptor dopants occupying Ga substitutional sites, a list that includes Li, Na, K, Be, Mg, Ca, Cu, Au, and Zn, result in deep acceptor levels and do not produce p-type conductivity in β-Ga2O3. These dopants are found to introduce deep acceptor levels with ionization energies of more than 1 eV. The prediction that holes are self-trapped to form polarons limits the current range of possible devices to unipolar conductivity. There has been observation of p-type conductivity, due most likely to ionized Ga vacancies at elevated temperature. The ionization energy of the acceptor level was measured to be 1.1 eV above the valence band edge. At elevated temperatures, up to 650 k, the measured hole mobility was 4.2 cm2 V s 1, much larger than predicted from the self-trapped (bound polaron) hole assumption. The availability of high-quality, large-area native substrates offer a complete platform for various applications such as high-performance power switching, radio frequency (RF) amplifiers, and harsh environment signal processing. In this book, we did not cover bulk crystal growth and single crystals, because this has been the subject of a series of recent reviews in the special issues of Japanese Journal of Applied Physics, Vol. 55, Issue 12 (2016), entitled Gallium Oxide and Related Semiconductors, and the Electrochemical Society Journal of Solid State Science and Technology Focus Issue on Ultrawide Bandgap Materials and Devices, Vol. 6, Issue 2 (2017), as well as in a recent review on the progress in the growth of β-Ga2O3 for power electronics applications (Materials Science in Semiconductor Processing, Vol. 78, May 2018, pages 132–146, authored by Michele Baldini, Zbigniew Galazka, G€unter Wagner). Examples of bulk crystals grown by various solution growth methods are shown in Fig. 3 and for more details, the reader is referred to the special journal issues listed above. In the materials synthesis area, we have instead focused on the topics of epitaxy and thin film deposition, as well as progress in nanostructures of β-Ga2O. We then explore the state-of-the-art in doping and structural properties, currently dominated

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Fig. 3 Bulk β-Ga2O3 crystals grown from the melt by: (A) float zone (FZ), (B) EFG, (C) Czochralski (CZ), and (D) vertical Bridgman (VB). Reprinted with permission from Elsevier, Copyright 2016, Michele Baldini, Zbigniew Galazka, G€ unter Wagner, Recent progress in the growth of β-Ga2O3 for power electronics applications, Mater. Sci. Semicond. Process. 78 (2018) 132–146.

by extended and point defects, look at the properties of (In,Ga)2O3 alloys, which will be critical for heterostructure devices based on Ga2O3 and examine the theory and mechanisms for high and low-field transport properties as well as the study and identification of electronic states using electron paramagnetic resonance. Hydrogen has a strong influence on the electrical properties of transparent conducting oxides where it can give rise to shallow donors and can passivate deep compensating defects. A lot of new information is now being reported about the nature of the dominant hydrogen-related defects in Ga2O3. This is followed by the crucial device processing technologies, namely, fabrication of Ohmic contacts and Schottky barriers and plasma and wet etching. Wide-band-gap devices are normally assumed to be radiation hard, allowing their use in environments like space. We include a chapter discussing the relative radiation hardness of Ga2O3 in comparison with GaN. There is a critical need to understand band alignments of common dielectrics on Ga2O3 and this is covered in one of the chapters. Finally, we have chapters on key device technologies, including solar-blind UV detectors using both thin films and exfoliated flakes, power MOSFETs and diodes, photo-assisted degradation of environmentally aggressive chemicals using Ga2O3 and finally, gas sensors capable of operation under elevated temperature conditions. In summary, it is clear that Ga2O3 has an excellent combination of materials properties with the exception of thermal conductivity. The availability of large, highquality bulk substrates, makes it an attractive option for power switching devices

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and solar-blind UV photodetection. There are still numerous areas that need additional work, including (i) Improved Ohmic contacts through optimized surface cleaning, interface tailoring, and increased doping capabilities. (ii) Better understanding of radiation damage in Ga2O3 and related heterostructures. (iii) Interface state density mitigation processes for dielectrics for MOS devices, as well as how these are affected by process/patterning and contacting conditions. (iv) Experimental clarification of energy levels of native defects such as oxygen vacancies and their role in residual conductivity relative to extrinsic impurities. (v) Doping studies to obtain p-type conductivity. (vi) Epi growth of β-(AlxGa1 x)2O3/Ga2O3 and β-(InxGa1 x)2O3 heterostructures growth to establish the stability regimes and compositional limits. (vii) Novel thermal management approaches for power devices to exploit the much higher thermal conductivities available with heat-sink materials like Cu, SiC, or diamond.

This book provides a framework for understanding the current state-of-the-art and how advances might be made in the areas listed above. Stephen Pearton Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States Fan Ren Department of Chemical Engineering, University of Florida, Gainesville, FL, United States Michael Mastro Electronics and Science Technology Division, US Naval Research Laboratory, Washington DC, United States

Progress in MOVPE growth of Ga2O3

1

Roberto Fornari*,† *Department of Mathematical, Physical and Computer Sciences, University of Parma, Parma, Italy, †IMEM-CNR Institute, Parma, Italy

Chapter Outline 1.1 Introduction 3 1.2 Homoepitaxial deposition of β-Ga2O3 4 1.3 Heteroepitaxial deposition of β-Ga2O3 11 1.4 Heteroepitaxial deposition of ε-Ga2O3 15 References 28

1.1

Introduction

Semiconducting sesquioxides, especially Ga2O3, are known since decades [1–3]; however, it is only in the past few years that they are massively investigated. This is essentially due to the development of suitable technologies for growth of large single crystals [4–9] and homo- and heteroepitaxial layers [10–15]. The possibility of growing single crystals and films with relatively low defect density, opened the way to new application areas, in addition to the well-known transparent conducting oxides (TCO) electrodes, namely, (i) substrates for GaN epilayers and (ii) high-power transistors, and (iii) UV detectors. Although the homoepitaxial growth is supposed to provide layers of higher crystallographic perfection, the deposition of gallium oxide on hetero-substrates such as α-Al2O3 (sapphire) is still very popular. It has the advantage of directly providing electrical isolation of the film, which may turn out useful in view of the application of β-Ga2O3 to power field effect transistors (FETs). Furthermore, sapphire wafers are commercially available with diameters of up to 600 and very cheap compared with the still expensive Ga2O3 substrates. This chapter provides a survey of the recent advances in homo- and heteroepitaxial deposition of Ga2O3 by metal-organic vapor-phase epitaxy (MOVPE). This includes (i) MOVPE deposition of β-Ga2O3 on differently oriented β-Ga2O3 homo-substrates; (ii) MOVPE deposition of β-Ga2O3 on hetero-substrates, with special emphasis on sapphire; and (iii) MOVPE deposition of metastable crystalline phases of gallium oxide. Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00001-4 © 2019 Elsevier Inc. All rights reserved.

4

1.2

Gallium Oxide

Homoepitaxial deposition of β-Ga2O3

Most investigations on homoepitaxial β-Ga2O3 have been carried out by molecular beam epitaxy (MBE) and more rarely by MOVPE. The reason for this lies in the gap that for a long time existed between the relatively poor structural quality of β-Ga2O3 layers grown by MOVPE in comparison to MBE materials. Indeed, the maximum electron mobilities of MOVPE layers were limited to about 40 cm2/Vs, while in layers grown by MBE on (010)-oriented homosubstrates were >100 cm2/Vs. However, as the MOVPE technology is more suitable for large-scale production of semiconductors, in the last times great efforts were made in order to improve the quality of homoepitaxial β-Ga2O3 by MOVPE. Different types of β-Ga2O3 wafers were tested as substrates and two orientations were generally used: (100) and (010). Of greatest importance is the wafer surface preparation before initiating the epitaxial growth: the (100) β-Ga2O3 substrates were first cleaned with acetone and isopropylic alcohol and then annealed in oxygen ambient at 1000°C for at least 1 h. This was sufficient to remove the subsurface polishing damage and get a regular step-and-terrace surface as shown in Fig. 1.1. Of course, the terrace width is a function of the substrate disorientation. In the example of Fig. 1.1, the surface deviated by about 0.2 degrees from the perfect (100) and the corresponding width was about 70–100 nm. Trimethylgallium (TMGa) is normally taken as Ga precursor while oxygen is supplied either as a pure gas (O2) or via water dissociation in the growth reactor. Surprisingly, the initial MOVPE deposition on (100) substrates did not result in coherent layers, but rather in nanowire and nanoparticles deposition, when using oxygen gas as oxidation source [11]. This happened for a wide range of epitaxial parameters: substrate temperature between 750°C and 850°C, reactor pressure in the range 5–100 mbar, oxygen-to-gallium ratios between 1000 and 9000. In the discussion on the influence of the oxygen precursor on the growth mode, authors of Ref. [11] considered the different thermodynamic conditions connected with a large oxygen or hydrogen excess in the reaction chamber. From thermodynamic simulation, they concluded that when using pure oxygen the main reaction products are solid Ga2O3 and gaseous oxygen, water, and carbon dioxide, whereas if water is used then hydrogen molecules, H2, are present in high concentration as compared with formed Ga2O3. O2 molecules are not present at all, since they are fully consumed by reaction with the organic species. The influence of the oxygen precursor on the growth mode was hence qualitatively discussed as follows: when pure O2 is employed, the formation energy of oxygen vacancies is substantially increased which in turn lowers their concentration. At the same time, H2O and CO2 are present in comparable concentrations, which promotes the formation of Ga2(CO3)3 that may act as a mask through which whiskers grow. On the other hand, when H2O is employed, the oxygen partial pressure is considerably low and oxygen vacancies are expected to form spontaneously. This time however H2O is present at much higher concentrations than CO2 so that the adsorption of CO2, and related carbonate formation, is prevented. H2O simply dissociates at oxygen vacancy sites thus promoting growth. As additional mechanism the authors also considered that hydrogen may occupy oxygen vacancy sites forming GadH bonds and thus reduce the surface state density. Therefore, this would

Progress in MOVPE growth of Ga2O3

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Fig. 1.1 (A) AFM image of (100)-oriented β-Ga2O3 after thermal annealing at 1000°C in oxygen ambient; (B) Corresponding surface scan. Note the formation of regular steps (0.6 nm high) and terraces (70–90 nm wide). From G. Wagner et al., Homoepitaxial growth of β-Ga2O3 layers by metal-organic vapor phase epitaxy, Phys. Status Solidi Appl. Mater. Sci. 211 (2014) 27–33.

change the surface potential and the kinetics at the substrate surface, promoting the layer-by-layer growth. The initial decomposition of the substrate, with the formation of tiny Ga clusters, that catalyzes the nucleation and growth of nanostructures cannot however be ruled out. If that is the actual case, one is therefore led to conclude that atomic oxygen from water dissociation in the high-temperature reactor chamber is more effective in preventing dissociation of the (100) surface of β-Ga2O3. What is sure is that taking H2O as atomic oxygen source leads to layer-by-layer growth and smooth layers as reported in Refs. [11, 13, 16]. The growth of homoepitaxial single-phase β-Ga2O3 on (010) substrates by lowpressure CVD was reported in Ref. [17]. Authors investigated the surface morphology and crystal quality as a function of growth temperature dependences. The growth studies were carried out in a custom-designed CVD system on commercial (010) β-Ga2O3

6

Gallium Oxide

substrates synthesized by the EFG growth method. The dislocation density of the substrate was estimated to be 10 times higher than those on (100) substrates because of a higher reevaporation rate of the highly volatile Ga2O species from the (100) plane [20]. Attempts to increase the growth rate by applying higher TEGa and O2 flow rates, while keeping a constant molar ratio between the two precursors, resulted in progressively increasing roughness up to a RMS value around 11 nm. The epitaxial growth on (010) Fe-doped semiinsulating β-Ga2O3 substrates proved to be very suitable from the point of view of n-type doping. Actually, by using tetraethylorthosilicate (TEOS) and tetraethyltin (TESn) as metallorganic precursors for Si and Sn, it was possible to achieve very high carrier concentrations. When the flow rates of TEOS and TESn were varied in the range  1  1011–1  108 mol/min, the free carrier concentrations could be adjusted in the doping range 1  1017 and 8  1019 cm3 [19]. It seems however that Si has to be preferred as donor, because of its higher incorporation rate and no memory effects in the MOVPE reactor. In a series of recent papers, the group at IKZ Berlin explained why the doping is more successful for epitaxial growth on (010) substrates and also provided some hints on elimination of structural defects in (100)-oriented films as a method to improve the doping efficiency in these epilayers [21, 22]. In Ref. [21], the electrical transport properties of Si-doped β-Ga2O3 (100) homoepitaxial layers were studied as a function of substrate misorientation: the layers perfectly (100) oriented, or with small offorientation, exhibited low mobility values as well as low doping efficiency. Furthermore, the Hall mobility collapsed below a threshold electron concentration (about 1018 cm3), in that reproducing the behavior already observed for Sn-doped layers [13]. By using transmission electron microscopy (TEM), authors showed that these homoepitaxial layers contained a very high density of twin lamellae with two different atomic structures, namely, coherent boundaries parallel to (100) and incoherent ones parallel to (001). While the former type of twins preserves the atomic coordination, in case of the incoherent twin boundaries (ITBs) one dangling bond per unit cell is present. The dangling bonds are arranged along the ITBs of thin twin lamellae and may act as acceptors. This model of electrically active dislocations is similar to the one first suggested by Read [23] and applied more recently in order to explain unusual carrier mobilities in GaN with high dislocation density [24, 25]. Based on the density and geometry of the ITBs, estimated by TEM measurements, authors could quantitatively confirm that mobility reduction and collapse, as well as part of the compensation, are due to the presence of twin lamellae. Fig. 1.4 shows the presence and distribution of such defects in layers homoepitaxially grown on substrates with different mis-cuts between 0.1 degrees and 4 degrees. It is apparent that only substrates with relatively large misorientation yield layers free of laminar twins. On the other hand, one should

Progress in MOVPE growth of Ga2O3

9

Fig. 1.4 TEM study of twin lamellae in homoepitaxial (100) β-Ga2O3: (A) substrate with offorientation of 0.1 degrees; (B) of 0.2 degrees; (C) of 2.0 degrees; and (D) of 4.0 degrees. From R. Schewski et al., Evolution of planar defects during homoepitaxial growth of β-Ga2O3 layers on (100) substrates—a quantitative model, J. Appl. Phys. 120 (2016) 225308.

remember that the terrace width changes with the misorientation, that is, it is larger for small angles, which enhances the probability of two-dimensional (2D) island nucleation, as schematically reported in Fig. 1.5. The way Ga atoms may stick to the surface is clearly twofold; therefore, if they do not reach the step, due to large terraces or reduced diffusivity, the chance of laminar twinning becomes high. Based on this model, the authors concluded that homoepitaxial growth by MOVPE on (100) β-Ga2O3 is very challenging. The stacking faults, in form of laminar twins, are characterized by a c/2 glide reflection as twin relation. Coalescence of twinned and epitaxial 2D islands results in the formation of ITBs that actually present electrically active dangling bonds, and thus compensate and reduce the mobility of charge carriers. The limited surface diffusion of the Ga-adatoms and the small substrate miscut-angle are responsible for the high density of stacking-faults experimentally observed in homoepitaxial films. Therefore, one should look for growth conditions that yield step-flow. Possible approaches toward twin-free epilayers include higher growth temperatures, application of surfactants in order to enhance the species diffusion on the surface or reduction of terrace width by adopting substrates with high miscut-angles of about 6 degrees. The diffusion coefficient of Ga at the gas–solid interface at the growth temperature of 850°C was estimated to be 7  109 cm2 s1, that is, two orders of magnitude lower than the one experimentally found in GaAs but 6 orders of magnitude higher than for cubic GaN.

10

Gallium Oxide

Fig. 1.5 Schematic of monolayer growth on a terraced (100) surface of β-Ga2O3. (A) Stick-andball model illustrating two possible positions of Ga on the (100); dark green and light green balls correspond to tetrahedral (Ga I) and octahedral (Ga II) bound Gallium atoms, respectively. The little red, orange, and yellow balls correspond to different oxygen sites OI, OII, and OIII. (B) Schematic diffusion model showing step-flow growth or island nucleation in dependence of terrace width and/or atom surface diffusion coefficient. From R. Schewski et al., Evolution of planar defects during homoepitaxial growth of β-Ga2O3 layers on (100) substrates—a quantitative model, J. Appl. Phys. 120 (2016) 225308.

In another study on homo- and heteroepitaxial β-Ga2O3, it was pointed out that in spite of an heavy Si-doping the films were found to be electrically insulating [26]. The positron annihilation study of these samples evidenced a relatively high Ga vacancy concentration (>5  1017 cm3), higher than the silicon concentration of 1017 cm3. Theoretical calculations predict that the VGa should be in a negative charge state when Fermi levels is located in the upper half of the band gap, hence compensating for n-type doping [27]. The high resistivity of those Si-doped β-Ga2O3 epilayers was therefore attributed to the experimentally observed Ga vacancy concentrations in films.

Progress in MOVPE growth of Ga2O3

1.3

11

Heteroepitaxial deposition of β-Ga2O3

As for homoepitaxial Ga2O3, also in the case of heteroepitaxy, most work has till now been done with plasma-assisted MBE and relatively little with MOVPE. (0001)Oriented sapphire wafers were by far the most used hetero-substrates [10, 14, 28–31], although attempts were made also on GaAs [15], MgAl2O4 [12], and MgO [32] substrates. In Ref. [14], MOVPE heteroepitaxy of β-Ga2O3 was performed on both a-plane and c-plane sapphire substrates using a TMGa precursor (22.5 Lmol/min) and approximately 5:1 O2/TMGa flow rate ratio. A pressure level of 6.67 kPa (0.066 atm, 50 Torr) at a temperature range of 600–800°C was used to obtain crystalline films in the preferred phase. The as-grown films were characterized by high-resolution XRD and Raman spectroscopy. According to Raman investigation, the films were not phase pure but probably contained some α-Ga2O3 inclusions as well as other Ga-related phases. It is to be recalled here that small α-phase inclusions were usually detected in heteroepitaxial Ga2O3 by high-resolution TEM whichever the deposition technique [33]. Fig. 1.6 shows the diffraction pattern corresponding to epilayers on both faces of sapphire. Clear reflections of β-Ga2O3 with (201) orientation were observed for films grown on c-plane sapphire (Fig. 1.6A); while for growth on a-face sapphire other peaks with different orientations were also observed, suggesting a less oriented film. A peak was observed near 36 degrees that could be attributed to (101) α-Ga2O3 in the film. Least-squares refinement was performed to obtain the lattice constant values of the films grown on c-plane and a-plane substrates. For the c-plane sample, the lattice ˚ , b ¼ 3.0315(2) A ˚ , and c ¼ 5.8237(10) A ˚ , and for the constants were a ¼ 12.2648(20) A ˚ , b ¼ 3.0384(7) A ˚ , and c ¼ 5.8315(6) A ˚. a-plane sample they were a ¼ 12.2212(14) A The full-width at half-maximum (FWHM) from the rocking curve of the (402) reflection was 2.42 degrees, that is, higher than typical FWHM values obtained for homoepitaxial β-Ga2O3 [34]. This result indicates that MOVPE on foreign substrates requires significant development efforts to improve the Ga2O3 crystallographic quality. The X-ray reflectivity data reported in Ref. [14] were fitted by authors to get an estimate of the heteroepitaxial film density. This was found to be about 5.6 g/cm3, which is lower than reported theoretical values (6.44 g/cm3), hence suggesting the presence of a large concentration of vacancies in the film. In a study aiming at assessing the Sn doping in homo- and heteroepitaxial β-Ga2O3, Gogova et al. [16] reported that MOVPE deposition of (0001) sapphire always results in (201)-oriented monoclinic films. The X-ray peaks positioned at 2θ ¼ 18.9, 37.8, 58.7, and 80.7 degrees were assigned to the 201, 402, 603, and 804 Bragg reflections of the β-phase. No additional Bragg peaks from the layer were observed proving evidence that the layers contained only β-phase. It must however be noted that (201)-oriented β-Ga2O3 layers indeed grow epitaxially on sapphire but are not single crystalline since there is an in-plane rotation of 60 degrees between grains due to the difference in the crystal symmetry of the substrate (hexagonal) and the epilayer (monoclinic). Such grains were previously observed by Nakagomi et al. [35] and Chen et al. [30] and also detected by TEM investigation. In Ref. [35], thin films were

12

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Fig. 1.6 X-ray diffraction pattern of β-Ga2O3 deposited on c-oriented (A) and a-oriented sapphire (B). From M.J. Tadjer, et al., Structural, optical, and electrical characterization of monoclinic β-Ga2O3 grown by MOVPE on sapphire substrates, J. Electron. Mater. 45 (4) (2016) 2031–2037.

Progress in MOVPE growth of Ga2O3

13

deposited via reaction of Ga vapor with an oxygen plasma on c- and a-oriented sapphire. The authors concluded that both films on the (001) c-plane or (110) a-plane sapphire are (201) oriented. β-Ga2O3 films formed on the (001) c-plane sapphire consist of six types of β-Ga2O3 crystal rotated every 60 degrees from the [110] direction. Similarly, β-Ga2O3 thin films on the (110) a-plane sapphire consist of six types of β-Ga2O3 crystals rotated by multiples of 60 degrees from the [110] or [110] direction. This crystal orientation of β-Ga2O3 formed on the c-plane or the a-plane sapphire depends on the oxygen atom arrangement on the substrate surface. While some of the aspects of the atom arrangement on the (110) a-plane sapphire were discussed previously, more studies are needed to understand the general arrangement in detail. Fig. 1.7A provides one example of the β-Ga2O3 arrangement on c-oriented sapphire, and well explains the origin of domains in the heteroepitaxially grown films. In Ref. [30], β-Ga2O3 films were grown on c-plane sapphire by MOVPE. Using high-resolution XRD θ  2θ and Φ-scan measurements, the Bragg angles of β-Ga2O3 (201), (401), (111), and (111) planes were measured in order to estimate the lattice constants a, b, c, and angle β of the monoclinic film. Comparing with the standard value reported for β-Ga2O3 single crystals, the values of a and c appear to be smaller, while the angle β becomes larger, only the value of b was practically unvaried. According to the epitaxial relationship between β-Ga2O3 film and c-plane sapphire substrate, one is led to conclude that β-Ga2O3 lattice would suffer compressive strain along [102] direction. The lattice mismatch and the six different types of β-Ga2O3 domains on c-oriented Al2O3 are therefore a major obstacle toward high-quality β-Ga2O3 films on sapphire. As shown in Fig. 1.7B, indeed there are three different types of β-Ga2O3 crystal grains plus another three types that originate from the rotation of 180 degrees of the diagram. The β-Ga2O3 (201) plane is parallel to the (0001) sapphire surface, while the β-Ga2O3 [102] and [010] directions run parallel to the Al2O3 [110] and [110] directions, respectively. It is to be observed that both cited works assume that the (201) plane of β-Ga2O3 lays parallel to the c-surface and that there are six in-plane domains resulting from 60 degrees rotations about the [201] axis of gallium oxide. However, the description of the gallium oxide-to-sapphire arrangement differs by 60 degrees, as it becomes evident by comparing the two schematics of Fig. 1.7. This may be due to the different deposition methods or to a misinterpretation of the XRD data. Unfortunately, there are no additional and complete studies to make clear how the two lattices actually position with respect to each other. Alema et al. [28] succeeded to dramatically increase the β-Ga2O3 growth rate by applying a MOVPE close-coupled showerhead reactor, where the substrate-toshowerhead separation was about 10 mm. The growth was carried out on (0001) sapphire (2 in.) substrates using solid Ga(DPM)3 (DPM ¼ dipivaloylmethanate), and liquid TEGa as well as TMGa metal organic (MO) precursors as the sources for Ga. The film growth was conducted from the respective sources by sublimating Ga(DPM)3 at 155°C or evaporating the TEGa and TMGa sources, respectively, at 22°C and 5°C. The vaporized Ga sources were carried into the reactor by Ar gas. A pure molecular oxygen (5 N) was separately injected into the reactor to mix with the TEGa, TMGa, and Ga(DPM)3 precursors in the proximity of the substrate for oxidation with oxygen flow rates of 300, 800, and 800 sccm, respectively. β-Ga2O3 thin film growth was performed at a chamber pressure of 30–120 Torr and substrate temperatures between 750°C and 950°C.

Fig. 1.7 (A) Arrangement of monoclinic Ga2O3 lattice on the hexagonal c-oriented sapphire substrate according to Ref. [35]. (B) Origin of domains in heteroepitaxially grown (201)oriented β-Ga2O3 layers according to Ref. [30].

Progress in MOVPE growth of Ga2O3

15

All Ga2O3 thin films grown from each MO sources were visibly smooth and transparent, but electrically highly resistive (sheet resistance >109 Ω/square), as no measurable conductivity was detected by the Hall effect technique. All deposited layers were (201) oriented and exhibited good crystallographic characteristics as proved by narrow peaks in positions overlapping those of high-quality bulk Ga2O3 (see Fig. 1.8). Unfortunately, authors did not perform a study about the actual epitaxial relationships between sapphire and β-Ga2O3 film, which would have confirmed one of the two options discussed above. It is however remarkable that the close-coupled reactor allows for a very high deposition rate, which may reach 10 μm/h, that is about one order of magnitude higher than for standard horizontal or vertical reactors. Fig. 1.9 reports a comparison between growth rates as a function of chamber backpressure for standard vertical reactors (with either TMGa or TEGa as gallium precursors) and the close-coupled reactor used by the authors: the difference is apparent. A maximum growth rate of about 10 μm/h was achieved using the TMGa precursor and a substrate temperature of 900°C. This successful demonstration of high growth rate for β-Ga2O3 films using various Ga-precursors was achieved thanks to the separate injection of oxygen and MO precursors, as well as the close-coupled showerhead design of the MOVPE reactor which prevents the premature oxidation of the MO precursors. These results are particularly interesting as they offer a solution to the long-standing challenge of realizing fast growth rate of Ga2O3 by MOVPE. Actually, so far only H-VPE proved to be able to grow monoclinic Ga2O3 at a rate of several μm/h [36, 37].

1.4

Heteroepitaxial deposition of ε-Ga2O3

(–603)

Δ

Δ Sapphire (–402)

Fig. 1.8 XRD profiles of bulk (201)oriented monoclinic Ga2O3 (a), and epitaxial films grown by close-coupled MOVPE using Ga(DPM)3 (b), TEGa (c), and TMGa (d) as gallium precursors. The substrate was c-oriented sapphire, kept at 900°C, 940°C, and 900°C for the samples (b), (c), and (d), respectively. From F. Alema, B. Hertog, A. Osinsky, P. Mukhopadhyay, M. Toporkov, W.V. Schoenfeld, Fast growth rate of epitaxial β-Ga2O3 by close coupled showerhead MOCVD, J. Cryst. Growth 475 (2017) 77–82.

(–201)

In the two previous sections, we discussed the recent advances in homo-and heteroepitaxial β-Ga2O3 grown by MOVPE. From one side, it is clear that no truly single crystalline films can be achieved by hetero-epitaxy and that β-Ga2O3 on

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Chamber pressure (Torr) Fig. 1.9 Growth rate as a function of reactor pressure for β-Ga2O3 film grown using TEGa (black line and dots) or TMGa (red line and dots) in a standard vertical MOVPE reactor, with molar flow rates of 103 and 75 μmol/min, respectively. For both cases, the O2 flow rate was 1800 sccm and the substrate temperature was 750°C. The blue dots and line correspond to growth at 750 degrees in a close-coupled showerhead reactor, with O2 flow rate of 300 sccm and TEGa molar flow rate of 50 and 103 μmol/min. Note the much higher growth rate in the latter case. From F. Alema, B. Hertog, A. Osinsky, P. Mukhopadhyay, M. Toporkov, W.V. Schoenfeld, Fast growth rate of epitaxial β-Ga2O3 by close coupled showerhead MOCVD, J. Cryst. Growth 475 (2017) 77–82.

sapphire always contains rotational domains. From the other side, it appears that even homoepitaxial films present a large density of structural defects (stacking faults), which may prevent an effective n-type doping. To prevent the occurrence of such defects, one has to apply tricky growth procedures such as using highly misoriented homo-substrates. Furthermore, monoclinic Ga2O3 exhibits a certain anisotropy of physical properties and a strong tendency to cleavage. The intrinsic drawbacks of monoclinic β-Ga2O3 have triggered a number of investigations on alternative polymorphs. Among them, the hexagonal α and the orthorhombic ε have attracted attention because of their higher crystallographic symmetry, better matching to Al2O3 substrates and acceptable thermodynamic stability. The α-phase Ga2O3 has been extensively studied by the group at the University of Kyoto [38, 39] and is the object of a specific chapter of this volume, therefore in the following a thorough discussion about growth and properties of the less known ε-Ga2O3 will be reported. Among the polymorphs, the ε-phase is particularly interesting because of its higher symmetry and milder epitaxial growth conditions with respect to the more popular β. Furthermore, it shows a favorable matching to commercial sapphire as well as to other hexagonal or pseudo-hexagonal substrates [40–42]. Theoretical studies aiming at establishing structure and formation energy, i.e., thermodynamic stability, of the different polymorphs [43] suggested that this polymorph is orthorhombic, although initial

Progress in MOVPE growth of Ga2O3

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Fig. 1.10 SEM image of the nucleation islands at the surface of an ε-Ga2O3 layer. From F. Boschi, M. Bosi, T. Berzina, E. Buffagni, C. Ferrari, R. Fornari, Hetero-epitaxy of ε-Ga2O3 layers by MOCVD and ALD, J. Cryst. Growth 443 (2016) 25–30.

experimental studies considered it to be hexagonal with P63mc space group symmetry [40, 44, 45]. This was because the nucleation islands and the XRD patterns showed an evident hexagonal symmetry (see for example Fig.1.10). In a later work based on detailed TEM observations [46], the real structure of the ε-phase was found to contain 5–10 nm large 120 degrees-rotational twin domains, each of them with orthorhombic structure and Pna21 space group symmetry, analogous to κ-Al2O3. The simplest way to look at the crystallographic structure ε-Ga2O3 is to consider that it consists of a 4H ABAC oxygen close-packed stacking, where Ga atoms in between occupy octahedral and tetrahedral sites forming two types of polyhedral layers parallel to (001), as schematically represented in Fig. 1.11. The high-resolution TEM investigation of ε-Ga2O3 film cross section showed that they possess a columnar structure (see Fig. 1.12) and also showed that epitaxial islands of γ-Ga2O3 were present at the interface between the c-oriented sapphire and the Ga2O3 film [46]. The occurrence of the γ-phase at the interface is an unexpected result, and its presence was evident only after highresolution TEM investigations. Although the role of the γ-phase may be important for the ultimate growth of ε-Ga2O3 on α-Al2O3 at the used temperature, its origin and evolution are still unclear. The plane-view investigation of the same specimen clearly showed that the columnar structure corresponds to domains with lateral extension of 5–10 nm (see Fig. 1.13). Based on these results, one can conclude that in the ε-phase the Ga is ordered at the nanoscale, and that the crystal structure of each Ga2O3 domain is analogous to the kappa-alumina (κ-Al2O3). The edge-sharing octahedra and the corner-sharing tetrahedra form zig-zag ribbons along the [100] direction, which ultimately give rise to twins that separate the domains [46]. These twins are schematically depicted in Fig. 1.14, along with another typical defect found in such epilayers, namely, the antiphase

A B A C c

A

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Fig. 1.11 Schematic of the ε-Ga2O3 lattice cross section (the vertical direction is c). The red dots indicate the (001) oxygen planes; between them are alternate layers of Ga in purely octahedral and octahedral/tetrahedral positions. In order to maintain the stoichiometric ratio, in each layer only 2/3 of the possible Ga sites are actually occupied. Here three examples of possible random distribution of tetrahedra and octahedra are shown. From F. Mezzadri, G. Calestani, F. Boschi, D. Delmonte, M. Bosi, R. Fornari, Crystal structure and ferroelectric properties of ε-Ga2O3 films grown on (0001)-sapphire, Inorg. Chem. 55 (2016) 12079–12084.

[100] a-Al2O3 [100] Ga2O3

006 ¯ 014 012

002 060

Fig. 1.12 Bright-field (BF) image of the vertical cross-section of epitaxial Ga2O3 on α-Al2O3 (top), selected-area electron diffraction (SAED) pattern (bottom right) and high-resolutionTEM image with the corresponding indexed fast Fourier transforms (A and B, bottom left). At the film-substrate interface a secondary phase could be identified as γ-Ga2O3 (region and FFT marked B). Note that ε-Ga2O3 shows a columnar texture; the corresponding reflections could be indexed taking an orthorhombic unit cell (region A). From I. Cora et al., The real structure of ε-Ga2O3 and its relations to κ-phase, CrystEngComm 11 (2017) 1509–1516.

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19

Fig. 1.13 (A) Plane-view high-resolution TEM image of a sample in [001] projection; (B–E) corresponding fast Fourier transforms (FFTs) from 5 nm  5 nm large areas. Domains have orthorhombic structure (see B–D) and are rotated by 120 degrees with respect to each other. FFT of the whole image (F) shows a pseudo-hexagonal symmetry. From I. Cora, et al., The real structure of ε-Ga2O3 and its relation to κ-phase, CrystEngComm 11 (2017) 1509–1516.

boundaries. A detailed XRD analysis carried out on small ε-Ga2O3 fragments obtained ˚, from very thick layers provided the following lattice constants: a ¼ 5.0463(15) A ˚ ˚ b ¼ 8.7020(9) A, and c ¼ 9.2833(16) A. Furthermore, the structure of the ε-phase was confirmed to be noncentrosymmetric, in agreement with the characteristics of the Pna21 space group. The positive and negative charges do not mutually compensate along the z-direction, giving rise to nonzero electrical dipoles. Based on the structural information, this phase should be pyroelectric with an estimated electrical polarization of 0.18 μC/cm2. Additional dynamic hysteresis measurements demonstrated that this material is indeed ferroelectric [45]. The optical properties of the ε-phase were studied by Pavesi et al. by means of optical absorption in the far UV range and photoconductivity (PC) [47]. The PC curve of Fig. 1.15 shows a relatively mild onset at about 4.2 eV that becomes sharper at about

20

Gallium Oxide

(A)

(B)

Normalized absorbance/PCS (a.u.)

Fig. 1.14 (A) Polyhedral structural model of the 120 degrees rotational twin boundary. (B) Polyhedral structural model of antiphase boundary inside the domains. From I. Cora, et al., The real structure of ε-Ga2O3 and its relation to κ-phase, CrystEngComm 11 (2017) 1509–1516. Fig. 1.15 Spectral photocurrent and optical absorption edge of ε-Ga2O3 films on c-oriented sapphire (normalized curves). From M. Pavesi, et al., ε-Ga2O3 epilayers as a material for solarblind UV photodetectors, Mater. Chem. Phys. 205 (2018) 502–507.

1.0 0.8 Absorbance PCS

0.6 0.4 0.2 0.0 2.0

2.5

3.0

3.5

4.0

4.5

5.0

5.5

6.0

Energy (eV)

4.6 eV. The onset of PC occurs well below the estimated bandgap and is ascribed to the presence of energy levels connected with point defects or complexes, as already observed for β-Ga2O3. The optical absorption spectrum, recorded at room temperature, looks different as a nonnegligible absorption is present already at 2 eV, and it

Progress in MOVPE growth of Ga2O3

21

increases monotonically up to band edge. The long tail, extending from the absorption edge down to 2.1 eV, may include other components in addition to excitation of photocarriers from deep levels, like for example light absorption by defects at substrate/ film interface. Photocurrent values in the same spectral region, far from bandgap edge, are too low to be appreciable in the same linear plot of Fig. 1.15. It is interesting to note that the photogenerated current is low even when the absorbance is nonnegligible. Considering that the bandgap width is a most crucial parameter for both electronic and opto-electronic applications, authors made an assessment of the bandgap of ε-Ga2O3 by carrying out absorbance measurements on a large set of films with thicknesses in the range 175 nm to 2 μm. In all cases, the onset due to band-band absorption started at about 4.6 eV and the measured absorption coefficients were in good agreement with those extrapolated from the data of Oshima et al. [42]. The analysis based on the set of films with variable thickness showed that the optical bandgap of all tested ε-Ga2O3 films converges between 4.6 and 4.7 eV, hence close to that of β-Ga2O3. However, extensive theoretical and experimental investigations on the band structure of the ε polymorph are still necessary in order to clarify whether the band is direct or indirect. The as-grown ε-Ga2O3 always exhibited a very high resistivity, larger than 106 Ωcm. Measurements of the dark current as a function of temperature in the range 20–300°C showed that the film conductance increases by three orders of magnitude with the temperature. The corresponding Arrhenius plot was linear and presented an activation energy of 0.695 eV. This Arrhenius slope is probably related to the thermal ionization of deep electronic levels located at about 0.7 eV below the conduction band. Similar electronic states were previously detected by DLTS in β-Ga2O3 single crystals [48], although their nature was not clarified. Cathodo-luminescence spectroscopy measurements on ε-Ga2O3 films gave no band–band or excitonic emission, which is common also for other Ga2O3 polymorphs. A broad, structured emission band peaked at about 2.6 eV was generally found. The Gaussian deconvolution of the spectra provided three narrower bands centered at 2.34 eV (530 nm), 2.67 eV (464 nm), and 2.75 eV (450 nm), respectively. The origin of these emission bands is not yet known, but the comparison with literature data for the more studied β-Ga2O3 suggests that the three main bands could be connected with intra-gap donor-acceptor pairs due to point defects [49, 50]. The full set of experimental results (dark current activation energy, tail of optical absorption, and CL) is quite complex. However, it is possible to reconcile them assuming the presence within the forbidden bandgap of a family of deep donors located around 0.7 eV below the conduction band, which accounts for the dark current increase with temperature, plus three deeper levels below midgap that account for the CL radiative emission, and the tails of the absorption and PC spectra. The thermal stability of ε-Ga2O3 is an important issue. Preliminary annealing experiments of MOCVD films suggested that this polymorph is thermally stable at least up to 800°C [41]. More recently, the thermal stability of the ε-phase was accurately investigated by complementary methods after thermal annealing at temperatures in the range 700–1000°C [51]. XRD, TEM, and differential scanning calorimetry (DSC) were used to assess the phase transformations following

DSC (mW)

22

Gallium Oxide Onset: 622°C

↑exo

Peak: 642°C

Inflection: 824°C

10 K/min

Inflection: 891°C Area: 550 mJ

1st derivative

10

6 Inflection: 799°C

5 K/min

Inflection: 879°C

2 200

300

400

500

600

700

800

900 Temperature (°C)

Fig. 1.16 DSC curves of two ε-Ga2O3 layers on sapphire between 200°C and 1000°C, recorded at different heating rates (10 or 5 K/min). The first derivative (DDSC) of the 10 K/min is also shown as a dashed line. Note the mild endothermic change above 600°C whose cause is still unknown, and the sharp exothermic peak at 880–900°C corresponding to the ε- to β-phase transition. From R. Fornari, et al., Thermal stability of ε-Ga2O3 polymorph, Acta Mater. 140 (2017) 411–416.

the high-temperature treatment. The results clearly indicate that ε-Ga2O3 initiates modifying its crystallographic structure above 650°C as demonstrated by a mild endothermic bent of the DSC curves. However, the effective transition to β-phase occurs quite suddenly at 880–900°C, depending of the DSC heating rate (see Fig. 1.16). XRD and TEM results confirm this evidence. TEM in particular shows that after annealing at 1000°C and rapid cooling the film is uniquely made of β-Ga2O3 grains, most of them with orientation (310) with respect to the sapphire substrate [51]. However, if the cooling rate is substantially reduced, the converted β-Ga2O3 layer tends to assume the standard orientation (201) jj to (0001) of the substrate. The authors of this study concluded that ε-Ga2O3 may actually be used for device fabrication provided that all fabrication steps are carried out at temperatures below 700°C. After this survey of the physical and thermodynamic properties of the ε-phase, let us turn the attention to the epitaxial deposition of such material. As already mentioned, the first report about epitaxial ε-Ga2O3 was presented in 2015 [42] and regarded the H-VPE growth at 550°C on GaN (0001), AlN (0001), and β-Ga2O3 (201) by using gallium chloride and pure O2 as precursors. A growth rate as high as 20 μm/h was achieved in this way. X-ray measurements and polar maps proved that the ε-Ga2O3 films were grown epitaxially on the three kinds of substrates, and mostly (0001)oriented, although some misoriented grains were also observed. The optical bandgap of this HVPE ε-Ga2O3 was determined to be 4.9 eV. Authors applied X-ray diffraction to find out the epitaxial relationships for planes and directions, which was seen to be: ε-Ga2O3 (0001)//GaN (0001) and ε-Ga2O3 [1010]//GaN [1010]. The relations to

Progress in MOVPE growth of Ga2O3

23

AlN (0001) templates were the same, although the in-plane lattice mismatches are quite different, respectively, 8.8% and 6.6% with GaN (0001) and AlN(0001). Differently from films on GaN and AlN substrates, additional diffraction peaks of ε-Ga2O3 were detected probably due to the larger volume fraction of misoriented threedimensional (3D)-grains. The epitaxial relationships between the c-plane ε-Ga2O3 and the substrate were found to be ε-Ga2O3 (0001)//β-Ga2O3 (201) and ε-Ga2O3 [1010]//β-Ga2O3 [102]. This reproduces the results observed in the case of the reciprocal growth of hexagonal GaN on the β-Ga2O3 [52, 53]. Further, the in-plane atomic arrangement of Ga or O in both ε-Ga2O3 (0001) and β-Ga2O3 (201) basically follow triangular lattices, and the mismatch is as small as 1.1%. The authors of Ref. [42] claimed that the use of c-oriented GaN and AlN substrates, whose space group is the same as that of ε-Ga2O3, is the first condition to deposit ε-Ga2O3 epitaxial layers, the second being the relatively low deposition temperature. It was thanks to the deposition temperature much lower than for standard CVD growth of oxides that ε-Ga2O3 could be heteroepitaxially grown even on monoclinic β-Ga2O3 (201), eventually with smaller FWHM than in films grown on GaN and AlN. It is worth noting that authors of [42] still defined the crystal structure of ε-Ga2O3 as hexagonal with space group P63mc, by virtue of the misleading sixfold symmetry observed in growth islands and X-ray maps. The first MOVPE growth of ε-Ga2O3 was reported in 2016 by Boschi et al. [40]. The epilayers were deposited on: (i) c-oriented sapphire, (ii) c-oriented GaN templates, (iii) (111)- and (001)-oriented 3C-SiC. TMGa and ultrapure water were used as precursors. Ga2O3 layers were deposited within the temperature range 450–715°C, at a fixed pressure of 100 mbar. Water and TMGa were delivered to the growth chamber by means of two separate gas lines, using palladium-purified H2 as carrier. TMGa was delivered through a standard double-dilution line, in order to operate in waterexcess conditions: precursors flows were set to get a H2O-to-TMGa partial pressure ratio of about 103. For the operation conditions of 650°C and 100 mbar, employed for most of the MOVPE experiments, as discussed below, the growth rate was about 500 nm/h. For growth temperature of 550°C or lower, no crystalline layers were obtained. Despite their very smooth surfaces, the deposited films were invariably amorphous, and X-ray profile scans did not exhibit any peaks related to whichever crystalline phase of gallium oxide, but only typical α-Al2O3 reflections from the substrate (see the upper X-ray scan of Fig. 1.17B). For deposition temperature of 715°C, weak peaks corresponding to (201) planes of β-Ga2O3 were observed, in addition to those of sapphire (see lower X-ray scan of Fig. 1.17B). The lattice parameters extracted from these XRD profiles matched well ˚ , b ¼ 3.04 A ˚, those of the (201) plane of bulk material [1, 25], namely, a ¼ 12.23 A ˚ c ¼ 5.80 A, and angle β ¼ 103.7 degrees. The peaks were quite broad and had low intensity, thus indicating a poor crystalline quality, also reflected by an irregular surface morphology. This result was not surprising, as the crystalline β-phase is usually obtained at much higher temperatures and often requires some technological measures, such as pretreatment of the substrate to inhibit the formation of planar domains and improve the overall crystalline quality [11, 54].

24

Gallium Oxide

550°C

4

10

3

10

Sapphire 0012

5

10

Sapphire 006

550°C

2

10

1 µm

1

2

10

ε-Ga2O3 010

3

10

650°C

ε-Ga2O3 008

4

10

ε-Ga2O3 006

5

10

ε-Ga2O3 004

650°C

ε-Ga2O3 002

Intensity (cps)

10

1

2

10

β-Ga2O3–1005

3

10

β-Ga2O3–804 ?

715°C

715°C

β-Ga2O3–402

4

10

β-Ga2O3–201

5

10

β-Ga2O3–603

10

1 µm

1

10

10

1 µm

(B)

20

30

40

50

60

70

80

90 100 110 120

2θ (degrees)

(A) Fig. 1.17 Ga2O3 samples grown by MOVPE at different temperatures on c-plane sapphire: (A) SEM images, and (B) corresponding high-resolution XRD scan profiles. From F. Boschi, M. Bosi, T. Berzina, E. Buffagni, C. Ferrari, R. Fornari, Hetero-epitaxy of ε-Ga2O3 layers by MOCVD and ALD, J. Cryst. Growth 443 (2016) 25–30.

Single-phase ε-Ga2O3 was obtained on c-sapphire only when growing at intermediate temperatures. Films deposited at 650°C were indeed smooth, with very intense and narrow X-ray diffraction peaks corresponding to (002) planes of ε-Ga2O3 (see intermediate X-ray scan in Fig. 1.17B), and roughness below 2 nm. X-ray polar maps showed that the Ga2O3 lattice is rotated by 30 degrees about the c-axis, with respect to the sapphire lattice. Thanks to this layer-substrate arrangement, the effective lattice mismatch reduces to +4.8% (compressive stress) between ε-Ga2O3 and (0001) sapphire. This situation resembles the one always found in the case of GaN epitaxy on c-oriented sapphire. Even though the presence of a substrate with an in-plane hexagonal arrangement of the atoms appeared to be more favorable for nucleation and coalescence of ε-Ga2O3 hexagonal islands, the results obtained on different substrates clearly showed that the growth temperature is the crucial factor that ultimately decides the crystal structure of Ga2O3 epilayers. The best ε-Ga2O3 films were grown at 650°C on c-oriented GaN templates; however, films of relatively high quality and acceptable morphology were deposited even on (111) 3C-SiC template, despite the cubic symmetry of this material. This can be explained by taking into consideration the hexagonal arrangement

Progress in MOVPE growth of Ga2O3

25

(pseudo-ternary symmetry) of the atoms on this specific plane of 3C-SiC. On the contrary, when growth was performed on the (001) plane of 3C-SiC, the morphological and structural disorder of the films increased dramatically. Nevertheless, it must be noted that under the growth conditions described above (moderate temperature, high water-to-TMGa ratio, low reactor pressure), the ε-phase remained dominant, as demonstrated by XRD and SEM analyses [40]. These experimental observations demonstrate that Ga2O3 films tend to grow as ε-phase when the deposition temperature remains below a limit of about 650°C, in that confirming that it is the second most stable phase after the β one. In the study of Xia et al. [41], ε-Ga2O3 films were deposited on 6H-SiC substrates by MOCVD at 500°C. The 6H-SiC was selected for its good electrical and thermal conductivity, but also because it has a smaller a-axis lattice mismatch (3.3%) with respect to ε-Ga2O3 films, compared with other commonly used hexagonal substrates. High-purity O2 gas and TEGa (kept at 5°C and 800 Torr) were used as oxygen and gallium sources, respectively. Authors decided to use TEGa because of its relatively low pyrolysis temperature and lower pre-reaction intensity with oxygen with respect to common TMGa sources. High-purity Ar gas was used as the carrier gas for TEGa. Both O2 and Ar flows were kept at 50 sccm. The [O2]/[TEGa] molar ratio was calculated to be 390, while the pressure in the growth chamber was maintained at 3500 Pa. The chamber pressure was defined as critical because above 5000 Pa pre-reaction between the precursors took place and nucleation in gas phase occurred, which resulted in formation of β-Ga2O3. It is actually surprising that ε-Ga2O3 (or even β-Ga2O3) could be obtained at a reactor temperature of 500°C. The question of phase formation in MOVPE deposition of Ga2O3 is still open and there seem to be some uncertainty about the conditions that may grant either pure ε- or β-Ga2O3. In a recent study, Zhuo et al. attempted to build a phase diagram for the MOCVD growth of Ga2O3 and concluded that the stabilization of metastable ε-Ga2O3 is primarily controlled by the growth temperature and VI/III ratio [55]. They performed the growth of a number of Ga2O3 films on c-plane sapphire substrates, utilizing TEGa and high-purity oxygen (O2) as precursors, and argon as carrier gas. During the growth, the chamber pressure was maintained at 9.1 Torr, but the flow rates were varied over a wide range: the flow rate of TEGa between 22.3 and 67.4 μmol/ min and that of O2 between 13.4 and 120.5 mmol/min, respectively. The growth temperature ranged between 450°C and 570°C. For growth temperatures above 535°C, and flow rate of TEGa and VI/III ratio set at 67.4 μmol/min and 596, respectively, a β-Ga2O3 polycrystalline thin film was obtained, as indicated by X-ray diffraction. Furthermore, SEM analysis showed that films were polycrystalline with very fine grain structures. Upon decreasing the growth temperature to 505°C, the films exhibited new diffraction peaks that were indexed as 0002, 0004, and 0006 diffractions of ε-Ga2O3 (again considered as hexagonal, in the following the peak notations of the original Ref. [55] are maintained for enabling readers to follow text and figures of the cited work). Therefore, the temperature decrease resulted into a mixture of ε- and β-Ga2O3. The full-width at half-maximum (FWHM) value of the diffraction peak of β-Ga2O3 (0.678 degrees) was more than three times larger than that of the ε-0006 diffraction peak (0.200 degrees), implying

26

Gallium Oxide

that the ε-Ga2O3 crystal fraction had better quality than β-Ga2O3 for those given growth conditions. Consistently with the X-ray results, a remarkable change in surface morphology was detected by SEM: small β-Ga2O3 was distributed among larger ε-Ga2O3 grains with (0001) orientation and flat tops. Authors found that metastable ε-Ga2O3 tends to stabilize at lower growth temperature. By keeping a constant TEGa flow and changing the O2 flow rates, authors also investigated the effect of the VI/III ratio. For VI/III ratios of 993 at 505°C, no ε-Ga2O3 was detected (see Fig. 1.18), and the samples were polycrystalline β, while in samples grown at various temperatures for a constant VI/III ratio of 1788 the ε-phase emerged again. According to Ref. [55], the nucleation layer was generally constituted of a mixture of β- and ε-Ga2O3 grains, but the ε-phase took progressively over as the growth proceeded up to forming a closed ε layer under certain growth conditions. This is schematically represented in the “phase diagram” of Fig. 1.18: the microstructure and phase of the films were plotted as a function of deposition temperature and VI/III ration, for a constant TEGa flow of 67.4 μmol/min. Lower VI/III ratios seem to be beneficial for the stabilization of ε-Ga2O3, provided that the temperature remains above 500°C. By reducing the growth temperature or VI/III ratio, the nucleation of Ga2O3 transformed from pure β form into a mixture of β- and ε-phases. Mixed-phase Ga2O3 acted as seed layer upon which an evident structural evolution with the layer thickness was observed. Authors concluded that the formation of ε-Ga2O3 is energetically favored during the first step of growth and maintained if growth conditions close to thermodynamic equilibrium are chosen. Actually, they reported that by reducing the growth rate they were able to grow a pure ε epilayer.

570

VI/III ratio: 1788 535°C 505°C

58

(A)

Temperature (°C)

505°C

450°C

60

2q (degrees)

β-Ga2O3

540 510 480

β+ε

450

Microcrystalline

480°C

β−603

Intensity [a.u. (log scale)]

VI/III ratio: 993

500

62

(B)

1500

2500

VI/III ratio

Fig. 1.18 (A) XRD patterns of Ga2O3 samples grown at various temperatures with higher VI/III ratios of 993 and 1788; (B) phase diagram of Ga2O3 films grown at constant TEGa flow rate of 67.4 μmol/min. From Y. Zhuo, Z. Chen, W. Tu, X. Ma, Y. Pei, G. Wang, β-Ga2O3 versus ε-Ga2O3: control of the crystal phase composition of gallium oxide thin film prepared by metal-organic chemical vapor deposition, Appl. Surf. Sci. 420 (2017) 802–807.

Progress in MOVPE growth of Ga2O3

27

By comparing Refs. [40, 41, 55], it appears that the question of phase control in MOVPE growth of Ga2O3 is still rather controversial; there are a number of incongruences regarding the preparation of the ε polymorph, beginning from the suitable temperature window. In very recent publication, it was shown that three different phases of Ga2O3 (α, β, and ε) could be successfully grown on c-plane sapphire by only tuning the flow rate of HCl along with other precursors in an MOCVD reactor [56]. A threefold increase in the growth rate of pure β-Ga2O3 was achieved by just introducing 5 sccm of HCl into reactor flow. With continuously increased HCl flow, a mixture of β- and ε-Ga2O3 was observed, until the Ga2O3 film became a pure ε-Ga2O3 with a smooth surface and the highest growth rate (1 μm/h) at a HCl flow rate of 30 sccm. By further increasing the HCl flow up to 60 sccm, the films tended to have a mixture of α- and ε-Ga2O3 with a dominant α-Ga2O3, while the growth rate dropped significantly (0.4 μm/h). The film appeared to be rough as a result of the mixed phases since the growth rate of ε-Ga2O3 is much higher than α-Ga2O3. In this HClenhanced MOCVD mode, the Cl impurity concentration was practically identical in the investigated samples, despite the different injected HCL amount. Authors calculated (density functional theory) that the relative energy between β-, ε-, and α-Ga2O3 became smaller thus inducing the phase change by increasing the HCl flow in the reactor. They concluded that HCl acted as a catalyst for the phase transformation process, so that the HCl-enhanced MOCVD approach might be the right way to achieve well-controlled heteroepitaxy of Ga2O3 films with different phases. This results confirms the findings of Yao et al. [57] who obtained films of β-Ga2O3 on (0001) sapphire substrates using MOCVD within the temperature range of 650–850°C, but they got the metastable phases α-and ε-Ga2O3 when growing in the same temperature range by HVPE. The authors attributed the HVPE growth of the metastable phases to a combination of high growth rates and low growth temperature. One reason for the nonequilibrium growth during HVPE that is common for all these samples is indeed the significantly faster growth rate, which probably does not enable the adatoms to rearrange and to form the thermodynamically stable β-phase. This would imply that the diffusing adatoms initially bonded to the oxygen and aluminum sites of the (0001) substrate surface, tend to follow the surface crystallography (α-Ga2O3 also possesses the corundum structure), and are unable to rearrange as βGa2O3 before the next layer of adatoms covers the surface. Each layer of α-Ga2O3 is thus constrained or buried under the subsequent layers. However, as also discussed in Ref. [56], it is the presence of Cl that most apparently affects the onset of formation and continued growth of either ε-phase or α-phase. Actually, as revealed by SIMS investigations, all HVPE samples showed a very high Cl concentration at the Ga2O3/Al2O3 interface. The presence of Cl at the outset of growth probably leads to the formation of very stable AldCl bonds at the Al2O3 (0001) surface that ultimately change the interfacial energy between the two oxides and strongly influences the initial film nucleation. The dominant factor (rapid growth rate, similarity in crystal structure and favorable lattice matching, high concentrations of Cl) that controls the formation of the metastable polymorphs however was not determined unequivocally, and further investigations are necessary in order to fully understand the different polymorphism aspects.

28

Gallium Oxide

References [1] R. Roy, V.G. Hill, E.F. Osborn, Polymorphism of Ga2O3 and the system Ga2O3-H2O, J. Am. Chem. Soc. 74 (3) (1952) 719–722. [2] S. Geller, Crystal structure of β-Ga2O3, J. Chem. Phys. 33 (1960) 676–684. [3] M.R. Lorenz, J.F. Woods, R.J. Gambino, Some electrical properties of the semiconductor β-Ga2O3, J. Phys. Chem. Solids 28 (1967) 403–404. [4] Y. Tomm, P. Reiche, D. Klimm, T. Fukuda, Czochralski grown Ga2O3 crystals, J. Cryst. Growth 220 (4) (2000) 510–514. [5] Y. Tomm, J. Ko, A. Yoshikawa, T. Fukuda, Floating zone growth of β-Ga2O3: a new window material for optoelectronic device applications, Sol. Energy Mater. Sol. Cells 66 (1–4) (2001) 369–374. [6] H. Aida, K. Nishiguchi, H. Takeda, N. Aota, K. Sunakawa, Y. Yaguchi, Growth of β-Ga2O3 single crystals by the edge-defined, film fed growth method, Jpn. J. Appl. Phys. 47 (11) (2008) 8506–8509. [7] A. Kuramata, K. Koshi, S. Watanabe, Y. Yamaoka, T. Masui, S. Yamakoshi, High-quality β-Ga2O3 single crystals grown by edge-defined film-fed growth, Jpn. J. Appl. Phys. 55 (12) (2016) 1202A2. [8] Z. Galazka, et al., Czochralski growth and characterization of β-Ga2O3 single crystals, Cryst. Res. Technol. 45 (12) (2010) 1229–1236. [9] Z. Galazka, et al., Scaling-up of bulk β-Ga2O3 single crystals by the Czochralski method, ECS J. Solid State Sci. Technol. 6 (2) (2017) Q3007–Q3011. [10] N.M. Sbrockey, T. Salagaj, E. Coleman, G.S. Tompa, Y. Moon, M.S. Kim, Large-area MOCVD growth of Ga2O3 in a rotating disc reactor, J. Electron. Mater. 44 (5) (2015) 1357–1360. [11] G. Wagner, et al., Homoepitaxial growth of β-Ga2O3 layers by metal-organic vapor phase epitaxy, Phys. Status Solidi Appl. Mater. Sci. 211 (1) (2014) 27–33. [12] W. Mi, J. Ma, C. Luan, H. Xiao, Structural and optical properties of β-Ga2O3 films deposited on MgAl2O4 (100) substrates by metal–organic chemical vapor deposition, J. Lumin. 146 (8) (2014) 1–5. [13] M. Baldini, et al., Semiconducting Sn-doped β-Ga2O3 homoepitaxial layers grown by metal organic vapour-phase epitaxy, J. Mater. Sci. 51 (7) (2016) 3650–3656. [14] M.J. Tadjer, et al., Structural, optical, and electrical characterization of monoclinic β-Ga2O3 grown by MOVPE on sapphire substrates, J. Electron. Mater. 45 (4) (2016) 2031–2037. [15] V. Gottschalch, et al., Growth of β-Ga2O3 on Al2O3 and GaAs using metal-organic vaporphase epitaxy, Phys. Status Solidi 206 (2) (2009) 243–249. [16] D. Gogova, M. Schmidbauer, A. Kwasniewski, Homo- and heteroepitaxial growth of Sn-doped β-Ga2O3 layers by MOVPE, CrystEngComm 17 (35) (2015) 6744–6752. [17] S. Rafique, L. Han, M.J. Tadjer, J.A. Freitas, N.A. Mahadik, H. Zhao, Homoepitaxial growth of β-Ga2O3 thin films by low pressure chemical vapor deposition, Appl. Phys. Lett. 108 (18) (2016) 182105. [18] M. Handwerg, R. Mitdank, Z. Galazka, S.F. Fischer, Temperature-dependent thermal conductivity in Mg-doped and undoped β-Ga2O3 bulk-crystals, Semicond. Sci. Technol. 30 (2) (2015) 24006. [19] M. Baldini, M. Albrecht, A. Fiedler, K. Irmscher, R. Schewski, G. Wagner, Si- and Sn-doped homoepitaxial β-Ga2O3 layers grown by MOVPE on (010)-oriented substrates, ECS J. Solid State Sci. Technol. 6 (2) (2017) Q3040–Q3044.

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[20] K. Sasaki, A. Kuramata, T. Masui, E.G. Vı´llora, K. Shimamura, S. Yamakoshi, Devicequality β-Ga2O3 epitaxial films fabricated by ozone molecular beam epitaxy, Appl. Phys. Express 5 (3) (2012) 35502. [21] A. Fiedler, et al., Influence of incoherent twin boundaries on the electrical properties of β-Ga2O3 layers homoepitaxially grown by metal-organic vapor phase epitaxy, J. Appl. Phys. 122 (16) (2017) 165701. [22] R. Schewski, et al., Evolution of planar defects during homoepitaxial growth of β-Ga2O3 layers on (100) substrates—a quantitative model, J. Appl. Phys. 120 (22) (2016) 225308. [23] W.T. Read, XVI. Scattering of electrons by charged dislocations in semiconductors, London, Edinburgh, Dublin Philos. Mag. J. Sci. 46 (373) (1955) 111–131. [24] D.C. Look, J.R. Sizelove, Dislocation scattering in GaN, Phys. Rev. Lett. 82 (6) (1999) 1237–1240. [25] J.-L. Farvacque, Z. Bougrioua, I. Moerman, Free-carrier mobility in GaN in the presence of dislocation walls, Phys. Rev. B 63 (11) (2001) 115202. [26] E. Korhonen, et al., Electrical compensation by Ga vacancies in Ga2O3 thin films, Appl. Phys. Lett. 106 (24) (2015) 242103. [27] J.B. Varley, H. Peelaers, A. Janotti, C.G. Van de Walle, Hydrogenated cation vacancies in semiconducting oxides, J. Phys. Condens. Matter 23 (33) (2011) 334212. [28] F. Alema, B. Hertog, A. Osinsky, P. Mukhopadhyay, M. Toporkov, W.V. Schoenfeld, Fast growth rate of epitaxial β-Ga2O3 by close coupled showerhead MOCVD, J. Cryst. Growth 475 (2017) 77–82. [29] S. Rafique, et al., Heteroepitaxy of N-type β-Ga2O3 thin films on sapphire substrate by low pressure chemical vapor deposition, Appl. Phys. Lett. 109 (2016) 132103. [30] Y. Chen, et al., The lattice distortion of β-Ga2O3 film grown on c-plane sapphire, J. Mater. Sci. Mater. Electron. 26 (5) (2015) 3231–3235. [31] Y. Lv, J. Ma, W. Mi, C. Luan, Z. Zhu, H. Xiao, Characterization of β-Ga2O3 thin films on sapphire (0001) using metal-organic chemical vapor deposition technique, Vacuum 86 (12) (2012) 1850–1854. [32] L. Kong, J. Ma, C. Luan, Z. Zhu, Structural and optical properties of Ga2O3: in films deposited on MgO (100) substrates by MOCVD, J. Solid State Chem. 184 (8) (2011) 1946–1950. [33] R. Schewski, et al., Epitaxial stabilization of pseudomorphic α-Ga2O3 on sapphire (0001), Appl. Phys. Express 8 (1) (2015) 11101. [34] H. Murakami, et al., Homoepitaxial growth of β-Ga2O3 layers by halide vapor phase epitaxy, Appl. Phys. Express 8 (1) (2015) 15503. [35] S. Nakagomi, Y. Kokubun, Crystal orientation of β-Ga2O3 thin films formed on c-plane and a-plane sapphire substrate, J. Cryst. Growth 349 (1) (2012) 12–18. [36] Q.T. Thieu, et al., Preparation of 2-in.-diameter (001) β-Ga2O3 homoepitaxial wafers by halide vapor phase epitaxy, Jpn. J. Appl. Phys. 56 (11) (2017) 110310. [37] K. Nomura, et al., Thermodynamic study of β-Ga2O3 growth by halide vapor phase epitaxy, J. Cryst. Growth 405 (2014) 19–22. [38] D. Shinohara, S. Fujita, Heteroepitaxy of corundum-structured α-Ga2O3 thin films on α-Al2O3 substrates by ultrasonic mist chemical vapor deposition, Jpn. J. Appl. Phys. 47 (2008) 7311–7313. [39] S. Fujita, K. Kaneko, Epitaxial growth of corundum-structured wide band gap III-oxide semiconductor thin films, J. Cryst. Growth 401 (2014) 588–592. [40] F. Boschi, M. Bosi, T. Berzina, E. Buffagni, C. Ferrari, R. Fornari, Hetero-epitaxy of ε-Ga2O3 layers by MOCVD and ALD, J. Cryst. Growth 443 (2016) 25–30.

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[41] X. Xia, et al., Hexagonal phase-pure wide band gap ε-Ga2O3 films grown on 6H-SiC substrates by metal organic chemical vapor deposition, Appl. Phys. Lett. 108 (20) (2016) 202103. [42] Y. Oshima, E.G. Vı´llora, Y. Matsushita, S. Yamamoto, K. Shimamura, Epitaxial growth of phase-pure ε-Ga2O3 by halide vapor phase epitaxy, J. Appl. Phys. 118 (8) (2015) 85301. [43] S. Yoshioka, H. Hayashi, A. Kuwabara, F. Oba, K. Matsunaga, I. Tanaka, Structures and energetics of Ga2O3 polymorphs, J. Phys. Condens. Matter 19 (34) (2007) 346211. [44] H.Y. Playford, A.C. Hannon, E.R. Barney, R.I. Walton, Structures of uncharacterised polymorphs of gallium oxide from total neutron diffraction, Chem. Eur. J. 19 (8) (2013) 2803–2813. [45] F. Mezzadri, G. Calestani, F. Boschi, D. Delmonte, M. Bosi, R. Fornari, Crystal structure and ferroelectric properties of ε-Ga2O3 films grown on (0001)-sapphire, Inorg. Chem. 55 (1) (2016) 12079–12084. [46] I. Cora, et al., The real structure of ε-Ga2O3 and its relation to κ-phase, CrystEngComm 11 (2017) 1509–1516. [47] M. Pavesi, et al., ε-Ga2O3 epilayers as a material for solar-blind UV photodetectors, Mater. Chem. Phys. 205 (2018) 502–507. [48] K. Irmscher, Z. Galazka, M. Pietsch, R. Uecker, R. Fornari, Electrical properties of β-Ga2O3 single crystals grown by the Czochralski method, J. Appl. Phys. 110 (6) (2011) 063720. [49] T. Onuma, et al., Correlation between blue luminescence intensity and resistivity in β-Ga2O3 single crystals, Appl. Phys. Lett. 103 (4) (2013) 31–34. [50] L. Binet, D. Gourier, Origin of the blue luminescence of β-Ga2O3, J. Phys. Chem. Solids 59 (8) (1998) 1241–1249. [51] R. Fornari, et al., Thermal stability of ε-Ga2O3 polymorph, Acta Mater. 140 (2017) 411–416. [52] E.G. Vı´llora, K. Shimamura, K. Kitamura, K. Aoki, T. Ujiie, Epitaxial relationship between wurtzite GaN and β-Ga2O3, Appl. Phys. Lett. 90 (23) (2007)234102. [53] M.M. Muhammed, et al., High optical and structural quality of GaN epilayers grown on (2 01) β-Ga2O3, Appl. Phys. Lett. 105 (4) (2014) 42112. [54] D. Gogova, et al., Structural properties of Si-doped β-Ga2O3 layers grown by MOVPE, J. Cryst. Growth 401 (2014) 665–669. [55] Y. Zhuo, Z. Chen, W. Tu, X. Ma, Y. Pei, G. Wang, β-Ga2O3 versus ε-Ga2O3: control of the crystal phase composition of gallium oxide thin film prepared by metal-organic chemical vapor deposition, Appl. Surf. Sci. 420 (2017) 802–807. [56] H. Sun, et al., HCl flow-induced phase change of α-, β- and ε-Ga2O3 films grown by MOCVD, Cryst. Growth Des. 18 (2018) 2370–2376. [57] Y. Yao, et al., Growth and characterization of α-, β-, and E-phases of Ga2O3 using MOCVD and HVPE techniques, Mater. Res. Lett. 6 (5) (2018) 268–275.

MBE growth and characterization of gallium oxide

2

Neeraj Nepal, D. Scott Katzer, David J. Meyer U.S. Naval Research Laboratory, Washington, DC, United States

Chapter Outline 2.1 Introduction 31 2.2 MBE growth of Ga2O3 and materials characterization 2.2.1 2.2.2 2.2.3 2.2.4 2.2.5

2.3 Current status and future prospects 2.4 Summary 44 Acknowledgments 44 References 44

2.1

33

Oxide MBE equipment considerations 33 Amorphous Ga2O3 for gate dielectrics 33 Heteroepitaxy of Ga2O3 33 Homoepitaxy of Ga2O3 34 Stabilizing metastable phases 39

43

Introduction

Gallium oxide (Ga2O3) is an ultrawide bandgap (UWBG) oxide semiconductor with an indirect bandgap of 4.5–5.2 eV [1–4]. The beta-phase (β-Ga2O3) is the commonly regarded as most stable of the several crystalline phases (or polymorphs) of Ga2O3. Because of its wide bandgap, it is transparent from ultraviolet to visible wavelengths. It had first been widely explored as a transparent conductive oxide (TCO) [1, 5] for optical devices such as light-emitting diodes. Also, it has been used as a gate dielectric in metal oxide semiconductor (MOS) structures in GaAs [6–9]. The β-Ga2O3 can be synthesized by melt growth techniques such as Czochraski [10, 11], floating zone (FZ) [12–15], and edge-defined film-fed growth (EFG) [16–18] at atmospheric pressure which can provide inexpensive large area bulk substrates. Kuramata Akito et al. has already demonstrated 4-in. diameter substrate by EFG and Tamura Corporation has commercialized 2-in. diameter wafers [16]. The commercial availability of large area Ga2O3 substrates is an important advantage over GaN and similar group III-N compound semiconductors in many potential electrical and optical device applications. Besides material benefit of UWBG (see Table 2.1), these substrates provide a highquality crystalline platform for power electronics devices that require higher crystalline quality, low-defect density material with precise doping control capabilities. As such, Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00002-6 © 2019 Elsevier Inc. All rights reserved.

32

Gallium Oxide

Table 2.1 Materials properties ultrawide band gap semiconductors compared to GaN for high-power, high-frequency applications Property

GaN

β-Ga2O3

Diamond

AlN

Bandgap Eg (eV) Dielectric constant εr Electric breakdown field Ec (MV/cm) Electron mobility μn (cm2/V s) Hole mobility μp (cm2/V s) Thermal conductivity (W/cm K) Saturated electron velocity (107 cm/s) Baliga’s figure of merit (BFOM) Johnson figure of merit (JFM)

3.4 9 3.3 1250 850 2.1 2.5 1 1

4.9 10 8 300 10 0.27 2.4 3.8 5.5

5.5 5.5 10 2200 850 10 2.7 29.4 10.7

6.2 8.5 17 300 14 2.9 2.2 28.3 19.3

Ga2O3 has attracted wide interest for such power electronics applications [4, 19]. Homo- and heteroepitaxial growth of Ga2O3 films has been carried out by molecular beam epitaxy (MBE) [5, 20, 21], metal–organic chemical vapor deposition (MOCVD) [22], chemical vapor deposition (CVD) [23], mist-CVD [24–27], pulsed laser deposition (PLD) [28], GaN oxidation [29], sol–gel [30], and halide vapor-phase epitaxy (HVPE) [31, 32] techniques. Among them, MBE has the advantages of enabling high-purity film growth with abrupt interfaces, controllable doping over a wide range, and in situ electron diffraction diagnostic tools, and is one of the most popular growth techniques for research into the epitaxial film growth of Ga2O3. There are several different crystalline structures of Ga2O3 that occur in nature [25, 27, 33–35]. The alpha phase, α-Ga2O3, is the structural analog of corundum or α-Al2O3. It is a member of the R-3c space group with a ¼ 0.49825 nm, c ¼ 1.3433 nm. It has been heteroepitaxially grown on α-Al2O3 substrates [14, 20]. β-Ga2O3 is the most stable polymorph of Ga2O3 (all other phases are reported to be metastable) and as has been mentioned is commercially available in large substrate form. It is a member of the C2/m space group with a ¼ 1.223 nm, b ¼ 0.304 nm, c ¼ 5.80 nm, and crystal angle β ¼ 103.7 degrees. γ-Ga2O3 has a structure analogous to MgAl2O4 or the spinel-type structure. It is a member of the Fd-3m space group with a ¼ 0.822 nm. Finally, ε-Ga2O3 is a hexagonal phase believed to be the next most stable phase of Ga2O3 after the beta phase. It is a member of the P63mc space group with a ¼ 0.2904 nm, c ¼ 0.9255 nm. The epsilon phase of Ga2O3 is potentially appealing for integration on heterostructures as its noncentrosymmetric polar crystal structure is expected to give rise to spontaneous and piezoelectric polarization charges at crystal discontinuities in analogy to those in the AlGaN/GaN system and thus potentially enabling polarization-engineered devices such as high electron mobility transistors (HEMTs). Note that Cora and colleagues recently found that nanoscale grains of what has been identified as ε-Ga2O3 is actually κ-Ga2O3 as a result of ordering of Ga atoms on the sublattice [36]. It is a member of the orthorhombic Pna21 space group with a ¼ 0.5046 nm, b ¼ 0.8702 nm, and c ¼ 0.9283 nm [36].

MBE growth and characterization of gallium oxide

33

In this chapter, we present growth and characterization of homoepitaxy and metastable phase stabilization of Ga2O3 by MBE. The cited MBE publications are representative of the best publicly available to the best of the authors’ knowledge.

2.2

MBE growth of Ga2O3 and materials characterization

2.2.1 Oxide MBE equipment considerations Many aspects of MBE growth of Ga2O3 are similar to those of more conventional III– V MBE growth—the use of effusion cells, the use of an ultrahigh vacuum (UHV) chamber, the use of ultrahigh-purity source materials when possible, and the use of reflection high-energy electron diffraction (RHEED) for in situ characterization and growth optimization. But the use of oxygen in the MBE system introduces additional complications above those of conventional MBE. In particular, while gallium (Ga) is not as reactive with oxygen as strontium and many other materials [37], Ga in the effusion source will react with oxygen in the MBE environment to some extent (up to 0.2 at.% oxygen solubility in liquid Ga at 1250°C) [38]. The extent of the oxidation of the source material depends on the various dimensions and operating characteristics of the MBE system [39], so further improvement in the technology of Ga2O3 MBE can be expected over time. Oxygen will also react with the RHEED electron gun cathode filament and with the heater filaments in effusion cells, potentially substantially reducing their lifetime (e.g., 1 year for RHEED filaments and 3 years for Ga effusion cell filaments [5]). Another important consideration is that molybdenum is the most commonly used metal for III–V substrate wafer holders in MBE. But molybdenum easily forms MoO3—a high vapor pressure compound—when exposed to oxygen [40]. To minimize the chance of contamination of the growing Ga2O3 oxide film with molybdenum, it is important to use molybdenum-free effusion cells and wafer mounts.

2.2.2 Amorphous Ga2O3 for gate dielectrics Much of the early work on the MBE growth of Ga2O3 were investigations of its utility as a gate dielectric for III–V MOS devices. Work by Callegari et al. [9] and Passlack et al. [7, 8, 41] showed the promise of amorphous Ga2O3 for these electron device applications and laid the foundation for subsequent epitaxial Ga2O3 growth investigations by MBE. A review of the status of MBE Ga2O3 gate dielectrics as of 2005 can be found in Ref. [6].

2.2.3 Heteroepitaxy of Ga2O3 Heteroepitaxial growth of Ga2O3 thin films to date has most commonly used sapphire (Al2O3) [19, 42–44]. Other substrates such as SiC [45, 46], MgO [25, 47], and MgAl2O4 [48] substrates have also been used. Some of these substrate materials have the advantage of being commercially available in reasonably large sizes, high quality, and relatively low cost. The surface preparation and cleaning of Al2O3, MgO, and

34

Gallium Oxide

MgAl2O4 is well understood, making them appealing for the MBE growth of new oxide materials. In addition, varying only the cations in the substrate to the Ga across the interface in the growing film reduces many of the complexities that would be present when growing on a crystalline substrate with a different anion (e.g., SiC, AlN, Si, GaAs, etc.). Even with a common anion, though, the crystal structure of monoclinic β-Ga2O3 on hexagonal c-plane sapphire is not ideal for obtaining perfect epitaxy. β-Ga2O3 grows with the {20-1} planes parallel to the c-Al2O3 (0001) surface, but has three rotational domains that have a similar oxygen arrangement between the film and the substrate— namely h102i β-Ga2O3 jj [11–20] Al2O3 [20, 44, 47]. Cubic MgO (001) has also been used for heteroepitaxial growth of β-Ga2O3. In that case, the Ga2O3 grows with the {100} orientation with h021i β-Ga2O3 jj [010] MgO [47]. Cubic spinel MgAl2O4 (001) has been used for heteroepitaxial growth of β-Ga2O3 by MOCVD [48] and γ-Ga2O3 films by PLD [48a]. In the former case, the epitaxial relationship was β-Ga2O3 (1 0 0) jj MgAl2O4 (1 0 0) with β-Ga2O3 [0 0 1] jj MgAl2O4 h0 1 1i, while in the latter case the film was single-domain with the expected cube-on-cube epitaxial arrangement with the spinel substrate. While the large bandgap of β-Ga2O3 offers promise for voltage handling capability in future UWBG semiconductor high-power/temperature electronics devices, its low thermal conductivity and experimentally observed breakdown voltages that are less than theoretically predicted present impediments to immediate application. For the optimal performance of Ga2O3-based devices, careful heterostructure design is required. Heteroepitaxy on high thermal conductivity substrates such as Si, SiC, and diamond are preferred to address the thermal limitations of Ga2O3. The (010) plane of β-Ga2O3 has a lattice mismatch of about 20.8% and 1.3% with Si(111) and SiC, respectively. Thus, the growth of Ga2O3 on SiC may result in a lower density of threading dislocations than on the other potential high-thermal-conductivity substrates. The thermal expansion coefficient of β-Ga2O3 along this direction is 3.37  106 K1 [49] and its thermal mismatch is about 30% and 25% with Si and 4H-SiC, respectively.

2.2.4 Homoepitaxy of Ga2O3 Homoepitaxy of Ga2O3 on native Ga2O3 substrates has many advantages in terms of the potential for achieving high crystalline quality of the growing film. Homoepitaxy eliminates complications involved with differing crystal structure, thermal expansion mismatch, lattice mismatch, and others which are unavoidably introduced in heteroepitaxy. It is for these reasons that there has been increasing interest in the MBE of Ga2O3 on commercially available native β-Ga2O3 substrates. Although β-Ga2O3 was first explored as TCO, it has better figures of merit for power and radio frequency (rf ) device applications than many other materials. For electronic application, high-quality homoepitaxy is one of the best options for obtaining the highest quality Ga2O3 films, and Ga2O3 has an important advantage in that single-crystal substrates can be grown by melt growth methods at atmospheric pressure. Standard melt growth techniques enable the mass production of large-

MBE growth and characterization of gallium oxide

35

diameter single-crystal wafers required to grow low defect density homoepitaxial films and vertical device structures needed for high-voltage and high-current power and rf devices. Ever since the development and availability of high-quality bulk Ga2O3 substrates, homoepitaxy has been one of the preferred growth approach in Ga2O3 research. Homoepitaxy has been carried out by various growth techniques such as MOCVD, PLD, CVD, and MBE.

2.2.4.1 Substrate selection β-Ga2O3 is single axis monoclinic structure that had two-fold symmetry. All three axes are unequal in length. It is quite common in epitaxy for many properties of the growing film to depend on the surface orientation of the substrate, and that is also the case in Ga2O3 growth. In fact, the substrate orientation selection is critical in homoepitaxial growth of Ga2O3. Sasaki et al. [50, 51] has shown that the Ga2O3 growth rate is highly dependent on the substrate orientation for growth by ozone MBE. Fig. 2.1 shows a relationship between the surface orientation and homoepitaxial β-Ga2O3 growth rate by MBE [51]. The horizontal axis in the figure is the angle between the substrate surface and the (100) plane. An undoped Ga2O3 layer was grown on Si-doped n-type Ga2O3 substrate samples of various orientations. The epitaxial growth rate on the (100) plane was about 10 nm/h, whereas the rate on the (010) and (310) planes was about 125 nm/h. Thus, it was found that the epitaxial growth rate can be increased by more than ten times by changing from the (100) plane to the (010) or (310) plane. The growth rate on the (100) plane was much lower than on other orientations because the adhesion energy on (100) terraces is lower than on other planes, thus the reevaporation rate of atoms supplied to the (100) terraces is higher than on other planes.

Fig. 2.1 Relationship between surface orientation of β-Ga2O3 substrate and homoepitaxial growth rate. The horizontal axis is the angle between the substrate surface and (100) plane. Adapted from K. Sasaki, M. Higashiwaki, A. Kuramata, T. Masui, S. Yamakoshi, J. Cryst. Growth 378 (2013) 591.

36

Gallium Oxide

2.2.4.2 Substrate cleaning Preparing the substrate surface for the homoepitaxial growth is an important consideration for resulting film quality [52, 53]. The growth surface can be prepared ex situ or in situ or combination of both for high-quality epitaxy. Sasaki et al. performed wet chemical treatment of the substrate before loading into the MBE chamber. Prior to loading the substrates into an MBE growth chamber, they were cleaned with organic solvent (acetone and methanol), acid [HF (46%) and H2SO4 + H2O2], and ultrapure water, and then bonded to a Si wafer with indium metal. Okumura et al. recovered a streaky RHEED pattern (indicative of a clean, well-ordered, smooth surface) from the as received surface through the use of in situ oxygen plasma treatment [53]. The as-received substrate did not have clear streaky RHEED streaks as shown in Fig. 2.2A, which becomes streaky (Fig. 2.2B) after plasma treatment. They also observed more than seven Pendell€ osung fringes for the homoepitaxial growth along (010) on plasma treated surface by plasma-assisted MBE, which indicate a high-quality interface between the substrate and epitaxial layer (Fig. 2.2C).

2.2.4.3 MBE growth optimization Substrate temperature One of the key growth parameters is the substrate temperature during growth (Tg). Sasaki et al. studied its effect on the structural and electrical properties of Sn-doped Ga2O3 homoepitaxial films grown by MBE on single-crystal β-Ga2O3 (010) substrates [54]. They obtained atomically smooth and flat surfaces for the samples grown in the range of Tg ¼ 550–650°C. Fig. 2.3A shows the Tg dependence of the growth rate of the unintentionally doped Ga2O3 epitaxial film. The growth rate decreased a little above Tg ¼ 700°C as a result of reevaporation of Ga metal from the growth surface. During Ga2O3 MBE growth,

Fig. 2.2 RHEED patterns of (A) as-received and (B) plasma-cleaned β-Ga2O3(010) substrate surfaces at an oxygen flux of 1.2  105 Torr at 600°C for 30 min. (C) (020) ω – 2θ scan diffraction peak of 135-nm-thick β-Ga2O3(010) film grown under slightly Ga-rich conditions (ϕGa  ϕO* ¼ 0.3) with an oxygen flux of 1.2  105 Torr (BEP). Adapted from O. Hironori, K. Masao, S. Kohei, K. Akito, H. Masataka, S.S. James, Appl. Phys. Express 7 (2014) 095501.

MBE growth and characterization of gallium oxide

37

Fig. 2.3 (A) Growth rates of Ga2O3 epitaxial films as a function of Tg, (B) surface RMS roughnesses of Ga2O3 epitaxial films as a function of Tg, (C) RHEED patterns, and (D) surface AFM images of Ga2O3 epitaxial films grown at 600°C. Adapted from K. Sasaki, M. Higashiwaki, A. Kuramata, T. Masui, S. Yamakoshi, J. Cryst. Growth 392 (2014) 30.

the growth rate decreased about 25% at Tg ¼ 800°C. Thus, it turned out that the Tg dependence of the growth rate is relatively small. Fig. 2.3B plots the surface RMS roughness of the Ga2O3 epitaxial films as a function of Tg (the roughness values were obtained from the AFM images). The surface RMS roughness was as small as 0.4 nm for the sample grown at Tg ¼ 600°C as shown in Fig. 2.3D. The RHEED patterns are presented in Fig. 2.3C. The sample grown at Tg ¼ 550–650°C had a smooth surface.

Ga and oxygen flux In MBE, Ga2O3 is mostly grown by the oxidation of evaporated Ga flux by using either ozone or excited (or active) oxygen such as from a plasma. Growth can be carried out either at oxygen or Ga rich conditions and the film growth strongly depends on the ratio of O/Ga. Patrick Vogt and Oliver Bierwagen presented an important growth parameter space phase diagram for the growth of Ga2O3 on c-sapphire by MBE

38

Gallium Oxide

[43]. Although quantitative growth parameter values would be different for homoepitaxy, qualitative comparisons can be made and the growth fundamentals should be same for homo- and heteroepitaxy. Desorption of the higher vapor pressure Ga suboxide (Ga2O) under Ga-rich condition decreases the growth rate, so to counteract this effect the growth is performed under oxygen-rich conditions. To increase growth rate Ga2O desorption need to be suppressed by providing additional oxygen to convert it to Ga2O3. This may be one of the reasons that ozone MBE provides higher growth rate compared to plasma MBE, as O3 is more reactive than excited O2. On the other hand, to improve material quality Ga-rich growth is preferred to suppress point defect such as Ga-vacancies. Okumura et al. studied the dependence of the Ga2O3 growth rate on Ga and active O fluxes (ϕGa and ϕO*) [53]. As shown in Fig. 2.4A, for Ga-limited (oxygen-excess) growth the maximum growth rate was 1 nm/min at 700°C and remained at 1 nm/min for the Ga flux up to 20 nm/min without any Ga droplet formation on the surface, however, the growth rate decreased with increasing Ga flux at T ¼ 500°C. It indicates that excess Ga forms volatile Ga2O at lower temperature. Excess Ga adatoms may have desorbed from the Ga2O3 surface at higher temperature (Tg > 600°C). As is illustrated in Fig. 2.4B, for Tg of 650–750°C, the growth rate was constant at 1 nm/min. For Tg > 750°C, the growth rate decreased with Tg and was independent of the Ga/O* ratio, apparently as a result of thermal decomposition of the Ga2O3 (010) films. For Tg < 650°C, the growth rate decreased with Tg, yielding a low growth rate at 500°C under slightly Ga-rich conditions. The RHEED image intensity decreased during growth under Ga-rich conditions for Tg < 600°C, indicating excesses Ga accumulation on the surface. Excess Ga on the surface could reduce oxygen incorporation and could result in the formation of

Fig. 2.4 (A) Relationship of the growth rate of β-Ga2O3 (010) to the impinging Ga flux ϕGa at 500°C (black circles) and 700°C (white circles). The oxygen plasma power and flux were 200 W and 1.5  105 Torr, respectively, corresponding to an O* flux of 1 nm/min. (B) Growth rate of β-Ga2O3 (010) under slightly Ga-rich conditions (ϕGa  ϕO* ¼ 0.3 nm/min) as a function of growth temperature. The oxygen flux was 1.5  105 Torr. Adapted from O. Hironori, K. Masao, S. Kohei, K. Akito, H. Masataka, S.S. James, Appl. Phys. Express 7 (2014) 095501.

MBE growth and characterization of gallium oxide

39

Ga2O. Growth temperatures above 600°C can promote the desorption of excess Ga and reduce the formation of Ga2O.

2.2.5 Stabilizing metastable phases Vogt et al. carried out systematic study of the reaction kinetics involved in the MBE growth of Ga2O3 [43, 55]. To understand the growth kinetics, they systematically varied the Ga flux, Ga-to-O ratio, and growth temperature, and developed a growth diagram (GD) as shown in Fig. 2.5. The GD has three distinct regions: (i) complete Ga incorporation O-rich, (ii) O-rich with Ga2O desorption-limited-growth regime with plateau of growth rate, and (iii) Ga-rich with a decreasing growth rate at increasing Ga flux. In each case, the difference between the Ga flux and the growth rate is related to the formation and desorption of the volatile Ga2O, which is catalyzed by the Ga2O3 surface. Growth rates lower than the Ga flux were either caused by the growth in the Ga-rich regime or by a higher growth temperature. Their results provide guidance for further increasing the maximum growth rate and crystal quality. In order to improve crystal quality, Ga-rich growth has to be established. A Ga-rich growth regime in the plateau region may suppress the formation of Ga vacancies. Table 2.2 summarizes the five known phases of Ga2O3. Single-crystal Ga2O3 exists in α, β, δ, ε, and γ, polymorphs with the free energy of formation in the order of β < ε < α < δ < γ [34, 35]. These metastable phases have unique properties that can be exploited for device applications. β-Structure is the most stable phase and α-Ga2O3 crystallizes in the corundum structure. Υ and δ crystallize in cubic structures

Fig. 2.5 Ga2O3 MBE growth diagram (GD) showing the variation of ϕGa/ϕO with growth temperature. It shows two major O-rich and Ga-rich regimes. The GD illustrates regimes of complete (i), partial ((ii), (iii), and (iii*)), and no Ga-incorporation (iv) as a function of TG and rGa. Adapted from P. Vogt, O. Bierwagen, Appl. Phys. Lett. 108 (2016) 072101.

Structure

Monoclinic (C2/m) Wurtzite (P63mc) Corundum (R 3 c) Cubic (I a 3) Cubic (F d 3 m)

β-Ga2O3 ε-Ga2O3 α-Ga2O3 δ-Ga2O3 Υ-Ga2O3 4.9 4.9 5.2 N/A 5.3

Eg (eV) 10 10 N/A N/A N/A

εr 12.214 2.9067 4.983 9.52 8.238

˚) a (A 3.037 N/A N/A N/A N/A

˚) b (A

Materials parameters of polymorphs of Ga2O3 [19, 31, 34, 35, 56–58]

Ga2O3 polymorphs

Table 2.2

5.798 9.255 13.433 N/A N/A

˚) c (A

103.83 N/A N/A N/A N/A

β (degree)

40 Gallium Oxide

MBE growth and characterization of gallium oxide

41

and have different space group. The ε-phase of Ga2O3, while, slightly energetically disfavored compared to the β-phase, has a hexagonal wurtzite crystal structure and is the only polar phase of Ga2O3. As has been demonstrated in GaN and related ternaries polarization-induced charges can result in two-dimensional carrier gases with higher mobility transport compared to conventional impurity doping, which can further increase the device power density by reducing the on-resistance of the device. The in-plane lattice mismatch of ε-phase Ga2O3 has been experimentally shown to be 8.8%, 6.6%, and 5.7% with (0001) oriented GaN, AlN, and hexagonal SiC, respectively [44]. The calculated polarization strength of ε-phase Ga2O3 is 10 and 3 times larger than that of GaN and AlN, respectively [56]. Different polymorphs of Ga2O3 can be grown epitaxially [25, 27, 33, 36, 59–61]. Besides β-Ga2O3, the metastable phases are of interest because of their promising material properties such as band gap versus lattice constant engineering (ε- and α-phases), wider band gap than III-nitrides and polarization engineered heterostructure design (ε-phase). Maccioni and Fiorentini predicted polarization-induced two-dimensional electron gases with high sheet carrier densities up to 1.6  1014 cm2 in heterostructures of ε-Ga2O3 with other materials such as GaN due to a high spontaneous polarization along (0001) direction [56]. Growth techniques such as mist-CVD have been used to crystallize these metastable phases [25–27, 62]. However, MBE has generally created the most stable β-phase unless carried out with the assistance of impurity dopants. The presence of impurities such as tin (Sn) [42] and indium (In) [63] at the doping level has been shown to stabilize the ε-phase by forming an intermediate impurity oxide and exchanging oxygen to Ga2O in the Ga-rich growth regime. It also prevents Ga2O desorption and increases the growth rate by suppressed etching due to suboxide formation.

2.2.5.1 Tin-assisted ε-Ga2O3 Using Sn as a dopant Kracht et al. synthesized β- and ε-Ga2O3 by MBE on c-plane sapphire substrates [42]. The Sn-dopant expanded the growth window of β-Ga2O3 in the Ga-rich regime, suppressed the Ga suboxides formation and stabilized phase-pure ε-Ga2O3. They proposed a growth model as shown in Fig. 2.6 based on the oxidation of Ga suboxide by reduction of an intermediate sacrificial tin oxide. Fig. 2.6 shows the Ga2O3 growth rates vs Ga beam equivalent pressure (BEPGa) for two different series with different Sn BEP (BEPSn) (the black squares, BEPSn ¼ 0) and (the red triangles, BEPSn ¼ 1.17  1011). Initially, the GR increased with Ga flux is observed in the Ga-limited oxygen-rich growth regime. Without Sn, the growth rate reaches a maximum with transition to Ga-rich conditions around BEPGa ¼ 5  108 mbar and it decreases to zero for a higher BEPGa due to Ga2O formation and desorption. However, introducing Sn at the doping level during growth increases the growth rate by five times for higher BEPGa value of 2.1  107 mbar indicating that the Sn suppresses Ga2O formation and desorption. As shown in Fig. 2.5, volatile Ga suboxide Ga2O is formed and desorbs from the sample surface when metal-rich growth conditions are applied. The presence of Sn on the surface suppresses the desorption process. Based on Gibbs energies calculations, Kracht et al. explain this finding by concluding that formed SnO2 and SnO in the

42

Gallium Oxide

Fig. 2.6 Comparison of the growth rates of series B and A (with and without tin) as a function of BEPGa. An attenuation of etching and an expansion of the growth window with the presence of tin are observed. Adapted from M. Kracht, et al., Phys. Rev. Appl. 8 (2017) 054002.

presence of tin adatoms react with Ga2O and growth rate increases by stabilizing more Ga2O3. The reduced Sn might be reoxidized to SnO and SnO2 that can further oxidize Ga2O and hence attenuate the etching process. Thus, the Sn impurity helps to overcome the growth limitations caused by the formation of volatile suboxides. Also the presence of Sn above a critical concentration (BEPSn of 1011 mbar) in metal-rich growth conditions results in the formation of ε-Ga2O3.

2.2.5.2 Indium-assisted ε-Ga2O3 Ga2O3 growth rate in plasma-assisted MBE was enhanced by an additional In flux [63]. The enhancement results from a catalytic effect by formation of In2O3 followed by an In-Ga interatomic exchange to form Ga2O3. Vogt and colleagues derived a simple model that quantitatively describes process as well as its consequences on the formation rate of Ga2O3. They also demonstrated the catalytic action of In2O3 that allows the synthesis of the metastable ε-Ga2O3. Desorption of Ga2O and In was measured using quadrupole mass spectroscopy (QMS). Fig. 2.7A shows the desorption flux of Ga2O (ϕGa2O) and In (ϕIn) from a β-Ga2O3 template at 700°C. The surface was exposed with O first (regime i) so there is no ϕGa2O and ϕIn flux. In regime (ii), when the surface was exposed to Ga flux, ϕGa2O is equal to supplied Ga flux. In regime (iii), ϕIn is supplied in addition to ϕGa and ϕO. After a short delay, the Ga2O desorption flux decreases by 40%, and laser reflectometry shows the growth of Ga2O3. Also, all supplied In to the surface desorbs resulting pure Ga2O3 film. As shown in region (iv), when the Ga shutter is closed and the In shutter is open, there is no In desorption and In2O3 nucleates. Ga2O3 grows when the Ga shutter is opened again [in time interval (v)]. This metal-exchange catalysis results the growth of ε-Ga2O3. Metal-exchange catalysis could apply to all materials whose binary constituents exhibit analogous kinetic and thermodynamic properties to In2O3 and Ga2O3.

MBE growth and characterization of gallium oxide

43

Fig. 2.7 (A) Desorption of Ga2O and In from β-Ga2O3 (201) surface at 700°C. ϕGa ¼ 6.5 nm–2 s–1 and ϕIn ¼ 4.0. The metals supplied in the time intervals (i)–(v) are indicated by the bars at the top. (B) XRD scan of In2O3-catalyzed ε-Ga2O3 film grown on a β-Ga2O3 on c-sapphire at 700°C, ϕGa ¼ 6.5 nm–2 s–1, and ϕIn ¼ 5.4 nm–2 s–1. Adapted from P. Vogt, O. Brandt, H. Riechert, J. Lohnemann, O. Bierwagen, Phys. Rev. Lett. 119 (2017) 196001.

For example, the oxidation efficiency of Sn (ηSn) is even larger than that of In, which, in turn, is larger than that of Ga, that is, ηSn > ηIn > ηGa [55]. This mechanism opens a new path for the epitaxial synthesis of transparent semiconducting oxides and their metastable phases. Vogt et al. used this catalysis technique to synthesis the metastable hexagonal ε-phase of Ga2O3 as shown in Fig. 2.7B [63].

2.3

Current status and future prospects

Substantial progress has been made in the epitaxial growth of several phases of Ga2O3 in recent years and the pace of progress is accelerating. Understanding of the properties and growth of the various technologically important Ga2O3 phases is increasing and there have not yet been any fundamental problems in epitaxy that MBE growth of these materials cannot address. However, the MBE growth of Ga2O3 is still in its infancy. In many respects, the MBE growth of Ga2O3 is at a stage similar to the early MBE investigations of GaN growth. Growth rates are relatively low and investigations of ways to increase the growth rate are underway. Early MBE GaN growth rates were often as low as 50 nm/h, but now growth rates exceeding 8 μm/h have been reported [64, 65]. Similarly, until recently, reported Ga2O3 growth rates were often 77 K, the energies of E0 and E00 show a gradually red shifting with the increasing

Fig. 4.9 Temperature-dependent TR spectra of the β-Ga2O3 nanowires between 30 and 320 K.

Synthesis, optical characterization, and environmental applications of β-Ga2O3 nanowires

81

temperature such as the general semiconducting behavior. It is noticed that an additional defect feature of ED1 presents only in the TR spectra of temperature range from 100 to 260 K. The ED1 feature cannot be clearly detected at low temperatures of 30–77 K because it may come from the VGa defect acceptor level, which is initially empty of electron. For T ¼ 100–260 K, the ED1 feature is detectable and its energy is nearly invariant to a value of
4.232 eV. This result indicated that the ED1 originates from a defect-to-defect transition. The ED1 can be assigned as the transition of VGa to the highest defect state in DB by either a VO0 or a Gai located near the conduction-band edge. For T  300 K, The VO0 or Gai defect state is nearly ionized and merged with the CB that reduces the TR amplitude of the ED1 feature as shown in Fig. 4.9. Temperature dependence of the EDB, E0, and E00 features obtained by the temperature-dependent TR spectra is depicted in Fig. 4.10. The solid lines are fitted to a Bose-Einstein expression:

EiðTÞ ¼ EiB  aiB 

8 > > < > > :

1+ 

9 > > =

2   > θiB ; 1 > exp T 

(4.3)

where i is the respective transition feature, aiB represents the strength of the electron– phonon interaction, and θiB corresponds to the average phonon temperature. The fitted values of EiB, aiB, and θiB for the β-Ga2O3 nanowires are 5.209  0.102 eV, 230  103 meV, and 560  100 K for E00 transition, 4.950  0.100 eV, 224  100 meV, and 560  100 K for E0 transition, and 3.528  0.101 eV, 115  55 meV, and 560  100 K for the EDB transition, respectively. The values of aiB for the E0 and E00 (band-to-band transition) are about doubled to that of EDB with VB to DB transition. It indicates that the variation speed of temperature-energy shift

Fig. 4.10 Temperature dependences of transition energies of E00 , E0, and EDB transitions in β-Ga2O3 nanowires.

82

Gallium Oxide

for E0 and E00 is faster than that of EDB as those in Fig. 4.10. The average phonon temperature θiB for the EDB, E0, and E00 transitions is 560  100 K. The phonon energy value (by kT) agrees well with the phonon replicas of
43 meV that induced by the vibration mode in the β-Ga2O3 nanowires as shown in Fig. 4.8B. Based on the experimental results, the near-band-edge structure of β-Ga2O3 nanowires is constructed and depicted in Fig. 4.11. The top of VB (Γ V) is consisted of oxygen 2p orbital and the lowest of CB (Γ C) constructed by Ga 4s electrons to render a direct gap E0 of 4.726 eV at 30 K. The valence-band top separates into higher Γ V and lower Γ V by crystal-field splitting with the energy of
0.252 eV. The E00 transition comes from the lower split Γ V ! Γ C. The defect donor band (DB) is mainly consisted by oxygen-vacancy levels VO as well as the highest defect donor of VO0 or Gai. The acceptors of β-Ga2O3 nanowires consist Ga vacancies VGa and VGa0 to form an acceptor band and VGa-VO vacancy-pair clusters to form near 1D quantum-well structure with size of L
2.2 nm. Experimental TR and PL results indicated that there is an excitonic level with a binding energy of
40–50 meV located below the DB band. All the transition assignments and transition energies of the ED3, EW1, EW2, EW3, ED2, Eex DB, EDB, ED1, E0, and E00 transitions are indicated distinctly in Fig. 4.11. An overall gapstate and near-band-edge picture for the β-Ga2O3 nanowires is hence being realized by the experimental study. In summary, we present the optical characterization of β-Ga2O3 nanowires by using temperature-dependent TR and PL measurements. The experimental results verified and confirmed the abundance of gap-state and near-band-edge transition exists in β-Ga2O3 nanowires. The origin of transition, transitional energy, optical-phonon

Conduction band

Ga 4s VO

Highest Vo' or Gai

Defect donor band

E0 ED3 ED2

EDB

E0'

ex EDB

ED1

VGa' Acceptors

W1 W2

VGa

W3

VGa

O 2p Valence band L ~ 2.2 nm

0.252 eV

VGa–VO pair clusters

Fig. 4.11 The gap-state and near-band-edge transitions in β-Ga2O3 nanowires.

Synthesis, optical characterization, and environmental applications of β-Ga2O3 nanowires

83

behavior, and temperature dependence of gap-state and near-band-edge transitions in β-Ga2O3 nanowires is discussed. The whole band-edge structure below and above the band gap of β-Ga2O3 nanowires was thus constructed.

4.4

Photocatalytic property of β-Ga2O3 nanowires

In the present section, we evaluated the correlations between defect states and photocatalytic activity of β-Ga2O3 nanowires. For this purpose, samples synthesized with different growth ambient conditions were prepared. Their structural and optical properties were characterized in the previous sections. The decomposition of two different dyes: rhodamine B (RB) and methyl blue (MB) with β-Ga2O3 nanowires under UV light illumination was used to evaluate their photocatalytic activity. The results show that the photocatalytic activity of β-Ga2O3 nanowires is greatly enhanced and mainly attributed to the increasing numbers of acceptor states associated with gallium defects. The highest photocatalytic activity was achieved from the sample with the largest numbers of defect states. A commercial UV lamp (Philips, TUV) with a wavelength centered at 254 nm was employed for UV irradiation. The photocatalytic activity was evaluated by the decomposition of MB and RB in a deionized water solution with an initial concentration of 1  105 M. A sample size of 1.5 cm2 was added to 10 mL of dye solution in a beaker covered with a quartz plate. Before irradiation, the samples were put in the dark for 10 min to ensure stable adsorption. After desired time intervals, the concentration of solution was analyzed by recording the characteristic absorption of MB (664 nm) and RhB (554 nm) using an UV–visible spectrometer. For comparison, the photocatalytic activities of commercial β-Ga2O3 powder (99.995%, Alfa Aesar) were tested under the same reaction conditions. The particle size of commercial β-Ga2O3 powder was ranging from 120 to 350 nm in diameter. The amount of β-Ga2O3 powder (0.0026 g) was similar to the estimated weight of β-Ga2O3 nanowires. The photocatalytic activities of β-Ga2O3 nanowires were evaluated by the photocatalytic degradation of MB and RB under UV illumination. Fig. 4.12A and C shows the progressive spectral change and degradation curve of MB in the presence of β-Ga2O3 nanowires upon UV light irradiation. MB shows the characteristic absorption band at 664 nm and decreased rapidly after illumination. Fig. 4.12B and D shows the absorption spectrum of RB solution in the presence of β-Ga2O3 nanowires prepared under 1% and 20% of ambient oxygen. RB shows a major absorption band at 554 nm and the absorbance decreased after illumination. The photocatalytic activity of the samples is significantly increased with the growth ambient oxygen. In addition, compared with β-Ga2O3 powder, the β-Ga2O3 nanowires show a higher photocatalytic activity. The photocatalytic performance of 20% β-Ga2O3 nanowires is 2.56 times better than that of β-Ga2O3 powder. One can see that the trend of the photocatalytic activities of RB is comparable with MB results [Fig. 4.12E and F]. The photocatalytic activities of β-Ga2O3 nanowires are slightly higher toward MB over RB. This may due to kinetics differences in photo-degradation reaction between RB and MB. Generally, the photocatalytic decomposition follows the Langmuir-Hinshelwood kinetic

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Fig. 4.12 Absorption spectrum of methyl blue (MB) and rhodamine B (RB) solution in the presence of 1% (A) and (B) and 20% (C) and (D) β-Ga2O3 nanowires sample under exposure to ultraviolet (UV) light for various durations. (E) and (F) The photocatalytic performance of different β-Ga2O3 nanowires samples.

model. The first-order order reaction rate constant can be calculated by the plots of the ln(C/C0) versus irradiation time (t). The obtained rate law may be present as lnðC=C0 Þ ¼ kt

(4.4)

where C is the concentration of dye, Co the initial concentration of dye, k the reaction constant, and t the irradiation time. The reaction constants for the samples prepared

Synthesis, optical characterization, and environmental applications of β-Ga2O3 nanowires

85

Table 4.1 Comparison of specific reaction rates for the photodegradation of rhodamine B (RB) and methyl blue (MB) over β-Ga2O3 nanowires grown under different amounts of ambient oxygen, and β-Ga2O3 powder Sample

kRB (min21)

kMB (min21)

β-Ga2O3 powder

0.0034

0.0053

1% O2

0.0054

0.0082

5% O2

0.0072

0.0104

10% O2

0.0078

0.0129

20% O2

0.0087

0.0142

under different ambient oxygen are listed in Table 4.1. The results confirm that the decomposition rates are much higher when β-Ga2O3 nanowires present, especially for those prepared under high ambient oxygen. The origin of enhanced photocatalytic activity will be discussed later. For sustainability concern, β-Ga2O3 is known with an excellent chemical and thermal stability. Here, a 20% β-Ga2O3 nanowires sample was examined for their antiphotocorrosion properties. The cyclic photodegradation of RB and MB under UV light irradiation in the presence of 20% β-Ga2O3 nanowires sample is shown in Fig. 4.13. The β-Ga2O3 nanowires exhibited an excellent photo-stability and showed no noticeable loss of photocatalytic activity after 10 cycles. The results demonstrate that nanostructured β-Ga2O3 is a promising green photocatalyst, particularly useful when be used under harsh environment. The photocatalytic activity is known to be dependent on the crystallinity, surface area, and morphology. The structural characterization shows the β-Ga2O3 nanowires samples have comparable surface area, and crystallinity. As a result, the different photocatalytic activity may be mainly attributed to their differences in the optical properties. In principle, when a photon with energy of hν exceeds the bandgap of materials absorbed, the electron will be promoted to CB leaving a hole in VB. While electron–hole pair can further proceed radiative (as light emission), nonradiative (as heat or lattice vibrations) recombination or react with electron donors and electron acceptors absorbed on the surface. The photocatalytic reactions are listed as follows: β  Ga2 O3 + hν ! β  Ga2 O3 ðeCB  + hVB + Þ eCB  + O2 ! O2  2eCB  + 2H + + O2 ! H2 O2 H2 O2 + eCB  ! OH +  OH hVB + + H2 O ! H + +  OH

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Gallium Oxide

Fig. 4.13 Cyclic photodegradation of (A) RB and (B) MB under UV light irradiation for 10 cycles using β-Ga2O3 nanowires grown under 20% ambient oxygen.

OH + organic dye ! CO2 + H2 O The photocatalysis activity is believed to be associated with highly reactive species of % peroxide (O 2 ) and hydroxyl radical ( OH) generated by electron and hole on the surface with water. If surface defect states exist, it may be able to trap the electron or hole,

Synthesis, optical characterization, and environmental applications of β-Ga2O3 nanowires

87

e– ECB

2.67 eV

3.02 eV

2.39 eV

–

O2

3.45 eV

O2 + e–

Eg = 4.7 eV

VO

VGa

h+ + OH–

•OH

(VGa, VO) EVB h+

Fig. 4.14 The band structure and near surface charge transfer process of β-Ga2O3 nanowires with defects.

the recombination is prevented and the rate of oxidation reduction reactions may be increased. In our case, the large populations of gallium defects have been found in samples prepared under high ambient oxygen. The large numbers of defects (VGa and VGa–VO) consist of robust acceptor states in the band gap, trapping the holes and preventing recombination. As a result, the photocatalytic activity is increased. The proposed schematic band structure and near surface charge transfer process of the β-Ga2O3 nanowires are illustrated in Fig. 4.14. Various defect bands (centered from 3.45 to 2.39 eV) promote the electron–hole pair separation rate. The UV luminescence has been reported to be independent of sample preparation methods and impurity concentrations, and was attributed to an intrinsic transistion [21]. The numbers of donor levels (Vo) does not change with growth ambient, which is evident in PL results. Thus, the oxygen vacancies (Vo) should not be responsible for the enhanced photocatalysis activity. As a result, the enhanced photocatalytic activity mainly came from the large numbers of acceptor states associate with gallium defects. The acceptor states not only extend the near UV light absorption edge but also slow down the electron–hole pair recombination rate. Both these two effects contribute to enhanced photocatalytic activity. Moreover, the 2.37 eV emission band dominates in the 20% sample, suggesting a large population of gallium vacancy are responsible for the enhanced photocatalytic activity. Therefore, the gallium vacancies can be considered to be the active sites of β-Ga2O3 nanowires. Thus, a strong correlation between optical property and photocatalytic activity was found in β-Ga2O3 nanowires. β-Ga2O3 nanowires were synthesized by an ambient controlled vapor transport process. Optical characterization results show a large number of acceptor states associated with gallium defects significantly extended the absorption edge. For photocatalysis application, the sample with a high concentration of gallium defects exhibits stable photocatalytic activity toward RB and MB degradation under UV irradiation. The gallium-associated defects are concluded to be active sites of β-Ga2O3 nanowires. The large numbers of acceptor states are believed to enhance photocatalytic activity in β-Ga2O3 nanowires.

88

4.5

Gallium Oxide

Summary

In this chapter, we present the synthesis, optical characterization, and environmental applications of the β-Ga2O3 nanowires. The gap-state and near-band-edge transitions of β-Ga2O3 nanowires were identified and studied. The defects states play an important role in their optical emission and photocatalytic property. Owing to its various interesting properties such as wide bandgap, chemical and thermal stability, robust defect states, large surface to volume ratios, β-Ga2O3 nanowires are very promising in potential applications in optoelectronic, environmental applications, and fundamental research in the future.

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Growth, properties, and applications of β-Ga2O3 nanostructures

5

Mukesh Kumar*,†, Sudheer Kumar*, Vikram Kumar†,‡, R. Singh*,† *Department of Physics, Indian Institute of Technology Delhi, New Delhi, India, †Nanoscale Research Facility, Indian Institute of Technology Delhi, New Delhi, India, ‡Solid State Physical Laboratory, New Delhi, India

Chapter Outline 5.1 Introduction 91 5.2 Study of β-Ga2O3 nanostructures 5.2.1 5.2.2 5.2.3 5.2.4

93

β-Ga2O3 nanostructures using the CVD technique 93 β-Ga2O3 nanowires: Morphological and structural properties 94 Optical properties of β-Ga2O3 nanostructures 99 Application of β-Ga2O3 nanostructures 101

5.3 Functional nanowires based on β-Ga2O3

103

5.3.1 Coaxial β-Ga2O3/GaN nanowires through ammonification of β-Ga2O3 nanowires 103 5.3.2 ZnGa2O4 nanowires through coaxial ZnO/β-Ga2O3 nanowires 107 5.3.3 β-Ga2O3 nanowires template mediated by high-quality ultralong GaN nanowires 108

5.4 Conclusions and future perspective Acknowledgments 110 References 110

5.1

109

Introduction

For the successful establishment of nanoscale device platforms and their integration in complex nanosystems, semiconductor nanostructures are potential building blocks for next-generation high-performance devices. They could serve as the fundamental component of various devices such as photodetectors, field effect transistors, sensors, lasers, light-emitting-diodes, photovoltaics, waveguides, thermoelectrics, batteries, nanoelectromechanical systems, and nanostructure-biological cell interfaces [1–6]. Wide bandgap semiconductor nanostructures also provide advantageous properties such as the ability to operate at high temperatures, low leakage current at room temperature, and radiation and chemical stability. Nanostructures exhibit the large surface-to-volume ratios, fewer defects, and less strain compared to their bulk counterpart, which could be utilized for realization of efficient nanoscale devices. Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00005-1 © 2019 Elsevier Inc. All rights reserved.

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Moreover, the physical dimension tuning capability of these nanostructures offers the advantage of tuned fundamental electronic and optical properties of nanoscale devices. Recently, beta gallium oxide (β-Ga2O3) with monoclinic crystal structure has emerged as a potential semiconductor, having outstanding properties such as large bandgap (
4.7–4.9 eV at room temperature (RT)), high breakdown field (8 MV cm1), and thermal and chemical stability at high temperatures [7, 8]. β-Ga2O3 based on nanostructures has shown promising applications in nanoscale devices such as photodetectors [9–15], gas sensors [15–25], and field effect transistors (FETs) [26–28]. In recent times, novel nanoscale devices have been demonstrated based on β-Ga2O3 nanostructures. Reddy et al. [29] showed high photocatalytic activity of β-Ga2O3 nanorods. Ling et al. [30] fabricated pH sensors based on Ga2O3 nanowires. Li et al. [28] demonstrated the FET based on β-Ga2O3 nanowires. Liu et al. [16] have demonstrated β-Ga2O3 nanowires based on gas sensors, and Li et al. [13] showed a solar-blind photodetector based on bridged Ga2O3 nanowires. Mechanisms for nanostructures growth are generally categorized as top-down or bottom-up approaches. The top-down approach involves etching of bulk material using lithographic technologies to form nanostructures. However, this approach has limitations when trying to decrease the size of nanostructures. The bottom-up approach is more common for nanostructure growth. This approach includes direct deposition methods (such as selective area epitaxy and screw dislocation driven growth), template directed methods, and metal-nanoparticle-mediated methods (such as vapor-liquid-solid (VLS) and vapor-solid-solid (VSS)). The most promising and widely adopted nanostructure for a variety of applications is nanowire, which is generally grown by the well-established VLS mechanism. Wagner and Ellis [31] proposed this mechanism for growth of Si whiskers in 1964. In VLS, metal or metal alloy nanoparticles are used as a catalyst for nanowire growth. This catalyst, at growth temperature, has a large accommodation coefficient compared to the surrounding surface of the substrate, and serves as a preferential site for nucleation and subsequent nanowire growth. Nanostructures of β-Ga2O3 such as nanowires, nanosheets, nanobelts, nanorods, and nanotubes have been synthesized using a number of growth techniques such as physical evaporation [32–34], arc-discharge [35, 36], laser ablation [37], carbothermal reduction [38], microwave plasma [39–41], metal-organic chemical vapor deposition [42, 43], and the most widely used technique, chemical vapor deposition (CVD) [44–64]. In this chapter, various growth techniques for β-Ga2O3 nanostructures are outlined, and the most widely used CVD method and the well-established VLS growth mechanism are described. Morphological, structural, and optical properties of β-Ga2O3 nanostructures are described, and investigations on growth aspect such as selfcatalytic growth, nanowire diameter control, and different substrate effects have been included. Novel applications of β-Ga2O3 nanostructures have been highlighted. Further, discussion on functional nanowires based on β-Ga2O3 nanowires such as coaxial GaN/β-Ga2O3 nanowires, ZnGa2O4 nanowires, and β-Ga2O3 nanowires mediated by high-quality long GaN nanowires are included for potential high-performance next-generation nanoscale devices.

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5.2

93

Study of β-Ga2O3 nanostructures

In the last two decades, β-Ga2O3 nanostructures including nanowires, nanosheets, nanobelts, nanorods, and nanotubes have been synthesized by various techniques which are outlined as follows. Using the physical evaporation technique, β-Ga2O nanostructures have been synthesized by a number of groups [32–34]. GaN powder [33] or Ga metal [34] as source materials, In [33] or Au [34] as catalysts, and sapphire [33] or Si [34] as substrate were employed to grow β-Ga2O nanostructures by this method. In the arc-discharge method for obtaining β-Ga2O nanostructures on the cathode, GaN powder and transition metals were placed into a small hole formed in the anode, and DC current was then applied between the electrodes of the anode and cathode [35, 36]. In the laser ablation technique, a laser beam is usually used to ablate the target [37]. Hu et al. [37] have synthesized β-Ga2O nanostructures on an alumina substrate at a suitable temperature (800°C) by ablating the Ga2O3 target using a KrF excimer pulsed laser. Using a mixture of Ga2O3 powders with graphite, β-Ga2O3 nanostructures have also been synthesized using the carbothermal reduction method [38]. Using microwave plasma [39–41], Sharma and Sunkara [39] demonstrated bulk synthesis of highly crystalline β-Ga2O3 nanotubes, nanowires, and nanobrushes where a thin film of molten Ga was exposed to a microwave plasma having a mixture of hydrogen (H2)/methane (CH4)/O2. In the metal-organic chemical vapor deposition method [42, 43], generally an organometallic precursor (such as gallium acetylacetonate (CH3COCHCOCH3)3Ga, having a low decomposition temperature (
196°C)) and O2 gas are used to grow β-Ga2O3 nanostructures. Other methods such as molecular beam epitaxy and pulsed laser deposition techniques are also very attractive for controlled doping and heterostructure formation based on β-Ga2O3 nanostructures. In addition to these techniques, one of the most widely used methods for β-Ga2O3 nanostructure growth is CVD [44–64], which is described in the next section.

5.2.1 β-Ga2O3 nanostructures using the CVD technique The CVD method is attractive due to a high deposition rate, capable of producing pure materials, reproducibility of synthesis and ability to control the morphology of nanostructures by controlling process parameters, and capability of producing materials on an industrial scale. Many groups have studied CVD growth of β-Ga2O3 nanostructures including nanowires, nanosheets, and nanobelts [44–64]. The morphology of β-Ga2O3 nanostructures depends on various experimental parameters in CVD such as precursors, growth temperature, growth duration, separation distance between metal source and substrate (with and without catalyst nanoparticles), type of catalyst nanoparticles, and source gases flow rates. For example, metallic Ga and oxygen as source materials have been widely used to grow β-Ga2O3 nanostructures [53, 62]. Different precursors like a mixture of Ga and Ga2O3 have also been adopted to grow β-Ga2O3 nanostructures with and without catalysts [47, 48, 65]. Au nanoparticles are commonly used as a catalyst to grow β-Ga2O3 nanostructures. Other catalysts such as Ni [65], Fe [60], and spin-coated Ga2O3 films as a catalyst

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[58] have also been reported. Commonly used substrates for β-Ga2O3 nanostructures include silicon [44, 46] and alumina [45]. β-Ga2O3 nanostructures such as nanoribbons and nanorods have been shown by Auer et al. [53] and suggested that the vapor-solid (VS) mechanism is responsible for the obtained nanostructures. Chang et al. [5] have grown β-Ga2O3 nanowires using Ga vapors and H2O as precursors. The reaction temperature was kept in the range of 700–950°C. Grown nanowires exhibited a VLS growth mechanism and had diameters in the order of a few tens of nanometers with lengths of a few microns. β-Ga2O3 nanostructures such as nanosheets, nanowires, and nanobelts have been observed in self-catalytic nanostructure growth using the CVD technique [58], where spin-coated Ga2O3 films were used as a catalyst to initiate the growth of β-Ga2O3 nanostructures. The resulting rectangular nanosheets of β-Ga2O3 have widths and lengths in the range of 0.5–1.2 and 5–15 μm, respectively. On the other hand, nanowires grown by the self-catalytic method have diameters and lengths of about 50–150 nm and a few tens of micrometers, respectively. In various kinds of nanostructures, nanowires are one of the more promising nanostructures of β-Ga2O3, and show potential for nanoscale device applications.

5.2.2

β-Ga2O3 nanowires: Morphological and structural properties

Nanowires are desirable due to their capability to interface with other nano-micro scale systems, realization of nano-bio devices due to their size regime being similar to biological macromolecules, and single nanowire-based device integration. They also have a higher surface-to-volume ratio with long axial morphology for applications in light detection, sensors, and catalytic functionalities. As outlined above, VLS is the well-established mechanism for nanowire growth. The VLS mechanism incorporates distinct stages for nanowire growth as alloying, nucleation, and growth. A model for the typical VLS growth mechanism is shown in Fig. 5.1. At growth temperature, source vapor atoms dissolve into the catalyst and form the catalyst alloy. Supersaturation in this alloy induces the precipitation of a solid phase, which leads to nanowire growth. However, several basic aspects of VLS growth are still unclear due to rapid kinetics under conventional growth conditions, and it is difficult to probe the VLS growth events. In the case of the VLS process, mainly two general models have been described [66, 67]. In one model, catalyst alloy nanoparticles are at the Incoming vapor atoms

Nanowire Alloy droplet Catalytic particle or

substrate

Fig. 5.1 Model for VLS mechanism for nanowire growth.

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bottoms of nanowires (root mode); in the other model, catalyst alloy nanoparticles are on the tips of nanowires (tip mode). However, their physical or chemical origins are still unclear. Nucleation of nanowires on a catalyst coated substrate depends on various parameters such as interfacial parameters, contact angle of catalyst-nucleus and nucleus-substrate, and nanowire material-catalyst phase diagram. Therefore, catalyst droplet morphology in addition to growth conditions is crucial for preference of one of the two nanowire growth modes (root mode and tip mode). In general, surface diffusion transport and direct nucleation of species from the vapor phase play a crucial role in the catalyst-mediated root mode, which leads to the absence of the catalyst alloy on the nanowire tips. In the tip mode, the catalyst alloy tends to float on nucleated nanowires, resulting in the presence of the catalyst alloy on the nanowire tips. Nanowire structure has potential as a component in building blocks of nanoscale devices through adjustment of fundamental optical and electronic properties by tuning its nanowire diameter. Diameter tuning of β-Ga2O3 depends on various experimental conditions in CVD, such as separation distance between metal source and substrate, gases flow rates, growth temperatures, and combination of catalyst nanoparticles’ size with growth temperatures [62]. The model shown in Fig. 5.2 illustrates diameter tuning of β-Ga2O3 nanowires at various growth conditions in CVD. Nanowire diameter Ar/O2 gases

Ga2O3 nanowire

Ga metal

900⬚C

Catalyst alloy Substrate

Decreasing catalyst nanoparticle size 800⬚C

Higher flow rate at 900⬚C

850⬚C Increasing growth temperature to 850⬚C

Decreasing temperature

Increasing distance between metal source to substrate

Fig. 5.2 Growth model for diameter tuning of β-Ga2O3 nanowires at various growth conditions in CVD system. Reproduced with permission from M. Kumar, V. Kumar, R. Singh, Diameter tuning of betaGa2O3 nanowires using chemical vapor deposition technique, Nanoscale Res. Lett. 12 (2017) 184 with author’s copyright.

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can be controlled by varying separation distance between the metal source and substrate. The diameter decreases with an increase in separation distance between the metal source and substrate due to depletion of reactant vapor species and an increase in developed boundary layer thickness in CVD further downstream. This has been demonstrated in the top panel of the model. Growth temperature is one of the crucial parameters to obtain β-Ga2O3 nanowires, and the diameters of these nanowires are strongly affected by the selection of growth temperature and catalyst nanoparticles’ size. It is usually observed that a higher growth temperature (such as 900°C) for β-Ga2O3 nanowires induces larger diameters, compared to a lower growth temperature (such as 800°C). This is demonstrated on the left-hand side of model by various growth temperatures where temperature-dependent vapor pressure and adatom mobility play a crucial role to control the nanowire diameter. To obtain smaller nanowire diameters, a lower growth temperature and smaller catalyst size must be used. As an example, when using a lower growth temperature such as 800°C and smaller catalyst nanoparticles such as 20 nm, β-Ga2O3 nanowires below 50 nm can be obtained [62]. This has been illustrated in the bottom part of the model. Further, a higher flow rate of gases can induce larger nanowire diameter, though the surface morphology of nanowires gets deteriorated as shown on the right-hand side of the model. β-Ga2O3 nanowires with different diameters are shown in Fig. 5.3, where FESEM images (A–C) represent the diameter control by varying separation distances between the metal

Fig. 5.3 FESEM images of β-Ga2O3 nanowires with different diameters. (A–C) nanowires grown at 900°C for 30 min using catalyst Au nanoparticles (50 nm) with different distances between source to substrate (A) 2 cm, (B) 4 cm, and (C) 6 cm. Image (D) nanowires grown at 800°C for 30 min using catalyst Au nanoparticles (20 nm) with 4 cm source to substrate distance. Reproduced with permission from M. Kumar, V. Kumar, R. Singh, Diameter tuning of betaGa2O3 nanowires using chemical vapor deposition technique, Nanoscale Res. Lett. 12 (2017) 184 with author’s copyright and from M. Kumar, S. Kumar, N. Chauhan, D.S. Kumar, V. Kumar, R. Singh, Study of GaN nanowires converted from beta-Ga2O3 and photoconduction in a single nanowire, Semicond. Sci. Technol. 32 (2017) 085012 with permission of IOP Publishing Ltd.

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source and the substrate, and FESEM image (D) represents the smaller diameter of nanowires obtained through selecting smaller temperature (such as 800°C) and smaller catalyst nanoparticle size (such as 20 nm). Conclusively, diameters of β-Ga2O3 nanowires can be tuned from several hundreds of nanometers to a few tens of nanometers by exploiting growth conditions in the CVD reactor. Further, the lengths of nanowires grown with a catalyst are usually greater than nanowires grown without a catalyst. Xiang’s group [65] has shown β-Ga2O3 nanowires grown by a mixture of Ga and Ga2O3 (ratio 4:1) in the presence of oxygen at high temperatures (1000–1200°C for 1 h) with and without a catalyst. Nanowires grown without a catalyst had lengths of 10–20 μm, whereas the lengths of nanowires grown with a catalyst (obtained from NiCl2H2O/ethanol solution) were up to 500 μm. Kumar et al. [56] investigated the effect of growth time on structural properties of synthesized β-Ga2O3 nanowires where VLS is the dominant growth mechanism. Average diameters of as-grown nanowires were 55 nm, 62 nm, and 81 nm for growth durations of 1, 2, and 4 h, respectively. Structural properties of β-Ga2O3 nanowires having a monoclinic crystal structure are shown by high-resolution TEM measurements (Fig. 5.4). It was observed that the crystalline quality of nanowires improved with increasing growth time. HRTEM investigation of a single β-Ga2O3 nanowire indicates high crystalline   quality with the 392 direction of the wire axis. Chang et al. [45] showed β-Ga2O3 nanowires with diameters and lengths in the range of 20–40 and 0.4–2.0 μm, respectively, where TEM study indicated the single-crystalline structure with lattice spacing
of
0.46 nm corresponding to 201 planes of monoclinic β-Ga2O3. Xiang et al. [65] showed β-Ga2O3 nanowire (diameters and lengths in the range of 30–80 nm and 10–20 μm, respectively) with a monoclinic single-crystalline structure grown uniformly in the [010] direction. Tien et al. [55] showed β-Ga2O3 nanowires where HRTEM study suggested monoclinic single crystals structure of nanowires grown in the [110] direction. Further, Raman spectroscopy is one of the other key characterization techniques to understand structural properties and various phonon modes of β-Ga2O3 nanostructures. In the case of β-Ga2O3, 27 optical modes at Г-point (Brillouin zone center) with irreducible representation [68] are represented as follows: Г opt ¼ 10Ag + 5Bg + 4Au + 8Bu

(5.1)

The Au and Bu symmetry are infrared active modes whereas Ag and Bg symmetry are Raman active modes. A typical room-temperature Raman spectrum obtained from β-Ga2O3 nanowire is shown in Fig. 5.5A. The Raman spectrum consisted of 11 Raman active modes around 143.7 (Bg), 168.4 (Ag), 198.9 (Ag), 319.3 (Ag), 345.6 (Ag), 415.4 (Ag), 474.3 (Bg), 628.7 (Ag), 651.2 (Bg), 657.2 (Ag), and 765.4 (Ag) cm1. These active modes can be classified into three categories: low (below 200 cm1), mid (300–500 cm1), and high (above 500 cm1) frequency modes. Low frequency modes (143.7, 168.4, and 198.9 cm1) are associated with libration and translation of tetrahedra-octahedra chains based on the octahedral (Ga2O6) and tetrahedral (GaO4) subunits in β-Ga2O3. Mid frequency modes (319.3, 345.6, 415.4, and 474.3 cm1)

Fig. 5.4 TEM images (A–C) showing the morphology of Au nanoparticle catalyzed β-Ga2O3 nanowires indicating VLS growth mechanism. (D) HRTEM image of a single β-Ga2O3     nanowire recorded along 136 zone axis indicating direction of the wire axis 392 shown by a white arrow. (E) Fast Fourier-Transform (FFT) of the single β-Ga2O3 nanowire in (D) specifying the orientation of the wire axis by the streaks. Reproduced with permission from S. Kumar, V. Kumar, T. Singh, A. Hahnel, R. Singh, The effect of deposition time on the structural and optical properties of beta-Ga2O3 nanowires grown using CVD technique, J. Nanopart. Res. 16 (2013) with permission of Springer Nature.

Fig. 5.5 Raman spectra of (A) β-Ga2O3 nanowire and (B) reference single crystal of β-Ga2O3. Reproduced with permission from S. Kumar, G. Sarau, C. Tessarek, M.Y. Bashouti, A. Hahnel, S. Christiansen, R. Singh, Study of iron-catalysed growth of beta-Ga2O3 nanowires and their detailed characterization using TEM, Raman and cathodoluminescence techniques, J. Phys. D. Appl. Phys. 47 (2014) 435101 with permission of IOP Publishing Ltd.

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are assigned to deformation of Ga2O6 octahedra. The high frequency modes (628.7, 651.2, 657.2, and 765.4 cm1) are attributed to stretching and bending of GaO4 tetrahedra. Choi et al. [36] and Zang et al. [32] have suggested that Raman peaks of β-Ga2O3 nanorods and β-Ga2O3 powder have approximately similar positions. However, Gao et al. [69] and Rao et al. [70] suggested that comparative Raman spectra of Ga2O3 nanostructures and Ga2O3 powder have large Raman shift (red-shift [69] and blue shift [70] in the range of several cm1). This anomaly has been clarified by comparative Raman spectra of β-Ga2O3 nanowire and reference bulk β-Ga2O3 single crystal in a recent study [60], as represented in Fig. 5.5A and B. This indicates that Raman peak positions of β-Ga2O3 nanowires and their bulk counterpart do not differ much due to size reduction effects (with the diameter well above the quantum confinement range). One unique advantage of nanowires over their bulk counterpart is fewer structural defects. In the case of bulk films, structural defects generally resulted due to the large lattice mismatch between the growing film and substrate. On the other hand, nanowires usually have negligible strain, dislocations, and other structural defects owing to the nature of free-standing nanowire growth [4]. Kumar et al. [57] showed the growth of β-Ga2O3 nanowires on different substrates such as Si, GaN/sapphire, and sapphire. A bandgap of β-Ga2O3 nanowires grown on these different substrates estimated by UV-VIS-NIR spectrophotometer is shown in Fig. 5.6A–C (plots between [F(R)hν]2 vs photon energy). In the bandgap estimation, the generated reflectance curve from a UV-VIS-NIR spectrophotometer is converted to Kubelka-Munk (K-M) function F(R) and then the bandgap (Eg) is correlated with the K-M function F(R) by the well-known Eq. [71] as follows: ½FðRÞhυ2 ¼ B hυ  Eg




(5.2)

where B is a constant. The estimated bandgaps from curves are in the order of 4.76–4.83 eV for β-Ga2O3 nanowires on different substrates (Si, GaN/sapphire, and sapphire) [57], indicating fairly comparable bandgap values.

5.2.3 Optical properties of β-Ga2O3 nanostructures Optical emission of β-Ga2O3 nanostructures contains band-to-band emission and usually sub-bandgap emission having blue, green, and red emission bands. The photoluminescence (PL) spectrum of β-Ga2O3 nanostructures at room temperature has been reported by Li et al. [13] showing intrinsic emission at 265 nm (
4.7 eV) and 278 nm (
4.5 eV) nm under an excitation wavelength of 230 nm. Anisotropy of the monoclinic phase is responsible for emission of these two peaks. In the case of sub-bandgap emission, it is suggested in a number of reports that the blue emission is related to recombination of electron and hole (due to the presence of donor states resulting from oxygen vacancies and acceptor states resulting from gallium vacancy or gallium-oxygen vacancy), the green emission due to the presence of impurities (such as Be, Ge, Sn, Li, Zr, and Si), and the red emission due to the presence of nitrogen impurity in β-Ga2O3 nanostructures [38, 72, 73].

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Fig. 5.6 A plot between [F(R)hν]2 versus photon energy (hν) for the determination of direct bandgap of as-grown β-Ga2O3 nanowires grown on (A) Si, (B) GaN/sapphire, and (C) sapphire substrates, respectively. Reproduced with permission from S. Kumar, C. Tessarek, S. Christiansen, R. Singh, A comparative study of beta-Ga2O3 nanowires grown on different substrates using CVD technique, J. Alloys Compd. 587 (2014) 812–818 with permission of Elsevier.

Cathodoluminescence (CL) is also very useful to investigate the optical properties of β-Ga2O3 nanostructures. A CL spectroscopy of β-Ga2O3 nanostructures containing nanowires/nanosheets in UV-visible range is shown in Fig. 5.7A. The CL spectrum showed a strong broad UV-blue emission centered around 2.64 eV, deconvoluted into three Gaussian peaks at 2.44, 2.60, and 3.01, respectively. Energy peaks at 2.60 eV and 3.01 eV were related to the blue band whereas an energy peak at 2.44 eV was related to the green emission which results from the defect states present inside the bandgap of β-Ga2O3 nanostructures. Fig. 5.7(B) is an energy band diagram, demonstrating the UV-blue and red emissions corresponding to β-Ga2O3 nanostructures. Guzman-Navarro’s group [74] studied UV-blue emission (at 3.31 eV) in β-Ga2O3 nanostructures and suggested that the thermal treatment of samples enhance the UV emission with quenching of the blue band due to elimination of point defects in β-Ga2O3 nanostructures. Nogales et al. [75] observed a UV-blue emission along

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Fig. 5.7 (A) CL spectrum of self-catalyzed β-Ga2O3 nanostructures on spin coated Ga2O3 film/ sapphire substrate showing sub-bandgap emission with deconvoluted emission bands. (B) Energy band diagram for β-Ga2O3 explaining the UV-blue and red emissions. Reproduced with permission from S. Kumar, V. Kumar, T. Singh, A. Hahnel, R. Singh, The effect of deposition time on the structural and optical properties of beta-Ga2O3 nanowires grown using CVD technique, J. Nanopart. Res. 16 (2013) with permission of Springer Nature.

with a red emission at 1.73 eV in β-Ga2O3 nanowires. In the case of isoelectronic (In, Al) doped β-Ga2O3 nanowires, a blue shift in optical emission for In as well as Al doping has been demonstrated using comparisons with undoped β-Ga2O3 nanowires [76].

5.2.4 Application of β-Ga2O3 nanostructures In the past few years, various applications of β-Ga2O3 nanostructures such as photodetectors [9–15], gas sensors [15–25], and FETs [26–28] have been demonstrated as follows. The high performance of bridged Ga2O3 nanowires for solar-blind photodetection has been shown by Li et al. [13]. The high photocatalytic activity of β-Ga2O3 nanorods under UV irradiation was reported by Reddy et al. [29]. Catalytic activity of thermosensitive Ga2O3 nanorods was presented by Kumar et al. [77]. The pH sensor based on Ga2O3 nanowires has been shown by Ling et al. [30], and they suggested possibilities for improvement in performance of pH sensors by nanowire dimensions related to sensing capabilities. FETs based on β-Ga2O3 nanowires have been demonstrated in the reports [26–28]. Villora et al. [27] observed an increase in conductance with an increase in back gate voltage, and a high on-off current ratio (Ion/Ioff  105) has been achieved at gate voltage from 0 to 7.5 V with drain voltage of 0.2 V. The measured electron mobility in their case was
65.4 cm2/Vs. Gas sensors of β-Ga2O3 nanostructures have been developed [15–25], illustrating there sensing capabilities. Fig. 5.8A represents the fabricated gas sensor based on β-Ga2O3 nanowires with platinum electrodes; this sensor is used for effective sensing of oxygen (O2) and carbon monoxide (CO) gases. Fig. 5.8B showed the response of O2 gas by an as-fabricated gas sensor device operated at 300°C; the resistance of device increased with exposure to O2. Response of the CO gas pulse by a sensor device

Fig. 5.8 (A) Schematic diagram of Ga2O3 multiple nanowire-based gas sensor, (B) dynamic response of fabricated gas sensor to O2 gas pulse at 300° C, and (C) dynamic response of fabricated gas sensor to CO gas pulse at 100°C. Reproduced with permission from Z.F. Liu, T. Yamazaki, Y. Shen, T. Kikuta, N. Nakatani, Y.X. Li, O-2 and CO sensing of Ga2O3 multiple nanowire gas sensors, Sensors Actuators B Chem. 129 (2008) 666–670 with permission of Elsevier.

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operated at 100°C is shown in Fig. 5.8C. The O2 and CO gases exhibit the opposite response detected by a β-Ga2O3 multiple nanowire-based gas sensor operated at a low temperature range (100–500°C) [16]. Photodetectors based on β-Ga2O3 nanowires have been demonstrated by a number of groups [9–15]. Feng and co-workers [10] have observed the photoresponse in a fabricated photodetector when the device was exposed to illumination by an UV lamp. The measured dark current was in the order of pico-amperes (pA), which was enhanced to several nano amperes under illumination. Response time, recovery time, and sensitivity were about 0.22 s, 0.09 s, and
1000, respectively. Lopez et al. [78] studied an ultraviolet light selective frequency photodetector based on β-Ga2O3 nanowires, and estimated the maximum responsivity
1.2 mA/W at 8 V for Sn doped β-Ga2O3 nanowires. Kumar et al. [11] have also studied photoconduction in a single β-Ga2O3 nanowire based on metal-semiconductor-metal photodetector and a photocurrent three orders higher than the dark current was observed. Du et al. [79] have studied a deep ultraviolet photodetector based on the β-Ga2O3 nanowire network. Li et al. [13] demonstrated a solar-blind photodetector based on bridged β-Ga2O3 nanowires. Photoconduction properties of the bridged β-Ga2O3 nanowires are shown in Fig. 5.9. Four orders of photocurrent to dark current ratio under UV illumination (
2 mW cm2 at 254 nm with period of 60 s at bias of 50 V) were shown (Fig. 5.9A). The photocurrent decay process is depicted in Fig. 5.9B. The measured dark current was below pA (Fig. 5.9C). The optical response wavelength region of bridged β-Ga2O3 nanowires (Fig. 5.9D) is well below the lowest wavelength of solar spectrum on Earth, confirming that the device is blind to solar light.

5.3

Functional nanowires based on β-Ga2O3

To realize the nanoscale device applications based on functional nanowires of β-Ga2O3 nanowires such as coaxial GaN/β-Ga2O3 nanowires, ZnGa2O4 nanowires, and β-Ga2O3 nanowires mediated by high-quality long GaN nanowires, further discussion on functional nanowires is presented below.

5.3.1 Coaxial β-Ga2O3/GaN nanowires through ammonification of β-Ga2O3 nanowires One of the approaches for obtaining coaxial β-Ga2O3/GaN nanowires involves postsynthesis phase transformation processes [80], where thermal nitridation of β-Ga2O3 in presence of ammonia (NH3) gas is useful to obtain phase transformations. Phase transformation of β-Ga2O3 to GaN can be realized by chemical reactions [81]: Ga2 O3 ðsÞ + 2H2 ðgÞ ! Ga2 OðsÞ + 2H2 OðgÞ

(5.3)

Ga2 OðsÞ + 2NH3 ðgÞ ! 2GaNðsÞ + 2H2 OðgÞ + 2H2 ðgÞ

(5.4)

Fig. 5.9 See legend on next page.

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2 Ammonia gas can be decomposed into NH radical, and H2 at temper2 radical, NH atures above 800°C, which provides the active source of nitrogen atoms. In the pres ence of H2, NH2 2 , and NH , there is a thermodynamical tendency of the Ga2O formation by reduction of β-Ga2O3 [82]. Therefore, the GaN formation from β-Ga2O3 to obtain coaxial β-Ga2O3/GaN nanowires can progress through two-step chemical reactions as shown by Eqs. (5.3) and (5.4). Ammonification temperature is one of the important parameters for obtaining coaxial β-Ga2O3/GaN nanowires through the conversion process. Ammonification temperature provides the required energy for diffusion of thermally activated atoms through a solid where atoms jump from lattice site to lattice site. Control in ammonification temperature for conversion from β-Ga2O3 nanowires to GaN nanowires offers the advantage of coaxial structure formation. Ammonification of β-Ga2O3 nanowires at temperature ranges from 900°C to 1050° C for 1 h [83], as shown in Fig. 5.10(I: A–D) and corresponding XRD pattern in Fig. 5.10(II: A–D). The onset of decomposition of β-Ga2O3 at the surface can be initiated from 350°C in the presence of H2 and N2 gases [84]. Therefore, a higher ammonification temperature (900–1050°C) increases the probability of gallium desorption from the surface of the nanowire, which is responsible for nanowire surface alterations due to ammonification. However, ammonified nanowires retain their nanowire-like shape. It must be noted that the nanowire’s diameter plays a crucial role in retaining a nanowire-like shape [85]. The progressive conversion of β-Ga2O3 to the GaN phase can be obtained when the ammonification temperature is increased to an appropriate range (such as 900–1050°C), and composition of coaxial β-Ga2O3/GaN nanowires can be controlled by selecting an appropriate ammonification temperature. The progressive conversion of β-Ga2O3 to the GaN phase has been demonstrated through XRD, as shown in Fig. 5.10(II: A–D). XRD peaks related to both β-Ga2O3 as well as GaN at lower ammonification temperature of 900°C indicate the coaxial β-Ga2O3/GaN nanowires formation, whereas full conversion to GaN nanowires can be obtained at higher temperatures (such as 1050°C) through the ammonification process.

Fig. 5.9 Photoconduction properties of the bridged β-Ga2O3 nanowires. (A) Time-dependent photoresponse of the bridged β-Ga2O3 nanowires measured in dry air under illumination (
2 mW cm2 at 254 nm) with a period of 60 s at a bias voltage of 50 V. The photocurrent to dark current ratio was
3  104. (B) Photocurrent decay process of the device. The inset shows the enlarged decay edge. (C) IdV characteristics of the bridged β-Ga2O3 nanowires in dark (squares), under 365 nm light (triangles), and under 254 nm light (circles). IdV curve measured under 254 nm light is plotted on a linear scale in the inset. (D) Spectral response of bridged β-Ga2O3 nanowires revealing that the device is blind to solar light. The dashed line indicates the lowest wavelength of the solar spectrum on Earth. Reproduced with permission from Y.B. Li, T. Tokizono, M.Y. Liao, M.A. Zhong, Y. Koide, I. Yamada, J.J. Delaunay, Efficient assembly of bridged beta-Ga2O3 nanowires for solar-blind Photodetection, Adv. Funct. Mater. 20 (2010) 3972–3978 with permission of John Wiley and Sons.

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(II) ∗(-313) ∗ (-603)

0.0 57 60

(-603)

∗ (-311)

0.2

0.2

0.1

0.1 35

40

0.0 57 60

0.3

0.3

0.2

0.2

0.1

0.1 35

40

0.0 57 60

0.3

0.3

0.2

0.2

0.1

0.1

0.0 30

30

40

0.3

0.0 30 (d)

35

0.1

0.3

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(c)

0.2

∗

0.0 30

0.3

(101)

0.1

(100)

(002)

0.2

(101)

(400) (100) ∗

Intensity (a.u.)

(b)

(002)

0.3

(a)

∗(-311)

(I)

35

40

0.0 57 60

40 50 60 2θ (degree)

70

Fig. 5.10 (I) FESEM images and (II) XRD patterns of GaN/β-Ga2O3 nanowires formed at ammonification temperatures of (A) 900°C, (B) 950°C, (C) 1000°C, and (D) 1050°C. Insets of (II): zoom segments with normalized intensity; Symbol “*” represents residual β-Ga2O3. Reproduced with permission from M. Kumar, G. Sarau, M. Heilmann, S. Christiansen, V. Kumar, R. Singh, Effect of ammonification temperature on the formation of coaxial GaN/Ga2O3 nanowires, J. Phys. D. Appl. Phys. 50 (2017) 035302 with permission of IOP Publishing Ltd.

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5.3.2 ZnGa2O4 nanowires through coaxial ZnO/β-Ga2O3 nanowires Zinc gallate (ZnGa2O4) is a semiconductor material having a wide bandgap of around 4.4 eV [86]. It has various applications including transparent conducting oxide, photocatalyst, and flat panel display as a blue phosphor [87, 88]. To obtain ZnGa2O4 nanowires through multistep processes [59], first β-Ga2O3 nanowires (diameters
150–300 nm, lengths up to tens of micrometers) are obtained using catalyst-mediated growth in CVD. In the second step, coaxial β-Ga2O3/ZnO nanowires are obtained using the atomic layer deposition technique by conformal deposition of the coaxial ZnO layer on β-Ga2O3 nanowires (diameters
300–500 nm). Finally, the as-obtained template is annealed at an appropriate temperature under oxygen ambient to activate an interfacial solid state reaction for the formation of ZnGa2O4 nanowires. The comparative Raman spectra containing peak positions related to different materials stages during ZnGa2O4 nanowire formation [59] are shown in Fig. 5.11. Raman spectra consist of ZnO film/sapphire (for comparison, Fig. 5.11A), β-Ga2O3 nanowires (Fig. 5.11B), coaxial β-Ga2O3/ZnO nanowires (Fig. 5.11C), and ZnGa2O4 nanowires (Fig. 5.11D). The presence of both ZnO and β-Ga2O3 phase-related peaks in Raman spectrum shown in Fig. 5.11C indicates the formation of coaxial β-Ga2O3/ZnO nanowires. Subsequently, high temperature annealing (1000°C for 1 h) of these coaxial

Fig. 5.11 Comparison of representative Raman spectra of (A) ZnO film on sapphire, (B) β-Ga2O3 nanowires, (C) coaxial ZnO/β-Ga2O3, and (D) converted ZnGa2O4 nanowires after annealing at 1000°C for 1 h. Reproduced with permission from S. Kumar, G. Sarau, C. Tessarek, M. Gobelt, S. Christiansen, R. Singh, Study of high quality spinel zinc gallate nanowires grown using CVD and ALD techniques, Nanotechnology 26 (2015) 335603.

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β-Ga2O3/ZnO nanowires results in significant change in the Raman spectra. Raman peaks related to β-Ga2O3 nanowires were eliminated, whereas new Raman peaks (Fig. 5.11D) around 330, 380, 410, 536, 580, 610, 670, and 710 cm1 were evolved. This indicated the formation of ZnGa2O4 nanowires through a process of interdiffusion. Further, ZnGa2O4 nanoparticles have been formed by Can et al. [89], and hollow spinel ZnAl2O4 nanotubes obtained through ZnO-Al2O3 core-shell nanowires have been shown by Fan and his coworkers [90].

5.3.3

β-Ga2O3 nanowires template mediated high-quality ultralong GaN nanowires

β-Ga2O3 nanowires can act as a potential template to obtain high-quality ultralong GaN nanowires as illustrated recently by Kumar et al. [91]. In this approach, each β-Ga2O3 nanowire acts as a substrate and in the presence of gallium metal, ammonia, and nitrogen gases at high temperatures (such as 1050°C), these β-Ga2O3 nanowires result in high-quality GaN nanowires. In this process, surface of β-Ga2O3 nanowires ammonified initially, resulting in GaN nucleation sites and therefore epitaxially GaN growths on nanowire surface are possible besides the phase conversion from β-Ga2O3 to GaN. One of the unique features of β-Ga2O3 nanowire for obtaining GaN nanowires is their long length, which is a well-known issue in VLS-grown GaN nanowires. Therefore, using β-Ga2O3 nanowires, very high-quality ultralong GaN nanowires with millimeter-length scales can be obtained. A photoconductor device based on an as-obtained high-quality single nanowire shows very high gain and responsivity. Current-voltage characteristic of a photoconductor device in the dark and under UV light illumination is shown in Fig. 5.12A. The performance metric of photoconduction is given by responsivity (R), gain (G), and normalized gain (GN) estimated using the following equations [92, 93]:
R ¼ Ilight  Idark =ðAl ∗Li Þ

(5.5)

G ¼ R  ðhν=ηqÞ

(5.6)

GN ¼ G

 2 L η ¼ μτη V

(5.7)

The terms are denoted as light receiving area (Al), light intensity (Li), quantum efficiency (η), electronic charge (q), and travel length of carriers between electrodes (L). The physical meaning of normalized gain is the optical to electrical conversion (η) with photocarrier transport (τμ), and avoids the dependence of voltage and feature size on R and G. The high values of R ¼ 5.1  103 A/W, G ¼ 2.5  104 (assuming η ¼ 1) and GN ¼ 0.68 cm2 V1 at bias voltage of 0.5 V are shown. In the case of nanowires, spatial separation of photogenerated electron-hole pair due to upward bending of electronic bands in the vicinity of nanowire surface causes the increased carrier life time [94].

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Fig. 5.12 (A) Current-voltage characteristics of single ultralong GaN nanowire obtained from β-Ga2O3 nanowires in the dark and in the presence of UV light. (B) Calculated responsivity and gain for the single nanowire photoconduction. Inset of (B) shows FESEM image of fabricated device under test. The Cr/Au contacts at the ends of nanowire are used for the photoconduction study. The other contacts (Ni/Au) on the device under test were meant for Schottky diode investigations and not used in this work. Reproduced with permission from M. Kumar, V. Kumar, R. Singh, Formation of ultralong GaN nanowires up to millimeter length scale and photoconduction study in single nanowire, Scr. Mater. 138 (2017) 75–78 with permission of Elsevier.

Therefore, prolonged life time and short carrier transit time (τt) between electrodes can lead to high gain and responsivity, as given by the following equation: G¼

τ τt

(5.8)

where τ is carrier lifetime and τt is carrier transit time between electrodes. In other words, carriers (electrons) are going through the photoconductor in a number of loops during carrier life time, resulting in high gain and responsivity.

5.4

Conclusions and future perspective

This chapter provided a brief overview for β-Ga2O3 nanostructures from a growth aspect to device applications. The outstanding properties of β-Ga2O3 such as large bandgap, high breakdown field, thermal and chemical stability, along with advantageous properties due to its nanostructures morphology such as large surface-to-volume ratio, fewer defects, and less strain makes it a potential material for development of high-performance nanoscale devices. β-Ga2O3 nanostructures have shown great promise for nanoscale devices such as deep-UV photodetectors, gas sensors, and FETs. In addition, functional nanowires based on β-Ga2O3 nanostructures can also be utilized for establishing the nanoscale device platform. However, research in β-Ga2O3

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nanostructures is not yet fully developed, and still requires a number of solutions for important issues as follows: l

l

l

The controlled doping in β-Ga2O3 nanostructures is one of the desirable requirements for their device application. Techniques such as MOCVD, MBE, and PLD could be most prominent for the doping aspect in β-Ga2O3 nanostructures during growth. Schottky diodes and field effect transistors based on β-Ga2O3 nanostructures need to be investigated or improved through studying current transport properties. Large-scale devices based on uniform assembly of β-Ga2O3 nanostructures for industrialscale applications are required for integration into the circuit with high scalability in the near future.

Acknowledgments The authors acknowledge the Department of Electronics and Information Technology, Government of India for partial financial support for this work and the Nanoscale Research Facility of IIT Delhi for providing various research facilities. Mukesh Kumar acknowledges the IIT Delhi for providing a research fellowship.

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[71] A.E. Morales, E.S. Mora, U. Pal, Use of diffuse reflectance spectroscopy for optical characterization of un-supported nanostructures, Rev. Mex. Fis. 53 (2007) 18–22. [72] L. Binet, D. Gourier, Origin of the blue luminescence of beta-Ga2O3, J. Phys. Chem. Solids 59 (1998) 1241–1249. [73] L.W. Chang, T.Y. Lu, Y.L. Chen, J.W. Yeh, H.C. Shih, Effect of the doped nitrogen on the optical properties of beta-Ga2O3 nanowires, Mater. Lett. 65 (2011) 2281–2283. [74] G. Guzman-Navarro, M. Herrera-Zaldivar, J. Valenzuela-Benavides, D. Maestre, CL study of blue and UV emissions in beta-Ga2O3 nanowires grown by thermal evaporation of GaN, J. Appl. Phys. 110 (2011) 034315. [75] E. Nogales, B. Mendez, J. Piqueras, Cathodoluminescence from beta-Ga2O3 nanowires, Appl. Phys. Lett. 86 (2005) 113112. [76] E. Nogales, B. Sanchez, B. Mendez, J. Piqueras, Cathodoluminescence study of isoelectronic doping of gallium oxide nanowires, Superlattice. Microst. 45 (2009) 156–160. [77] V.B. Kumar, R.K. Mishra, I.N. Pulidindi, Z. Porat, J.H.T. Luong, A. Gedanken, Preparation and catalytic activity of thermosensitive Ga2O3 Nanorods, Energy Fuel 30 (2016) 7419–7427. [78] I. Lopez, A. Castaldini, A. Cavallini, E. Nogales, B. Mendez, J. Piqueras, Beta-Ga2O3 nanowires for an ultraviolet light selective frequency photodetector, J. Phys. D. Appl. Phys. 47 (2014) 415101. [79] J.Y. Du, J. Xing, C. Ge, H. Liu, P.Y. Liu, H.Y. Hao, J.J. Dong, Z.Y. Zheng, H. Gao, Highly sensitive and ultrafast deep UV photodetector based on a beta-Ga2O3 nanowire network grown by CVD, J. Phys. D. Appl. Phys. 49 (2016) 425105. [80] J. Thangala, Z.Q. Chen, A. Chin, C.Z. Ning, M.K. Sunkara, Phase transformation studies of metal oxide nanowires, Cryst. Growth Des. 9 (2009) 3177–3182. [81] Z.H. Lan, C.H. Liang, C.W. Hsu, C.T. Wu, H.M. Lin, S. Dhara, K.H. Chen, L.C. Chen, C. C. Chen, Nanohomojunction (GaN) and nanoheterojunction (InN) nanorods on onedimensional GaN nanowire substrates, Adv. Funct. Mater. 14 (2004) 233–237. [82] A. Othonos, M. Zervos, C. Christofides, A systematic investigation into the conversion of beta-Ga2O3 to GaN nanowires using NH3 and H-2: Effects on the photoluminescence properties, J. Appl. Phys. 108 (2010) 124319. [83] M. Kumar, G. Sarau, M. Heilmann, S. Christiansen, V. Kumar, R. Singh, Effect of ammonification temperature on the formation of coaxial GaN/Ga2O3 nanowires, J. Phys. D. Appl. Phys. 50 (2017) 035302. [84] R. Togashi, K. Nomura, C. Eguchi, T. Fukizawa, K. Goto, Q.T. Thieu, H. Murakami, Y. Kumagai, A. Kuramata, S. Yamakoshi, B. Monemar, A. Koukitu, Thermal stability of beta-Ga2O3 in mixed flows of H-2 and N-2, Jpn. J. Appl. Phys. 54 (2015) 041102. [85] M. Kumar, S. Kumar, N. Chauhan, D.S. Kumar, V. Kumar, R. Singh, Study of GaN nanowires converted from beta-Ga2O3 and photoconduction in a single nanowire, Semicond. Sci. Technol. 32 (2017) 085012. [86] T. Omata, N. Ueda, K. Ueda, H. Kawazoe, New ultraviolet-transport Electroconductive oxide, ZnGa2O4 spinel, Appl. Phys. Lett. 64 (1994) 1077–1078. [87] Q. Liu, D. Wu, Y. Zhou, H.B. Su, R. Wang, C.F. Zhang, S.C. Yan, M. Xiao, Z.G. Zou, Single-crystalline, ultrathin ZnGa2O4 Nanosheet scaffolds to promote photocatalytic activity in CO2 reduction into methane, ACS Appl. Mater. Interfaces 6 (2014) 2356–2361. [88] S.H. Wu, H.C. Cheng, Preparation and characterization of nanosized ZnGa2O4 phosphors, J. Electrochem. Soc. 151 (2004) H159–H163. [89] M.M. Can, G.H. Jaffari, S. Aksoy, S.I. Shah, T. Firat, Synthesis and characterization of ZnGa2O4 particles prepared by solid state reaction, J. Alloys Compd. 549 (2013) 303–307.

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Properties of (In,Ga)2O3 alloys Holger von Wenckstern Universit€at Leipzig, Felix-Bloch-Institut f€ur Festk€orperphysik, Halbleiterphysik, Linnestraße 5, Leipzig, Germany

6

Chapter Outline 6.1 Introduction 119 6.2 Overview on crystal structures observed in (In,Ga)2O3 6.3 Lattice parameters of bulk material 122 6.3.1 6.3.2 6.3.3 6.3.4

120

Rhombohedral (InxGa1
x)2O3 122 Monoclinic β-(InxGa1
x)2O3 123 Cubic (GaxIn1
x)2O3 123 Hexagonal InGaO3 123

6.4 Thin film growth

125

6.4.1 Growth of (In,Ga)2O3 thin films with lateral composition spread by PLD 126 6.4.2 Growth and phase formation of (In,Ga)2O3 thin films 127

6.5 Deep-UV absorption and band gap engineering 130 6.6 Phonon modes 131 6.7 Dielectric function and index of refraction 134 6.8 Schottky barrier diodes 135 6.9 Photodetectors 137 6.10 Summary and outlook 139 References 141 Further reading 148

6.1

Introduction

Band gap engineering by alloying of different semiconducting materials is utilized for charge carrier localization within high-electron mobility transistors or quantum well structures or for tuning of emission energy and absorption onset in light emitters and detectors, respectively. Further, alloying allows for strain management in epitaxial growth. The miscibility of solid solutions depends on the equilibrium crystallographic structures of the binaries and is typically higher for isostructural solid solutions. If the binaries have different equilibrium crystal structures, phase separation will occur. In the case of (In,Ga)2O3 various polymorphs were reported for the binaries [1–3]. The equilibrium crystal structure of indium oxide is the cubic bixbyite structure [1]. At elevated temperatures and high pressure, a phase transition to α-In2O3 having rhombohedral corundum structure occurs [2]. A first comprehensive description of the Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00006-3 © 2019 Elsevier Inc. All rights reserved.

120

Gallium Oxide

polymorphism and transformation relations of gallium oxide was reported by Roy et al. [3]. α-Ga2O3 has rhombohedral corundum structure similar to α-In2O3. The monoclinic beta-gallia structure is the equilibrium crystal structure of gallium oxide, γ-Ga2O3 is a defective spinel structure. Roy et al. proposed a cubic bixbyite structure labeled δ-Ga2O3 [3], however, Playford et al. [4] demonstrated that it is not a distinct polymorph but rather a nanocrystalline modification of the ε-phase having an orthorhombic unit cell with space group Pna21 as suggested by Yoshioka et al. [5]. However, experimentally the symmetry of this ε-phase appears often hexagonal with space group P63mc [4]. The polymorphism of the group-III sesquioxides impacts alloy formation and results will not only depend on composition but also growth conditions and substrates used may lead to the stabilization of the one or the other polymorph. So far, alloys of indium and gallium oxide were reported for the rhombohedral α-phase, the cubic bixbyite phase, and the monoclinic β-phase. Further, a hexagonal high-pressure phase labeled InGaO3 II and being similar to hexagonal YAlO3 was reported by Shannon and Prewitt for powder samples [6] and observed later in thin films by Kranert et al. [7].

6.2

Overview on crystal structures observed in (In,Ga)2O3

Within the corundum structure (space group R3c), the cations are sixfold coordinated as shown in Fig. 6.1A, the Wyckoff notation of these lattice sites with site symmetry C3 is 12c [9, 10]. There are six formula units in the unit cell (two formula units per primitive unit cell) [9, 11]. The crystal structure has optically uniaxial behavior. The irreducible representation of optical phonon modes is Γ ¼ 2A1g + 2A1u + 3A2g + 2A2u + 5Eg + 4Eu [10]. Raman active modes are the A1g and the Eg modes. Binary thin films of α-In2O3 [12] and α-Ga2O3 [13] were realized on α-Al2O3 rendering the rhombohedral modification especially promising for band gap engineering. The monoclinic β-gallia structure, depicted in Fig. 6.1B, has space group C2/m. It contains four formula units per unit cell (two within the primitive unit cell) [14]. There exist two equally portioned distinct cation sites, one of them is fourfold (Ga(1)) and one of them is sixfold (Ga(2)) coordinated. Further, the oxygen atoms occupy three different lattice sites labeled O(1), O(2), and O(3). O(1) and O(2) are threefold coordinated, O(3) bonds to four cations. The MO4 tetrahedra and the strongly distorted MO6 octahedra form chains along the crystallographic b-direction that have been claimed as source for electrical anisotropy of the material [15]. However, measurements on single-crystalline bulk material revealed isotropic electric transport properties [16, 17]. The monoclinic β-gallia structure is biaxial and resulting optically anisotropic properties were described in detail in Refs. [18–21]. Optical phonon modes have the irreducible representation Γ ¼ 10Ag + 5Bg + 8Bu + 4Au, the 15 g-modes are Raman active [7]. The unit cell of the bixbyite structure, shown in Fig. 6.1C, belongs to space group I3a, contains 16 formula units corresponding to 80 atoms [22] (8 formula units are within the primitive unit cell [23]). Two different sixfold coordinated cation sites exits, 8 are at (b)-site [labeled In(1)] with S6 site symmetry and 24 are at the (d)-site [labeled

Properties of (In,Ga)2O3 alloys

121

c

b

a c

Ga(1) a b c

c b

a

a b

O(1)

O(2)

c

a b

O(3)

Ga(2)

(A)

(B) 8b

24d c a

b

c c c a

(C)

b

a

c b

a

a

c b

a

b

b

(D)

Fig. 6.1 (A) Rhombohedral crystal structure of corundum-like α-M2O3 (M: In, Ga; shown in bluish color) and nearest-neighbor configuration of the MO6 octahedron. (B) Monoclinic β-gallia crystal structure and nearest-neighbor configuration of the Ga(1)O4 tetrahedron, the Ga(2)O6 octahedron, of the threefold coordinated O(1) and O(2) and the fourfold coordinated O(3) lattice sites. (C) Cubic bixbyite lattice structure, the different cation positions are highlighted in blue (Wyckoff position 8b) and green (Wyckoff position 24d), respectively. The oxygen environment of the different cation sites is illustrated by the corner positions on a cube, the actual atomic positions are shown below. (D) Crystal structure of hexagonal InGaO3 II and bonding configuration of Ga (green) and In (purple). Ga is fivefold and In sixfold coordinated. Black lines indicate the unit cell for each crystal structure. The representations were created using VESTA [8].

In(2)] with site symmetry C2 [14]. For both sites, the surrounding oxygen atoms are approximately located at corners of a cube with the cation in its center. For In(1) two corners on opposite sides of the body diagonal of the cube are unoccupied and for In(2) two corners on opposite sides of a face diagonal are empty [14]. Due to the cubic crystal structure material properties are isotropic. The large number of atoms within the primitive unit cell results in a large number of optical phonon modes with representation Γ ¼ 4Ag + 4Eg + 14Tg + 5Au + 5Eu + 16Tu [23]. The 22 g-modes are Raman active. In the hexagonal InGaO3 II phase [Fig. 6.1D, space group P63/mmc], the indium atoms are sixfold coordinated. The surrounding oxygen atoms form a trigonal antiprism. Ga is fivefold coordinated; the surrounding oxygen atoms form a trigonal bipyramid. Along the c-direction, these cation polyhedra form a layer structure that might result in anisotropic properties and possibly a net polarization in strained layers. For indium content between x ¼ 0.45 and x ¼ 0.55 calculations of Maccioni and Fiorentini resulted in zero polarization for unstrained material [24]. Optical phonon modes have the irreducible representation Γ ¼ 1A1g + 3A2u + 1E1g + 3E1u + 3E2g + 2E2u, the g-modes are Raman active [25].

122

Gallium Oxide

In summary, the crystal structures reported for (In,Ga)2O3 have fourfold, fivefold, ˚ and sixfold coordinated cation sites. Indium is with its ionic radius of rIn3+ ¼ 0.79 A ˚ much larger than Ga with rGa3+ ¼ 0.62 A and will be preferentially incorporated at octahedral lattice site. Further, threefold and fourfold coordinated anion sites exist. The various cation and anion environments provide means to stabilize the one or the other phase by doping/alloying with cations and anions, respectively, of appropriate size and valency. So far, the rhombohedral corundum structure was stabilized for both binaries and in principle offers isostructural alloying of indium gallium oxide. Further, the stabilization of (In,Ga)2O3 and (Al,Ga)2O3 thin films in the polar ε-phase appears highly interesting [26–30]. Formation of a two-dimensional (2D) electron gas with potentially high carrier density can be induced by the polarization difference at ε-(In,Ga)2O3/ε-(Al,Ga)2O3 heterointerfaces presenting another highly interesting field of research.

6.3

Lattice parameters of bulk material

Early studies on phase relationships of sesquioxides date back to the 1920s of the last century and were carried out in the laboratories of V.M Goldschmidt at University of Oslo [8]. Schneider et al. [31] predicted subsolidus binary-phase diagram for the In2O3-Ga2O3 system and proposed solid solution within the cubic bixbyite structure for In2O3 with Ga2O3 admixtures up to about 15 mol%. Solid solutions with the monoclinic lattice type were predicted for Ga2O3 with In2O3 admixtures up to about 13 mol % [31]. Further, an unknown structure being similar to κ-alumina was observed for Ga2O3 mole fractions between about 57% and 72% [31]. Higher (lower) mole fractions lead to the occurrence of the beta-gallia (bixbyite) structure. MacDonald et al. [32] observed two monoclinic type structures in mixed crystals and suggested that these represent two polymorphs with a 1:1 composition ratio. These studies provided first insights in the phase diagram of the In2O3-Ga2O3 system triggering more sophisticated and systematic studies on phase formation, stability, and the dependence of lattice parameters on alloy composition.

6.3.1 Rhombohedral (InxGa12 x)2O3 Data on bulk (InxGa1
x)2O3 in rhombohedral modification are not available and literature on heteroepitaxial films is scarce. The incorporation of In2O3 into α-Ga2O3 layers grown heteroepitaxially on α-Al2O3 was reported up to 8 at.% and leads to a ˚ , howdecrease of the a-lattice constant [33]. A linear fit yields a(x) ¼ (5.35
4.85x)A ever, it has to be noted that only three data points are available and hence the fit provides a first estimate only. Further, the influence of strain and thickness of these heteroepitaxial layers needs to be determined. Phase-pure rhombohedral samples were also obtained for x  0.67. Films with 0.08 < x < 0.67 showed phase separation (occurrence of the β-gallia structure).

Properties of (In,Ga)2O3 alloys

123

6.3.2 Monoclinic β-(InxGa12 x)2O3 The admixture of In2O3 to monoclinic β-Ga2O3 was investigated by Shannon and Prewitt [6], Edwards et al. [34, 35], and Kranert et al. [7] using powder samples. The incorporation of indium leads to a linear increase of the a, b, and c lattice parameter, while the lattice parameter β decreases linearly. Shannon and Prewitt reasoned that indium is incorporated at octahedral lattice site up to the composition GaInO3 [6]. It was Shannon and Prewitt [6] as well that grew solid solutions using the flux method with either an approximate 1:1 composition or low indium admixture of 5 mol%. The millimeter-sized samples crystallized in the monoclinic beta-gallia structure. Singlecrystalline solid solutions were obtained by Vasyltsiv et al. [36, 37] and Tomm et al. using the floating zone method [38]. Phase-pure monoclinic samples were obtained for x < 35%; for InGaO3 transparent reddish crystals with unidentified structure are obtained. The diffraction pattern was reported to be similar to that of the κ-alumina structure [38]. Using structural data acquired by X-ray diffraction (XRD) Vasyltsiv et al. and later Edwars et al. confirmed that In3+ preferentially occupies octahedral lattice site [35, 37], validated by perturbed angular correlation measurements of Pasquevich et al. [39]. However, in several studies, it was found that phase separation occurs before all of the Ga(2)-lattice sites are occupied by indium [35, 37]. The solubility limit was determined for the powder samples to be about 44 mol% for a temperature of 1000°C. It decreases to about 41.4 mol% for 1400°C. The variation of lattice parameters is collected in Fig. 6.2 and includes literature data from X-ray studies on powder as well as bulk single-crystalline samples. All parameters exhibit a linear dependence on the alloy composition; the dependencies are provided in the figure and in Table 6.1. The relative change of lattice constants is strongest along the a- and b-direction. For (In0.4Ga0.6)2O3 a and b increased about 4.5% whereas c increased by about 2.7%. Deviations from the linear behavior would be expected for indium admixtures higher than 50% (not realized experimentally) due to the occupation of tetrahedral sites by indium.

6.3.3 Cubic (GaxIn12 x)2O3 Incorporation of Ga2O3 into cubic In2O3 ceramics was investigated by Edwards et al. and Regoutz et al. and XRD revealed a decrease of the lattice constant with increasing Ga content [35, 40]. Results of the two studies are depicted in Fig. 6.3 together with a linear fit of the data. The solubility of Ga2O3 into cubic In2O3 is much lower than that of In2O3 into monoclinic Ga2O3; the solubility limit was reached between 6 [40] and 7.5 at.% [35].

6.3.4 Hexagonal InGaO3 High-pressure synthesis of hexagonal samples with a 1:1 In:Ga composition was reported by Shannon and Prewitt [6]. The source material was heated at 1200°C at a pressure of ˚ and c ¼ 12.039 A ˚ . The density of hexag65 kbar. The lattice parameters are a ¼ 3.31 A
3
3 onal InGaO3 is with 6.756 g cm higher than the 6.447 g cm of monoclinic InGaO3 explaining the necessity of high pressure during synthesis of that polymorph.

Fig. 6.2 See legend on next page.

124 Gallium Oxide

Properties of (In,Ga)2O3 alloys

125

Dependence of the monoclinic lattice parameters on indium incorporation in (InxGa12 x)2O3 (rows 1–4)

Table 6.1

˚ a(x) ¼ (12.2247 + 1.4161  x)A ˚ b(x) ¼ (3.0422 + 0.3447  x)A ˚ c(x) ¼ (5.8014 + 0.3889  x)A β ¼ (103.876 – 3.3873  x) ˚ a(x) ¼ (10.111 – 0.7781  x)A The solubility limit is reached for x  0.43. Dependence of the a-lattice constant on gallium incorporation in cubic (GaxIn1
x)2O3 (bottommost column). The solubility limit is reached at x  0.07.

6.4

Thin film growth

Thin films of (InxGa1
x)2O3 have been obtained by molecular beam epitaxy (MBE), metal organic chemical vapor deposition (MOCVD), the sol-gel method (SGM), pulsed-laser deposition (PLD), the molecular precursor method (MPM), mist chemical vapor deposition (MIST-CVD), and magnetron sputtering (MS). However, most of the results discussed in the following chapters were obtained on thin films grown by continuous composition spread PLD (CCS-PLD) [41]. Up to today, this method is used in the laboratories of the semiconductor physics group of Universit€at Leipzig, only. Therefore, we provide a short introduction to combinatorial material research and CCS-PLD.

Fig. 6.2 Dependence of the a-, b-, c-, and β-lattice parameters on indium content in monoclinic (InxGa1
x)2O3 powder samples (Shannon, Edwards, Kranert) and single crystals grown by the floating zone method (Vasyltsiv, Tomm). The solid lines are linear fits to the complete data set. The fit parameters are provided in each subfigure. Dashed lines indicate saturation values of lattice parameters. Data from R.D. Shannon, C.T. Prewitt, Synthesis and structure of phases in the In2O3-Ga2O3 system, J. Inorg. Nucl. Chem. 30 (1968) 1389–1398, doi:10.1016/0022-1902(68)80277-5; V.I. Vasyltsiv, Y.I. Rym, Y.M. Zakharko, Optical absorption and photoconductivity at the band edge of β-Ga2
xInxO3, Phys. Stat. Sol. (B). 195 (1996) 653–658, doi:10.1002/pssb.2221950232; D.D. Edwards, T.O. Mason, Subsolidus phase relations in the Ga2O3-In2O3-SnO2 system, J. Am. Ceram. Soc. 81 (1998) 3285–3292, doi:10.1111/j.1151-2916.1998.tb02769.x; Y. Tomm, J.M. Ko, A. Yoshikawa, T. Fukuda, Floating zone growth of beta-Ga2O3: a new window material for optoelectronic device applications, Sol. Energy Mater. Sol. Cells. 66 (2001) 369–374; C. Kranert, J. Lenzner, M. Jenderka, M. Lorenz, H. von Wenckstern, R. Schmidt-Grund, et al., Lattice parameters and Raman-active phonon modes of (InxGa1–x)2O3 for x < 0.4, J. Appl. Phys. 116 (2014) 013505, https://doi.org/10.1063/1. 4886895.

126

Gallium Oxide

Fig. 6.3 Dependence of the a-lattice constant of cubic (GaxIn1 2 x)2O3 on gallium admixture. The dashed line presents a linear fit to the data; fitting parameters are shown in the figure. Data from D.D. Edwards, T.O. Mason, Subsolidus phase relations in the Ga2O3-In2O3-SnO2 system, J. Am. Ceram. Soc. 81 (1998) 3285–3292, doi:10.1111/j.1151-2916.1998.tb02769.x; A. Regoutz, R.G. Egdell, D.J. Morgan, R.G. Palgrave, H. Tellez, S.J. Skinner, et al., Electronic and surface properties of Ga-doped In2O3 ceramics, Appl. Surf. Sci. 349 (2015) 970–982, https://doi.org/10.1016/j.apsusc.2015.04.106.

6.4.1 Growth of (In,Ga)2O3 thin films with lateral composition spread by PLD The creation and analysis of material libraries has proven highly efficient in the exploration and optimization of novel functional materials [41–52]. Concerning research on solid-state thin films combinatorial material synthesis was reported for evaporation [53], MBE [51], MS [42, 44], PLD [41, 54, 55], and so on. Especially sputtering is highly convenient for creating thin films with CCS since state-of-the-art systems are equipped with multiple magnetrons to be used for co-sputtering. For PLD, the ablation of multiple targets is possible by beam splitting [56, 57] or laser line scanning [58]. In other combinatorial PLD systems, targets have to be exchanged and movable masks or apertures, behind which the substrate is partially shadowed, are used to provide lateral control of the particle flux on the substrate [54, 55]. In 2013, von Wenckstern et al. established a facile CCS-PLD approach for the creation of material libraries relying on the ablation of segmented PLD targets. This CCS-PLD method can be used in existing off-set PLD system without any modification of the hardware [41]. A spatial offset ε between the substrate center and the centerline of the expanding plasma plume and a synchronized rotation of substrate and target are the necessary conditions for obtaining a CCS. The distribution of the different elements originating from different target segments on the substrate depends besides thermodynamical aspects on the background pressure, the target-to-substrate distance, and the offset ε [41]. As example we depict the lateral variation of the indium content of a

Properties of (In,Ga)2O3 alloys

127

Fig. 6.4 (A) False-color representation of the indium concentration within a 2 in. in diameter (In1 2 xGax)2O3 thin film on (100)MgO. Black dots indicate sites of energy-dispersive X-ray (EDX) spectroscopy measurements, the false-color data are an interpolation of the EDX results. The white arrow indicates the direction of the compositional gradient and the direction of a more detailed EDX line scan depicted in (B). Data from C. Kranert, J. Lenzner, M. Jenderka, M. Lorenz, H. von Wenckstern, R. Schmidt-Grund, et al., Lattice parameters and Raman-active phonon modes of (InxGa1–x) 2O3 for x < 0.4, J. Appl. Phys. 116, 2014, 013505, https://doi.org/10.1063/1.4886895.

(InxGa1
x)2O3 thin film grown by CCS-PLD using a target consisting of semicircular In2O3 and Ga2O3 segments in Fig. 6.4. The sample was deposited at a growth temperature of 650°C, an oxygen pressure 0.08 mbar, a target-to-substrate distance of 90 mm, and an offset ε ¼ 18 mm on a 2 in. in diameter (100)MgO substrate [7]. Energy-dispersive X-ray (EDX) spectroscopy was used for spatially resolved chemical analysis. The indium concentration varies between 0.1 and 0.8 and shows a slight S-shaped dependence along the gradient direction [cf. Fig. 6.4B] being in agreement with calculations [41]. Along lines perpendicular to the gradient direction, the indium concentration is in principle constant. (In,Ga)2O3 thin films grown at various conditions by CCS-PLD on (100)MgO or (00.1)Al2O3 substrates were used for investigations of phase formation, phonon mode properties, absorption edge, Schottky barrier diodes (SBDs) and photodetectors discussed in the following sections of this chapter with respect to literature results.

6.4.2 Growth and phase formation of (In,Ga)2O3 thin films Thin films of (InxGa1
x)2O3 in monoclinic modification have been obtained by MBE [59–61], MOCVD [62–66], SGM [67], PLD [7, 65, 68–76], MPM [77], MIST-CVD [33], and MS [78]. The most commonly used substrates are (00.1)Al2O3, (100)MgO, and (100)ZrO2. The growth direction is along [-201] for growth on (00.1)Al2O3 and along [100] for growth on (100)MgO or (100)ZrO2. Growth on amorphous glass results in polycrystalline layers [78]. All of these heteroepitaxial thin films contain rotational domains preventing investigations on anisotropic material properties of monoclinic (InxGa1
x)2O3.

128

Gallium Oxide

Fig. 6.5 (A) Indium content determined by EDX spectroscopy and (B) growth rate calculated from transmission data along the gradient direction of (InxGa1 2 x)2O3 thin films grown by CCS-PLD on 2 in. in diameter (00.1)Al2O3. The growth temperature was about 650°C and the oxygen growth pressure as labeled.

The incorporation of indium strongly depends on growth temperature and oxygen pressure (and availability of reactive oxygen species). A detailed analysis of thermodynamics and kinetics during MBE growth was presented by Vogt and Bierwagen [60]. In principle, the stronger GadO bond with respect to IndO favors incorporation of gallium, however, volatile Ga2O suboxide may form resulting in low growth rate especially at high growth temperature and low oxygen flux/pressure. To illustrate these effects, we compare in Fig. 6.5A the cation distribution of CCS-PLD thin films grown on (00.1)Al2O3 at a temperature of about 650°C but different oxygen pressure of 8  10
2 and 3  10
4 mbar, respectively. For an oxygen pressure of 8  10
2 mbar, the lateral variation of the indium content corresponds to results for the thin film on (100)MgO grown under the same growth conditions (cf. Fig. 6.4) and to calculations [41], implying that the stoichiometry of the incident cation flux is transferred into the thin film. For a growth pressure of 3  10
4 mbar, the indium content is for most positions z along the gradient clearly lower than for the sample grown at higher oxygen pressure. This is due to the fact that the higher III/VI ratio favors formation of MedO bonds for the cation with higher bonding strength (here GadO bond) and hence oxidation and incorporation of indium is less probable than for higher oxygen pressure. However, Ga2O suboxide formation and desorption of the suboxide results in low incorporation of gallium as well. Hence, we observe very low growth rates on the Ga-rich side of the wafer as the comparison in Fig. 6.5B reveals. The growth rate increases by a factor of five for In-rich conditions (large positive z-values in Fig. 6.5) compared with Ga-rich conditions [79]. For the sample grown at higher oxygen pressure, the growth rate is independent of the composition of the cation flux as it is the case for the film on (100)MgO (cf. Fig. 6.4). For these films, the growth rates are similar with values of 8.4 and 8.3 pm/pulse, respectively. Growth studies of β-(InxGa1
x)2O3 by MOCVD showed that indium acts as surfactant promoting step-flow growth which was explained referring to a model derived for the growth of (In,Ga)N [66, 80].

Properties of (In,Ga)2O3 alloys

129

Phase-pure monoclinic (InxGa1
x)2O3 was observed for MBE-grown thin films at 600°C on (00.1)Al2O3 up to x ¼ 0.35 [59], but independent of the growth method occurrence of the cubic phase is typically reported for x > 0.3 [59, 62, 64, 67]. On (100)MgO substrate, phase-pure films were obtained by MOCVD for x < 0.1 [65] and by PLD for x < 0.3 [7]. Cubic thin films of (GaxIn1
x)2O3 were obtained by MOCVD at 550°C on (100) YSZ (yttria-stabilized ZrO2) for x  0.5 [63]. The monoclinic structure was not observed even for Ga-rich (x > 0.7) thin films, which were X-ray amorphous. On (00.1)Al2O3 phase-pure cubic thin films were reported for x  0.5 by MOCVD at 700°C [62, 64], for x  0.2 by SGM (highest annealing temperature 900°C for 1 h) [67], and for x  0.44 by PLD at 500°C [69]. For growth on (100)MgO by PLD at 650°C alloyed films crystallize in the bixbyite phase only for x  0.2 [7]. Rhombohedral (InxGa1
x)2O3 thin films were realized by MIST-CVD at 500°C on (00.1)Al2O3 with a α-Ga2O3 buffer layer. Phase-pure films were reported for x  0.08 and x  0.67 [33]. The hexagonal polymorph InGaO3 II, discovered by Shannon and Prewitt for high-pressure synthesis, was recently observed in thin films grown by CCS-PLD [7, 70, 73, 79]. The hexagonal phase is not only observed for sample positions for which the cation ratio is 1:1 but also in an extended composition range. In Fig. 6.6A 51 θ–2θ XRD diffractograms acquired along a line parallel to the compositional gradient, highlighted by the arrow in Fig. 6.6B, are shown in false colors. The sample was grown at 650°C and an oxygen pressure of 3  10
4 mbar (cf. Fig. 6.5). For

Fig. 6.6 (A) False-color representation of 51 θ–2θ X-ray diffractograms acquired along the compositional gradient (white arrow in B) of a (InxGa1 2 x)2O3 CCS-PLD sample on (00.1) Al2O3 deposited at about 650°C [79]. (B) False-color representation of the indium distribution obtained by EDX spectroscopy of the 2 in. in diameter CCS-PLD sample, black dots indicate measuring spots, other data were interpolated. The white arrow points from low to high z-values. (C) θ–2θ X-ray diffractogram taken at the position indicated by the solid line in (A). Peaks of highest intensity belong to the sapphire substrate. Next highest intensity belongs to (004) reflection of the hexagonal phase. Reflexes belonging to the cubic and monoclinic side phase are labeled by c and m, respectively.

130

Gallium Oxide

sample positions between 7 and 10 mm, the hexagonal phase is predominant as visible in the diffractogram of Fig. 6.6C; its growth direction is along the c-axis. However, for these thin films, the hexagonal phase was so far not reported without side phase, typically alongside the cubic modification.

6.5

Deep-UV absorption and band gap engineering

Band gap engineering is one important issue in the study of ternary alloys. For the case of (InxGa1
x)2O3, we will again consider the polymorphism and present published results for the monoclinic, bixbyite, and the rhombohedral phase. Most of the studies available investigated the reduction of the absorption edge of monoclinic (InxGa1
x)2O3 with increasing indium incorporation. Monoclinic Ga2O3 is actually an indirect semiconductor, however, due to the small energetic difference between the indirect and direct transition it essentially behaves like a direct semiconductor [14]. The optical anisotropy of monoclinic (InxGa1
x)2O3 was investigated by Vasyltsiv et al. using single-crystalline bulk material for polarization-dependent absorption measurements [37]. Similar to binary Ga2O3 [81, 82], the absorption edge is highest if the electric field vector E of the incident light is parallel to the crystallographic b-axis (E jj b). If the electric field vector is perpendicular to the b-axis (E ? b), the absorption edge is about 300 meV lower. Results of transmission measurements of a heteroepitaxial thin film grown by CCSPLD [79] are depicted as function of the alloy composition in Fig. 6.7. For x  0.15, the sample has monoclinic structure without side phases. The absorption edge shifts monotonously to lower energy with increasing indium content. This is also observed in regions of the sample having mixed phases. The spectral dependence of the square of the absorption coefficient, depicted in Fig. 6.7A, was used to determine the absorption edge shown in Fig. 6.7B. Within the monoclinic phase, the bandgap decreases linearly with increasing indium content.

Fig. 6.7 (A) Square of absorption coefficient and (B) bandgap of a heteroepitaxial (InxGa1 2 x)2O3 thin film grown by CCS-PLD on (00.1)Al2O3 as a function of the indium content x. Modified from H. von Wenckstern, D. Splith, M. Purf€urst, Z. Zhang, C. Kranert, S. M€ uller, M. Lorenz, M. Grundmann, Structural and optical properties of (In,Ga)2O3 thin films and characteristics of Schottky contacts thereon, Semicond. Sci. Technol. 30 (2) (2015) 024005.

Properties of (In,Ga)2O3 alloys

131

Fig. 6.8 Dependence of band gap or absorption edge energy on alloy composition of various polymorphs of (InxGa1 2 x)2O3. Upward triangles, circles, squares, and diamond-shaped markers represent data of monoclinic (m) [37, 67, 79, 83–86], hexagonal (h) [83], cubic (c) [62, 63, 69, 83, 84], and rhombohedral (rh) [87] (InxGa1 2 x)2O3. Experimental data were obtained from transmission measurements at room temperature. Data of Vasyltsiv et al. were acquired by photoconductivity measurements at room temperature. The dashed line represents a fit to the data for monoclinic (InxGa1 2 x)2O3 of Kokubun et al. [67] and von Wenckstern et al. [79] and cubic (InxGa1 2 x)2O3 of Yang et al. [62], respectively. The growth method and the substrate used are provided in the legend.

Literature data on variation of the bandgap of (InxGa1
x)2O3 are compiled in Fig. 6.8 for phase-pure monoclinic, bixbyite and rhombohedral thin films, and monoclinic bulk crystals. Further, results of first-principles calculation for monoclinic, bixbyite, and hexagonal (InxGa1
x)2O3 are included [83, 84]. Theory and most of the experimental results confirm the linear variation of the bandgap of monoclinic (InxGa1
x)2O3 with x. A linear band gap variation was also found for the bixbyite and the hexagonal phase within the composition range considered. Calculated results for cubic (InxGa1
x)2O3 of Peelaers et al. [84] showed band gap bowing with a bowing parameter b ¼ 1.69 eV. Parameters derived from fits of the band gap variation are summarized in Table 6.2. Data on rhombohedral (InxGa1
x)2O3 are scarce. Only two/three data points for In-rich/Ga-rich films were published [33]. Hence, the fitting parameters of Table 6.2 provide rough estimates only. The different a-values obtained indicate that more data points are required and that bowing might not be neglected for the rhombohedral polymorph. Further strain, induced by the α-Ga2O3 buffer layer, likely influences lattice parameters and phase formation and stability.

6.6

Phonon modes

Phonon mode properties of cubic as well as rhombohedral In2O3 are described in the literature [23, 88–90]. For monoclinic Ga2O3 detailed investigations of vibrational properties exist [18, 19, 82, 91]. Further, first studies on heteroepitaxial Ga2O3 thin

132

Gallium Oxide

Table 6.2 Parameters of absorption edge variation for various polymorphs of the ternary alloy (InxGa12 x)2O3 Crystal structure

Eg(0) (eV)

a (eV)

Composition range

Monoclinic Hexagonal Bixbyite

4.72 4.97 4.16 3.48 4.60 5.25 3.73 4.91


2.48
2.03
0.69 0.69
1.825
3.52 [2.01]
0.49 [2.01]
2.377

0–0.13 0.47–0.55 0.88–1 0–0.12 0–1 0–1 0–1 0–0.32

4.855 3.695 5.30


1.16 1.16
4.17

0.5–1 0–0.5 0–0.08

0.92

0–0.33

Monoclinic Bixbyite Monoclinic Cubic Corundum

3.70

Comment

References

Theory, T ¼0K

[83]

Theory, T ¼0K

[84]

Fit of Fig. 6.8 Fit of Fig. 6.8 Three data points Two data points

[67, 79] [62, 64] [33]

For the bixbyite modification upper lines are with reference to (InxGa1
x)2O3 (parameters corresponding to the fit shown in Fig. 6.4), lower lines with respect to (GaxIn1
x)2O3. The parameters were derived by linear fitting of experimental or calculated data according to Eg(x) ¼ Eg(x ¼ 0) + a  x. Values in square brackets denote the bowing parameter b (Eg(x) ¼ Eg(x ¼ 0) + ax + bx2 ) in case linear fitting was not adequate. For the data of corundum structure only three and two data points were available for the Ga-rich and the In-rich side, respectively, of the phase diagram.

films with α- and ε-phase are available [27, 92]. For the ternary alloy, data for the cubic and monoclinic polymorph are published. Vibrational properties of (GaxIn1
x)2O3 ceramics with cubic modification were investigated by Regoutz et al. [40]. Due to the lower mass of Ga Raman mode frequencies show a monotonic increase up to a Ga content of about 6%. For higher gallium admixture phonon mode frequencies saturate and additional modes at 171 and 429 cm
1 appear indicating phase separation. This is in agreement with XRD results (cf. Fig. 6.3) that revealed a saturation of the lattice constant for Ga contents of about 7% and higher [40]. Fig. 6.9 depicts the composition dependence of the 307 cm
1 Raman Ag mode. The phonon mode frequency increases linearly and saturates for x  0.07 at 309.05 cm
1. Phonon modes of a PLD grown (Ga0.17In0.83)2O3 cubic thin films were reported by Wang et al. and show strong broadening compared with a binary In2O3 PLD thin film [93]. For measuring temperatures between 77 and 500 K, a linear shift of the phonon mode frequencies was observed with a temperature coefficient of 0.017 and (2) (2) 0.024 cm
1 K
1 for the A(1) g - and the Ag /Tg -mode, respectively. The temperature coefficients are higher than values of 0.006 and 0.009 cm
1 K
1 for the A(1) g - and (2) the A(2) g /Tg -mode, respectively, determined for a binary cubic In2O3 thin film [93]. The higher broadening and the stronger temperature dependence of the modes was attributed to lattice defects and structural disorder introduced by Ga incorporation.

Properties of (In,Ga)2O3 alloys

133

Fig. 6.9 Dependence of the Ag Raman phonon mode on alloy composition of (InxGa1 2 x)2O3 solid solutions. The dashed line is a linear fit to the data with parameters as labeled. Data from A. Regoutz, R.G. Egdell, D.J. Morgan, R.G. Palgrave, H. Tellez, S.J. Skinner, et al., Electronic and surface properties of Ga-doped In2O3 ceramics, Appl. Surf. Sci. 349 (2015) 970–982, https://doi.org/10.1016/j.apsusc.2015.04.106.

Phonon mode properties of monoclinic (InxGa1
x)2O3 solid solutions were reported by Vigreaux et al. and Kranert et al. up to x ¼ 0.4 [7, 94]. Since indium is heavier than gallium, phonon modes shift to lower frequencies with increasing x. For phonon mode ν1, a two-mode behavior was observed [94] and related to the large difference of the cation masses. The phonon mode frequency jumps between 0.06 < x < 0.1 from ν1(x ¼ 0.06) ¼ 114 cm
1 to ν1(x ¼ 0.1) ¼ 90 cm
1. For other modes, a two-mode behavior was not reported. The cation displacement has higher impact on low-frequency modes [95] and hence these show a stronger relative shift compared with higher frequency modes. For indium admixture above 0.4 the phonon mode frequencies saturate, the solubility limit is reached. Further, modes corresponding to cubic indium oxide appear and indicate mixed phases. The composition dependence of Raman modes reported for (InxGa1
x)2O3 solid solutions is summarized in Table 6.3. Kranert et al. investigated thin films with lateral variation of composition on (100) MgO (cf. Fig. 6.4) and (00.1)Al2O3 (cf. Fig. 6.6) grown by CCS-PLD [7]. The Raman mode frequencies monotonically decrease with indium concentration, mode splitting was not observed. Higher-order Raman modes (e.g., ν10/11) were found to be insensitive to strain and have a linear x-dependence. These modes can be used for a nondestructive determination of the alloy composition [7]. Lower-order modes showed a linear x-dependence up to about 0.2. For higher indium admixture bending occurred and a stronger x-dependence was found and attributed to lattice strain in accordance with investigations of the lattice parameters by XRD revealing a decrease in slope of the x-dependence of the out-of-plane a-lattice constant for the film on (100)MgO

134

Gallium Oxide

Phonon modes and their composition dependence in solid solutions of monoclinic (In,Ga)2O3

Table 6.3

dω/dx Mode

Symmetry

ω0 (cm21)

Vigreaux et al.

ν1a ν5 ν6 ν9 ν10/11 ν12 ν13 ν14 ν15

Ag Ag Ag Ag Ag/Bg Ag Bg Ag Ag

113 200 319 416 475 629 651 658 767


56
75
93


82

Kranert et al.
64
125
79
101
131
98
45
66

This line shows a two-mode behavior, data are valid for x  0.08. The table was compiled using data from Vigreaux et al. [94] and Kranert et al. [7]. Two subscripts of a mode indicate that an unambiguous identification of the respective phonon mode was not possible.

a

for x > 0.2. For intermediate alloy composition, a phonon mode at 260 cm
1 appears that can neither be attributed to the β-gallia nor to the bixbyite structure. Since XRD data shown in Fig. 6.6 suggest the formation of the hexagonal InGaO3 II phase this phonon mode was attributed to lattice vibrations within this polymorph [7].

6.7

Dielectric function and index of refraction

The dielectric function of cubic In2O3 is well described in literature [96–99], for rhombohedral In2O3 results of first-principles calculations are available [99]. For binary Ga2O3, the dielectric function for the monoclinic phase and its anisotropy was investigated in detail [19, 82, 100–102], for other polymorphs calculations of the dielectric function exist [101] and an experimental study for α-Ga2O3 was reported [103]. For the (InxGa1
x)2O3 alloy, little data have been published so far. Schmidt-Grund et al. investigated CCS-PLD grown (InxGa1
x)2O3 thin films on (100)MgO (cf. Fig. 6.4) and (00.1)Al2O3 (cf. Fig. 6.6) with continuous lateral composition spread by spectroscopic ellipsometry measurements [53]. Again, the anisotropic optical properties of the β-gallia structure were not accessible due to the rotational domains within these heteroepitaxial thin films [14]. Schmidt-Grund et al. [70] analyzed the dielectric function, the absorption coefficient, and the refractive index for 0.02  x  0.61 within the spectral range from 0.5 to 8.5 eV. Concerning the dielectric function, different line shapes are observed for regions with predominately monoclinic and cubic structure, respectively. Within the predominately monoclinic part of the samples, the features within the dielectric function show in general a red shift with increasing x. The signal of a feature between 6 and 7 eV increases with increasing indium content. In the mixed phase region of the samples, the line shape of the dielectric function changes from

Properties of (In,Ga)2O3 alloys

135

β-Ga2O3-like to a line shape typical for cubic In2O3. Independent of the substrate, the electronic properties generally correspond to β-(InxGa1
x)2O3 for x < 0.3 and to cubic (InxGa1
x)2O3 for x > 0.4. A peculiar observation was reported for a dielectric function acquired in the mixed phase regime at x  0.33. The spectral dependence has similarities to the dielectric function predicted for rhombohedral In2O3 [99], however, the X-ray and Raman investigations discussed above, revealed the occurrence of the hexagonal InGaO3 II structure for which a dielectric function was not reported so far. Schmidt-Grund et al. argued that the results suggest similarities in the electronic structure of rhombohedral In2O3 and hexagonal InGaO3 II. Using the Cauchy model, the energy dispersion of the refractive index was evaluated in the transparency regime. The high-frequency refractive index n∞ varies from 1.9 to 1.99, these values correspond well to the refractive indices of the binaries β-Ga2O3 and cubic In2O3 [53].

6.8

Schottky barrier diodes

Fabrication of SBDs on binary Ga2O3 is straightforward. Schottky diodes with remarkable properties were reported for both rhombohedral [104] and monoclinic Ga2O3 [17, 105–108] and stimulated their exploitation within high-power devices [14, 109–111]. In contrast, realization of Schottky barrier contacts on In2O3 remains challenging. SBDs were achieved on cubic thin films and bulk crystals [112–115]. The key for obtaining highly rectifying diodes is (i) the removal of the surface electron accumulation layer by, for example, a remote oxygen plasma treatment [112, 116] and (ii) the deposition of the Schottky contact metal in an environment containing reactive oxygen species, for example, in a reactive sputtering process [14, 113]. This recipe enables fabrication of highly rectifying SBDs to other oxide semiconductors (crystalline or amorphous) as well [117–119]. The literature reporting on properties of SBDs on ternary (In,Ga)2O3 alloys is scarce. SBDs were realized on (In,Ga)2O3 and (In,Ga)2O3:Si CCS-PLD thin films by reactive sputtering of Pt [73, 79]. Schottky contact properties were reported for regions with β-gallia structure and for parts of the thin films with mixed phase; we will restrict the following discussion to sample regions with predominately monoclinic crystal structure. The series resistance of diodes on the undoped (InxGa1
x)2O3 thin film depends strongly on the indium content and decreases with increasing In admixture [79]. It limits the forward diode current, hence hinders the determination of Schottky diode parameters by curve fitting and has strong impact on the overall current rectification. Further, the reverse diode current increases, as expected from properties of SBDs on the binaries, with increasing indium content. The maximal rectification of about 4.5 orders of magnitude was reported for x  0.06. Much better performance was reported for (InxGa1
x)2O3:Si (Si content is about 0.5 at.%) based Pt SBDs. Similar to the diodes on undoped β-(InxGa1
x)2O3, we observe a decrease of series and parallel resistance for increasing indium incorporation [73]. In Fig. 6.10, we compare room temperature current density-voltage characteristics of reactively sputtered SBDs on β-(InxGa1
x)2O3 and β-(InxGa1
x)2O3:Si. Independent of the alloy composition the forward current density in the saturation

136

Gallium Oxide

Fig. 6.10 Room temperature current-density voltage characteristics of reactively sputtered Pt Schottky barrier contacts on β-(InxGa1 2 x)2O3 (dashed lines) and β-(InxGa1 2 x)2O3:Si (solid lines, Si content in thin film is about 0.5 at.%) thin films grown by CCS-PLD on (00.1)Al2O3 substrate. The characteristics are shown for positive bias sweep direction (from
2V to +2V). The In-content is as labeled in the figure. Data from H. von Wenckstern, D. Splith, A. Werner, S. M€uller, M. Lorenz, M. Grundmann, Properties of Schottky barrier diodes on (InxGa1–x)2O3 for 0.01 x 0.85 determined by a combinatorial approach, ACS Comb. Sci. 17 (2015) 710–715, doi:10.1021/ acscombsci.5b00084; H. von Wenckstern, D. Splith, M. Purf€ urst, Z. Zhang, C. Kranert, S. M€ uller, et al., Structural and optical properties of (In,Ga)2O3 thin films and characteristics of Schottky contacts thereon, Semicond. Sci. Technol. 30 (2015) 024005, https://doi.org/10.1088/ 0268-1242/30/2/024005.

regime is higher for diodes on the Si-doped layer. Further, the reverse current density is more influenced by the indium incorporation than by the Si doping. In general, for diodes on the Ga-rich side of the wafers, the reverse currents are lowest and a splitting of the zero crossing for positive and negative bias sweep is visible. Such behavior was already observed and modeled for SBDs on binary β-Ga2O3:Si PLD thin films [120] and bulk single crystals [121]. Largest effective Schottky barrier heights of about 0.94 eV are obtained for lowest indium contents. With increasing indium content, the effective Schottky barrier height decreases to about 0.7 eV for x ¼ 0.15. In contrast to diodes on the undoped layer, the rectification decreases with increasing indium content similar to the trend of the effective barrier height. The ideality factor is with about 2 considerably higher than for similar SBDs on binary β-Ga2O3 PLD thin films and bulk crystals exhibiting ideality factors below 1.5 [107] and 1.1 [121], respectively. Reactively sputtered Pt SBDs were also realized on regions of the (InxGa1
x)2O3 CCS-PLD sample having predominately bixbyite structure (the hexagonal InGaO3 II structure was observed as side phase) [73]. The rectification saturates just below 10 due to high reverse leakage currents. The effective barrier height is about 0.5 eV.

Properties of (In,Ga)2O3 alloys

137

Fig. 6.11 Room temperature current-density voltage characteristics of reactively sputtered PtOδ/(InxGa1 2 x)2O3:Si Schottky barrier diodes for indium contents as labeled. The characteristics are shown for diodes in as-deposited state (solid line) and after being subjected to a high temperature of 400°C for 10 min in nitrogen ambient. Data from H. von Wenckstern, D. Splith, A. Werner, S. M€ uller, M. Lorenz, M. Grundmann, Properties of Schottky barrier diodes on (InxGa1–x)2O3 for 0.01 x  0.85 determined by a combinatorial approach, ACS Comb. Sci. 17 (2015) 710–715, https://doi.org/10.1021/ acscombsci.5b00084.

The properties of these SBDs are inferior compared with similar diodes on heteroepitaxial binary In2O3 thin films [113, 115]. Investigations of SBDs on phase-pure cubic (InxGa1
x)2O3 thin films are required to allocate the potential of this material for, e.g., metal-semiconductor field-effect transistors. SBDs on the Ga-rich side (predominately monoclinic) are very stable in hightemperature environment. In Fig. 6.11, current density-voltage characteristics of as-deposited SBDs are compared with characteristics acquired on the same diodes after 10 min annealing at 400°C in nitrogen. For lower Ga-content the forward current density increased after the annealing; the reverse current density showed little effect such that the rectification overall improved due to the annealing.

6.9

Photodetectors

First (InxGa1
x)2O3-based photoconductors were realized by Vasyltsiv et al. using monoclinic single crystals [37] and allowed polarization-dependent investigations of the photoconductivity for incident light configurations with E jj b or E ? b. Similar to the transmission experiments discussed above, the absorption edge is consistently higher for the configuration E jj b. A systematic redshift of the absorption edge with increasing In content was observed for both configurations. For all samples investigated a sub-band gap absorption starting at about 3.7 eV was noted and its intensity

138

Gallium Oxide

increased with increasing indium content. It was assigned to a charge transfer from oxygen 2p levels to indium 5s levels within an InO6 octahedron [37]. Interestingly, it was found that the dark conductivity as well as the photoconductivity decrease with increasing indium concentration [37]; for thin films grown by CCS-PLD, the opposite is observed [79]. First thin film photoconductive detectors based on (InxGa1
x)2O3 grown by SGM (0  x  0.2) by Kokubun et al. had a dark resistivity of about 107 Ω cm independent of the indium concentration [67]. The responsivity of the detectors is highest for the highest indium content. The long wavelength threshold expectedly shifts to lower energy with increasing x. For detectors with x ¼ 0.05 and x ¼ 0.2, the responsivity decreases for λ < 250 nm. Photoconductive detectors based on PLD-grown monoclinic (InxGa1
x)2O3 with x ¼ 0.25, 0.36, and 0.49 by Zhang et al. exhibited highest responsivity for the highest In content [75] similar to the findings of Kokubun et al. The turn-on wavelength shifts linearly to higher wavelength with increasing indium concentration [75]. The rise and fall times have multiple components; the slower process has a rise (fall) time of 7.8s (1.0s) which is attributed to emission and capture of electrons by oxygen vacancies. Z. Zhang used the (InxGa1
x)2O3:Si CCS-PLD thin film already discussed in the chapter on SBDs to fabricate (InxGa1
x)2O3:Si based metal-semiconductor-metal photodetectors with reactively sputtered Pt as electrodes [74]. Most of the detectors were realized on a sample stripe being cut in a direction parallel to the alloy composition gradient. For such detectors, a systematic shift of the onset of the photo response was observed. In Fig. 6.12, the spectral responsivity of devices fabricated along the gradient is depicted. Similar to results of Kokubun et al. and F. Zhang et al. the responsivity increases with increasing indium concentration. For x > 0.25 gain is observed, even though the sample has mixed phase for this composition. The gain mechanism was investigated by means of current-voltage measurements under

Fig. 6.12 Spectral dependence of the room temperature photo responsivity of β-(InxGa1 2 x)2O3: Si metal-semiconductor-metal photodetectors with alloy composition as labeled. The detectors were fabricated along the gradient line of a CCS-PLD thin film on c-plane sapphire. The dashed line indicates quantum efficiency of 1. The external bias was for each measurement 4 V. Data from Z. Zhang, H. von Wenckstern, J. Lenzner, M. Lorenz, M. Grundmann, Visible-blind and solar-blind ultraviolet photodiodes based on (InxGa1
x)2O3, Appl. Phys. Lett. 108 (2016) 123503, https://doi.org/10.1063/1.4944860.

Properties of (In,Ga)2O3 alloys

139

Fig. 6.13 Spectral dependence of the room temperature photo responsivity of β-(InxGa1
x)2O3: Si metal-semiconductor-metal photodetectors with alloy composition as labeled. The detectors were fabricated along an iso-composition line of a CCS-PLD thin film on c-plane sapphire. The dashed line indicates quantum efficiency of 1. The external bias was for each measurement 4 V. Data from Z. Zhang, H. von Wenckstern, J. Lenzner, M. Lorenz, M. Grundmann, Visible-blind and solar-blind ultraviolet photodiodes based on (InxGa1
x)2O3, Appl. Phys. Lett. 108 (2016) 123503, https://doi.org/10.1063/1.4944860.

illumination in dependence on the excitation wavelength [74]. For excitation above the absorption edge, the Schottky barrier height decreases indicating an accumulation of holes within the space charge region. For detectors based on binary β-Ga2O3, the gain mechanism was ascribed by Armstrong et al. to the formation of self-trapped holes that lead to the accumulation of positive charges within the space-charge region and with that to a barrier lowering [122]. Further, despite the mixed phases for x > 0.2, the absorption onset Eabs(x) of the detectors shifts linearly to lower energy following Eabs(x) ¼ 4.84eV
(1.87  x)eV. Overall, these detectors cover the spectral range from UV-A over UV-B into the UV-C. Further, detectors were prepared for a direction being perpendicular to the compositional gradient (these detectors have similar cation composition and should have similar absorption onset energy) and were selected from the composition line with x  0.11. The responsivity of detectors is depicted in Fig. 6.13. Despite small variations in the maximum value of the responsivity, the spectral shape of the photo response is essentially the same for all devices investigated and substantiates the usefulness of the CCS-PLD approach discussed throughout this chapter for basic material research and fabrication of demonstrator devices.

6.10

Summary and outlook

We have reviewed literature results on material properties of different polymorphs observed in (InxGa1
x)2O3 bulk and thin film samples. The schematic of Fig. 6.14

140

Gallium Oxide

Fig. 6.14 Schematic miscibility diagram of (InxGa1
x)2O3 for three crystallographic modifications. The bluish (greenish) bars indicate data obtained for heteroepitaxial thin films (ceramic bulk material). The range of band gap energies accessible within the given composition range is provide above each bar.

summarizes miscibility of (InxGa1
x)2O3 for the rhombohedral, the monoclinic, and the bixbyite structure as experimentally determined for thin film and ceramic bulk samples. For heteroepitaxial thin films with rhombohedral structure, a miscibility gap between about 0.08  x  0.67 was reported [33, 123], despite the fact that both binaries were stabilized in the rhombohedral α-phase [87, 124]. Several basic properties of ternary α-(InxGa1
x)2O3 need to be determined to extend the field of application beyond the field of high-power applications. For monoclinic (InxGa1
x)2O3 phase-pure ceramic samples were obtained up to x ¼ 0.43. For heteroepitaxial thin films phase separation typically occurs for x ¼ 0.35 or below, depending on growth condition such as substrate, growth temperature, or doping. The change of lattice parameters and band gap energy with alloy composition is linear; the same holds true for phonon mode frequencies. Higher frequency modes were found to shift independent of epitaxial strain and can be used to determine the alloy composition nondestructively. First demonstrator devices in the form of SBDs and metalsemiconductor-metal photodetectors were discussed. For ceramic (InxGa1
x)2O3 samples with cubic bixbyite structure phase separation occurred for x  0.92 (corresponding to a Ga admixture of 0.08 or higher). The lattice constant a decreases and phonon mode frequencies increase with increasing Ga content [35, 40]. Cubic heteroepitaxial thin film samples were obtained up to about x ¼ 0.5. The optical band gap and the absorption edge increases linearly with increasing Ga admixture. The hexagonal InGaO3 II phase was observed in CCS-PLD thin films, however, systematic experimental investigation of its properties is still to come. Independent of the polymorph of (InxGa1
x)2O3 the determination of crucial material properties such as electric and thermal transport properties, mechanical properties and thermal expansion coefficients, doping limits, and defect formation needs further investigations for the complete range of alloy composition to allow fabrication of

Properties of (In,Ga)2O3 alloys

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ternary thin film samples with tailored functionality. Further, the physical mechanisms behind phase formation have to be understood to provide a tool box for the growth of (InxGa1
x)2O3 with material properties in the polymorph suited best for a desired application. This specifically applies to ε-(InxGa1
x)2O3 which offers due to its large polarization an additional feature for device design. For that and equally important for the other polymorphs, the growth and investigation of high-quality heterostructures is demanded.

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x)2O3 alloys for transparent electronics, Phys. Rev. B. 92 (2015) 085206. F. Zhang, K. Saito, T. Tanaka, M. Nishio, M. Arita, Q. Guo, Wide bandgap engineering of (AlGa)2O3 films. Appl. Phys. Lett. 105 (2014) 162107, https://doi.org/10.1063/ 1.4900522. S. Fujita, T. Oshima, MBE Growth and Characterization of Ga2O3 and (InGa)2O3 Films for Ultraviolet Applications, Meet. Abstr. MA2007-01. (2007) 769. N. Suzuki, K. Kaneko, S. Fujita, Growth of corundum-structured In2O3 thin films on sapphire substrates with Fe2O3 buffer layers. J. Cryst. Growth 364 (2013) 30–33, https://doi.org/10.1016/j.jcrysgro.2012.11.065. B. Garcia-Domene, H.M. Ortiz, O. Gomis, J.A. Sans, F.J. Manjo´n, A. Munoz, et al., High-pressure lattice dynamical study of bulk and nanocrystalline In2O3. J. Appl. Phys. 112 (2012) 123511, https://doi.org/10.1063/1.4769747. C. Kranert, R. Schmidt-Grund, Raman active phonon modes of cubic In2O3. Phys. Status Solidi Rapid Res. Lett. 8 (2014) 554–559, https://doi.org/10.1002/ pssr.201409004. J.A. Sans, R. Vilaplana, D. Errandonea, V.P. Cuenca-Gotor, B. Garcia-Domene, C. Popescu, et al., Structural and vibrational properties of corundum-type In2O3 nanocrystals under compression. Nanotechnology 28 (2017) 205701 https://doi.org/10.1088/ 1361-6528/aa6a3f. C. Kranert, C. Sturm, R. Schmidt Grund, M. Grundmann, Raman tensor formalism for optically anisotropic crystals. Phys. Rev. Lett. 116 (2016) 127401, https://doi.org/10.1103/ PhysRevLett.116.127401. R. Cusco´, N. Dome`nech-Amador, T. Hatakeyama, T. Yamaguchi, T. Honda, L. Artu´s, Lattice dynamics of a mist-chemical vapor deposition-grown corundum-like Ga2O3 single crystal. J. Appl. Phys. 117 (2015) 185706, https://doi.org/10.1063/1.4921060. V. Wang, W. Xiao, D.M. Ma, R.-J. Liu, C.M. Yang, Structural, electronic, and optical properties of GaInO3: a hybrid density functional study. J. Appl. Phys. 115 (2014) 043708, https://doi.org/10.1063/1.4863210. C. Vigreux, L. Binet, D. Gourier, B. Piriou, Formation by laser impact of conducting β-Ga2O3–In2O3 solid solutions with composition gradients. J. Solid State Chem. 157 (2001) 94–101, https://doi.org/10.1006/jssc.2000.9043. D. Dohy, G. Lucazeau, A. Revcolevschi, Raman spectra and valence force field of singlecrystalline β Ga2O3. J. Solid State Chem. 45 (1982) 180–192, https://doi.org/10.1016/ 0022-4596(82)90274-2. R. Schmidt-Grund, H. Krauß, C. Kranert, M. Bonholzer, M. Grundmann, Temperature dependence of the dielectric function in the spectral range (0.5–8.5) eV of an In2O3 thin film. Appl. Phys. Lett. 105 (2014) 111906, https://doi.org/10.1063/1.4896321.

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[97] J.B. Varley, A. Schleife, Bethe–Salpeter calculation of optical-absorption spectra of In2O3 and Ga2O3. Semicond. Sci. Technol. 30 (2015) 024010, https://doi.org/ 10.1088/0268-1242/30/2/024010. [98] M. Feneberg, J. Nixdorf, C. Lidig, R. Goldhahn, Z. Galazka, O. Bierwagen, et al., Manyelectron effects on the dielectric function of cubic In2O3: effective electron mass, band nonparabolicity, band gap renormalization, and Burstein-Moss shift. Phys. Rev. B. 93 (2016) 045203, https://doi.org/10.1103/PhysRevB.93.045203. [99] F. Fuchs, F. Bechstedt, Indium-oxide polymorphs from first principles: quasiparticle electronic states. Phys. Rev. B. 77 (2008) 155107, https://doi.org/10.1103/PhysRevB.77. 155107. [100] C. Sturm, J. Furthm€uller, F. Bechstedt, R. Schmidt-Grund, M. Grundmann, Dielectric tensor of monoclinic Ga2O3 single crystals in the spectral range 0.5–8.5 eV. APL Mater. 3 (2015) 106106, https://doi.org/10.1063/1.4934705. [101] J. Furthm€uller, F. Bechstedt, Quasiparticle bands and spectra of Ga2O3 polymorphs. Phys. Rev. B. 93 (2016) 115204–115216, https://doi.org/10.1103/PhysRevB.93.115204. [102] T. Onuma, S. Saito, K. Sasaki, T. Masui, T. Yamaguchi, T. Honda, et al., Spectroscopic ellipsometry studies on β-Ga2O3 films and single crystal. Jpn. J. Appl. Phys. 55 (2016) 1202B2, https://doi.org/10.7567/JJAP.55.1202B2. [103] A. Segura, L. Artu´s, R. Cusco´, R. Goldhahn, M. Feneberg, Band gap of corundum like α
Ga2O3 determined by absorption and ellipsometry. Phys. Rev. Mater. 1 (2017) 024604, https://doi.org/10.1103/PhysRevMaterials.1.024604. [104] M. Oda, R. Tokuda, H. Kambara, T. Tanikawa, T. Sasaki, T. Hitora, Schottky barrier diodes of corundum-structured gallium oxide showing on-resistance of 0.1 mΩ cm2 grown by MIST EPITAXY®. Appl. Phys. Express. 9 (2016) 021101, https://doi.org/ 10.7567/APEX.9.021101. [105] M. Higashiwaki, K. Konishi, K. Sasaki, K. Goto, K. Nomura, Q.T. Thieu, et al., Temperature-dependent capacitance–voltage and current–voltage characteristics of Pt/Ga2O3 (001) Schottky barrier diodes fabricated on n––Ga2O3 drift layers grown by halide vapor phase epitaxy. Appl. Phys. Lett. 108 (2016) 133503, https://doi.org/10.1063/1.4945267. [106] K. Konishi, K. Goto, H. Murakami, Y. Kumagai, A. Kuramata, S. Yamakoshi, et al., 1-kV vertical Ga2O3 field-plated Schottky barrier diodes. Appl. Phys. Lett. 110 (2017) 103506, https://doi.org/10.1063/1.4977857. [107] S. M€uller, H.V. Wenckstern, F. Schmidt, D. Splith, H. Frenzel, M. Grundmann, Method of choice for the fabrication of high-quality β-gallium oxide-based Schottky diodes. Semicond. Sci. Technol. 32 (2017) 065013, https://doi.org/10.1088/1361-6641/aa6a8d. [108] K. Sasaki, D. Wakimoto, Q.T. Thieu, Y. Koishikawa, A. Kuramata, M. Higashiwaki, et al., First demonstration of Ga2O3 trench MOS-type Schottky barrier diodes. IEEE Electron Device Lett. 38 (2017) 783–785, https://doi.org/10.1109/LED.2017.2696986. [109] M. Higashiwaki, K. Sasaki, H. Murakami, Y. Kumagai, A. Koukitu, A. Kuramata, et al., Recent progress in Ga2O3 power devices. Semicond. Sci. Technol. 31 (2016) 034001, https://doi.org/10.1088/0268-1242/31/3/034001. [110] S.J. Pearton, J. Yang, P.H. CaryIV, F. Ren, J. Kim, M.J. Tadjer, et al., A review of Ga2O3 materials, processing, and devices. Appl. Phys. Rev. 5 (2018) 011301, https://doi.org/ 10.1063/1.5006941. [111] S. Fujita, M. Oda, K. Kaneko, T. Hitora, Evolution of corundum-structured III-oxide semiconductors: growth, properties, and devices. Jpn. J. Appl. Phys. 55 (2016) 1202A3, https:// doi.org/10.7567/JJAP.55.1202A3. [112] O. Bierwagen, J.S. Speck, T. Nagata, T. Chikyow, Y. Yamashita, H. Yoshikawa, et al., Depletion of the In2O3(001) and (111) surface electron accumulation by an oxygen plasma surface treatment. Appl. Phys. Lett. 98 (2011) 172101, https://doi.org/10.1063/ 1.3583446.

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[113] H. von Wenckstern, D. Splith, F. Schmidt, M. Grundmann, O. Bierwagen, J.S. Speck, Schottky contacts to In2O3. APL Mater. 2 (2014) 046104, https://doi.org/10.1063/ 1.4870536. [114] M. Nazarzadehmoafi, S. Machulik, F. Neske, V. Scherer, Schottky contact by Ag on In2O3 (111) single crystals. Appl. Phys. Lett. 105 (2014) 162104, https://doi.org/ 10.1063/1.4899143. [115] F. Schmidt, D. Splith, S. M€uller, H. von Wenckstern, M. Grundmann, Electronic defects in In2O3 and In2O3:Mg thin films on r-plane sapphire. Phys. Stat. Sol. (B) 252 (2015) 2304–2308, https://doi.org/10.1002/pssb.201552328. [116] O. Bierwagen, Indium oxide—a transparent, wide-band gap semiconductor for (opto) electronic applications. Semicond. Sci. Technol. 30 (2015) 024001, https://doi.org/ 10.1088/0268-1242/30/2/024001. [117] M.W. Allen, S.M. Durbin, J.B. Metson, Silver oxide Schottky contacts on n-type ZnO. Appl. Phys. Lett. 91 (2007) 053512, https://doi.org/10.1063/1.2768028. [118] A. Lajn, H.V. Wenckstern, Z. Zhang, C. Czekalla, G. Biehne, J. Lenzner, et al., Properties of reactively sputtered Ag, Au, Pd, and Pt Schottky contacts on n-type ZnO. J. Vac. Sci. Technol. B 27 (2009) 1769, https://doi.org/10.1116/1.3086718. [119] S. Bitter, P. Schlupp, H. von Wenckstern, M. Grundmann, Vital role of oxygen for the formation of highly rectifying Schottky barrier diodes on amorphous zinc–tin–oxide with various cation compositions. ACS Appl. Mater. Interfaces 9 (2017) 26574–26581, https://doi.org/10.1021/acsami.7b06836. [120] D. Splith, S. M€uller, F. Schmidt, H. von Wenckstern, J.J. van Rensburg, W.E. Meyer, et al., Determination of the mean and the homogeneous barrier height of Cu Schottky contacts on heteroepitaxial β-Ga2O3 thin films grown by pulsed laser deposition. Phys. Stat. Sol. (A). 211 (2014) 40–47, https://doi.org/10.1002/pssa.201330088. [121] S. M€uller, H. von Wenckstern, F. Schmidt, D. Splith, F.L. Schein, H. Frenzel, et al., Comparison of Schottky contacts on β-gallium oxide thin films and bulk crystals. Appl. Phys. Express. 8 (2015) 121102–121104, https://doi.org/10.7567/APEX.8.121102. [122] A.M. Armstrong, M.H. Crawford, A. Jayawardena, A. Ahyi, S. Dhar, Role of self-trapped holes in the photoconductive gain of β-gallium oxide Schottky diodes. J. Appl. Phys. 119 (2016) 103102, https://doi.org/10.1063/1.4943261. [123] S. Fujita, K. Kaneko, Epitaxial growth of corundum-structured wide band gap III-oxide semiconductor thin films. J. Cryst. Growth 401 (2014) 588–592, https://doi.org/10.1016/ j.jcrysgro.2014.02.032. [124] D. Shinohara, S. Fujita, Heteroepitaxy of corundum-structured α-Ga2O3 thin films on α-Al2O3 substrates by ultrasonic mist chemical vapor deposition. Jpn. J. Appl. Phys. 47 (2008) 7311–7313, https://doi.org/10.1143/JJAP.47.7311.

Further reading [125] K. Momma, F. Izumi, IUCr, VESTA 3 for three-dimensional visualization of crystal, volumetric and morphology data. J. Appl. Crystallogr. 44 (2011) 1272–1276, https://doi.org/ 10.1107/S0021889811038970.

Low-field and high-field transport in β-Ga2O3

7

Krishnendu Ghosh, Uttam Singisetti Electrical Engineering Department, University at Buffalo, Buffalo, NY, United States

Chapter Outline 7.1 Introduction 149 7.2 Electron-phonon interaction in β-Ga2O3

150

7.2.1 Electron-LO phonon coupling 151 7.2.2 Electron-LO phonon-plasmon coupling 153 7.2.3 Short-range (nonpolar) electron-phonon coupling 156

7.3 Electron mobility in β-Ga2O3

158

7.3.1 Bulk electron mobility 158 7.3.2 2DEG mobility 162

7.4 Velocity-field curves in β-Ga2O3 7.5 Summary 166 Acknowledgments 166 References 166

7.1

163

Introduction

β-Ga2O3 has recently emerged as a novel wide-bandgap semiconductor with immense potential for applications in power electronics and optoelectronics. Experimental advancements in the past 5 years have been significant toward realizing commercial β-Ga2O3 devices in the near future [1–7]. Matured crystal growth and processing techniques make the material further promising [8–10]. In terms of power electronic applications, MOSFETs based on this material have been demonstrated that could withstand record high voltages [11, 12]. The accuracy of n-type doping and the difficulty of p-type doping make electrons the primary charge carriers in β-Ga2O3. Although β-Ga2O3 has lower electron mobility compared to other wide-bandgap semiconductors, it is found to have a superior Baliga’s figure of merit that jointly accounts for on-state resistance and breakdown voltage [4]. So it is important to investigate in rigor the fundamentals behind β-Ga2O3 material properties that could be beneficial to gain an understanding on the causes that control mobility and breakdown voltage. There are theoretical reports on fundamental materials aspects including electronic structure [13] and optical properties [14], lattice dynamical and dielectric properties [15], and thermal properties [16, 17] as well. The primary physics behind both Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00007-5 © 2019 Elsevier Inc. All rights reserved.

150

Gallium Oxide

mobility (and hence the device on resistance) and breakdown voltage lies in the electron transport phenomenon. There have been a few experimental reports that try to characterize the electron transport and scattering mechanisms in β-Ga2O3 with Hall measurements being reported a few times to predict temperature dependence and also crystal orientation dependence of the electron mobility [18, 19]. On the other hand, we are making a systemic study on the theoretical understanding of electron transport in β-Ga2O3 starting from the first principles [20–22]. The main idea is to follow a bottom-up approach in order to develop an understanding of the near-equilibrium and far-from-equilibrium electron dynamics in β-Ga2O3. This is unique compared to conventional semiconductors in a way that β-Ga2O3 has a low-symmetry crystal structure and a fairly large primitive unit cell that gives rise to many phonon modes. On several occasions, the traditional notions of electron transport that are applicable to Si and GaAs actually do not quite hold well in the case of β-Ga2O3. In this chapter, we attempt to provide a comprehensive picture of electron transport in β-Ga2O3 under low and moderately high electric fields based on our work in the recent years. In the following, first we discuss the nature of the electron-phonon coupling in β-Ga2O3 taking into account several important aspects including the plasmon-phonon hybridization and dynamical screening. Long-range [polar optical phonon (POP)] and short-range (nonpolar optical) couplings are separately described as they play different roles based on the regime (strength of the applied field) of transport. Next, we turn to the electron mobility aspects that are revealed by incorporating the different scattering mechanisms in a semiclassical Boltzmann transport formalism. Interesting features such as the anisotropy of electron mobility is analyzed with a thorough understanding of the underlying physics that is unique in case of β-Ga2O3. Also, an electron mobility in a two-dimensional electron gas (2DEG) is predicted in this context. In the last section, we focus on high-field transport properties in β-Ga2O3 that includes the velocity-field curves.

7.2

Electron-phonon interaction in β-Ga2O3

Electron-phonon interaction (EPI) plays a pivotal role in controlling the intrinsic mobility of a material. The vibrational motion of the lattice adds a perturbation to the periodic crystal potential that couples with the electronic motion and results in scattering of the electrons. This becomes important in the diffusive regime of transport, which is generally the case for most power semiconductor devices where the transit time of electrons from source to drain is much higher compared to typical time intervals between scattering events. On the other hand, in materials with partially ionic bonds, for example, polar semiconductors including β-Ga2O3, an additional EPI arises from the vibration of the net dipole moment of the ionic bond. This interaction is Columbic in nature, and hence long ranged. In most polar semiconductors, under low electric field, this long-range EPI controls the intrinsic electron transport and β-Ga2O3 is no exception. This is the same dipole moment that produces splitting of the transverse optical (TO) and longitudinal optical (LO) vibrational modes in the long-wavelength limit. The LO modes gain a slightly higher energy compared to its TO counterparts. The long-range EPI can be viewed as a coupling between the electron and the LO mode. This coupling is also referred to as polar coupling, and the resulting scattering is referred to as POP scattering.

Low-field and high-field transport in β-Ga2O3

151

7.2.1 Electron-LO phonon coupling β-Ga2O3 has a base-centered monoclinic lattice (C2m space group) with a 10 atom primitive unit cell. The crystal structure is shown in Fig. 7.1A with the green balls denoting Ga atoms and the red balls denoting the O atoms. The Cartesian directions are also shown as a reference and the same direction convention is followed throughout this chapter. Fig. 7.1B shows the first Brillouin zone (BZ) along with the three reciprocal vectors. The black dots show the high-symmetry points. The lattice parameters as known from experiments and structural relaxation calcula˚ , b ¼ 3.05 A ˚ , and c ¼ 5.82 A ˚ . Before going into the details tions [13] are a ¼ 12.27 A of the EPI calculation, a few important points are to be discussed regarding the electronic structure and the lattice dynamical properties in β-Ga2O3. Plane-wave pseudopotential-based density functional theory (DFT) calculations reveal an isotropic electronic band-structure near the BZ center (Γ point) with an approximate electronic effective mass of 0.3 times the free electron mass. For low-field transport, the primary concern regarding the electronic band is its behavior near the Γ point. β-Ga2O3 has a spherically symmetric parabolic band near the Γ point, and hence the effective mass approximation-based transport theories are expected to be well applicable if the scattering rates are appropriately formulated including the microscopic details.

K3

Z

N1

a b

x

K1

y

(A)

N

z c

K2

(B)

Fig. 7.1 (A) The crystal structure of β-Ga2O3 along with the Cartesian direction convention used in this chapter. (B) The Brillouin zone and the reciprocal lattice vectors. (A) Courtesy K. Ghosh, U. Singisetti, Ab initio velocity-field curves in monoclinic β-Ga2O3. J. Appl. Phys. 122 (3) (2017) p. 035702, with permissions from AIP Publishing LLC. (B) Courtesy K. Ghosh, U. Singisetti, Ab initio calculation of electron–phonon coupling in monoclinic β-Ga2O3 crystal. Appl. Phys. Lett. 109 (7) (2016) p. 072102, with permissions from AIP Publishing LLC.

152

Gallium Oxide

Lattice dynamical computations based on density functional perturbation theory (DFPT) reveal several aspects important for EPI. The essential quantities are the Born effective charge tensor, the high-frequency dielectric tensor, the phonon energies, and their displacement patterns. In β-Ga2O3, there are a total of 30 phonon modes arising from the 10 atoms. Out of those, 12 are infrared (IR) active as they possess a net dipole oscillation. These IR active modes are responsible for the long-range EPI. Based on their displacement patterns, the IR active modes could be classified as Au modes and Bu modes. The four Au modes vibrate along the Cartesian y direction, while the eight Bu modes vibrate on the x–z plane. The most important point to note here is that since the projection of the net dipole oscillation strength of the Bu modes are different along the x direction and the z direction, the IR strengths and the LO-TO splitting are highly anisotropic. This is a signature of the low symmetry of the crystal and this anisotropy reflects in the POP scattering as well. The anisotropy of the IR activity is experimentally observed [15] and its trend is consistent with theoretical reports. Table 7.1 shows the DFPT calculated direction dependent LO-TO splitting. As could be seen the Bu 1 mode has a strong splitting along the z direction, it has a negligible splitting along the x direction. The four Au modes also have different splitting. Overall, one cannot postulate a generic strength for LO-TO splitting (and hence long-range EPI) in β-Ga2O3 since different modes show varied range of splitting and the strengths are also dependent on the direction of the phonon wave vector. Hence, each mode needs to be treated independently. Next, we focus on the features of long-range EPI in β-Ga2O3. The phonon modewise trend of the EPI strength is expected to follow the trend of the LO-TO splitting since both originate from the net dipole oscillation. However, an additional aspect in the long-range EPI strength is the eigenenergy of the phonon mode. Unlike LO-TO splitting, the long-range EPI strength depends on the amplitude of the vibration also. As described by the Vogl model [23], the electron-LO phonon coupling strength is inversely proportional to the square root of the phonon mode energy. The overall POP scattering rate is a combined result of electron-LO phonon coupling strength, electron energy, and temperature. The relative contribution of the 12 different phonon

Table 7.1

Direction dependent LO-TO splitting in β-Ga2O3

x direction

y direction

z direction

TO (meV) LO (meV)

TO (meV) LO (meV)

TO (meV) LO (meV)

21.30 29.30 34.20 43.00 52.00 70.20 83.70 90.10

13.02

13.15

36.30

41.63

55.90

67.61

82.00

93.14

21.30 29.30 34.20 43.00 52.00 70.20 83.70 90.10

21.68 32.68 34.98 44.64 59.18 77.13 88.92 92.76

28.49 29.43 34.23 43.13 57.84 79.99 84.87 96.15

Low-field and high-field transport in β-Ga2O3

153

´ 1013 T = 300 K

B1u

Ek = 60 meV

POP scattering rate (/s)

5

Ek = 100 meV 4 Ek = 170 meV 3 B6u 2 B2u

1 A1u 14

29

B8u B4u

B3u 21

B5u

A2u

37

43

A4u

A3u 53

71

83

B7u 91

wqn (meV)

Fig. 7.2 The rates by the different POP modes for three different electron energies. Courtesy K. Ghosh, U. Singisetti, Ab initio calculation of electron–phonon coupling in monoclinic β-Ga2O3 crystal. Appl. Phys. Lett. 109 (7) (2016) p. 072102, with permissions from AIP Publishing LLC.

modes in causing POP scattering is shown in the bar plots of Fig. 7.2. The three bar plots are shown at room temperature for three different electronic energies relevant under low electric field. At low electron energies, the most dominant POP mode is Bu 1 . The scattering of a higher energy electron mediated by the Bu 1 mode is lower due to the long-range nature of the coupling. On the other hand, for relatively higher energy electrons, the dominant mode is Bu 6 . Low-energy electrons are unable to emit Bu 6 , and hence their corresponding scattering rate is lower. This picture reveals the interesting trend of the scattering strength with the coupling strength and the electron energy in a many phonon system. For the sake of qualitative arguments and simpler transport calculations, under low electric field one could broadly say that the Bu 1 mode controls the low-field electron mobility in β-Ga2O3.

7.2.2 Electron-LO phonon-plasmon coupling A plasmon is a longitudinal mode of collective oscillation of electrons. Plasmons can couple with the LO phonons and give rise to hybridized LO-plasmon coupled (LOPC) modes. Certain interesting features arise when an electron interacts with such coupled modes and this becomes even more interesting for a material such as β-Ga2O3, which has several LO modes and each LO mode can couple in a different way with the plasmon mode. Before talking about the properties of electron-LOPC coupling in β-Ga2O3, let us clarify when electron-LOPC coupling is important to be considered

154

Gallium Oxide

and when a pure electron  LO coupling is good enough for transport calculations. The energy of the plasmon mode increases with increasing free electron density. The coupling between an LO phonon mode and the plasmon mode is significant if their energies are similar. In β-Ga2O3, the different LO phonon energies vary from about 13 to 96 meV. On the other hand, for doping levels ranging from 1017 to 1019/cm3, the plasmon energy in β-Ga2O3 sweeps from about 10 to 100 meV considering an isotropic effective mass of 0.3 and an average isotropic high-frequency dielectric constant of 4.3 (calculated from DFPT). So, as seen in the previous section, typically the lower energy LO modes will contribute to controlling the low-field mobility, it would be judicious to say for doping levels from 1017/cm3 to moderately high levels such as 5  1018/cm3, one needs to be concerned about LOPC modes and outside this doping window a pure electron-LO mode interaction will be good for low-field transport calculations. In the following, we describe different features of LOPC modes. The LOPC energies are found from the zeros of the net frequency-dependent dielectric constant for a given phonon wave vector. The net frequency-dependent dielectric constant is formed under a generalized Lyddane-Sachs-Teller formalism (for lattice contribution) and plasmon-pole model (for plasmon contribution). The number of LOPC modes depends on the direction of the wave vector. For wave vectors along the y direction, there are 5 LOPC modes (Au symmetry), whereas for x–z plane there are 9 LOPC modes (Bu symmetry). Otherwise, for any arbitrary wave vector there are 13 LOPC modes (mixed symmetry). The LOPC mode energies are plotted for different electron concentrations (ns) and are shown in Fig. 7.3A–C. The pure plasmon energy is shown in blacked dashed curve. At low ns, higher energy modes [say, the blue line in Fig. 7.3A or the cyan line in Fig. 7.3B] have an LO phonon energy. However, at higher ns, as the plasmon energy rises, modes that are below the pure plasmon mode [say, the green lines in both Fig. 7.3A and B] have TO phonon energy. The reason behind this is the fact that since the pure plasmon has higher energy compared to those modes, it could efficiently screen the net dipole moment that is responsible for producing the splitting. This phenomenon cannot be ignored while estimating the electronic scattering rates. Fig. 7.3A considers the Bu symmetry LOPC modes when the wave vector is along the z direction. So modes that have a higher dipole projection along the x-axis than that along the z couples less with the plasmon and hence their energy remains more or less flat [like the red line in Fig. 7.3A]. This is of course a signature of the low-symmetry crystal. Next, we turn to discuss the coupling strength between such LOPC modes and an individual electron. Under the same spirit as Frolich electron-LO phonon interaction model, the electron-LOPC interaction can be formulated mode wise by considering the difference of dielectric elements formed by keeping a particular mode “frozen” and then keeping the same in “full-response.” For a detail of such calculation refer to Ref. [24] and to our recent work [21]. Note that this formulation automatically takes into account frequency-dependent dynamical screening under long wavelength limit from both lattice modes and the plasmon. Here we report the main concepts of the electron-LOPC interaction strength in β-Ga2O3. Fig. 7.4A shows the screened oscillator strengths with the same color scheme as in Fig. 7.3A. Although the following discussion is done for the Bu symmetry LOPC modes, the same idea is valid for

Low-field and high-field transport in β-Ga2O3 0.14

0.14

qˆ = zˆ

Bu symmetry

0.10 0.08 0.06

LOPC5

LOPC4

0.04 0.02 0.00

(A)

0.4 0.6 ns (/cm3)

0.8 1.0 ´1019

0.14

0.10 0.08 0.06 0.04 0.00

0.2

(B)

0.4 0.6 ns (/cm3)

0.8 1.0 ´1019

Mixed symmetry qˆ = yˆ + zˆ

0.12 wnLOPC (eV)

qˆ = yˆ

0.02

LOPC3

0.2

Au symmetry

0.12 wnLOPC (eV)

wnLOPC (eV)

0.12

155

0.10 0.08 0.06 0.04 0.02 0.00

0.2

(C)

0.4 0.6 ns (/cm3)

0.8 1.0 ´1019

(A)

108 107 106 105 104 103 102 101 100

5 Bu symmetry

LOPC5

LOPC4

´1014 T = 300 K

qˆ = zˆ Scattering rate (/s)

KnLOPC

Fig. 7.3 The LO-plasmon coupled mode energies with (A) Bu symmetry, (B) Au symmetry, and (C) mixed symmetry. The pure plasmon mode is shown in black dashed line. Courtesy K. Ghosh, U. Singisetti, Electron mobility in monoclinic β-Ga2O3—Effect of plasmon-phonon coupling, anisotropy, and confinement. J. Mater. Res. (2017) 1–11, with permissions from Cambridge University Press.

LOPC3

4 3 2 1

0.2

0.4

0.6

ns (/cm3)

0.8

1.0

´1019

ns = 3 ´ 1018 /cm2

0.00

(B)

ns = 3 ´ 1017 /cm2 0.05 0.10 0.15 Electron energy (eV)

0.20

Fig. 7.4 (A) The dynamically screened coupling strength of the LOPC modes with Bu symmetry. (B) The scattering rates mediated by the LOPC modes for two different electron concentrations. Courtesy K. Ghosh, U. Singisetti, Electron mobility in monoclinic β-Ga2O3—Effect of plasmon-phonon coupling, anisotropy, and confinement. J. Mater. Res. (2017) 1–11, with permissions from Cambridge University Press.

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the Au modes. The essential idea to understand here is the effect of dynamic screening on the interaction strengths. Let us consider the modes LOPC3, LOPC4, and LOPC5 as marked in Fig. 7.4A. With increasing electron density each coupled mode goes through three stages—first, it has a phonon-like character with LO energy, then it becomes plasmon like, and the finally it regains its phonon character with the energy of a TO mode. Around a density of 2  1018/cm3, LOPC3 becomes plasmon like and provides strong antiscreening to LOPC4 since LOPC4 has a slightly higher energy. With increasing ns, the plasmon character shifts to LOPC4 and it starts antiscreening LOPC5 while screening LOPC3. After this, when the plasmon character shifts to LOPC5, it starts screening LOPC4 since the former has a higher energy than LOPC4. Finally, the plasmon character shifts to further high-energy modes and LOPC5 also gets screened. Hence, in a many LO mode material such as β-Ga2O3 the plasmon character propagates from one coupled mode to the next and the screening of the modes change accordingly. The electronic scattering rates are computed using the Fermi-Golden rule. The LOPC mode-mediated scattering rates in β-Ga2O3 are shown in Fig. 7.4B for two different ns at room temperature. As can be seen at lower electronic energies, the scattering rate is higher under a low ns (say at 3  1017/cm3), which follows from the fact that the low-energy LOPC modes are antiscreened at such ns. As ns increases, the scattering rate for the low-energy electron gets suppressed since the low-energy LOPC modes get well screened and the low-energy electrons are not able to emit any high-energy LOPC mode due to energy conservation. As ns (say at 3  1018/cm3) increases the scattering rate of the high-energy electrons shoots up since the higher energy LOPC modes get antiscreened. This type of behavior of the scattering rates has a significant role in determining the trend of electron mobility with varying ns.

7.2.3 Short-range (nonpolar) electron-phonon coupling As mentioned before electrons can couple with the perturbations of the periodic crystal potential that arise from the lattice vibrations. This coupling is nonpolar (nonColumbic) in nature and present in both ionic and covalent materials. Also, this coupling is typically short ranged in real-space which means the corresponding scattering mechanisms can provide large momentum relaxation to electrons. This short-range EPI is primarily responsible for effects such as velocity saturation that occurs under high-electric field. Here we show the scattering rates calculated from first-principles computation of short-range EPI in β-Ga2O3. Due to many phonon modes and low crystal symmetry, the calculation is somewhat complex and intensive. The details of the calculation are not given here for brevity and interested readers can refer to Ref. [22]. The nonpolar electron-phonon scattering in β-Ga2O3 could be broadly classified due to three types of phonons—zone-center (ZC) acoustic, zone-edge (ZE) acoustic, and optical. We provide here a qualitative understanding of three different types of scatterings along with corresponding deformation potentials and effective phonon mode energies. Fig. 7.5A shows the scattering rate due to all the 27 optical phonons. The dashed line shows the analytical fitting to the actually computed rate from first

Low-field and high-field transport in β-Ga2O3

157

1015 Scattering rate (/s)

300 K 10

14

1013 1012 1011

Sn = 4-30Sn(E) DP optical 0

(A)

1

2

3

Electron energy (eV)

Scattering rate (/s)

1015 300 K 1014 1013 Sn = 1-3Sn(E)

1012 1011 0

(B)

DP (ZC acoustic) DP (ZE acoustic) 3 1 2 Electron energy (eV)

Fig. 7.5 The nonpolar scattering rate by the (A) optical modes, and (B) acoustic modes. Courtesy K. Ghosh, U. Singisetti, Ab initio velocity-field curves in monoclinic β-Ga2O3. J. Appl. Phys. 122 (3) (2017) p. 035702, with permissions from AIP Publishing LLC.

principles, which is shown by the solid line. Due to the nice parabolic behavior of the electronic bands at low electron energies, the scattering rate follows a typical square root law at low energies and then starts deviating due to band nonparabolicties. The corresponding deformation potential and effective phonon mode energy are shown in Table 7.2. However, it is very important to keep in mind that the effective phonon energies do not represent the existence of any actual physical mode at that energies. They are just meant to carry out compact scattering rate calculations using analytical equations. On the other hand, Fig. 7.5B shows the scattering rates corresponding to ZC and ZE acoustic modes. The need for treating the two types of acoustic modes Table 7.2 The fitted deformation potential and phonon energies for β-Ga2O3

Zone-center acoustic Zone-edge acoustic Optical

Deformation potential (DA or Do)

Phonon energy (ω0)

3 eV 5  107 eV/cm 6  108 eV/cm

– 10 meV 50 meV

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separately arises from the different dispersion type of the modes at ZC and at ZEs. The conventional analytical acoustic deformation potential models are good only for ZC acoustic modes. The ZE acoustic modes have dispersion more like optical modes and hence the deformation potential is provided in units of eV like the other optical modes. The essential idea to note from here is that at low electron energies the polar scattering rate as shown in Fig. 7.4B is almost an order of magnitude higher than that of the nonpolar scattering rates. Hence, in β-Ga2O3, the intrinsic electron mobility is limited by the POPs.

7.3

Electron mobility in β-Ga2O3

Next, we turn to talk about the electron mobility and its behaviors with varying temperature, carrier concentration, and crystal orientation. The electron mobility is calculated by solving Boltzmann transport equation using Rode’s iterative technique. We discuss the mobility in two different forms—first the bulk mobility and next the mobility of 2DEG formed at heterojunctions such as AlxGa2xO3/Ga2O3.

7.3.1 Bulk electron mobility For discussing the temperature dependence, we consider a doping of about 1017/cm3 and LOPC is disregarded for this particular study of temperature dependence. Scattering due to ionized impurities is included based on a simple Brook’s-Herring model [25]. The impurity scattering is significant only at low temperatures where POP scattering is weak. The resulting trend of the electron mobility with respect to temperature is shown in Fig. 7.6. The room temperature mobility is about 115 cm2/Vs. The trend of the mobility with respect to temperature is similar to that observed in other conventional semiconductors such as GaAs. The initial rise in mobility is due to the fact the ionized impurity scattering decreases with increasing electron energy and with increasing temperature the tail of the Fermi-Dirac distribution spreads out in energy.

400 Mobility (cm2/Vs)

Fig. 7.6 The temperature dependence of electron mobility under low doping. Experimental data [18] are shown for comparison. Courtesy K. Ghosh, U. Singisetti, Ab initio calculation of electron–phonon coupling in monoclinic β-Ga2O3 crystal. Appl. Phys. Lett. 109 (7) (2016) p. 072102, with permissions from AIP Publishing LLC.

ND = 1.1 ´ 1017/cm3

300

200

100

0 101

Parisini et al. (measured data) This work 102 Temperature (K)

103

Low-field and high-field transport in β-Ga2O3

159

Hence, the effective scattering rate decreases resulting to a gradual increase in mobility. However, with further increase of temperature POP scattering tries to dominate and the rate of which increases with increasing temperature due to an increasing Bose occupation number. Hence, the electron mobility starts dropping off. For comparison the experimentally measured Hall mobility is also shown [18]. The agreement is fairly good. Deviations at higher temperature could be due to inaccuracies in the POP scattering rates since the DFPT-computed POP modes are at zero temperature. On the other hand, the discrepancy near 80 K is possibly due to uncertainty in the doping of the sample. To compare with other wide-bandgap semiconductors, the room temperature electron mobility is significantly lower in β-Ga2O3. However, there is a lot of room toward improving the mobility by device engineering such as forming heterojunctions. The key idea is to screen the POP modes efficiently to suppress the corresponding scattering. This requires a higher carrier concentration and that is where considering LOPC becomes important to accurately probe the theoretical limits of the mobility. In the next paragraph, we shade some light on how increasing the carrier concentration can help modulating the electron mobility. To probe the dependence of the mobility owing to screening by the carriers at room temperature, the mobility is shown in Fig. 7.7 under two conditions—without ionized impurity scattering and including the same. The triangles show the mobility without considering impurity scattering. Initially with increasing carrier concentration the mobility drops. This is due to the antiscreening of the low-energy LOPC modes that results to enhanced scattering rate. However, beyond 108/cm3 carrier concentration the mobility starts improving drastically due to strong screening of the LOPC modes. Including the dynamic nature of screening enables to capture the nonmonotonic nature

mn (cm2/Vs)

500 400

T = 300 K y direction

300

z direction

200 100

1018

1019

1020

ns (/cm ) 3

Fig. 7.7 The carrier concentration dependence of the room temperature mobility for two different crystal directions. The triangles show the case when the mobility is computed without taking into account ionized impurity scattering while the circles show the case when ionized impurity scattering is present. Courtesy K. Ghosh, U. Singisetti, Electron mobility in monoclinic β-Ga 2 O 3—Effect of plasmon-phonon coupling, anisotropy, and confinement. J. Mater. Res. (2017) 1–11, with permissions from Cambridge University Press.

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of the dependence of mobility. Under high enough carrier concentration, where the effects of antiscreening are absent, a static Thomas-Fermi screening model is expected to yield similar results. It could be noted that the mobility along the x direction is somewhat higher than that along the z direction. The reasons behind such anisotropy is discussed later on. Now we look at the mobility corresponding to the calculation including ionized impurity scattering. The initial trend is similar to the case without considering the impurity scattering. However, beyond 1018/cm3 the rise in mobility is not as sharp as the previous case. The reason is rather obvious—the ionized impurities provide a higher scattering rate while the enhanced carrier concentration tries to improve the mobility through screening of LOPC modes. Overall the mobility increases slowly. Around 5  1019/cm3 the mobility reaches around 200 cm2/Vs beyond which it is expected to drop due to very strong impurity screening. It is important to mention here that the ionized impurity concentration in this analysis (Fig. 7.7) is taken to be the same as the electron concentration. Next, we discuss the reasons behind the anisotropy in low-field electron mobility. Such small anisotropy (10%–15%) in the mobility is observed experimentally as well [19]. However, it is known the electronic bands near the Γ point are highly isotropic, so the question that arises is what causes the anisotropy in the mobility. Here, we try to investigate that with a Monte Carlo simulation of the electron transport carried out under a small electric field applied along two different directions. The simulation does take care of an ab initio band structure calculated from DFT. The previously discussed POP scattering is included along with nonpolar scatterings (although negligible under low electric field). LOPC is not considered in this particular case because within first order the plasmon energy is expected to be isotropic as the effective mass is isotropic. Hence, we trace the anisotropy in mobility from the anisotropy of electron-LO coupling strength. Before discussing the observations from the Monte Carlo simulation, first an important qualitative picture is clarified which is shown in Fig. 7.8A. Consider an electron with wave vector k that emits a POP with wave vector q and ends up on a 0 wave vector k . Under low electric field the electron energy is comparable to the LO energy and hence the scattering is significantly inelastic; hence, the event of such scattering in energy space looks like the diagram shown in Fig. 7.8A. Note that the allowed phonon wave vectors are inclined toward the initial wave vector of the electron meaning that low-energy electrons are more likely to interact with phonons that have aligned wave vectors. Now, if we look at Table 7.1, we see that the strongest LO-TO splitting among the low-energy phonons occur for a wave vector along the z direction for the mode Bu 1 . The splitting is much lower in the x direction for Bu 1 and also any other low-energy mode does not show such high splitting in any directions. So the electron mobility is also expected to be low along the z direction, which is what is seen in Fig. 7.7 and also observed experimentally [19]. Now, we look at the results of the Monte Carlo simulation to have a further clear understanding. Fig. 7.8B and C shows the POP emission rates with respect to time that begins when the electric field is turned on. While the electric field is applied along the z direction in both cases, the difference is the magnitude of the electric field. In Fig. 7.8B, we see that Bu 1 has a much higher scattering rate compared to Bu 6 and Au 2 . This is because

Fz q

k¢

k

(A)

nph (per electron per ps)

Low-field and high-field transport in β-Ga2O3

20

161

Bu1

300 K

15

Fz = 5.0E6 V/m

10

Au2

Bu6 5 0 0

0.2

(B)

0.4

0.6

Time (ps)

nph (per electron per ps)

30 300 K Bu1 20 Bu6

Fz = 2.0E7 V/m

10 Au2 0

(C)

0

0.2

0.4

0.6

Time (ps)

Fig. 7.8 (A) The schematic representation to show the momentum conservation relation between electron and phonon wave vectors during scattering of a low-energy electron. (B–C) The emission rates of three different LO modes for electric fields of 5  106 V/m and 2  107 V/m, respectively. Courtesy K. Ghosh, U. Singisetti, Electron mobility in monoclinic β-Ga 2 O 3—Effect of plasmon-phonon coupling, anisotropy, and confinement. J. Mater. Res. (2017) 1–11, with permissions from Cambridge University Press.

Bu 6 has a higher energy than Bu 1 and hence under the given field electrons do not have enough energy to emit Bu 6 phonons. On the other hand, Au 2 is polarized along the y direction, and hence does not couple well with electrons propagating along the z direction for reasons described in the previous paragraph. Fig. 7.8C shows a good increase in the emission rate of Bu 6 since a higher electric field boosts up the electron energy enabling the electrons to emit Bu 6 . Whereas the emission rate of Au 2 only slightly increases although it has a lower energy than Bu 6 . The reason is again due to the poor coupling with electrons propagating along the z direction. So the results of the Monte Carlo simulation validate the qualitative picture described in the previous paragraph and also provide a solid argument to support experimentally the observed anisotropy in electron mobility. Overall, this anisotropic mobility in β-Ga2O3 is unique in a sense that, unlike conventional semiconductors, it does not arise from the anisotropy of the band structure rather it occurs due to the anisotropy of the electron-LO phonon coupling, which in turn is a result of low crystal symmetry.

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Gallium Oxide

7.3.2 2DEG mobility Next, we turn to discuss the mobility trend in 2DEG formed at heterojunctions such as AlxGa2xO3/Ga2O3. Such heterojunctions help separate the impurities from conducting electrons thereby eliminating ionized impurity scattering while the high density of the electrons help screen the LOPC modes. Both effects together provide a significant enhancement in mobility. The inset in Fig. 7.9 shows the typical structure of the heterojunction along with the position of the dopant impurities referred to as the remote impurity (RI) center. Usually the RI center is a few nanometers away from the conducting channel. Further, the lower center would be the interaction of the conducting electrons with the Coulombic potential due to the dopant impurities. However, for a few nanometers the interaction is not negligible due to the long-range nature of the Coulombic interactions. Now, a qualitatively talk is made about the way the scattering rate due to such RI centers is treated. The dopant impurities are essentially taken as a sheet of charge density (like a δ doping). Although such assumption could be raised by employing a Poisson solver, we do not go into that here in order to avoid unnecessary complexity. The decaying Columbic potential interacts with the confined 2DEG electron wave function. For calculating the matrix elements for such Coulomb potential, the in-plane 2DEG wave function is taken to be a plane wave, while the outof-plane envelope function has a Fang-Howard form [26]. The screening of the matrix elements involves static dielectric elements with proper wave-vector dependence. For 2D systems, the corresponding Thomas-Fermi wave vector is independent of electron concentration and is simply given by 2/aB where aB is the effective Born radius [26]. Now, we discuss the LOPC-mediated scattering in the 2DEG. The idea is exactly same as that in the bulk case except the expression related to the plasmon dispersion is modified for the 2D case. Also since the energies of the strong LOPC modes are low enough, intersubband scatterings of the 2DEG are ignored. In the present case, the direction of confinement is taken to be along the y direction. So only Bu symmetry

mn (cm2/Vs)

104 d

Dopants

AlGO y

No RI

2DEG GO z

3

10

d = 5 nm d=0

102

0

2

T = 300 K 4

ns (/cm2)

d = 10 nm

6 ´ 1012

Fig. 7.9 Mobility of β-Ga2O3 2DEG for different distances of the RI center from the 2DEG channel. (Inset) shows a schematic of the heterojunction that helps form the 2DEG. Courtesy K. Ghosh, U. Singisetti, Electron mobility in monoclinic β-Ga2O3—Effect of plasmon-phonon coupling, anisotropy, and confinement. J. Mater. Res. (2017) 1–11, with permissions from Cambridge University Press.

Low-field and high-field transport in β-Ga2O3

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modes can cause scattering and Au or mixed symmetry modes remain inactive in scattering due to momentum conservation. Fig. 7.9 shows the computed mobility along the z direction for the 2DEG. With increasing electron density, the mobility improves due to enhanced screening. The antiscreening behavior is not observed since that occurs at a lower 2DEG density. Also Fig. 7.9 shows the trend of the 2DEG mobility as the RI center moves away from the interface. When the RI center is right at the interface, the 2DEG mobility for a density of 5  1012/cm2 is about 418 cm2/Vs. The red dashed line gives the upper limit of the mobility if the RI center is turned off from the calculation. Usually the center is within 5 nm from the 2DEG channel, so the mobility is expected to be around 1000 cm2/Vs for a 2DEG density of 5  1012/cm2 at room temperature. For the sake of simplicity, here the dopant density at the RI center is taken to be the same as the 2DEG density.

7.4

Velocity-field curves in β-Ga2O3

Velocity-field curve is an important property of any semiconducting material. This curve provides the performance limit of electronic devices made from that material. Here, we show the intrinsic velocity-field curves of bulk β-Ga2O3 at room temperature calculated using Monte Carlo simulation. The electron concentration is assumed to be low enough and hence no LOPC effect is included. Instead of going into any mathematical details, we focus on the qualitative picture of the simulation and also discuss the results. Unlike low-field calculations, high-field calculations cannot be performed using Rode’s iterative technique because the electron distribution goes far from equilibrium. So Monte Carlo simulation is a traditional approach to perform high-field transport calculations. While for conventional semiconductors one can use several approximations such as the deformation potential-based scattering rates, no dispersion of the phonon modes, Frolich model for POP scatterings, these approximations are not expected to hold well for β-Ga2O3 due to low symmetry and multiple phonons. Hence, the complete wave vector (both electron and phonon) dependence of all EPI needs to be taken into account. This is what makes the calculation complex and challenging but interesting at the same time. The essential idea of the simulation is given in the following. For more details please see Ref. [22]. First, the DFT computed electronic band structure, DFPT computed phonon eigenvalues, electron-POP matrix elements and scattering rates, and short-range (nonpolar) matrix elements and scattering rates are computed and stored. A renormalized scattering table is formed to estimate the relative strength of the different scattering mechanisms. The notion of a scattering mechanism includes the final electronic band after scattering, the index of the phonon mode that caused the scattering, whether the scattering is polar or nonpolar in nature, and whether the scattering involves absorption or emission of a phonon. For example, a simulation that takes into account two electronic bands and all 30 phonon modes there are total 180 possible scattering mechanisms. It is assumed here that POP scattering cannot cause interband transitions since the matrix elements are long range in nature and for a given electronic wave vector the Bloch states are orthogonal. So the renormalized scattering table contains the relative

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scattering strengths (on a scale of 0–1) of the 180 mechanisms. Next the electron distribution is initialized thermodynamically using a Boltzmann distribution and the electric field is turned on. The trajectory of the electrons in the reciprocal space is tracked with respect to time. The free flight time of an electron is stochastically determined based on the maximum scattering rate from the band at which the electron resides. Once a scattering event occurs, a mechanism is selected stochastically based on the previously formed renormalized scattering table. This provides information on final band index and final energy of the electron. Following this, the task is to get a suitable final wave vector of the electron that satisfies both momentum and energy conservation. This is the most intense part, and we skip this here for brevity. The details on this can be found in Refs. [22, 27]. Once a final wave vector is selected the electron is allowed to drift again until the next scattering event occurs. This interplay of drifting and scattering occurs until the electron distribution reaches a steady state. Once a steady state is reached, the average drift velocity of the electron ensemble is extracted. The simulation is repeated for multiple electric fields along various Cartesian directions to obtain the velocity-field curves. Fig. 7.10A shows the transient response of the velocity for different electric fields. Initially, the velocity increases sharply with time as the electrons get accelerated by the external field. At this regime, the transport is primarily controlled by POP scattering due to low electron energy. Since POP matrix elements are long range in nature, they provide small momentum relaxation rate compared to the actual scattering rate and hence the net drift velocity shoots up rapidly. But as the electron energy increases nonpolar scattering becomes significant. Nonpolar matrix elements being short range in nature can provide large change in momentum. Hence, the velocity drops due to higher momentum randomization. This phenomenon, also observed in other semiconductors, is known as the velocity overshoot. This finds application in ultrascaled transistor devices. If the transistor gate length is so small that the electrons can reach from source to drain over a timescale comparable to the one at which velocity overshoot ´105

´105

Drift velocity (m/s)

3.5 2.5 2.0 1.5 1.0 0.5 0.0 0

(A)

1.5

2

4 3 Time (s)

5 6 ´ 10−13

b-Ga2O3

1 0.5

Fz = 50 kV/cm

1

300 K

2

Increasing electric field with 50 kV/cm increment

3.0

Drift velocity (m/s)

4.0

0

(B)

0

1 2 3 Electric field (V/m)

x y z 4 ´ 107

Fig. 7.10 (A) The transient velocity of the electrons for different electric fields. (B) The velocity-field curves of electrons in β-Ga2O3 for three different Cartesian directions. Courtesy K. Ghosh, U. Singisetti, Ab initio velocity-field curves in monoclinic β-Ga2O3. J. Appl. Phys. 122 (3) (2017) p. 035702, with permissions from AIP Publishing LLC.

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165

occurs then the electrons can potentially avoid the higher momentum randomization and provide a higher net drift velocity making the transistor faster. However, to utilize this in β-Ga2O3 transistors, the gate length needs to be below 20 nm even for a field of 450 kV/cm [the green curve in Fig. 7.10A]. This is challenging in terms of fabrication but feasible using state-of-art techniques. On the other hand, Fig. 7.10B shows the velocity-field profile for three different Cartesian directions. The velocity increases for all three directions up to an electric field of 200 kV/cm followed by which negative differential conductivity (NDC) starts appearing. The NDC effect in β-Ga2O3 is less than that observed in GaAs, and the reasons behind NDC in β-Ga2O3 are fundamentally different as revealed by the Monte Carlo simulations. Usually NDC happens due to the transition of electrons to remote satellite valleys, which have a higher effective mass than that in the lowest conduction band valley. However, in β-Ga2O3 such satellite valleys occur at a much higher energy (about 2.4 eV) compared to that in GaAs (about 0.3 eV). So it is unlikely that a significant number of electrons can reach such high-energy valleys at moderately high field. This is also verified by histogram plots shown and discussed in Ref. [22]. The NDC effect in β-Ga2O3 is a result of short-range intravalley scattering within the Γ valley at higher energies. Also, the nonparabolic effect of the bands far from the Γ point reduces the increase of group velocity of the electrons with increasing wave vector. Overall, the velocity drops after a critical field of about 250 kV/cm. Next we come to the anisotropy of the velocity-field curves. At lower electric fields, the anisotropy is a result of the anisotropy in electron-LO phonon interaction as shown by the anisotropic phonon emission picture in Fig. 7.8B and C. This makes the velocity along the z direction lower compared to that along the other directions. On the other hand, at higher fields, the drift velocity is higher along the y direction. This is due to a higher group velocity in that direction compared to the other two directions. In order to facilitate device simulation, we provide here fitting parameters for the velocity-field curves based on the Barnes model [28] of NDC in Table 7.3. Here, μ0 is the low-field mobility, vsat, Fc, and γ are adjustable fitting parameters. This model is often used by the commercial device simulators to model device performance including parallel field-dependent mobility. However, it is very important to note that, the fitting parameters are good only up to an electric field of about 500 kV/cm beyond which other high-field phenomenon such as impact ionization needs to be considered for an accurate description. Table 7.3 Computed fitting parameters that can be used on the Barnes model [28] for device simulation including velocity saturation and NDC

μ0 (cm2/Vs) vsat (cm/s) Fc (V/cm) γ

x

y

z

140 107 2.25  105 2.84

140 1.5  107 1.54  105 2.47

112 107 2.63  105 3.35

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7.5

Gallium Oxide

Summary

In summary, this chapter reviewed some essential concepts related to low-field and high-field transport properties in β-Ga2O3. Electron-phonon coupling is described with significant details and very few assumptions are made in computing the matrix elements. The temperature dependence of the mobility is discussed and compared with experimental reports. To explore the potential improvement of mobility by free carrier screening, the coupling between LO phonons and plasmon is studied including dynamical screening effects under long wavelength. Mobility values are predicted for 2DEG formed at heterojunctions such as AlxGa2xO3/Ga2O3. A 2DEG mobility of up to 1000 cm2/Vs under an electron density of 5  1012/cm2 is estimated for RI centers located optimally about 5 nm from the 2DEG channel. Anisotropy of electron mobility is traced microscopically from the anisotropy of electron-POP coupling and also supported by Monte Carlo simulation results. For high-field transport, velocity-field curves in β-Ga2O3 at room temperature computed from Monte Carlo simulation are analyzed with some insights on effects such as NDC. A critical electric field of 250 kV/cm is estimated beyond which NDC starts appearing. Fitting parameters for scattering rates and field-dependent mobility are provided to guide simpler transport calculations and device simulations.

Acknowledgments The electronic structure and lattice dynamical calculations are carried out using Quantum ESPRESSO [29]. The Kohn-Sham eigenvalues are interpolated using Wannier90 [30]. The EPI matrix element calculations are done using a modified version of the EPW code [31, 32]. This work is supported by National Science Foundation (NSF) grant (ECCS 1607833) monitored by Dr. Dimitris Pavilidis. The authors also acknowledge the high-performance computing clusters provided by the Center for Computational Research at the University at Buffalo.

References [1] M. Higashiwaki, et al., Ga2O3 Schottky Barrier Diodes with n--Ga2O3 Drift Layers Grown by HVPE, IEEE Device Research Conference, 2015. [2] M. Higashiwaki, et al., Depletion-mode Ga2O3 metal-oxide-semiconductor field-effect transistors on β-Ga2O3 (010) substrates and temperature dependence of their device characteristics, Appl. Phys. Lett. 103 (12) (2013) 123511. [3] M. Higashiwaki, et al., Gallium oxide (Ga2O3) metal-semiconductor field-effect transistors on single-crystal β-Ga2O3 (010) substrates, Appl. Phys. Lett. 100 (1) (2012) 013504. [4] M. Higashiwaki, et al., Recent progress in Ga2O3power devices, Semicond. Sci. Technol. 31 (3) (2016) 034001. [5] K. Sasaki, et al., Ga2O3 Schottky barrier diodes fabricated by using single-crystal β-Ga2O3 (010) substrates, IEEE Electron Device Lett. 34 (4) (2013) 493–495. [6] T. Oishi, et al., High-mobility β-Ga2O3(201) single crystals grown by edge-defined filmfed growth method and their Schottky barrier diodes with Ni contact, Appl. Phys. Express 8 (3) (2015) 031101.

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[7] K. Zeng, et al., Ga2O3 MOSFETs using spin-on-glass source/drain doping technology, IEEE Electron Device Lett. 38 (4) (2017) 513–516. [8] Z. Galazka, et al., Czochralski growth and characterization of β-Ga2O3 single crystals, Cryst. Res. Technol. 45 (12) (2010) 1229–1236. [9] K. Irmscher, et al., Electrical properties of β-Ga2O3 single crystals grown by the Czochralski method, J. Appl. Phys. 110 (6) (2011) 063720. [10] Y. Tomm, et al., Czochralski grown Ga2O3 crystals, J. Cryst. Growth 220 (4) (2000) 510–514. [11] K.D. Chabak, et al., Enhancement-mode Ga2O3 wrap-gate fin field-effect transistors on native (100) β-Ga2O3 substrate with high breakdown voltage, Appl. Phys. Lett. 109 (21) (2016) 213501. [12] M.H. Wong, et al., Field-plated Ga 2 O 3 MOSFETs with a breakdown voltage of over 750 V, IEEE Electron Device Lett. 37 (2) (2016) 212–215. [13] H. Peelaers, C.G. Van de Walle, Brillouin zone and band structure of β-Ga2O3, Phys. Status Solidi B 252 (4) (2015) 828–832. [14] K.A. Mengle, et al., First-principles calculations of the near-edge optical properties of β-Ga2O3, Appl. Phys. Lett. 109 (21) (2016) 212104. [15] M. Schubert, et al., Anisotropy, phonon modes, and free charge carrier parameters in monoclinicβ-gallium oxide single crystals, Phys. Rev. B 93 (12) (2016) 125209. [16] Z. Guo, et al., Anisotropic thermal conductivity in single crystal β-gallium oxide, Appl. Phys. Lett. 106 (11) (2015) 111909. [17] B. Liu, et al., Lattice dynamical, dielectric, and thermodynamic properties of β-Ga[sub 2] O[sub 3] from first principles, Appl. Phys. Lett. 91 (17) (2007) 172102. [18] A. Parisini, R. Fornari, Analysis of the scattering mechanisms controlling electron mobility inβ-Ga2O3crystals, Semicond. Sci. Technol. 31 (3) (2016) 035023. [19] M.H. Wong, et al., Electron channel mobility in silicon-doped Ga2O3 MOSFETs with a resistive buffer layer, Jpn. J. Appl. Phys. 55 (12) (2016) 1202B9. [20] K. Ghosh, U. Singisetti, Ab initio calculation of electron–phonon coupling in monoclinic β-Ga2O3 crystal, Appl. Phys. Lett. 109 (7) (2016) 072102. [21] K. Ghosh, U. Singisetti, Electron mobility in monoclinic β-Ga2O3—Effect of plasmonphonon coupling, anisotropy, and confinement, J. Mater. Res. 4142-4152 (2017) 1–11. [22] K. Ghosh, U. Singisetti, Ab initio velocity-field curves in monoclinic β-Ga2O3, J. Appl. Phys. 122 (3) (2017) 035702. [23] C. Verdi, F. Giustino, Frohlich Electron-phonon vertex from first principles, Phys. Rev. Lett. 115 (17) (2015) 176401. [24] M.V. Fischetti, et al., Effective electron mobility in Si inversion layers in metal–oxide– semiconductor systems with a high-κ insulator: The role of remote phonon scattering, J. Appl. Phys. 90 (9) (2001) 4587–4608. [25] D. Chattopadhyay, H. Queisser, Electron scattering by ionized impurities in semiconductors, Rev. Mod. Phys. 53 (4) (1981) 745. [26] T. Ando, et al., Electronic properties of two-dimensional systems, Rev. Mod. Phys. 54 (2) (1982) 437–672. [27] T. Kunikiyo, et al., A Monte Carlo simulation of anisotropic electron transport in silicon including full band structure and anisotropic impact-ionization model, J. Appl. Phys. 75 (1) (1994) 297. [28] J.J. Barnes, et al., Finite-element simulation of GaAs MESFET’s with lateral doping profiles and submicron gates, IEEE Trans. Electron Devices 23 (9) (1976) 1042–1048. [29] P. Giannozzi, et al., Quantum Espresso: A modular and open-source software project for quantum simulations of materials, J. Phys. Condens. Matter. 21 (39) (2009) 395502.

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[30] A.A. Mostofi, et al., wannier90: A tool for obtaining maximally-localised Wannier functions, Comput. Phys. Commun. 178 (9) (2008) 685–699. [31] J. Noffsinger, et al., EPW: A program for calculating the electron–phonon coupling using maximally localized Wannier functions, Comput. Phys. Commun. 181 (12) (2010) 2140–2148. [32] F. Giustino, et al., Electron-phonon interaction using Wannier functions, Phys. Rev. B 76 (16) (2007) 015003.

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

8

N.C. Giles*, L.E. Halliburton† *Department of Engineering Physics, Air Force Institute of Technology, Wright-Patterson Air Force Base, Dayton, OH, United States., †Department of Physics and Astronomy, West Virginia University, Morgantown, WV, United States

Chapter Outline 8.1 8.2 8.3 8.4

Introduction 169 Crystal structure of β-Ga2O3 171 Shallow donors and conduction electrons Acceptors and self-trapped holes 173

172

8.4.1 Doubly ionized gallium vacancies 174 8.4.2 Neutral Mg acceptors 178 8.4.3 Self-trapped holes 181

8.5 Transition-metal and rare-earth ions 8.5.1 8.5.2 8.5.3 8.5.4 8.5.5

Cr3+ ions Fe3+ ions Mn2+ ions Ti3+ ions Er3+ ions

184

184 185 186 186 186

8.6 Oxygen vacancies 187 Acknowledgments 188 References 188

8.1

Introduction

Electron paramagnetic resonance (EPR) is often used to identify and characterize defects and impurities having unpaired spins in semiconductors and insulators [1–4]. Examples include donors, acceptors, vacancies, interstitials, antisites, small polarons, transition-metal ions, and rare-earth ions. Simply stated, an EPR spectrum represents the absorption of microwaves at discrete values of magnetic field. This allows the energy-level scheme of the unpaired spin system to be determined. The importance of EPR arises from its high sensitivity and high resolution. Depending on the line widths, concentrations of defects as low as a few parts per billion can be easily observed. Each paramagnetic defect has a unique EPR spectrum that reflects its g matrix and hyperfine matrices, as well as the zero-field splittings when S is greater than 1/2. Once the identity of the responsible defect has been established, a spectrum can be used to Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00008-7 © 2019 Elsevier Inc. All rights reserved.

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compare concentrations of the defect in different samples and also to monitor changes in the concentration of the defect during photoexcitations and thermal anneals. A word of caution is in order when using EPR to observe defects in thin films. The strength of an EPR signal depends on the absolute number of unpaired spins in a sample, and not simply the concentration of spins. Thus, the observed intensity of a particular spectrum depends on the volume of the sample (and hence the number of spins), assuming the spins are uniformly distributed. With their smaller volumes, detecting EPR signals from thin films is often challenging unless the concentration of the responsible defect is large. For this reason, nearly all of the EPR spectra reported for β-Ga2O3 have been obtained from bulk crystals. Most EPR spectrometers in use today are commercial instruments produced by Bruker, a German-based company having offices worldwide. The source of microwaves is often a solid-state Gunn diode operating near 9.5 GHz. A microwave cavity, either rectangular or cylindrical, with built-in 100-kHz modulation coils is used. The magnet is usually water-cooled and can produce fields up to 1.4 T. In an EPR experiment, a sample containing an observable concentration of a paramagnetic defect is placed in the static magnetic field and a Zeeman splitting of the energy levels associated with the spin system occurs. The microwave magnetic field in the resonant cavity drives transitions between these levels. To obtain an EPR spectrum, the microwave frequency remains constant and the magnetic field is slowly swept through the region of interest. Because of the 100-kHz magnetic field modulation superimposed on the slowly varying magnetic field, the EPR signals appear as first derivatives of the absorption lines. To increase sensitivity and resolution, EPR spectrometers are available that operate at microwave frequencies of 90 GHz or higher with superconducting magnets providing the needed fields. An important experimental condition in most EPR experiments is the sample temperature. The signal-to-noise ratio of a spectrum often increases as the temperature is lowered because of the increasing population difference between spin-up and spindown levels. There are, however, many spin systems with spin-lattice relaxation times that vary rapidly with temperature. In some cases, an EPR line may be broadened beyond recognition at higher temperatures because of very short relaxation times. In other cases, the intensity of an EPR line is reduced at low temperature because long relaxation times lead to microwave power saturation of the spin transitions. The sample temperature can be varied from below 10 K to near room temperature in most spectrometers by using a helium-gas flow system or a closed-cycle helium system. This allows the optimum temperature to be determined for observing a particular EPR spectrum. For a simple isotropic S ¼ 1/2 spin system with only an electron Zeeman interaction, the selection rule ΔMS ¼  1 gives a single EPR line with hν ¼ gβB as the resonance condition. Here, ν is the microwave frequency, β is the Bohr magneton, and B is the magnetic field. The g value of the defect can then be calculated if experimental values are known for ν and B. In many cases of interest, however, there are multiple lines (possibly unresolved) in an EPR spectrum. The spin Hamiltonian describing a defect in a single crystal often has an anisotropic g matrix and anisotropic hyperfine matrices involving adjacent magnetic nuclei (i.e., nuclei with I  1/2) and perhaps a

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

171

central magnetic nucleus. These anisotropies produce angular dependence in the EPR spectrum of the defect (i.e., a shifting and possible splitting of lines as the direction of the magnetic field is rotated in various crystal planes). The principal values and the directions of the principal axes of these anisotropic matrices are extracted, via a fitting process, from the angular dependence of the spectrum. If S is greater than 1/2, zero-field splittings often cause very large angular-dependent shifts in the EPR line positions [5, 6].

8.2

Crystal structure of β-Ga2O3

The structure of the β-Ga2O3 crystals is monoclinic and the space group is C2/m (C32h). ˚ , b ¼ 3.0371 A ˚ , c ¼ 5.7981 A ˚ , and Lattice constants at 273 K are a ¼ 12.214 A β ¼ 103.83 degrees [7, 8]. The usual convention is followed where the crystallographic b axis is perpendicular to the crystal’s mirror plane. There are two inequivalent gallium sites and three inequivalent oxygen sites in this complex binary crystal. The Ga(I) ions have four oxygen neighbors, and the Ga(II) ions have six oxygen neighbors. The O(I) and O(II) ions have three gallium neighbors, and the O(III) ions have four gallium neighbors. A ball-and-stick model of the β-Ga2O3 crystal is shown in Fig. 8.1. In acquiring and presenting the angular dependence of EPR spectra in β-Ga2O3, a* and c* directions may be introduced where a* is perpendicular to b and c and c* is perpendicular to a and b. Fig. 8.1 Crystal structure of β-Ga2O3. Gallium ions are green, and oxygen ions are red. The two inequivalent gallium sites are Ga(I) and Ga(II), and the three inequivalent oxygen sites are O(I), O(II), and O(III). Reproduced from B.E. Kananen, L.E. Halliburton, E.M. Scherrer, K.T. Stevens, G.K. Foundos, K.B. Chang, N.C. Giles, Electron paramagnetic resonance study of neutral Mg acceptors in β-Ga2O3 crystals, Appl. Phys. Lett. 111 (2017) 072102 with permission from the American Institute of Physics.

a b c

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Gallium Oxide

Often, EPR investigators find it helpful to have an actual physical model of a crystal available when visualizing the principal-axis directions of a g matrix, the relative positions of neighboring ions that are responsible for hyperfine structure, and the nature of possible lattice relaxations. Small desk-top ball-and-stick crystal models containing up to several hundred ions can be obtained from Miramodus Ltd. [9].

8.3

Shallow donors and conduction electrons

A widely reported EPR signal in β-Ga2O3 crystals has been attributed to shallow donors and conduction-band electrons [10–22]. The spectrum consists of a single line with a g value near 1.96. Precise angular-dependent measurements show a slight anisotropy with gx ¼ 1.9606, gy ¼ 1.9577, and gz ¼ 1.9630, where z corresponds to the crystal’s b axis and x and y are in the b plane with x making an angle of 23 degrees with the a axis [22]. The shifts of the g values from the free spin value of 2.0023 are a result of the mixing of p-type valence band states with the s-type conduction band. Initially, the presence of this g ¼ 1.96 EPR signal in β-Ga2O3 was attributed to oxygen vacancies. It is now accepted that the oxygen vacancy is a deep donor in this material [23], and that n-type conductivity and the corresponding g ¼ 1.96 EPR signal are caused by impurities such as Si, Ge, and Sn serving as shallow donors [19–21]. Observations of hyperfine structure [from EPR or electron-nuclear double resonance (ENDOR) experiments] have not been reported for this signal, thus the chemical identity of the responsible donor is not immediately obtained from the spectrum. Electrical conductivity, infrared absorption, and EPR experiments provide indirect evidence that can be used to establish correlations with specific donor impurities. Son et al. [22] have recently completed a detailed study of the shallow donor in β-Ga2O3 crystals. A series of unintentionally doped samples were investigated before and after annealing treatments, and spectra were obtained over a wide temperature range. Activation of the donors was shown to occur when a crystal was heated in flowing N2 gas at temperatures between 1100°C and 1150°C. Using SIMS data, these investigators identified Si as the primary donor. Fig. 8.2 shows the EPR spectrum of the shallow donor taken at various temperatures with the magnetic field along the c axis [22]. Superhyperfine interactions with neighboring 69,71Ga nuclei broaden the signal at low temperature, and motional narrowing sharpens the signal at higher temperature. Based on the temperature dependence of the EPR signal, Son et al. [22] suggest that the shallow donor exhibits negative-U behavior in partially activated β-Ga2O3 crystals. Their analysis shows that the negative charge state DX is 16–20 meV below the neutral charge state of the donor and the neutral charge state is 28–29 meV below the conduction band minimum. In fully activated crystals with higher donor densities, the donor electrons form impurity bands and the donor activation energy reduces to 15–17 meV [22]. Because of the silicon’s small ionic radius and its proclivity for tetrahedral bonding, the most likely site in β-Ga2O3 for the active Si donor is the tetrahedral Ga(I) position. A question yet to be resolved, however, is the nature of the donor activation mechanism at the temperatures near 1100°C. Specifically, what is the location

EPR intensity (linear scale)

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

3427

173

UID b-Ga2O3 anneal 1100°C 20 min 140 K 96 K 54 K 170 K 200 K 240 K 260 K

3429

B|| 9.415 GHz 3431

3433

3435

Magnetic field (G)

Fig. 8.2 EPR spectra of the residual shallow donor measured at different temperatures in an unintentionally doped β-Ga2O3 substrate that was annealed at 1100°C in flowing N2 gas for 20 min to partially activate the donor. Reproduced from N.T. Son, K. Goto, K. Nomura, Q.T. Thieu, R. Togashi, H. Murakami, Y. Kumagai, A. Kuramata, M. Higashiwaki, A. Koukitu, S. Yamakoshi, B. Monemar, E. Janz
en, Electronic properties of the residual donor in unintentionally doped β-Ga2O3, J. Appl. Phys. 120 (2016) 235703 with permission from the American Institute of Physics.

and environment of the Si ions in the unintentionally doped samples before they become active? Son et al. [22] suggest that Si is initially passivated by other impurities or intrinsic defects, similar to Mg acceptors in GaN [24]. A passivation model with the Si on the Ga(I) site and an adjacent defect is most likely correct, but consideration should also be given to other possibilities such as Si initially at interstitial sites or forming Si2 molecules, perhaps in a split-interstitial configuration. Interesting results obtained in early experiments on conduction-band electrons in β-Ga2O3 revealed a magnetic bistability associated with the conduction-electron spins [10–14]. As described by Aubay and Gourier [11], hyperfine interactions between conduction-electron spins and nuclear spins of gallium are responsible for a strong dynamic nuclear polarization when the EPR of the conduction electrons is saturated (i.e., the Overhauser effect). The resulting nuclear field acting on the electron spins is bistable, thus causing a hysteresis of the resonance line.

8.4

Acceptors and self-trapped holes

Localization of holes in oxide crystals has been widely studied. Small polarons of the type described in detail by Schirmer [25, 26] and Stoneham et al. [27] are typically observed in EPR experiments, with acceptor-bound small polarons being the most common hole-like defects in the oxides. A common example is an acceptor at a cation site that converts from the singly ionized A charge state to the neutral A0 state when a

174

Gallium Oxide

hole is localized (i.e., trapped) on an adjacent oxygen ion. In this case, the participating acceptor may be either a cation vacancy or an impurity (with less positive charge than the constituent cation it replaces). As demonstrated in this section, holes in β-Ga2O3 crystals can be trapped on oxygen ions adjacent to Ga vacancies or adjacent to Mg ions at Ga sites. Also, at sufficiently low temperatures, holes can be self-trapped at oxygen ions in otherwise unperturbed regions of the crystal. Acceptors similar to Mg, such as Li, Cu, and Zn substituting for gallium ions, are expected to be studied with EPR in the future. Although their spectra also have not yet been reported, another class of acceptors that may become important in β-Ga2O3 involves Group V impurities (N, P, and As) substituting for oxygen ions.

8.4.1 Doubly ionized gallium vacancies The EPR spectrum of the doubly ionized gallium vacancy (V2 Ga ) has been observed by Kananen et al. [28] in a β-Ga2O3 crystal that was initially irradiated with high-energy neutrons. Before the neutron irradiation, the crystal was n-type because of the unintentional presence of Si shallow donors. The neutron irradiation then lowered the Fermi level in the crystal by producing a significant concentration of gallium-vacancy acceptors, as a result of the displacement of gallium ions by the high-energy neutrons. This allowed the V2 Ga charge state of the gallium vacancies to be seen without photoexcitation, as the Fermi level moved below the 2 /3  level of the isolated gallium vacancy. The EPR spectrum of the doubly ionized gallium vacancy (V2 Ga ) in the neutronirradiated β-Ga2O3 crystal is shown in Fig. 8.3A. This S ¼ 1/2 spectrum has a symmetrical set of partially resolved hyperfine lines caused by interactions with the 69Ga and 71 Ga nuclei at two adjacent Ga sites. These isotopes have I ¼ 3/2 nuclear spins, their natural abundances are 60.1% and 39.9%, respectively, and their magnetic moments are 69μ ¼ +2.0166βn and 71μ ¼ + 2.5623βn. As described in Ref. [28], there are three combinations of the two Ga isotopes that give rise to the observed hyperfine pattern: they are (i) two 69Ga nuclei, (ii) one 69Ga nucleus and one 71Ga nucleus, and (iii) two 71 Ga nuclei. The relative amounts of each combination are 36.1%, 48.0%, and 15.9%, respectively. Separate stick diagrams above the experimental spectrum in Fig. 8.3A represent each of these combinations. When the unpaired spin interacts equally with two 69Ga nuclei or two 71Ga nuclei, sets of seven lines (with relative intensities 1:2:3:4:3:2:1) are produced. In contrast, when one 69Ga nucleus and one 71Ga nucleus interact with the unpaired spin at the gallium vacancy, a set of 16 lines are produced, all having the same intensity. The lowest stick diagram in Fig. 8.3A is the sum of the three upper stick diagrams, and thus can be directly compared to the experimental spectrum. Experimental values of 1.28 and 1.63 mT for the 69Ga and 71Ga hyperfine parameters, respectively, were obtained from the EPR spectrum in Fig. 8.3A. A simulated spectrum using these values is shown in Fig. 8.3B. Spectra taken with the magnetic field along the b and c directions showed that the 69Ga and 71Ga hyperfine matrices are nearly isotropic. The angular dependence of the g matrix associated with the V2 Ga acceptor is shown in Fig. 8.4. Turning points occur along the b direction and near the a direction, thus

Fig. 8.3 (A) EPR spectrum from the S ¼ 1/2 doubly ionized gallium vacancy (V2 Ga ) in β-Ga2O3. The spectrum was taken at room temperature with the magnetic field along the a direction in the crystal. Above the spectrum, stick diagrams identify individual hyperfine lines arising from interactions with gallium nuclei at two neighboring gallium sites. (B) Simulated EPR spectrum using the Bruker SimFonia program. Reproduced from B.E. Kananen, L.E. Halliburton, K.T. Stevens, G.K. Foundos, N.C. Giles, Gallium vacancies in β-Ga2O3 crystals, Appl. Phys. Lett. 110 (2017) 202104 with permission from the American Institute of Physics.

2.032

34

347

2.020

348

2.014

349

2.008

350

2.002

351 a

30

60

b

c 30 60 Angle (degrees)

30

60

Magnetic field (mT)

g values

2.026

a*

Fig. 8.4 Angular dependence of the g matrix. The change in position of the center line of the V2 Ga acceptor is shown for rotations in three planes. Magnetic field values along the right vertical axis correspond to a microwave frequency of 9.835 GHz. Reproduced from B.E. Kananen, L.E. Halliburton, K.T. Stevens, G.K. Foundos, N.C. Giles, Gallium vacancies in β-Ga2O3 crystals, Appl. Phys. Lett. 110 (2017) 202104 with permission from the American Institute of Physics.

176

Gallium Oxide

placing the principal axes of the g-matrix approximately along the crystallographic a, b, and c directions. The corresponding principal values of the g matrix are ga ¼ 2.0034, gb ¼ 2.0322, and gc ¼ 2.0097. These small and positive g shifts suggest that the unpaired spin associated with the doubly ionized V2 Ga acceptor is localized in a pz orbital on an oxygen ion located adjacent to the gallium vacancy [29]. The O ion, with the hole, has a 2p5 configuration (2p2x 2p2y 2pz). Discrete energy levels, E1, E2, and E3, are formed when the local crystalline electric field removes the threefold orbital degeneracy of this 2P state (L ¼ 1, S ¼ 1/2). The ground state with the hole in the pz orbital is E1, and the excited states with the hole in the px and py orbitals of the ion are E2 and E3, respectively. The spin-orbit interaction mixes these excited states with the ground state and gives the following first-order expressions for the expected principal g values. g a ¼ ge

(8.1)

gb ¼ g e 

2λ E2  E 1

(8.2)

gc ¼ g e 

2λ E 3  E1

(8.3)

Here, ge (¼2.0023) is the g value for a “free” electron and λ (¼ 135 cm1) is the spinorbit coupling constant for the O ion. The negative sign for λ produces positive g shifts (i.e., g values greater than 2.0). The measured value of 2.0034 for ga is close to 2.0023, thus indicating that the pz orbital containing the unpaired electron spin is oriented approximately along the a direction in the crystal. The model proposed by Kananen et al. [28] for the doubly ionized gallium vacancy (V2 Ga ) is shown in Fig. 8.5. The hole is on a threefold oxygen ion at an O(II) site and the gallium vacancy is at the neighboring sixfold Ga(II) site. As suggested by the g matrix analysis, the oxygen pz orbital containing the unpaired spin (blue in Fig. 8.5) is oriented along the crystal’s a axis and points toward the Ga vacancy. This orientation of the pz orbital results when the large negative “effective” charge of the gallium vacancy strongly attracts the positive hole. Also, this orientation minimizes the total energy of the doubly ionized acceptor by allowing the hole to avoid its positive neighbors as much as possible. The resolved hyperfine lines seen in Fig. 8.3 are due to the interactions with the 69,71Ga nuclei at the two equivalent sites labeled GaA(I) and GaB(I) in Fig. 8.5. Kananen et al. [28] also observed singly ionized gallium vacancies (V Ga) in the neutron irradiated β-Ga2O3 crystal. The two holes in this acceptor are trapped at individual oxygen ions located on opposite sides of the gallium vacancy. They are weakly coupled and form a triplet S ¼ 1 state. Fig. 8.6 shows an EPR spectrum from the V Ga acceptor. The S ¼ 1 spectrum consists of two sets of lines located at lower and higher fields with the V Ga spectrum in the middle. These two spectra have very nearly the same g values and they both have resolved hyperfine patterns from neighboring 69,71 Ga nuclei. As expected, the separations between adjacent hyperfine lines in the

p

Vacancy

a

b

c

Fig. 8.5 Model of the doubly ionized gallium vacancy (V2 Ga ) in β-Ga2O3. An unpaired spin (the hole) is localized in a pz orbital on a threefold oxygen ion, O(II), adjacent to a gallium vacancy (dashed square) at a sixfold Ga(II) site. Resolved hyperfine interactions are with the two equivalent gallium ions labeled GaA(I) and GaB(I). Reproduced from B.E. Kananen, L.E. Halliburton, K.T. Stevens, G.K. Foundos, N.C. Giles, Gallium vacancies in β-Ga2O3 crystals, Appl. Phys. Lett. 110 (2017) 202104 with permission from the American Institute of Physics. − (S = 1) VGa

Fig. 8.6 EPR spectrum from the S ¼ 1 singly ionized gallium vacancy (V Ga) in β-Ga2O3, taken with a microwave frequency of 9.8306 GHz. The direction of the magnetic field is approximately midway between the a and c directions. A stick diagram above the spectrum identifies gallium hyperfine lines. Reproduced from B.E. Kananen, L.E. Halliburton, K.T. Stevens, G.K. Foundos, N.C. Giles, Gallium vacancies in β-Ga2O3 crystals, Appl. Phys. Lett. 110 (2017) 202104 with permission from the American Institute of Physics.

2– VGa

340

345 350 355 Magnetic field (mT)

360

178

Gallium Oxide

S ¼ 1 spectrum are approximately a factor of two less than those in the S ¼ 1/2 spec trum. In Fig. 8.5, the V Ga acceptor is formed when the VGa acceptor traps a second hole at the O(III) oxygen that is opposite the hole at O(II), the top oxygen ion.

8.4.2 Neutral Mg acceptors The EPR spectrum of the neutral Mg acceptor (Mg0Ga) has been investigated by Kananen et al. [30]. An Mg-doped crystal was irradiated near 77 K with 60-kV X-rays and then transferred to the EPR spectrometer without significant warming. The X-rays generated large concentrations of electrons and holes in the conduction and valence bands, respectively. Although many of these electrons and holes immediately recombined, an easily observed number of holes were trapped by the Mg ions and a corresponding number of electrons were trapped at unintentionally present Fe3+ and Cr3+ ions. The S ¼ 1/2 EPR spectrum of the neutral Mg acceptor (Mg0Ga), shown in Fig. 8.7A, contains partially resolved hyperfine lines from neighboring 69Ga and 71Ga nuclei. Spectra taken with the magnetic field along the b and c directions in the crystal were very similar, thus indicating that the hyperfine matrices are nearly isotropic. The Mg0Ga spectrum resembles the spectrum in Fig. 8.3A from the doubly ionized gallium vacancy (V2 Ga ) [28]. A comparison of the spectra in Figs. 8.3A and 8.7A, however, reveals a distinct difference. The spectrum in Fig. 8.3B has a line at the center of the symmetrical hyperfine pattern, whereas the spectrum in Fig. 8.7A does not have a center line. Kananen et al. [30] attributed this absence of a center line in the Mg0Ga spectrum to unequal interactions of the unpaired spin with the gallium nuclei at two neighboring cation sites. An analysis of the hyperfine pattern in Fig. 8.7A showed that the parameters for the two Ga neighbors differed by a factor of 2.2. Fig. 8.7 (A) EPR spectrum from the S ¼ 1/2 neutral magnesium acceptor (Mg0Ga) in Mg-doped β-Ga2O3, taken at 40 K with the magnetic field along the a direction and a microwave frequency of 9.3979 GHz. (B) Simulated EPR spectrum produced using EasySpin. Reproduced from B.E. Kananen, L.E. Halliburton, E.M. Scherrer, K.T. Stevens, G.K. Foundos, K.B. Chang, N.C. Giles, Electron paramagnetic resonance study of neutral Mg acceptors in β-Ga2O3 crystals, Appl. Phys. Lett. 111 (2017) 072102 with permission from the American Institute of Physics.

Experiment

(A)

Simulation

(B) 328

333

338

Magnetic field (mT)

343

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

179

A simulated spectrum of the Mg0Ga acceptor, created with EasySpin [31], is shown in Fig. 8.7B. This simulation was produced using principal hyperfine values of 2.61 and 3.32 mT for the 69Ga and 71Ga nuclei at one Ga neighbor and values of 1.18 and 1.49 mT for the 69Ga and 71Ga nuclei at the other Ga neighbor. As noted by Kananen et al. [30], hyperfine lines from 25Mg nuclei were not present in the experimental spectrum because of their small natural abundance (10%) and their expected small hyperfine parameters. The angular dependence associated with the g matrix of the Mg0Ga acceptor is shown in Fig. 8.8. Turning points appear near the a direction and near the c direction. The corresponding principal values of the g matrix are ga ¼ 2.0038, gb ¼ 2.0153, and gc ¼ 2.0371. From an analogy to the doubly ionized gallium vacancy (V2 Ga ) in β-Ga2O3, the small and positive g shifts observed for the neutral Mg0Ga acceptor indicate that the hole is located in a p orbital on an oxygen ion adjacent to the Mg ion [28]. The pz orbital containing the unpaired spin is aligned nearly along the a direction in the crystal since the value of 2.0038 for ga is close to the 2.0023 “free” spin value. The model proposed by Kananen et al. [30] for the neutral magnesium acceptor (Mg0Ga) is shown in Fig. 8.9. The hole is localized on a threefold oxygen ion at an O(I) site with the Mg ion at a neighboring sixfold Ga(II) site. In constructing this 2.044 329 2.038 330

g values

331 2.026 332 2.020 333 2.014

Magnetic field (mT)

a

2.032

334 2.008 335 2.002 a

30 60

b

30 60 c 30 60 a* 120 150 –c Angle (degrees)

Fig. 8.8 Angular dependence of the g matrix of the Mg0Ga acceptor. The center of the hyperfine pattern is plotted versus direction of the magnetic field for rotations in three planes. The discrete points are experimental results and the solid curves are computer-generated. Magnetic field values along the right vertical axis correspond to a microwave frequency of 9.398 GHz. Reproduced from B.E. Kananen, L.E. Halliburton, E.M. Scherrer, K.T. Stevens, G.K. Foundos, K.B. Chang, N.C. Giles, Electron paramagnetic resonance study of neutral Mg acceptors in β-Ga2O3 crystals, Appl. Phys. Lett. 111 (2017) 072102 with permission from the American Institute of Physics.

180

Gallium Oxide

O(III)

O(III) pz O(III)

O(III)

O(I)

Mg

O(III)

O(I)

O(I)

Ga(II)

Ga(I) O(II)

a*

O(II) c

b

Fig. 8.9 Model of the neutral magnesium acceptor (Mg0Ga) in a β-Ga2O3 crystal. The unpaired spin (the hole shown in blue) is localized in a nonbonding p orbital on a threefold oxygen ion, O(I), adjacent to the Mg ion at a sixfold Ga(II) site. The primary hyperfine interactions are with the Ga(I) and Ga(II) ions adjacent to the hole. Reproduced from B.E. Kananen, L.E. Halliburton, E.M. Scherrer, K.T. Stevens, G.K. Foundos, K.B. Chang, N.C. Giles, Electron paramagnetic resonance study of neutral Mg acceptors in β-Ga2O3 crystals, Appl. Phys. Lett. 111 (2017) 072102 with permission from the American Institute of Physics.

model, the following observations were used. Having resolved hyperfine from two, not three, neighboring Ga sites requires the hole to be localized on a threefoldcoordinated oxygen ion with the Mg ion at one of the three nearest-neighbor Ga positions, either an O(I) or an O(II) ion. The inequivalence of the two Ga hyperfine interactions strongly suggest that the hole must be on an O(I) ion, instead of an O(II) ion. As suggested by computational studies [32, 33], the Mg ion is at a sixfoldcoordinated gallium site in Fig. 8.9. The alignment of the hole’s p orbital along the a direction corresponds to a minimum energy configuration for the ground state of the neutral Mg acceptor. Specifically, the hole is oriented approximately perpendicular to the plane defined by the three cation neighbors (the Ga(I), Ga(II), and Mg ions) and thus lowers its energy by avoiding the positive Mg and Ga neighbors. It is appropriate to view the neutral Mg acceptor (Mg0Ga) as an acceptor-bound small polaron in the partially ionic β-Ga2O3 crystal [25]. In ionic terms, the Mg2+ ion at the Ga3+ site attracts a hole, with the hole residing on an adjacent oxygen ion instead of either forming an Mg3+ ion or having the hole delocalized but centered on the Mg ion. After producing the neutral Mg acceptors (Mg0Ga) during an X-ray irradiation near 77 K, Kananen et al. [30] performed a thermal pulse anneal study. The Mg0Ga acceptors became thermally unstable near 250 K as the hole “moves” away from the Mg ion via a small-polaron hopping process. It is also possible that the Mg0Ga acceptor decays because an electron is released from Fe2+ or Cr2+ ions and recombines with the hole at the Mg site. In either case, an approximate activation energy can be determined.

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

181

The approximation E  23kTm where Tm corresponds to the temperature where half of the acceptors have thermally decayed (Tm is 250 K for the Mg0Ga acceptor) gives an activation energy near 500 meV [34].

8.4.3 Self-trapped holes The self-trapped hole (STH) is an important intrinsic defect in β-Ga2O3. Computational studies [35, 36] have predicted the existence of STHs in β-Ga2O3, and Kananen et al. [37] have experimentally observed their EPR spectrum at temperatures below 80–100 K in single crystals. At room temperature, the short-lived STHs are the precursor to the self-trapped excitons (STEs) that are responsible for the dominant luminescence observed near 370–380 nm in β-Ga2O3 crystals [38]. As described in Ref. [37], the direct observation of the STHs with EPR requires a β-Ga2O3 sample having (1) a low Fermi level and (2) the presence of other defects that can serve as stable deep electron traps. In their study of the STHs, Kananen et al. [37] used a crystal that had been irradiated with fast neutrons. The neutron irradiation significantly lowered the Fermi level by introducing gallium vacancies [28] and allowed the STHs to be produced during a subsequent irradiation at 77 K with X-rays. An EPR spectrum from the STHs in the X-ray-irradiated β-Ga2O3 crystal is shown in Fig. 8.10. The STH spectrum has a distinctly different 69,71Ga hyperfine pattern when compared 0 to the gallium vacancy (V2 Ga ) and Mg acceptor (MgGa) spectra in Figs. 8.3 and 8.7. A different choice of sample to use for observing a stable STH with EPR would be an Mg-doped crystal, since these crystals also have low Fermi levels and usually contain trace amounts of Fe3+ and Cr3+ impurities that can act as stable electron traps. It should be possible to produce stable STHs during an irradiation of an Mg-doped β-Ga2O3 crystal at 77 K with X-rays. The STH spectrum, however, will be difficult to isolate in an EPR experiment because it will be completely overlapped by the Mg0Ga spectrum for all orientations of magnetic field (the two spectra have similar g values and the Ga hyperfine separations are smaller for the STH spectrum). Equivalent interactions with the 69Ga and 71Ga nuclei at two neighboring Ga sites are responsible for the symmetrical hyperfine pattern in Fig. 8.10. As described earlier for the doubly ionized gallium vacancy, there are three combinations of the two Ga isotopes that contribute to the observed hyperfine pattern. Instead of the 29 possible individual lines in the spectrum, only the “envelopes” of seven sets of these lines are resolved because of overlaps arising from large linewidths. Kananen et al. [37] show that the 69,71Ga hyperfine matrices are nearly isotropic and they extract values of 0.92 and 1.16 mT for the 69Ga and 71Ga hyperfine parameters, respectively, when the magnetic field is along the a direction. The Ga nuclei at the two adjacent Ga sites have the same hyperfine values. The angular dependence associated with the g matrix of the STH in β-Ga2O3 is shown in Fig. 8.11. Similar to the Mg0Ga acceptor, turning points occur near the a and c directions. This places the principal axes of the g matrix approximately along the a, b, and c directions in the crystal. The principal values of the g matrix for the STH are ga ¼ 2.0026, gb ¼ 2.0072, and gc ¼ 2.0461. These positive and small g shifts suggest that the STH has the hole localized in a pz orbital on one oxygen ion (very much

Fig. 8.10 EPR spectrum from the STH in β-Ga2O3 taken at 30 K immediately after an X-ray irradiation at 77 K (the crystal had been previously neutron irradiated). The magnetic field is along the a direction, and the microwave frequency is 9.3968 GHz. Stick diagrams above the spectrum refer to the hyperfine lines corresponding to different pairs of 69Ga and 71 Ga nuclei at the two neighboring gallium sites. Reproduced from B.E. Kananen, N.C. Giles, L.E. Halliburton, G.K. Foundos, K.B. Chang, K.T. Stevens, Self-trapped holes in β-Ga2O3 crystals, J. Appl. Phys. 122 (2017) 215703 with permission from the American Institute of Physics.

69

Ga,69Ga

69

Ga,71Ga

71

Ga,71Ga

Combined

Experiment

331

333

337 335 Magnetic field (mT)

339

328 2.044 329

a

330

2.032 g values

331 2.026 332

2.020

333

2.014

Magnetic field (mT)

2.038

334 2.008 335 2.002 0306090120150180210240270300330360 a 30 60 b 30 60 c 30 60 a* 120 150 −c Angle (degrees)

Fig. 8.11 Angular dependence associated with the g matrix of the STH in β-Ga2O3. The shifts in the middle position of the STH spectrum are shown for rotations in three planes. Discrete points are experimental and the solid curves are computer generated. Magnetic field values along the right vertical axis correspond to a microwave frequency of 9.3968 GHz. Reproduced from B.E. Kananen, N.C. Giles, L.E. Halliburton, G.K. Foundos, K.B. Chang, K.T. Stevens, Self-trapped holes in β-Ga2O3 crystals, J. Appl. Phys. 122 (2017) 215703 with permission from the American Institute of Physics.

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

O(III)

183

O(III) pz

O(III)

Ga(II)

O(III)

O(III) O(III)

O(I)

O(I)

O(I)

Ga(II)

a* Ga(I) O(II)

O(II)

c O(II) O(II)

b

Fig. 8.12 Model of the STH in β-Ga2O3. The unpaired spin (the hole) is localized in a pz orbital on a threefold oxygen ion O(I). The two Ga(II) ions are responsible for the resolved 69Ga and 71 Ga hyperfine interactions observed in the EPR spectra. In order to self-trap the hole, the Ga(I) ion relaxes away from the O(I) ion. Reproduced from B.E. Kananen, N.C. Giles, L.E. Halliburton, G.K. Foundos, K.B. Chang, K.T. Stevens, Self-trapped holes in β-Ga2O3 crystals, J. Appl. Phys. 122 (2017) 215703 with permission from the American Institute of Physics.

like the doubly ionized gallium vacancy [28] and the neutral Mg acceptor [30]). The value of 2.0026 for ga is close to 2.0023, which indicates that the pz orbital containing the unpaired spin is oriented nearly along the a direction in the crystal. The model proposed by Kananen et al. [37] for the STH in β-Ga2O3 is shown in Fig. 8.12. The hole is located at an O(I) ion, in a nonbonding p orbital aligned along the a direction. This nonbonding oxygen orbital is oriented approximately perpendicular to the plane defined by the three neighboring Ga ions, two Ga(II) and one Ga(I), and thus allows the hole (i.e., the unpaired spin) to avoid its positive Ga neighbors. As seen in Figs. 8.9 and 8.12, the STH and the neutral Mg acceptor (Mg0Ga) have very similar models. The primary difference is the Mg0Ga acceptor has one of the two Ga(II) neighbors replaced by the Mg ion. Kananen et al. [37] suggest that the shallow potential well needed to self-trap the hole is produced when the Ga(I) ion in Fig. 8.12 moves away from the O(I) ion (with the hole) and toward the plane containing its remaining three neighboring oxygen ions. The Ga(I) ion, the two O(II) ions, and the one O(III) ion form a more planar GaO3 unit. Nuclei at the two equivalent Ga(II) sites in Fig. 8.12 are responsible for the resolved hyperfine interactions seen in the EPR spectrum in Fig. 8.10. Nuclei at the relaxed Ga(I) site have a weaker, and experimentally unresolved, hyperfine interaction. Thermal stability is an important characteristic of the STH in β-Ga2O3. In a series of isochronal thermal anneals, Kananen et al. [37] determined that the STH becomes unstable in the 80–110 K region, after initially being produced at 77 K with X-rays. An estimate of the activation energy for this decay was obtained from the approximation E  23kTm, where Tm corresponds to the temperature where half of the STHs have disappeared [34]. Using Tm  97 K, Kananen et al. [37] found an approximate value

184

Gallium Oxide

of 190 meV for the thermal activation energy for decay of the STH. In a β-Ga2O3 crystal with a low Fermi level, the STHs do not decay in the 80–110 K region because an electron returns to an STH, but rather because the hole moves to a trapped electron. When there is sufficient thermal energy, the hole migrates through the lattice (i.e., by hopping from oxygen to oxygen) until a trapped electron is encountered. Thus, the observed STH decay in the 80–110 K region in our crystal is an intrinsic property of the STHs and does not depend on the specific identities of the participating electron traps [37].

8.5

Transition-metal and rare-earth ions

8.5.1 Cr3+ ions A detailed study of the EPR spectrum of Cr3+ ions in doped β-Ga2O3 crystals has been reported by Yeom et al. [39]. This work followed the initial observation of the Cr3+ spectrum by Tippens [40]. The Cr3+ ions have a 3d3 configuration with S ¼ 3/2. In this monoclinic crystal, the four spin states have a large zero-field splitting and their EPR spectrum is strongly angular dependent. Fig. 8.13 shows this angular dependence when the magnetic field is rotated in the plane perpendicular to the c axis. The EPR lines occur over a 1.3 T range for microwave frequencies near 9.5 GHz. The following spin Hamiltonian, where y is perpendicular to the crystal’s mirror plane, was used to describe the Cr3+ spectrum [39]. H ¼ βS g  B + B02 O02 + B22 O22 + B12 O12 Bqk are experimentally determined coefficients and Oqk are Stevens operators [5]. Yeom et al. [39], in fitting the angular dependence of the spectra in three crystal planes, Fig. 8.13 The stack plot of the Cr3+ EPR spectra, measured at 9.504 GHz in the a*b plane at room temperature. The maximum separation of the fine structure occurs when the static magnetic field is along the b axis. Reproduced from T.H. Yeom, I.G. Kim, S.H. Lee, S.H. Choh, Y.M. Yu, Electron paramagnetic resonance characterization of Cr3+ impurities in a β-Ga2O3 single crystal, J. Appl. Phys. 93 (2003) 3315 with permission from the American Institute of Physics.

b-axis

a* -axis

0.0

0.3

0.6

0.9

Magnetic field (T)

1.2

1.5

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

185

obtained 1.962, 1.964, 1.979 for the principal values of the g matrix and B02 ¼  1548, B22 ¼  1535, B12 ¼  2668 for the zero-field parameters (in units of 104 cm1). These results demonstrated that the Cr3+ ions are on Ga(II) sites with six oxygen neighbors, as opposed to the tetrahedral Ga(I) sites. Many β-Ga2O3 crystals contain trace amounts of Cr because of impure starting materials used in the growth. The majority of the Cr ions are expected to be in the Cr3+ charge state, but other charge states may also be present. In n-type crystals (with high Fermi levels), a significant portion of the unintentionally present Cr ions may be in the Cr2+ (3d4) charge state. In contrast, crystals with lower Fermi levels (e.g., those containing Mg acceptors or gallium vacancies) may have a portion of the Cr ions in the Cr4+ (3d2) charge state. Thus far, however, there have been no reports in the literature of EPR spectra from these Cr2+ or Cr4+ ions.

8.5.2 Fe3+ ions An EPR spectrum from Fe3+ ions in a β-Ga2O3 crystal has been reported by Meil’man [41]. The Fe3+ ions have a 3d5 configuration with S ¼ 5/2 and their spectrum exhibits a strong angular dependence for microwave frequencies near 9.5 GHz. The principal g values are 2.004, 2.002, and 2.007, and the primary zero-field parameters are B02 ¼ 2.212 GHz and B22 ¼ 2.090 GHz [41]. Effects of the monoclinic structure are readily seen in the angular dependence data. Although not established in Ref. [41], the Fe3+ ions most likely occupy the Ga(II) site. An EPR spectrum from Fe3+ ions in a recently grown Mg-doped β-Ga2O3 crystal is shown in Fig. 8.14. This spectrum was taken at room temperature with the magnetic Fe3+ in b-Ga2O3 RT, 9.4 GHz B along b axis

200

400

600

800

1000

1200

Magnetic field (mT)

Fig. 8.14 Room temperature EPR spectrum from Fe3+ ions in an Mg-doped β-Ga2O3 crystal. The four larger lines are from the Fe3+ ions, and the weak lines between 600 and 1200 mT are from O2 molecules from the air within the microwave cavity.

186

Gallium Oxide

field along the crystal’s b axis and a microwave frequency of 9.4 GHz. In Fig. 8.14, there are four prominent Fe3+ lines near 165, 285, 590, and 1290 mT and two lines from Cr3+ ions near 130 and 1300 mT. Although not yet observed in β-Ga2O3 crystals, future investigations are expected to reveal the presence of Fe2+ ions when the Fermi level is high and Fe4+ ions when the Fermi level is low.

8.5.3 Mn2+ ions Kim et al. [42] have reported a complete study of the EPR spectrum of Mn2+ ions in a β-Ga2O3 crystal. This spectrum has also been investigated by Yamaga et al. [21]. The Mn2+ ions have a 3d5 configuration and thus S ¼ 5/2. Resolved hyperfine lines are present in the spectrum from the 55Mn nuclei (100% abundant with I ¼ 5/2). The following spin Hamiltonian describes the Mn2+ EPR spectrum. H ¼ βS  g  B + I  A  S + B02 O02 + B22 O22 + B12 O12 + B04 O04 + B14 O14 + B24 O24 + B34 O34 + B44 O44

The principal g values are 2.014, 2.012, and 2.001, and the primary zero-field parameters are B02 ¼ 154.4, B22 ¼ 138.8, B12 ¼ 198 for the zero-field parameters (in units of 104 cm1) [41]. These results support having the Mn2+ ions at the Ga(II) sites with six oxygen neighbors.

8.5.4 Ti3+ ions EPR and ENDOR investigations of titanium ions in β-Ga2O3 crystals have been reported by Mentink-Vigier et al. [43–45]. Although Ti3+ (3d1) ions are expected in the titanium-doped crystals, large hyperfine interactions with at least eight Ga neighbors led these researchers to assign their observed spectra to a conduction electron trapped by a Ti4+ ion at a Ga(II) sixfold-coordinated site. Specifically, they describe the spin system as a diffuse (Ti4+-e) pair rather than a localized Ti3+ ion. The principal axes of the g matrix are along the crystal’s a, b, and c directions with principal values ga ¼ 1.949, gb ¼ 1.850, and gc ¼ 1.923. An EPR spectrum taken at 20 K with the magnetic field along the b direction showed numerous partially resolved hyperfine lines from neighboring 69Ga and 71Ga nuclei. Rich ENDOR spectra were also observed at 20 K from the surrounding Ga nuclei and the analysis of their angular dependence provided complete sets of hyperfine and nuclear electric quadrupole parameters. The authors [43] suggest that titanium in β-Ga2O3 (with its multiple Ga hyperfine interactions) is a candidate for a spin-bus system that may find use in quantum information processing.

8.5.5 Er3+ ions An EPR spectrum from Er3+ ions in β-Ga2O3 has been reported by Vincent et al. [46]. The electron configuration of Er3+ is 4f11. According to Hund’s rule, the ground state multiplet of this Kramers ion has orbital angular momentum L ¼ 6 and electron spin

Electron paramagnetic resonance (EPR) from β-Ga2O3 crystals

187

S ¼ 3/2, and the resulting values of J are 15/2, 13/2, 11/2, and 9/2. The four terms in this multiplet are 4I15/2, 4I13/2, 4I11/2, and 4I9/2, with 4I15/2 lying lowest in energy because the 4f shell is more than half full. Additional splittings, due to the partially shielded crystalline electric field, occur when the Er3+ ion is incorporated in a host lattice. The 2J + 1 states of the 4I15/2 term split into three quartets and two doublets in a cubic or tetragonal field. In the monoclinic structure of β-Ga2O3, all degeneracies are removed and eight Kramers doublets are formed. The lowest of these Kramers doublets gives rise to the observed EPR spectrum. The EPR spectra of the Er3+ ions were obtained at 6.5 K from β-Ga2O3 crystals with nominal Er dopings of 0.5 and 1.5 mol% [46]. An angular-dependence study showed that the principal-axis directions of the g matrix were close to the a, b, and c directions in the crystal and that the principal values were ga ¼ 5.75, gb ¼ 9.45, and gc ¼ 0.9. A well-resolved set of eight hyperfine lines from 167Er nuclei (I ¼ 7/2, 22.9% abundant) were observed when the magnetic field was along the b direction. The Er3+ ions were found to occupy only one crystallographic site in the crystal, with a slight ˚ ionic radius, the Er3+ ion is expected g-factor distribution. Because of its large 0.89 A to substitute for Ga at the Ga(II) octahedral site.

8.6

Oxygen vacancies

The EPR spectrum from shallow donors (Si, Sn, etc.) in β-Ga2O3 has been mistakenly assigned in several papers to the singly ionized oxygen vacancy. Although the assignments were reasonable assumptions in the early stages of research on this material, a recent computational study has shown that the oxygen vacancy is a deep donor in β-Ga2O3, and thus will not directly contribute to conductivity [23]. Of the three charge + 0 + states of an oxygen vacancy (V2+ O , VO, and VO), only the singly ionized VO state has an unpaired spin and thus can be observed with EPR. The singly ionized oxygen vacancy is expected to have a predominantly s-like nature due to the localization of the unpaired spin primarily on the nearest-neighbor Ga ions. Thus, instead of a single relatively sharp line similar to the hydrogenic shallow donors, the EPR spectrum of the V+O centers in β-Ga2O3 will be broad (tens of mT) with partially resolved hyperfine structure from the neighboring 69,71Ga nuclei. The EPR spectra of singly ionized oxygen vacancies in α-Al2O3, LiAlO2, and ZnO all have resolved hyperfine structure caused by the large interactions with their nearest-neighbor cations [47–50]. These previous studies of EPR signals from oxygen vacancies can provide guidance as to what spectral features are expected for the same defect in β-Ga2O3. In Al2O3 crystals, the EPR spectrum of the oxygen vacancy (V+O) consists of a set of lines regularly spaced at about 0.5 mT intervals and extending over nearly 70 mT [47]. Interactions with the 27Al nuclei (100% abundant, I ¼ 5/2) at two adjacent Al sites are resolved in the EPR spectrum. In LiAlO2 crystals, the EPR spectrum of the oxygen vacancy (V+O) consists of a large number of lines extending over nearly 100 mT [48]. This structure is caused by the hyperfine interactions with the two 27Al nuclei that are adjacent to the vacancy. A similar situation occurs in ZnO, with the difference being the 67Zn isotope is only 4.1% abundant with I ¼ 5/2. The low abundance of 67Zn

188

Gallium Oxide

produces a narrow I ¼ 0 line, with less intense hyperfine lines extending over a 13 mT region centered on the I ¼ 0 line [49,50]. Based on the insights gained from these other wide-band-gap oxides, the EPR spectrum from the oxygen vacancy in β-Ga2O3 is expected to have a resolved hyperfine pattern from the 69,71Ga nuclei (together the two Ga isotopes are 100% abundant) that extends over at least 15 mT. Because of the anticipated broad lines, a large concentration of V+O centers will be required to successfully observe their EPR spectrum.

Acknowledgments The authors wish to acknowledge in-depth discussions with B. E. Kananen, E. M. Scherrer, and C. A. Lenyk at the Air Force Institute of Technology (AFIT) concerning defects in β-Ga2O3. The views expressed in this chapter are those of the authors and do not necessarily reflect the official policy or position of the Air Force, the Department of Defense, or the United States Government.

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[36] P. Dea´k, Q.D. Ho, F. Seemann, B. Aradi, M. Lorke, T. Frauenheim, Choosing the correct hybrid for defect calculations: A case study on intrinsic carrier trapping in β-Ga2O3, Phys. Rev. B 95 (2017) 075208. [37] B.E. Kananen, N.C. Giles, L.E. Halliburton, G.K. Foundos, K.B. Chang, K.T. Stevens, Self-trapped holes in β-Ga2O3 crystals, J. Appl. Phys. 122 (2017) 215703. [38] S. Yamaoka, Y. Furukawa, M. Nakayama, Initial process of photoluminescence dynamics of self-trapped excitons in a β-Ga2O3 single crystal, Phys. Rev. B 95 (2017) 094304. [39] T.H. Yeom, I.G. Kim, S.H. Lee, S.H. Choh, Y.M. Yu, Electron paramagnetic resonance characterization of Cr3+ impurities in a β-Ga2O3 single crystal, J. Appl. Phys. 93 (2003) 3315. [40] H.H. Tippins, Optical and microwave properties of trivalent chromium in β-Ga2O3, Phys. Rev. 137 (1965) A865. [41] M.L. Meilman, EPR of Fe3+ ions in β-Ga2O3 crystals, Sov. Phys. Solid State 11 (1969) 1403. [42] I.G. Kim, T.H. Yeom, S.H. Lee, Y.M. Yu, H.W. Shin, S.H. Choh, Electron paramagnetic resonance studies of Mn2+ ions in β-Ga2O3 single crystal, J. Appl. Phys. 89 (2001) 4470. [43] F. Mentink-Vigier, L. Binet, G. Vignoles, D. Gourier, H. Vezin, Giant titanium electron wave function in gallium oxide: A potential electron-nuclear spin system for quantum information processing, Phys. Rev. B 82 (2010) 184414. [44] F. Mentink-Vigier, L. Binet, D. Gourier, Strong isotopic effect in the electron-mediated nuclear-nuclear interaction in solids, Phys. Rev. B 83 (2011) 214409. [45] F. Mentink-Vigier, L. Binet, D. Gourier, H. Vezin, Origin of the decoherence of the extended electron spin state in Ti-doped β-Ga2O3, J. Phys. Condens. Matter 25 (2013) 316002. [46] J. Vincent, O. Guillot-Noe¨l, L. Binet, P. Aschehoug, Y. Le Du, F. Beaudoux, P. Goldner, Electron paramagnetic resonance and optical spectroscopy of Er-doped β-Ga2O3, J. Appl. Phys. 104 (2008) 033519. [47] S.Y. La, R.H. Bartram, R.T. Cox, The F+ center in reactor-irradiated aluminum oxide, J. Phys. Chem. Solids 34 (1973) 1079. [48] M.S. Holston, I.P. Ferguson, J.W. McClory, N.C. Giles, L.E. Halliburton, Oxygen vacancies in LiAlO2 crystals, Phys. Rev. B 92 (2015) 144108. [49] J.M. Smith, W.E. Vehse, ESR of electron irradiated ZnO confirmation of F+ center, Phys. Lett. 31A (1970) 147. [50] C. Gonzales, D. Galland, A. Herve, Interactions hyperfines du centre F+ dans ZnO, Phys. Status Solidi B 72 (1975) 309.

Hydrogen in Ga2O3

9

Michael Stavola*, W. Beall Fowler*, Ying Qin*, Philip Weiser*, Stephen Pearton† *Department of Physics, Lehigh University, Bethlehem, PA, United States, †Department of Materials Science and Engineering, University of Florida, Gainesville, FL, United States

Chapter Outline 9.1 Introduction 191 9.2 Hydrogen in the transparent conducting oxides ZnO, SnO2, and In2O3 9.3 Hydrogen in β-Ga2O3 195 9.3.1 9.3.2 9.3.3 9.3.4

192

Theory 195 Thermal stability of deuterium in Ga2O3 196 Muon spin resonance 196 Vibrational properties of H in Ga2O3 196

9.4 Conclusion 205 Acknowlegments 206 References 206

9.1

Introduction

Semiconductors with bandgaps larger than the 3.4 eV bandgap of GaN are emerging as a new class of ultrawide-bandgap (UWBG) electronic materials [1–5]. In spite of the promising applications that are possible for UWBG materials, an understanding of their fundamental properties is at an early stage of development. The focus of this chapter is the hydrogen impurity and its interactions with other defects in β-Ga2O3, a transparent conducting oxide with an ultrawide bandgap of 4.9 eV [6–9]. (It is the most thermally stable monoclinic β phase of Ga2O3 to which we refer throughout this chapter.) The UWBG semiconductors show promise for device applications with dramatically improved performance. The large bandgap of Ga2O3 leads to a theoretical breakdown field of 8 MV/cm [3, 4, 9, 10]. The Baliga figure of merit for Ga2O3 (the figure of merit for power devices) is at least four times larger than those of GaN and 4H-SiC. An experimental breakdown field of 3.8 MV/cm for Ga2O3 has already been achieved in metal-oxide-semiconductor field-effect transistors and is higher than the critical field strengths of GaN and SiC [11]. The ultrawide bandgap of Ga2O3 also opens up opportunities for optoelectronic devices in the deep UV and for devices that can operate in harsh environments. An advantage of Ga2O3 over the wide-bandgap nitrides is the availability of native single-crystal substrates [4]. Bulk Ga2O3 can be grown by the floating zone [12], Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00009-9 © 2019 Elsevier Inc. All rights reserved.

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Czochralski [13], vertical Bridgeman [14], and edge-defined film-fed growth (EFG) methods [15]. With the EFG method, it has been possible to produce high-quality, 400 -diameter, Ga2O3 wafers. Ga2O3 shows unintentional n-type conductivity [16]. While the conventional wisdom has been that O vacancies can cause this conductivity, theory finds that O vacancies are not shallow donors in Ga2O3 and suggests that H can be a cause of n-type behavior [17]. This chapter is a survey of the properties of hydrogen in Ga2O3. These properties are discussed in the context of hydrogen’s behavior in other transparent conducting oxides.

9.2

Hydrogen in the transparent conducting oxides ZnO, SnO2, and In2O3

The passivation of dopants and defects in conventional semiconductors such as Si, GaAs, and GaN by hydrogen is well known and has an impact on technology that is widely recognized [18]. However, hydrogen in semiconducting oxides is dramatically different [19]. Recent theory has reinvigorated interest in hydrogen as a possible source of conductivity [17, 20–25] and has drawn attention to work performed in the 1950s where hydrogen in ZnO was found to give rise to shallow donors [26, 27]. The properties of hydrogen in ZnO, SnO2, and In2O3 are surveyed briefly here so that the behavior of hydrogen in Ga2O3 can be compared and contrasted with the behavior of hydrogen in other oxides that have been studied recently. Two defects have been predicted to be shallow donors in several oxide hosts: interstitial hydrogen, Hi, and hydrogen trapped at an oxygen vacancy, HO [17, 20–25]. Hi forms strong O-H bonds with stretching frequencies above 3000 cm1. HO gives rise to a novel multicenter bond with a much lower vibrational frequency. The muon is a positively charged particle with 1/9 the mass of a proton that mimics the properties of hydrogen. The spectroscopy of implanted muons has been widely used to investigate the properties of hydrogen in solids [28, 29]. Implanted muons have been found to form shallow donors in several oxide hosts, suggesting that the behavior of hydrogen will be similar. Hydrogen also interacts with cation vacancies in transparent conducting oxides [30] and can also form H2 molecules [31, 32]. ZnO: ZnO has the hexagonal wurtzite structure and a bandgap of 3.4 eV [19]. Hydrogen in ZnO gives rise to both Hi and HO shallow donors whose properties have been investigated extensively by theory and experiment. Muon spin resonance measurements showed that implanted muons form shallow donors in ZnO [33], so a similar result is expected for hydrogen. An O-H vibrational line that is polarized along the c-axis of ZnO was found at 3611 cm1 [34]. The 3611 cm1 line was found to be marginally stable at room temperature and was assigned to the Hi shallow donor [35–37]. When Hi is annealed away at temperatures near 150°C, hydrogen then forms interstitial H2 molecules that provide a reservoir of hydrogen in the sample that can be converted back to Hi by thermal annealing [31, 32].

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193

ZnO also contains a more thermally stable hydrogen shallow donor that is annealed away at  500°C [35–37] and that theory has predicted is due to HO [22]. The HO center in ZnO has been found to form a novel multicenter bond where H is weakly bonded to four Zn neighbors [22]. The vibrational mode predicted for HO in ZnO lies in a region where ZnO is strongly absorbing. In this case, photothermal ionization spectroscopy (PTIS) was used to detect the vibrational lines of HO at 742 and 792 cm1 that were assigned to the A1 and E modes of the C3V center [38]. The Hi and HO shallow donors in ZnO have electronic transitions that also have been studied by photoluminescence, Raman, and PTIS measurements [37]. Infrared (IR) lines at 3312.2 and 3349.6 cm1 have been assigned to the O-H stretching modes of the VZnH2 center in ZnO [34]. IR lines at 3303 and 3321 cm1 have been assigned to the doubly degenerate and singly degenerate O-H stretching modes, respectively, of the VZnH3 center in ZnO [39]. SnO2: SnO2 has the rutile structure and a direct bandgap of 3.6 eV [19]. Early experimental work showed that annealing SnO2 in a hydrogen-containing ambient gives rise to strong n-type conductivity [40]. Muon spin resonance experiments have found that muons implanted into SnO2 form shallow donors [41]. Interstitial hydrogen and hydrogen at an oxygen vacancy have been predicted by theory to be donors in SnO2 [23, 42]. Hi was predicted to be mobile near room temperature, whereas HO was predicted to be more thermally stable. Recent experiments confirm that annealing SnO2 single crystals in an H2 ambient increases their conductivity and also gives rise to several O-H vibrational lines with distinctive polarization properties [42, 43]. Two hydrogen shallow donors have been discovered, one that is marginally stable at room temperature and a second donor that is stable up to 650°C [43]. A vibrational line at 3156.1 cm1 was assigned to the less stable Hi center. The more stable donor has properties consistent with the HO center predicted by theory [23]. The vibrational modes of several additional O-H centers in SnO2 have also been discovered [42, 43]. Structures with H bound to a Sn vacancy (VSn) [42] or with H bound to a Sn interstitial (ISn) [44] have been proposed. The polarization properties of the vibrational lines determined the O-H bond angles for two of these additional O-H centers and support their assignment to structures with H bonded to an interstitial defect, possibly ISn [44]. Other lines remain unassigned and could be due to structures involving VSn. The electrically active, hydrogen-shallow-donor centers have been found to interact with the additional centers that involve hydrogen complexed with native defects [42, 43]. These defects can be interconverted from one to another by thermal treatments and can give rise to unexpected changes in the conductivity of SnO2 samples that contain H upon annealing. In2O3: In2O3 has the cubic bixbyite structure with a conventional unit cell that contains 80 atoms [19, 45]. The oxygen sites are all equivalent, and there are two inequivalent In sites [24]. Recent theoretical and experimental works find that hydrogen can be an important shallow donor in In2O3. Muon-spin-resonance experiments have found that implanted muons form shallow donors in In2O3 [41]. In2O3 thin films containing hydrogen show n-type conductivity with high mobility that has attracted

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attention for solar-cell applications [46, 47]. Furthermore, theory finds that Hi and HO are shallow donors [24]. Annealing In2O3 single crystals in an H2 ambient at temperatures near 450°C has been found to produce a thin conducting layer near the sample surface with a thickness near 100 μm and a carrier concentration of 1019 cm3 [48]. Infrared absorption experiments have found that substantial free-carrier absorption and several O-H vibrational lines are produced by hydrogenation treatments. The broad absorption due to free carriers that increases as the frequency decreases is shown in Fig. 9.1A. The O-H absorption lines produced by the hydrogenation of In2O3 are shown in Fig. 9.1B. In annealing experiments, the O-H vibrational line at 3306 cm1 was found to be correlated with the free-carrier absorption and was assigned to the interstitial hydrogen shallow-donor center that is responsible for the hydrogen-related conductivity [48]. Unlike ZnO and SnO2 where HO is the dominant hydrogen shallow donor, Hi has been found to be the dominant H shallow donor in In2O3. Several additional O-H lines near 3400 cm1 were produced by annealing In2O3 in a hydrogen ambient [48]. These lines have been suggested to be due to defect complexes with H trapped by indium vacancies (VIn).

600 2

0.3

as treated

500 Absorbance

300 1

0.2 400

400 500

0

0.1

300

600

as treated 0.0

2000

3000

4000

3300

3400

–1

Frequency (cm )

(A)

(B)

Fig. 9.1 A selection of IR absorption spectra (T ¼ 4.2 K) for an In2O3 sample that initially had been hydrogenated by an anneal (30 min) in an H2 ambient at 500°C. The sample was then annealed sequentially in flowing He at the temperatures shown in °C. (A) The absorption due to free carriers. (B) The IR absorption lines in the O-H stretching region. These spectra were baseline corrected to remove the contribution from free carriers.

Hydrogen in Ga2O3

195

Hydrogen in β-Ga2O3

9.3

β-Ga2O3 has a complex monoclinic crystal structure (Fig. 9.2A) [49, 50]. There are two inequivalent Ga sites and three inequivalent oxygen sites. Ga(1) and Ga(2) are tetrahedrally and octahedrally coordinated, respectively. O(1) and O(2) are threefold coordinated and O(3) is fourfold coordinated.

9.3.1 Theory Ga2O3 shows unintentional n-type conductivity [16]. While the conventional wisdom has been that O vacancies can cause this conductivity, theory finds that O vacancies are not shallow donors in Ga2O3 and suggests that H can be a cause of n-type behavior [17]. Both Hi and HO have been predicted to be shallow donors [17]. Given the number of possible Ga and O sites in Ga2O3, there are several possible configurations for Hi and HO centers. Hydrogen can affect the electrical properties of Ga2O3 by acting as a shallow donor or by compensating other defects, such as the gallium vacancy, that act as deep acceptors [17, 30]. The Ga vacancy has low formation energy and is predicted to interact with H to produce VGa-H defects with high thermal stability [30]. The interactions of H with VGa in β-Ga2O3 are predicted to give rise to large configurational relaxations of the VGa defect [30, 51].

[1 0 2] direction

a Light perpendicular to (–2 0 1) plane b

c

[0 1 0] direction out of plane

(A)

(B)

Fig. 9.2 (A) The unit cell of β-Ga2O3. The inequivalent sites in this and subsequent figures are color coded as follows: Ga(1), purple; Ga(2), dark green; O(1), red; O(2), yellow; O(3), light green. (B) Experimental orientation of the samples used in our experiments. These and subsequent figures were constructed using MOLDRAW (P. Ugliengo, Torino 2006, available at http://www.moldraw.unito.it/) and POV-Ray (http://povray.org).

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9.3.2 Thermal stability of deuterium in Ga2O3 Deuterium was introduced into the (201) face of β-Ga2O3 by plasma exposure (up to 270°C) and by ion implantation (100 keV, 1015 cm2) to investigate its thermal stability [52, 53]. Secondary ion mass spectrometry (SIMS) experiments showed that the out-diffusion of deuterium from samples treated in a plasma occurs at a lower temperature than for samples deuterated by ion implantation. [Deuterium was used in these experiments because the detection limit for D by SIMS is approximately three decades lower than for H because of the small natural isotopic abundance of D (0.0156%).] For samples deuterated by plasma exposure, the out-diffusion of D from the sample surface was proposed to be limited by the formation of D2 molecules and occurred at an annealing temperature of 500°C [53]. For samples deuterated by ion implantation, the out-diffusion of D was proposed to by limited by the interaction of D with implantation damage and occurred at an annealing temperature of 650°C [53]. Simulations of the in-diffusion and out-diffusion behaviors of D for both plasma-treated and ion-implanted samples produced satisfactory fits to the data.

9.3.3 Muon spin resonance When muons were implanted into powdered samples of β-Ga2O3, they were found to give rise to a shallow donor center with a binding energy between 15 and 30 meV and a ˚ [54]. These results are consistent with the muon forming a shallow Bohr radius of 20 A effective mass-like defect in β-Ga2O3. A subsequent report for muons implanted into single-crystal specimens of Ga2O3 below room temperature resolved two neutral centers, Mu1 and Mu2, with electron binding energies of 7 and 16 meV [55]. The earlier measurements [54] in powder samples were suggested to be an unresolved average of the Mu1 and Mu2 centers. At reduced temperatures, several metastable states for the muon were observed. At temperatures above 620 K, thermalized muons occupy a diffusively mobile state that shows ground state behavior [55]. An activation energy for Mu+ diffusion of 1.65 eV was inferred from the observed hop rate for this state.

9.3.4 Vibrational properties of H in Ga2O3 Hydrogen-containing defects in Ga2O3 have been investigated and microscopic properties have been determined through studies of their O-H vibrational modes [56]. Properties of the broad IR absorption that can arise from free carriers were investigated to determine whether hydrogen centers give rise to conductivity [57, 58]. Furthermore, the polarization properties of the hydrogen vibrational modes in ZnO and SnO2, for example, have provided valuable information about the structures of the defects that can form [34, 43, 44]. Similarly, the polarization properties of the vibration modes of O-H centers in Ga2O3, when combined with theory, help to identify defect structures [51].

Hydrogen in Ga2O3

197

9.3.4.1 Vibrational spectroscopy Fig. 9.2B shows the sample orientation used in these experiments. H and D were introduced into Ga2O3 by annealing n-type, single-crystal substrates purchased from MTI Corporation in an H2 or D2 ambient at temperatures above 800°C. Fig. 9.3A and B show that O-H and O-D vibrational lines are produced at 3437.0 and 2546.4 cm1, respectively, by the introduction of H and D. These lines are strongly polarized and are seen for the polarization with electric vector E // [102] but not for the polarization with electric vector E // [010]. Because this sample had a (201) face, the polarization with E // [201] could not be thoroughly investigated. These results do not preclude the possibility of absorption in the direction of the incident light (E // [201]), which would not be observed. Experiments are ongoing to pursue this further. The O-H and O-D lines shown in Fig. 9.3, along with a few additional lines, can also be produced by the implantation of protons or deuterons at room temperature [51]. Fig. 9.4 shows results for a Ga2O3 sample purchased from the Tamura Corporation that was implanted at room temperature with deuterons with multiple doses and energies up to 280 keV to produce a deuterated layer, 1200 nm thick, with a deuterium concentration of approximately 1020 cm3. Introducing deuterium into Ga2O3 by the implantation of deuterons produces the 2546 cm1 line that corresponds to the

Fig. 9.3 Polarized IR absorption spectra obtained at 10 K for the hydrogenated, deuterated, and codoped samples from MTI.

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Fig. 9.4 Polarized IR absorption spectra obtained at 10 K for a Ga2O3 sample from the Tamura Corp. that had been implanted at room temperature with deuterons. Effects of annealing at 400°C and 500°C are shown. Upper (blue) spectra were obtained with E // [102], the lower null spectra (red) with E // [010].

3437 cm1 O-H line along with several weaker lines at 2518, 2592, and 2632 cm1. All of these lines are seen for the polarization with E // [102]. Ga2O3 samples can be prepared that contain both H and D to test whether the defects that have been seen contain more than one hydrogen atom. (A defect that contains two H atoms, for example, will have an isotopic sibling that contains both H and D with distinctive vibrational properties that allow it to be identified.) The spectra shown in Fig. 9.3C and D reveal that Ga2O3 samples prepared by annealing in a mixture of both H2 and D2 do, in fact, show additional new vibrational lines at 3438.2 cm1 for O-H and 2547.1 cm1 for O-D [51]. These new lines are a signature of a defect that contains two identical H or D atoms. To explain the spectra shown in Fig. 9.3, the defect that contains two identical H (or D) atoms must have two coupled O-H (or O-D) modes, the first of which is IR active and the second of which is IR inactive. For a defect that contains both H and D, the O-H and O-D modes of the defect become dynamically decoupled and give rise to new vibrational lines for the decoupled oscillators. Furthermore, because the coupled and decoupled modes lie close in frequency, the coupling of the two O-H (or O-D) oscillators must be weak.

9.3.4.2 Evidence for a “hidden hydrogen” species An especially effective way to introduce H into Ga2O3 and to produce the 3437 cm1 center has been found to involve a two-step annealing process. First, a Ga2O3 sample was annealed in an H2 ambient at a temperature of 800°C or greater. Immediately following this annealing treatment, it was found that the 3437 cm1 line could be either weak or absent [as is shown in Fig. 9.5, spectrum (i)]. The second step was an

Hydrogen in Ga2O3

199

Fig. 9.5 IR absorbance spectra for a Ga2O3 sample that had received a two-step annealing treatment to introduce hydrogen. Spectrum (i) was measured for a sample annealed in an H2 ambient for 1 h at 900°C. Spectrum (ii) is for the same sample after a subsequent anneal at 400°C in flowing N2.

annealing treatment at a lower temperature in an inert ambient. Fig. 9.5, spectrum (ii), shows that a second annealing treatment at 400 °C in flowing N2 produced the 3437 cm1 O-H vibrational line. These results show that H is introduced into Ga2O3 by the first anneal at T > 800°C in a form that does not give rise to an observable O-H line. A second anneal at a temperature near 400°C transforms this hidden reservoir of H into the defect that gives rise to the 3437 cm1 center. (Samples that had been deuterated by annealing in a D2 ambient also showed a “hidden D” species that gave rise to a strong 2546 cm1 line in a similar two-step annealing process.) The first anneal in an H2 ambient at high temperature (T  800°C) charges the sample with hydrogen. The second anneal at 400°C produces the 3437 cm1 center. There are several possibilities for the hidden form of hydrogen that is produced by annealing in H2 gas at elevated temperatures. (i) There could be O-H centers with transition moments oriented primarily along the [201] direction. Such a defect would not be visible for the polarizations of the probing light that have been used in the experiments described above and, therefore, would not have been seen. (ii) Hydrogen in Ga2O3 could be introduced in the form of interstitial H2 molecules whose vibrational modes are not IR active. The introduction of hidden H2 was discovered to be important in ZnO where H2 acted as a source and sink for H in defect reaction [31, 32]. (iii) The hidden-hydrogen defect could be the HO shallow donor predicted by theory [17] that has a low vibrational frequency that appears in the spectral region where oxides are highly absorbing. An example of this possibility is provided by hydrogen trapped at an oxygen vacancy in ZnO [22, 38]. Ga2O3 samples hydrogenated (or deuterated) by the ion implantation of protons (or deuterons) also show interesting hydrogen reactions upon annealing at 400°C in an

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Gallium Oxide

inert ambient. Fig. 9.5 shows that annealing a sample at 400°C that had been implanted with D increases the strength of the 2546 cm1 line by roughly 4 times, suggesting the presence of the same hidden species that is formed in Ga2O3 samples annealed in an H2 ambient. Data shown in Fig. 9.6 support the suggestion that an H-related shallow donor can be a source of “hidden” hydrogen in Ga2O3 samples. Spectrum (i) in Fig. 9.6A shows the free-carrier absorption for a Ga2O3 sample that was annealed in an H2 ambient at 1000°C. This annealing treatment does not produce the 3437 cm1 O-H line. Spectrum (ii) shows that the free-carrier absorption is reduced by a subsequent anneal at 400°C in an inert ambient. Furthermore, this second annealing treatment also produces the 3437 cm1 line [spectrum (ii) in Fig. 9.6A and B]. These results suggest that an H-related shallow donor introduced by the anneal in H2 is converted into the 3437 cm1 O-H center by a subsequent anneal at 400°C in an inert ambient. Spectrum (iii) shows that further annealing at 1000°C in an inert ambient removes the 3437 cm1 center from the sample. The spectra labeled (iv) in Fig. 9.6A and B show differences of spectra (i) and (ii) in the same panels. These difference spectra emphasize contributions arising from H rather than from other donors (presumably Sn or Si) that can be present. The broad absorption shown in Fig. 9.6A, spectrum (iv), that increases at low frequency is due to the free-carrier absorption that is due to hydrogen in the sample. The downward going IR line shown in Fig. 9.6A and B, spectrum (iv), is due to the 3437 cm1 center that grows in during the second anneal at 400°C in an inert ambient.

Fig. 9.6 IR absorption spectra (T ¼ 77 K) for a Ga2O3 sample from Tamura Corp. that initially had been hydrogenated by annealing at 1000°C. This sample was subsequently annealed at in Ar at the temperatures indicated. (A) Shows the absorption due to free carriers and (B) shows the 3437 cm1 O-H vibrational line that corresponds to the free-carrier data shown in (A).

Hydrogen in Ga2O3

201

The data in Fig. 9.6 suggest that the Hi or HO shallow donors predicted by theory are good candidates for the reservoir of hidden hydrogen that can be produced by annealing a Ga2O3 sample in an H2 ambient. However, these results do not eliminate other possibilities (like H2) that may also provide a reservoir of hydrogen in our samples that are difficult to observe directly.

9.3.4.3 Theory of defect structures and their vibrational properties There are a variety of defects that can exist in this complex structure [59–62], and hydrogen can interact with many of them. Therefore, there are many possible configurations for interstitial H and for H complexed with native defects such as O or Ga vacancies. The total absence of any O-H vibrational absorption in the [010] direction is remarkable and provides severe constraints regarding the possible nature and structure of the observed defects. (Analysis of the dichroism associated with monoclinic Ga2O3 reveals [63, 64] that its effect will be negligibly small in these experiments.) Furthermore, as noted, the line shifts associated with codoping with H and D further narrow the possible candidates for the observed defects. While qualitative chemical arguments can be used to suggest approximate O-H axes, detailed calculations are needed to provide, with some confidence, the structures and properties of possible defect configurations. Such calculations have been carried out using the CRYSTAL06 code [65], choosing density functional theory (DFT) with a gradient-corrected approximation to the exchange-correlation functional (Becke’s B3LYP hybrid potential [66] with 20% exact exchange and Lee-Yang-Parr correlation [67]). This approach has been used successfully for defect calculations in a number of similar systems [43, 48, 68–73]. Calculations were carried out in fully relaxed periodic supercells containing 80, 120, or 160 host atoms, plus one or more H impurities, with lattice constants computed for the relaxed perfect crystal. A 2 2  2 k-point mesh of Monkhorst-Pack type [74] was used. The SCF convergence criterion was 107 Ha except for vibrational calculations, where 1010 Ha was used. Gaussian basis functions were of the type 311p(1) for H [75], 8411 for O [76], and 864111d(41) for Ga [77]. Charge states for the defect calculations are based on atomic ions; e.g., the charge state for one H at a Ga vacancy is (2). Harmonic and anharmonic vibrational frequencies may be calculated by CRYSTAL06. A number of potential defect configurations were investigated, beginning with interstitial H. Fig. 9.7 shows how Hi could attach to one of two opposing sites on O(1) or O(2), or one of four quasitetrahedral sites on O(3). Several of these have their O-H dipole perpendicular to [010] and therefore bear consideration as candidates for single O-H defects. A second candidate as a host for O-H dipoles is a Ga vacancy, for which there is both experimental [78] and theoretical [30] support. While there is theoretical [30, 51] evidence that Ga(1) vacancies are energetically favored over Ga(2) vacancies, both types are possible candidates for H traps. When atomic relaxation is taken into account, there are three inequivalent Ga(1) vacancy configurations [30, 62] to consider. One is the simple vacancy, with four “dangling bonds.” The other two arise from

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[1 0 2] direction

Incident light perpendicular to (–2 0 1) plane

[0 1 0] direction out of plane

Fig. 9.7 Possible hydrogen interstitial sites (in blue) in an unrelaxed Ga2O3 lattice.

a large shift of a neighboring Ga to a site midway between two resulting halfvacancies. These three configurations are shown and labeled in Fig. 9.8. Varley et al. [30] and Kyrtsos et al. [62] find a shifted configuration to be lower in energy than the unshifted, while Deak et al. [61] find the opposite. Recent results [51] as well favor the stability of the shifted configurations, with the Ga(1)21 configuration more stable than the Ga(1)23 configuration. A search for single-hydrogen sites associated with Ga vacancies such that the O-H axis has no [010] component leads to a number of possibilities. Two inequivalent H locations for the Ga(2) vacancy satisfy this condition [Fig. 9.9A]. There are two inequivalent H sites for the unshifted Ga(1) vacancy that could qualify [Fig. 9.9B] and also two pairs of equivalent Ga(1)21 sites [Fig. 9.10A and B]. Two equivalent H locations for the shifted Ga(1)23 site are candidates (Fig. 9.11).

Shifted Ga(1)23

g

g

Shifted Ga(1)21

Ga(1) vacancy

(A)

(B)

(C)

Fig. 9.8 Possible Ga(1) vacancy sites: (A), unrelaxed; (B), neighboring Ga(1) shifted to site with O(2) and O(3) neighbors [Ga(1)23]; (C) neighboring Ga(1) shifted to site with O(2) and O(1) neighbors [Ga(1)21].

[1 0 2] direction H

H

H H

Incident light perpendicular to (–2 0 1) plane

(A)

[0 1 0] direction out of plane

(B)

Fig. 9.9 (A) Ga(2) vacancy site plus two inequivalent H sites with no (010) O-H projection. (B) Unshifted Ga(1) vacancy site plus two inequivalent H sites with no (010) O-H projection.

[1 0 2 ] direction H H

H

H

Incident light perpendicular to (–2 0 1) plane

(A)

(B)

[0 1 0] direction out of plane

Fig. 9.10 (A, B) Two relaxed configurations of Ga(1)21, each with two equivalent H sites with no (010) O-H projection.

D dipole moment = 0

[1 0 2] direction

Incident light perpendicular to (–2 0 1) plane D dipole moment ¹ 0 [0 1 0] direction out of plane

Fig. 9.11 Relaxed configuration of Ga(1)23 plus two equivalent H sites with no (010) O-H projection. Shown also are the stretch modes for the corresponding two O-H defect, one with zero change in dipole moment and the other with nonzero change in dipole moment.

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The emergence of new lines in the region of the main O-D and O-H lines upon codoping with H and D is, as noted, the signature of a defect that involves two hydrogens. But because these two hydrogens generate only one observed line for H or D, one of the two O-H stretch modes is unobservable. Furthermore, the small shift associated with HD means that the individual O-H components are weakly coupled. In order for a stretch mode involving two O-H partners to have zero dipole moment, the two O-H’s must be symmetry equivalent and either parallel or antiparallel. The configurations shown in Figs. 9.10 and 9.11, which are associated with a shifted Ga(1) vacancy, may be reconsidered, now with both H sites occupied, in which case they do satisfy this condition. (Those shown in Fig. 9.9 are not symmetry equivalent and so are not candidates.) However, the configuration of Fig. 9.10A has only a modest (102) component and that of Fig. 9.10B lacks a (102) projection. Shifted Ga(1)23 plus two H, shown in Fig. 9.11, is the most likely candidate for the 3437 cm1 line. It has a large (102) component and also has the lowest energy (by 0.6 eV) of all configurations investigated, including a number that do not satisfy the experimental polarization constraints. Furthermore, the large distance between the individual O-H components leads to weak coupling between them.

9.3.4.4 Additional IR lines More recent experiments have revealed several additional IR lines that may be due to some of the other defect structures that have been predicted by theory. The focus is on O-D centers because their IR lines are detected with higher signal to noise ratio than for the corresponding O-H centers. The 2547 cm1 vibrational line is the dominant feature seen in a Ga2O3 sample deuterated in a D2 ambient. However, when such a sample was annealed at elevated temperature (T > 900 °C) to remove H or D and then retreated in D2, a few additional IR lines were produced. Fig. 9.12 shows O-D lines at 2584 and 2632 cm1 that were seen

Fig. 9.12 Polarized IR absorption spectra measured at 77 K for a Ga2O3 sample from the Tamura Corp. that had been repeatedly annealed in a D2 ambient.

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(in addition to the 2547 cm1 line) for the polarization with E // [102] for a Ga2O3 sample that had been repeatedly treated in D2 at 900°C. [The 2632 cm1 line was also seen in a sample implanted with deuterons (Fig. 9.4)]. It is not yet certain whether the additional lines that enter upon heat treatment involve more than one H. If not, one or more of the configurations shown in Figs. 9.8–9.11 with only one H, including hydrogen interstitials, may well be responsible for these lines. Work is ongoing by both experiment and theory that promises to identify additional hydrogen defect structures.

9.4

Conclusion

Properties of the hydrogen impurity in Ga2O3 have been discussed and compared with hydrogen’s behavior in other oxides. SIMS measurements show that hydrogen forms thermally stable defects in Ga2O3 [52, 53]. Muon spin resonance finds that implanted muons form shallow donors in Ga2O3 [17], and H is suggested to behave similarly [55]. Theory predicts that Hi and HO centers are shallow donors in Ga2O3 [17] and that the VGa deep acceptor traps H to produce a defect with low formation energy and high thermal stability [30]. Vibrational spectroscopy and its analysis by theory are being applied to Ga2O3 to provide insight into the structures and properties of the defects that hydrogen can form [51]. Interstitial hydrogen is prominent in ZnO, SnO2, and In2O3 single crystals that had been hydrogenated by annealing in an H2 ambient where it behaves as a shallow donor with a strong O-H stretching line with a vibrational frequency >3000 cm1 [34, 42, 43, 48]. On the contrary, no definitive evidence has yet been found for a Hi center in Ga2O3 that had been hydrogenated by annealing in H2 or by ion implantation [51]. The O-H vibrational lines of complexes of hydrogen with native defects have been observed for ZnO [34, 39], SnO2 [42, 43], and In2O3 [48]. The structures and vibrational properties of these defects have been predicted by theory [30]. Similarly, a strong O-H vibrational line at 3437 cm1 is produced in Ga2O3 by hydrogen introduced either by annealing in an H2 ambient or by proton implantation [51]. Theory has investigated the structures and vibrational properties of several possible VGa–Hn complexes in Ga2O3. The observed polarization properties of the 3437 cm1 line and its behavior in samples that contained both H and D suggest the assignment of the 3437 cm1 line to a particular VGa-2H structure (Fig. 9.11) [51]. ZnO [31, 32] and SnO2 [42, 43] have been found to contain “hidden” hydrogen that does not give rise to prominent O-H vibrational lines. Both interstitial H2 and HO are species that do not give rise to O-H absorption but that can provide a reservoir of H in a sample. When Ga2O3 is annealed in an H2 ambient, a hydrogen species not seen by IR absorption is produced. A subsequent anneal in an inert ambient at 400 °C converts this “hidden” species into the VGa-2H center with a strong O-H line at 3437 cm1. Both H2 and HO are candidates for the reservoir of hidden H in Ga2O3 as are hydrogen centers that could have their transition moments along a [201] direction. While progress has been made by experiment and theory toward determining the behavior of hydrogen in Ga2O3, much remains to be done before the properties of H in

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Ga2O3 are understood at a level that is typical of other semiconducting oxides such as ZnO, SnO2, and In2O3.

Acknowlegments The work at Lehigh University was supported by NSF Grant No. DMR 1160756 and the Sigma Xi Grants-in-Aid of Research program. The work at UF is partially supported by HDTRA1-17-1-0011. The project or effort depicted is sponsored by the Department of the Defense, Defense Threat Reduction Agency. The content of the information does not necessarily reflect the position or the policy of the federal government, and no official endorsement should be inferred.

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Ohmic contacts to gallium oxide Marko J. Tadjer U.S. Naval Research Laboratory, Washington, DC, United States

10

Chapter Outline 10.1 Introduction 211 10.2 Ohmic contacts and contact resistance 212 10.3 Ohmic contacts to gallium oxide 214 10.4 Development of Ohmic contacts for Ga2O3 microelectronics 10.5 Research opportunities for Ohmic contacts to Ga2O3 225 Acknowledgment 225 References 225

216

10.1 Introduction The deposition of a metal onto a semiconductor surface to provide low contact resistance, high-reliability electrical contacts without adversely affecting the device during the metallization process is one of the most important challenges in device fabrication. Consequently, a fundamental understanding of how contacts work is essential for successful device manufacturing and commercialization. The physics of carrier transport across the metal-semiconductor junction renders metal contacts either rectifying (a.k.a. Schottky) or nonrectifying. A nonrectifying contact whose relationship between current and voltage has a low interfacial contact resistance Rc, and is preferably linear, is referred to as an Ohmic contact. Achieving low contact resistance Rc (Ω mm) or contact resistivity ρc (Ω cm2) has required a great amount of investigation for every relevant semiconductor material in the past. Typically, the successful formation of an Ohmic contact has relied on three constituent requirements: highly or degenerately doped semiconductor, choice of metallization, and thermal annealing. In the case of silicon, for instance, diffusion processes have been the topic of much early work but ultimately the control and reproducibility of ion implantation have rendered it an industry standard. For compound semiconductor heterostructure devices based on GaAs or GaN, the presence of a two-dimensional electron gas (2DEG) has necessitated a multilayer metallization deposition and annealing scheme, the details of which took many years to optimize. Particularly in the case of III-nitride high electron mobility transistors (HEMTs), Ohmic contacts were relatively easy to make on heteroepitaxal GaN due to its high dislocation density as the barrier height was reduced through defect-assisted formation of metal-nitride alloys during the anneal. Subsequent breakthroughs in GaN crystal growth, however, resulted in several orders Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00010-5 © 2019 Elsevier Inc. All rights reserved.

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of magnitude lower dislocation density homoepitaxial GaN, and naturally the contact resistance obtained under identical process conditions was higher [1]. Regrowth techniques to provide n+-doped GaN have become commonplace as a result. Recent advances in bulk growth of single-crystal β-Ga2O3 substrates have changed how this material is regarded in the device community and have brought it into the spotlight as an emerging wide bandgap semiconductor with potential for radio frequency (rf ) and power electronics applications. Gallium oxide substrates have emerged as a device development platform since the commercial introduction of 2-in. β-Ga2O3 wafers by Tamura Corporation after only a few years of research and development after initial reports of the edge-defined film-fed growth method [2, 3]. Related early breakthroughs in epitaxy and doping control meant that Ga2O3 transistors were on the horizon, and the first report of a Ga2O3 transistor was published by Higashiwaki in 2012 [4–6]. In his report, Higashiwaki used a Ti/Au (denoting Au over Ti henceforth in this chapter) as the Ohmic contact metallization, preceded by a short BCl3 plasma etch in order to introduce defect states on the surface of the β-Ga2O3 film [6]. The academic study of forming Ohmic contacts to Ga2O3 has been a less pressing topic while even more fundamental issues in growth and devices have been garnering most of the attention. Still, it appears that there is a sufficient body of literature to date in order to establish a preliminary understanding of how Ohmic contacts to Ga2O3 work. In this chapter, we will review the fundamentals of Ohmic contact operation, review the existing body of literature on Ohmic contacts to Ga2O3, and will recommend future directions of research.

10.2 Ohmic contacts and contact resistance Thermionic-field emission theory defines a characteristic energy of a metalsemiconductor contact, E00 (eV), as follows [7]: vffiffiffiffiffiffiffiffiffiffiffiffiffiffi u qh u N E00 ¼ u m∗ 4π t ε0 εr m0

(10.1)

where q is the electron charge, h is the Planck’s constant, N is the doping density in the semiconductor, m* is the tunneling effective mass, εr is the permittivity of the semiconductor, and ε0 is the permittivity of free space. E00 is inversely proportional to the probability of an electron tunneling through the metal-semiconductor contact; thus, a lower value for E00 suggests that current due to thermionic emission would dominate the total current, whereas a higher value for E00 means that tunneling will dominate and the contact is more likely to be Ohmic [7]. The specific contact resistivity of a contact can be defined as follows [8]: 

∂J ρC  ∂V

1 (10.2) V¼0

Ohmic contacts to gallium oxide

213

The contact resistance RC is the most important parameter related to Ohmic contacts, and it is customary to report along with the rest of device performance data in the literature. Each research laboratory has a slightly different contact metallization process, and these slight variations in fabrication process mean that merely specifying the process conditions (i.e., metals used, annealing conditions, cleaning procedures, etc.) are not sufficient to judge or reproduce the performance of an Ohmic contact exactly. The measurement of contact resistance by a relatively straightforward technique, in most cases the transfer length method (TLM) with its linear (LTLM) and circular (CTLM) variations, is thus ordinarily reported particularly when three-terminal devices are being reported. The resistance data measured using the LTLM method are simpler to analyze because of the geometry of the test structure, which requires an isolation (mesa) etch to confine the current flow between adjacent pads only. For details on the LTLM and CTLM methods, the reader is referred to the text by Prof. Dieter Schroder [8]. Fig. 10.1 shows the dependence of E00 on the doping density calculated for β-Ga2O3 using published values for the anisotropic electron effective mass in the conduction band [9]. The effective mass values along the b- and c-axes are very similar, resulting in slightly lower values for E00. In all cases, it can be observed that to ensure that in the absence of surface states the contact to the β-Ga2O3 surface would have Ohmic behavior, the doping in the Ga2O3 would have to be >1020 cm3 (E00/kT ≫ 1). Thus, for the case of field emission (high doping density), the contact resistance is related to the metal-semiconductor barrier height and the characteristic energy in Eq. (10.3):  ρC ∝ exp

ΦB E00

 (10.3)

Fig. 10.1 Characteristic energy E00 of β-Ga2O3 as a function of doping density calculated using the anisotropic values for electron effective mass in the conduction band [8, 9].

214

Gallium Oxide

Substitution of Eq. (10.1) into Eq. (10.3) yields the following relationship between the contact resistivity and doping density:  pffiffiffiffiffiffiffiffiffiffi   2 ε S m ∗ ΦB pffiffiffiffi ρC ∝ exp ħ N

(10.4)

This dependence has made growth of highly doped n+ Ga2O3 films very relevant to Ohmic contact technology, and growth by pulsed-laser deposition (PLD) and lowpressure chemical vapor deposition(LPCVD) have been demonstrated [10, 11]. Most recently, Ga2O3 transistors fabricated on n+ Ga2O3 capped epi or molecular beam epitaxy (MBE)-regrown Ohmic contacts have been demonstrated as well [12, 13], and contact resistivity of 50 at a bias insufficient to cause breakdown, interpreted to be due to self-trapped hole formation near the Schottky contact. [8]

(-201)

EFG

3.9  1018

74 (300 K)

Solvent/ piranha/ BOE

Ni

0.9 (I-V) 1.49 (C-V)

1.4

Comment: Ohmic contact - In on top surface; I-V and C-V measurements conducted at room temperature. Thermionic emission dominated carrier transport below a forward bias of 0.75 V. At higher biases observed ideality factor of 1.4 is probably due thermionic-field emission arising from high effective carrier concentration, observed current density - 96.8 A/cm2 at a forward bias of 1.6 V, on-resistance – 7.4 mΩ.cm2. Continued

Table 11.1

Continued

Ref.

Orientation

Growth method

[34]

(-201)

–

Eff. Carr. Conc. (ND–NA)/cm3 UID 3  1017

Mobility cm2/V.s

Surface prep

Contact metal

Barrier height eV

Ideality factor

–

UV/ozone (10 min RT)

Pt TiN (ALD)

1.01 (30°C) 0.98 (30°C)

1.07 (30°C) 1.09 (30°C)

Comment: Ti/Au back contact – no anneal; measurements from 30°C to 210°C. Both Pt and TiN contacts were near ideal. Barrier heights and ideality factors were independent of temperature. At 30°C the Pt contact conducted higher forward and reverse currents than the TiN contact. Above 110°C the TiN contact showed higher forward and reverse currents; at 210°C reverse current for the Pt contact was about three orders of magnitude lower than the TiN contact. Degradation of the TiN contact was suspected at temperatures above 170°C as the ideality factor increased slightly. [9]

(-201)

EFG

3.6  1018

–

–

Ni Pt

1.07 (RT) 1.04 (C-V)

1.3 (RT) 1.28 (RT)

Comment: Back contact entire area Ti/Au; measurements at 25°C to 200°C; ideality factor decreased with temperature, barrier height increased, RON much lower for Ni (6.73 mΩ.cm2) than Pt (382 (mΩ.cm2). Reverse recovery times of 28 ns for the Ni contact and 26 ns for the Pt contact were observed at 25°C and these remained almost constant to 150°C. [10]

(-201)

EFG

5–8  1018

–

Organic/ HCl/ boiling H2O2

Ni

0.9 4 (I-V)

1.3

Comment: Studied several metals: W, Cu, Ni, Ir and Pt; weak correlation between Schottky barrier height and metal work function.

[13]

(-201)

PLD Si-doped hetero-epi on c-plane sapphire

1.6  1018

–

–

Cu (DC sputtered)

0.92 (I-V RT) 1.32 (I-V 500 K)

1.22 (RT) 1.03 (500 K)

Comment: Studied temperature dependence and inhomogeneity of barrier height. These contacts were stable to 550 K. Considering a Gaussian distribution of barrier heights, a mean barrier height of 1.32 eV was determined. This is consistent with barrier height of 1.32 eV determined at 550 K where the carrier transport is not influenced by inhomogeneity of the barrier and the contact behavior is close to ideal. [14]

(-201)

PLD Si-doped hetero-epi on c-plane sapphire

5  1017–1  1018

–

–

Pt (LT-RS, DC)

1.42 (I-V)

1.28

Comment: PtOx capped with Pt by long-throw (LT) DC magnetron sputtering. Eff. carr. conc., effective carrier concentration (ND – NA); FZ, floating zone growth method; BOE, buffered oxide etchant; CZ, Czochralski growth method; UID, unintentionally doped; as-fab, as fabricated; DLTS, deep-level transient spectroscopy; DLOS, deep level optical spectroscopy; UHV, ultra-high vacuum; PES, photoemission spectroscopy; HVPE, halide vapor phase epitaxy; CMP, chemical-mechanical polishing; RT, room temperature; RTA, rapid thermal annealing; EFG, edge-defined film-fed growth method; I-V-T, current-voltage as a function of temperature (Richardson plot); C-V-T, capacitance-voltage as a function of temperature; IPE, internal photoemission; PLD, pulsed laser deposition; ALD, atomic layer deposition.

244

Gallium Oxide

Dry etching of β-Ga2O3 can be conveniently achieved by reactive-ion etching (RIE) and inductively coupled plasma (ICP) etching using BCl3/Ar [27–29] employed for achieving ohmic contacts to Ga2O3 [30]. Removal of material by dry etching is an important process step for mesa formation and fabrication of device structures [27, 28]. However, removal of damage introduced by dry etching is essential when the etched structure forms an integral part of an active device. Yang et al. [31] used
Ni SBDs on 201 Ga2O3 to study the efficacy of thermal annealing for damage removal. They found that the barrier height, ideality factor, and reverse breakdown voltage could be restored to the level of unetched material following a 450°C anneal of the etched material when Ni was deposited following the anneal treatment. However, the reverse-bias characteristics degraded severely for Ni contacts deposited on the etched surface prior to the 450°C anneal treatment. Ohmic contacts are formed either on a selected area on the front surface or on the entire back surface of the wafer. A number of approaches and metallizations have been used [32]. A commonly employed procedure, first reported by Higashiwaki et al. [30] involves RIE with BCl3 followed by deposition of a Ti-Au film. Alternative procedures involved (i) painting the back surface with a film of Ga-In eutectic [16, 17], (ii) depositing a sputtered film of Ti on the entire back surface [23], (iii) e-beam evaporated Sn on the back surface and rapid thermal annealing (RTA) performed at 500°C [12], and (v) Ti-Au film deposition followed by anneal at 400°C or 450°C [10, 33]. In some device applications a multienergy Si-ion implantation and subsequent anneal treatment followed by deposition of Ti/Au film has been used to form low resistivity ohmic contacts to Ga2O3 [4]. Another approach that yielded low resistivity ohmic contacts was by Ti/Au metallization on a highly doped surface region produced by diffusing Sn from a spin-on source [34]. Lovejoy et al. [35] studied the surface point defects and band bending on the β-Ga2O3 (100). XPS show the bands bend upward by 0.5 eV at the surface after cleaving in air and annealing. Annealing leads to accumulation of negatively charge defects at surface, causing upward band bending. The defects are either Ga vacancies, interstitials, or adsorbed oxygen. This upward band bending should make it more difficult to form ohmic contacts during device fabrication. It is apparent from an examination of the entries in Table 11.1 that a wide range of processing conditions, crystal orientations, and crystal growth methods have been used to form Schottky contacts on β-Ga2O3, which makes comparisons difficult. However, there are some overarching trends that are noted. One observation is that most of the reported SBHs on β-Ga2O3 are between 1.0 and 1.5 eV. While the reported ideality factors typically varied between 1.01 and 1.30, it is interesting to note that barrier heights on (010)
material—typically 1.5 eV—tended to be higher than values reported on 201 , although there are fewer Schottky contact studies reported on the (010) orientation. In addition to crystal orientation, the underlying carrier concentration can have an effect on the ideality factor, with values near unity favored on bulk crystals and epitaxial films with low carrier concentrations: 1  1017/cm2 or lower. Higher ideality factors 1.3–1.4 were observed for Ga2O3 with carrier concentrations >1  1018/cm3 [8–10]. In materials with carrier concentrations >1  1018/cm3 carrier

Schottky contacts to β-Ga2O3

245

transport may be determined by thermionic-field emission, consequently yielding higher ideality factors [36]. However, it appears that e-beam evaporation yields higher quality (near unity ideality factor) Schottky diodes in comparison to sputter deposited contacts; sputter-induced surface damage may have a degrading effect on Schottky contacts. However, Muller et al. [14] using long-throw DC magnetron sputter deposition (distance between target and substrate 25 cm) obtained PtOx contacts with an ideality factor of 1.09. Annealing of the ohmic contact at a temperature 400–500°C also tends to yield higher ideality factors. A study of SBHs of metal contacts on β-Ga2O3 as a function of metal work function was reported Farzana et al. [21]. They studied contacts established with Au, Ni, Pd, and Pt on (010) oriented substrates from an UID EFG grown crystal with an effective carrier concentration of 1.1  1017/cm3. The contacts showed excellent rectifying behavior, as shown in Fig. 11.3, with ideality factors ranging from 1.03 to 1.09. Barrier heights determined from room temperature Is, using A* ¼ 41.1 A cm2 K2, were 1.62, 1.50, 1.29, and 1.53 eV for Au, Ni, Pd, and Pt, respectively. Barrier heights were also determined from C-V, I-V-T (Richardson plot, T ¼ 180–400 K), and IPE. The calculated barrier heights from all of the measurement techniques for each metal were in good agreement with each other. Near unity ideality factor and excellent linearity of modified Richardson plots, ln(Is/AT2) as a function (1/nkT), indicated that thermionic emission was the dominant transport process in Ni, Pd, and Pt contacts on the (010)-oriented β-Ga2O3 substrates. From these observations, Farzana et al. inferred that Fermi-level pinning or the presence of a dielectric layer at the interface did not dominate the Schottky barrier behavior formed with those three metals. In contrast, Au contacts behaved somewhat differently. Barrier height values obtained from the saturation current, C-V and IPE for Au were higher than those obtained for Pt, although Pt has a higher work function. The barrier height extracted for Au contacts from the modified Richardson plot was

Current density (A/cm2)

100

10–2

Pd Ni Pt Au

Fig. 11.3 J—V curves for Pd, Ni, Pt, and Au SBDs on (010) β-Ga2O3 single-crystal substrate with carrier concentration Nd ¼ 1.1  1017 cm3. Ideality factors (n) were between 1.03 and 1.09. Reproduced from Farzana, E., Zhang, Z., Paul, P.K., Arehart, A.R., Ringel, S.A., Appl. Phys. Lett. 110 (2017) 202102, with the permission of AIP Publishing.

n~1.09 n~1.03

10–4

n~1.04 n~1.05

10–6

10–8 –3

–2

–1 0 Voltage (V)

1

246

Gallium Oxide

significantly lower than that obtained from I-V, C-V, and IPE, and the Richardson’s constant extracted from the plot was much lower than the theoretical value of 41.41 A cm2 K2. It was therefore concluded that thermionic emission was not the dominant carrier transport process for the Au/Ga2O3 SBD. This complexity might be due to spatial barrier inhomogeneity.

11.5

Defects relevant to β-Ga2O3 Schottky contacts

11.5.1 Point defects and impurities Defects and impurities, such as vacancies and interstitials, unintended impurities, and even a high concentration of intentionally added impurities, during crystal growth, play important roles in gallium oxide Schottky diodes. For example, impurities may act as dopants and contribute to n-type and p-type conductivities, but p-type Ga2O3 has very low hole mobility, so it has very little practical use [37]. Hydrogen is a common impurity, which can occupy either interstitial (Hi) or substitutional sites (HO) in Ga2O3, due to its small size, acting as a shallow donor in both configurations. The low formation energy of Hi allows it to be easily incorporated as an unintentional impurity, but HO only has low formation energy under O-poor conditions; in contrast, the dissociation energy of H+O into H+i and V0O is low, indicating that both H+i and H+O can be removed by thermal annealing [37]. A number of impurities have been shown to produce n-type behavior in Ga2O3. Theoretical calculations predict n-type doping from Si, Ge, and Sn on Ga sites, as well as F and Cl on O sites [37]. Both Sn and Si have been commonly used as n-type dopants in both bulk crystals and epitaxial layers; whereas Ge is being investigated as an n-type dopant in MBE-grown epitaxial layers [38]. Silicon has also been found [16] to be an impurity background dopant in bulk crystals that originates from the source powder. Theoretical calculations of specific defect complexes of N-doped and Al-doped Ga2O3 using a pseudopotential plane-wave approach based on density functional theory (DFT), indicate that N in particular acts as a deep acceptor, thus is not suitable as an effective p-type dopant [39]. According to the calculated band structure, the valence band is nearly flat, indicating a large hole effective mass, which leads to low hole mobility that hinders the fabrication of p-type β-Ga2O3. Besides impurities, native point defects, such as vacancies and interstitials, can also behave as shallow or deep donors, create defect states, and affect formation of Schottky contacts. Because the O stoichiometry at the surface of Ga2O3 can change, depending on processing and annealing conditions, and therefore affect Schottky contact behavior, it is important to consider the electrical properties of O point defects. Oxygen vacancies were discussed in detail by Hajnal et al. [40]. Charged ionic O vacancies are formed by the following reaction: 1 Ga2 O3 ¼> Ga2 O3 : VO+ + + O2 + 2e : 2

Schottky contacts to β-Ga2O3

247

Calculations indicate the presence of three inequivalent VO sites [40]. Since the transition levels are more than 1 eV below the conduction band minimum, VO acts as a deep donor and should not contribute to n-type conductivity.

11.5.2 Extended crystallographic defects As of the point of this writing, there has been very little investigation into the effects on specific defects on the behavior of Ga2O3 Schottky diodes. In one study [41] Ni
SBDs were fabricated on the 010 surface of a β-Ga2O3 substrate grown by the EFG growth method. Using differential interference microscopy, etch pits attributed to either dislocations or voids were observed. The average densities for dislocation etch pits and void etch pits were 1.1  104 cm2 and 6  103 cm2, respectively. The authors measured the I-V characteristics of a large number of diodes and found that many of the diodes that had high densities of dislocations or voids also had high leakage currents (at 15 V). However, there was also a significant fraction of diodes that had both high leakage currents and low densities of dislocations and voids. Therefore, additional study is needed to make conclusions about how these specific defects affect
010 β-Ga2O3 Schottky diode behavior. In (001) substrates, etched grooves and etch-pits indicative of line-shaped voids and small defects were observed [42]. SBD characteristics (leakage currents, ideality factors, and barrier heights) did not show strong correlations with the densities of these defects. However, there was an apparent effect from mechanical damage resulting from the substrate preparation process. Current-voltage characteristics obtained from
SBDs on the (001)-oriented surface appear qualitatively similar to those on the 010 surface.
Three different types of etch pits were observed on 201 β-Ga2O3 substrates: (1) line-shaped pattern with average density of 5  102/cm2, (2) arrow-shaped etch pits, average density of 7  104/cm2, and (3) gourd-shaped etch pits, average density of 9  104/cm2 [43]. Although SBDs with high leakage currents were observed, no relationship between the incidence of these defects and high leakage was found. The general nature of the current-voltage characteristics was similar to those of diodes on the

010 -oriented surface. Dislocations in the 201 material run mainly along the [010]
direction and do not go through the 201 surface, and therefore do not act as path for
leakage currents. Although a higher density of etch pits was observed for the 201 oriented substrates, the number of SBDs with high leakage current was lower than the (010)-oriented substrates. Kasu et al. suggested that the forward leakage current may be associated with a low-energy barrier. A similar forward leakage was also observed
by Yao et al. [10] in their study of Ni SBDs on 201 -oriented substrates. They considered this effect to be related to surface or subsurface defects and modeled it as two diodes in parallel, one with a high barrier and the other with low barrier. In one study of crystal defects in Ga2O3 [44], surface defects were identified and classified from different types of etch pits on a β-Ga2O3 (010) single crystal, grown by EFG method. The surface was etched in H3PO4, and subsequently observed using

248

Gallium Oxide

SEM and AFM. The etch pits were classified into A-F types based on their shapes (Fig. 11.4), whereas groove-shaped pits on an unetched surface were classified as type G (Fig. 11.5). One method to study the effects of defects on SBD performance is to fabricate the device and measure the electrical properties, then remove the electrodes by etching, and observe etch pit densities in the semiconductor surface to determine whether correlations exist between electrical performance of the device and densities of certain defects. Using this method [41] void-related type-A etch pits and dislocation-related
type-F etch pits were observed on 010 β-Ga2O3 from SEM images. These authors
conducted similar characterization and analysis on 201 β-Ga2O3 [43]. According to
their previous study on 010 , dislocations are considered as a leakage path only along the [010] direction. Three types of etch pits were identified: (1) line-shaped etch pattern, (2) arrow-shaped etch pit, and (3) gourd-shaped etch pit. The origins of lineshaped etch patterns are reportedly void defects extending in the [010] direction, which were reported as well in previous papers [43, 45]. The origins of arrow-shaped and gourd-shaped etch pits were attributed to dislocations. Although etch pit densities
were higher on the 201 plane than the (010) plane [43], the number of SBDs on
201 with high leakage currents is less than that on the (010) plane. In other words,

Fig. 11.4 Different types of surface etch pits observed by Hanada et al. [44]. Republished with permission from Hanada, K., Moribayashi, T., Koshi, K., Sasaki, K., Kuramata, A., Ueda, O., Kasu, M., Jpn. J. Appl. Phys. 55, (2016) 1202BG.

Schottky contacts to β-Ga2O3

249

Fig. 11.5 Groove-type-shaped pits on unetched surface observed by Hanada et al. [44]. Republished with permission from Hanada, K., Moribayashi, T., Koshi, K., Sasaki, K., Kuramata, A., Ueda, O., Kasu, M., Jpn. J. Appl. Phys. 55 (2016) 1202BG.


no relationship between defect density and leakage current on the 201 plane was found. A double-barrier behavior can be identified from the I-V curve, which might arise from intermediate states introduced by defects.

11.5.3 Defect states in the band gap Ga2O3 SBDs were employed to investigate the concentrations of defect states and their locations within the bandgap, using DLOS (deep-level optical spectroscopy) and DLTS (deep-level transient spectroscopy) [20]. The measurements were performed on Ni/β-Ga2O3 SBDs on an (010) UID substrate, grown by the EFG method. From DLTS three deep-level states were identified: EC—0.62, 0.82, and 1.00 eV, with trap concentrations of 4.7  1014 cm3, 3.6  1016 cm3, and 3.7  1015 cm3, respectively. DLOS revealed two additional trap states: EC—2.16, and 4.40 eV, with trap concentrations of 1.0  1015 cm3 and 1.5  1016 cm3, respectively. The trap levels

250

Gallium Oxide

Ec—0.62, 0.82, and 1.00 eV are similar to traps found by Irmscher et al. [16] in β-Ga2O3 grown by CZ method, suggesting that they may share common physical sources. According to Zhang et al., EC—0.82 eV and EC—4.40 eV had the highest trap concentrations among all of the deep states. Calculation predicted that either SnGa or VO point defects may form deep traps near EC—0.82 eV, and VGa-related defects may form deep traps near EC—4.40 eV. Similar defect states were also found in unintentionally doped, single-crystal β-Ga2O3 grown by CZ method [16]. A Ni Schottky contact with barrier height of 1.1 eV was fabricated on the (100) surface of crystals with carrier concentrations in the range from 6  1016 cm3 to 8  1017 cm3. Three deep traps were detected at EC—0.55, 0.74, and 1.04 eV by DLTS. The deep trap at EC—0.74 eV has accepter-like characteristics. Fe and Co may be related to the 0.55 and 1.04 eV deep traps. Further research is needed to identify the physical sources of these states. An unintentional donor with energy 110 meV below the conduction band in UID Ga2O3 grown by EFG method was found to limit the breakdown voltage in Ga2O3 SBDs [46]. The on-resistance increased and the breakdown voltage decreased in diodes that contained this impurity level, compared with diodes that only contained shallow donors [46].

11.6

Nonideal and inhomogeneous Schottky barriers


Yao et al. [10] observed that I-V and C-V determined SBH values on 201 β-Ga2O3 showed weak correlation with the metal work functions, suggesting the presence of some degree of Fermi-level pinning (Fig. 11.6). It is suspected that a high density of

Schottky barrier height (eV)

3

– (201) Bulk lateral I-V – (201) Bulk vertical I-V (010) Epi lateral I-V – (201) Bulk lateral C-V – (201) Literature I-V – (201) Literature C-V (010) Literature I-V (010) Literature C-V Schottky-Mott

2.5

W 2

1.5

Ir

Pt Ni

Cu I-V

I-V 29

4.8

5.0

C-V 27 (1.18 eV)

C-V 23 C-V 28

C-V 26 (1.12 eV) C-V 25 (1.05 eV) I-V 26

1 4.6

27

I-V 23 C-V 22 I-V 22

I-V 25 (1.08 eV) I-V 28 5.2

5.4

5.6

Metal work function (eV)

Fig. 11.6 Schottky barrier heights on β-Ga2O3 vs. metal work function, from Yao et al. [10]. Republished with permission: From Yao, Y., Gangireddy, R., Kim, J., Das, K.K., Davis, R.F., Porter, L.M., J. Vac. Sci. Technol. B 35 (2016) 03d113.

Schottky contacts to β-Ga2O3

251


electrically active defects may account for this behavior; for example, 201 -oriented Ga2O3 is reported to contain twin defects [47]. Additionally, the high doping levels in the substrate could account for relatively high ideality factors (typically 1.3–1.6) on the
201 diodes in that study. One observation that is worth further investigation is that the SBH’s on (010) β-Ga2O3 was higher (by 0.08–0.3 eV) than their respective values on
201 . A couple of studies [20, 21] suggest that Fermi-level pinning is reduced on (010) Ga2O3. Farzana et al. [21] observed a correlation between the SBH and the work function of three metals (Pd, Ni, and Pt) on (010)-oriented β-Ga2O3 (moderately doped 1  1017 cm3). Interestingly, Au did not fit the Schottky-Mott trend. The contrary behavior of Au SBDs may be associated with SBH barrier inhomogeneity, which was evidenced by larger differences in the I-V and C-V measurements of the Au contacts. SBH inhomogeneity is a well-documented phenomenon that has been explored by other authors on different semiconductors [48–50].
Splith et al. studied Cu contacts established on heteroepitaxial 201 -oriented β-Ga2O3 films grown on sapphire substrates. These films were grown by PLD and doped during growth with Si with an effective carrier concentration of 8  1017/cm3. From room temperature I-V measurements, they determined an effective barrier height of 0.88–0.92 eV and ideality factor of 1.2–1.4, indicating that these Schottky diodes were nonideal. In this study diode characteristics were also obtained from 50 to 320 K. Subsequently using an analytical approach reported by Werner and G€ uttler, the difference between barrier heights obtained from C-V and I-V measurements was plotted as a function of 1/T [50]. The standard deviation was 126 meV, and the mean value of the barrier height was 1.32 eV. Interestingly from I-V measurements conducted at 550 K, they obtained the homogeneous barrier height as 1.32 eV (that is insensitive to the standard deviation) and an ideality factor of 1.03, indicating a near ideal behavior at that temperature. The mean barrier height from the Gaussian distribution and homogeneous barrier height at 550 K appear to be in good agreement. Werner and Guttler’s analysis also established that the barrier height from C-V measurements yields the mean value of the barrier as obtained from the Gaussian distribution.
Jayawardena et al. [23] studied Ni Schottky diodes on 201 -oriented Si-doped (2.8  1017 cm3) EFG grown substrates. At room temperature they obtained a barrier height and an ideality factor of 1.08 eV and 1.19, respectively. Using the same approach employed by Splith et al., they determined a mean barrier height of 1.26 eV and a standard deviation of 121 meV from measurements at temperatures in the 85–300 K range.
SBH inhomogeneity has also been reported on 201 -oriented β-Ga2O3 by Yao et al. [10]. They modeled the forward-bias I-V curves as two Schottky barriers in parallel as shown in Fig. 11.7. Calculations yielded SBH values of 1.00 and 0.80 eV for the high and low barriers, respectively. The second, lower barrier was attributed to electrically active defects in the bulk or near-surface regions. Cleaning/etching of the semiconductor surface using wet chemicals is typically used as a means to create a pristine surface and to unpin the Fermi-level of SBDs. Ohira et al. and Oshima et al. examined various acids and bases such as H3PO4, H2SO4, HCl, HNO3, HF, H2O2, KOH, and NaOH. They found that the two chemicals

252

100

Current (A)

Fig. 11.7 Log I vs. V plot of Ni/Ga2O3 showing two distinct linear regions and modeled as two Schottky barriers in parallel, observed by Yao et al. [10]. Republished with permission from Yao, Y., Gangireddy, R., Kim, J., Das, K.K., Davis, R.F., Porter, L.M., J. Vac. Sci. Technol. B 35 (2016) 03d113.

Gallium Oxide

10–5

10–10

10–15

Measurement Low barrier fit High barrier fit Total fit 0

0.5

1 Voltage (V)

1.5

2

that continuously etched the substrates were NaOH and H3PO4. Similarly, the etch rate and resulting rate constant in general are functions of temperature, which yields a tunable etch rate that can be tailored to different applications. Similarly, Yao et al.
[10] analyzed the effects that various wet chemical treatments have on Ni= 201 β-Ga2O3 SBDs. They used BOE(HF), BOE + H2O2, H2O2, HCl, HCl + H2O2, and an organic solvent clean before metal deposition to examine the effects on the SBD transport properties. It was found that the HCl + H2O2 treatment yielded the highest SBHs and lowest series resistances. Dry etching methods are also used in semiconductor processing and can be used to clean the semiconductor surface via a plasma. Hogan et al. and Shah et al. [27, 28] examined etching gases on β-Ga2O3 and found BCl3 to have the highest etch rate and thus most selectivity. The other gases examined were O2, Cl2, SF6, O2, CF4,
CHF3, and Ar. The 201 and (010) planes were found to have the same etch rate as a function of etching bias power; the etch rate for all orientations increased monotonically with both RF and ICP bias powers [27, 28]. Similarly, the roughness was relatively unchanged for BCl3/Ar, provided the BCl3 flow rate was greater than the Ar flow rate [29]. This reflects a balance between the physical and chemical etching components. However, when the Ar flow rate rose above the BCl3 flow rate, the roughness increased in part due to the perturbation of the equilibrium between the chemical and physical effects [20]. SBDs that were fabricated post-BCl3 etch showed an increase in ideality factor and a decrease in SBH with RIE bias power [51]. Although no chemical residue from the etching process was detected on the Ga2O3 surface using Auger electron spectroscopy (AES), they believe there may be effects from chemicals that are present below the detection limit. In another study [31] postdeposition annealing at 450°C reversed all effects that BCl3 etching had on the ideality factor and SBH of Ni/Au SBDs on
201 . These results suggest that at least some of the etching damage is reversible by annealing.

Schottky contacts to β-Ga2O3

11.7

253

β-Ga2O3 Schottky devices

11.7.1 SBDs as rectifiers Work on β-Ga2O3 device technology is at a very early stage of development. Early reports on breakdown voltage, on-resistance, forward current density, and reverse leakage currents using device structures without properly designed edge-termination or surface passivation point to the potential of Ga2O3 SBDs as high-voltage high-current rectifiers. It has been demonstrated by Konishi et al. [26] that breakdown voltage and reverse current of a Ga2O3 SBD can be improved substantially with the use of a field plate structure. Published studies that report breakdown voltage and on-resistance are summarized below. Sasaki et al. reported an n/n+ structure formed with a 1.4 μm thick Sn-doped homoepitaxial film. The epilayer (carrier concentration of 4  1016/cm3) was grown by ozone MBE on a highly doped (010)-oriented substrate (carrier concentration ¼ 1  1019/cm3) [33]. They used Pt as the Schottky contact and an ohmic contact was formed on the back surface with e-beam evaporated Ti/Au annealed at 450°C in N2 ambient. The Schottky diodes (diameter ¼ 200 μm) had a breakdown voltage of 100 V, an on-resistance of 2 mΩ cm2, and a forward current density of 200 A/cm2 at a bias of 1.7 V. Sasaki et al. [4] subsequently reported SBDs fabricated using Pt contacts on an UID (010)-oriented substrate with effective carrier concentration of 3–5  1016/cm2. A back contact was prepared by depositing Ti/Au metallization after etching using BCl3 RIE. These diodes had an ideality factor 1.05 and a barrier height of 1.52eV. A breakdown voltage of 150 V was observed for an effective carrier concentration of 3  1016/cm3 and on-resistance of 4.3 mΩ cm2; the breakdown voltage for an effective carrier concentration of 5  1016/cm3 was 115 V and the on-resistance was 7.85mΩ cm2, indicating that an n/n+ structure provides a lower on-resistance. Oh et al. [12] studied Ni contacts fabricated on an n/n + structure, where the UID 2-μm-thick epitaxial film was grown on a Sn-doped (010) substrate with an effective carrier concentration of  4.1  1018/cm3. The back contact was formed by evaporating a Sn film followed by RTA at 500°C for 30 s in N2 ambient. They reported a breakdown voltage of 210 V at room temperature. The Schottky contact properties included a high ideality factor (n ¼ 3.38) and a barrier height of 0.95 eV at room temperature. However, significant changes were observed at elevated temperature; the respective values were 1.21 and 1.01 eV at 225°C. The on-resistance values were 2582 Ω cm2 and 0.043 Ω cm2 at room temperature and 225°C, respectively.
In another study SBDs were fabricated with Ni and Pt contacts on 201 -oriented substrates having an effective carrier concentration of 3.6  1018/cm2 [9]. The back contact was formed by e-beam evaporated Ti/Au film over the entire area. For the Ni contact a barrier height of 1.07 eV and an ideality factor of 1.3 were reported; for the Pt contact the values were 1.04 and 1.28 eV, respectively. Interestingly, the on-resistance for Pt (382 mΩ cm2) was higher than that for Ni (6.73 mΩ cm2), and the room-temperature breakdown voltage for Pt (85 V) was lower than that for the Ni contact (154 V). Similar reverse recovery times of 28 and 26 ns were observed for the Ni and Pt diodes, respectively.

254

Gallium Oxide

A relatively high-breakdown (>800 V) vertical SBD without edge termination using a rectifying Ni/Au contact was reported [19]. The device structure is shown schematically in Fig. 11.8. The authors used an HVPE-grown, Si-doped (2  1016/cm3) epitaxial film grown on a Sn-doped β-Ga2O3 (001)-oriented substrate with an effective carrier concentration of 3.6  1018/cm3. As-grown 20-μm-thick epitaxial films were reduced to 10 μm using CMP to remove growth pits. Diode fabrication involved forming an ohmic contact by e-beam evaporation of Ti/Au (20/80 nm) on the entire back surface of the wafer. Using photolithography, e-beam deposition of Ni/Au (20/80 nm) and subsequent lift-off, 105- and 210-μm diameter SBDs were defined on the front surface. These diodes had breakdown voltages at room temperature in the 810–1016 V range. Average breakdown voltage for the 105-μm diameter diodes was 975  40 V, whereas the 210 μm diodes had an average breakdown voltage of 810  3 V. The forward current was determined by thermionic emission whereas the reverse current was dominated by thermionic-field emission, increasing by an order of magnitude for every 25°C rise in temperature. The on-resistance was 6 mΩ cm2, corresponding to a figure of merit, (V2BR/Ron), of 154 MW cm2. A 1-kV vertical Ga2O3 field-plated SBD with Pt as the Schottky contact, shown schematically in Fig. 11.9, was fabricated on HVPE epitaxial films grown on (001)-oriented EFG grown substrates [26]. The diode design was based on a twodimensional device simulation that specified an epitaxial film with effective carrier concentration of 1.0  1016/cm3; however, the actual effective carrier concentration was 1.8  1016/cm3. The fabricated device was characterized over a range of temperature from 22°C to 200°C, yielding an ideality factor of 1.03 over the entire temperature range. Linearity of the Richardson plot indicated transport by thermionic emission over the barrier. A barrier height of 1.46 eV was calculated for this fieldplated Pt/(001)Ga2O3 diode structure. In comparison, it is noted that for Pt contacts on (001)-oriented epitaxial films, a barrier height of 1.15 eV has generally been observed, as reported by Higashiwaki et al. [30]. The larger barrier height in the field-plated device shown in Fig. 11.9 is considered to be an effect of F ion incorporation, with an effective sheet concentration of 2  1013/cm2, at the metal-semiconductor interface due to processing with BOE. A result of the larger barrier height was a reduced tunneling component for Au (80 nm) Ni (20 nm) Si-doped epitaxial (~10 µm) Sn-doped bulk Ga2O3 (650 µm) Ti/Au (20 nm/80 nm)

Fig. 11.8 Schematic representation of vertical Schottky barrier diode on (001) β-Ga2O3, which had breakdown voltages ranging from 800 to 1000 V [19]. Reproduced from Yang, J., Ahn, S., Ren, F., Pearton, S.J., Jang, S., Kim, J., Kuramata, A., Appl. Phys. Lett. 110 (2017) 192101, with the permission of AIP Publishing.

Schottky contacts to β-Ga2O3

255

Fig. 11.9 1 kV vertical Ga2O3 field-plated SBD device schematic [26]. Reproduced from Konishi, K., Goto, K., Murakami, H., Kumagai, Y., Kuramata, A., Yamakoshi, S., Higashiwaki, M., Appl. Phys. Lett. 110 (2017) 103506, with the permission of AIP Publishing.

reverse bias; analysis of I-V-T plots (current-voltage measured as a function of temperature) indicated that the reverse current was dominated by thermionic emission. This observation is in contrast with diodes without field-plate edge-termination, where the reverse current is dominated by thermionic-field emission due to tunneling across the image-force-lowered barrier [11, 19, 23]. A reverse breakdown voltage of 1076 V was achieved for these diodes. Other room-temperature device performance characteristics included a forward current density of 300 A/cm2 at a forward bias of 3 V, a reverse current density of 1  104 A/cm2 at a reverse bias of 1000 V, and a specific on-resistance of 5.1 mΩ cm2 (including a substrate component of 2.0 mΩ cm2). Oda et al. [6] reported SBDs, fabricated on n/n+ freestanding α-Ga2O3 films separated from sapphire substrates, using a process shown schematically in Fig. 11.10. They were able to control the electron concentration in the range between 1  1017 Electrode (Pt/Ti/Au)

n- Ga2O3 +

n- Ga2O3 +

n Ga2O3

n Ga2O3

Sapphire

Sapphire

-

n- Ga2O3

+

n+ Ga2O3

n Ga2O3 n Ga2O3

Electrode (Ti/Au)

Fig. 11.10 Schematic representation of the fabrication process for α-Ga2O3 Schottky barrier diodes on a freestanding film. From Oda, M., et al., Schottky barrier diodes of corundum-structured gallium oxide showing on-resistance of 0.1 mΩcm2 grown by MIST EPITAXY®. Appl. Phys. Express, 9(2) (2016), p. 021101. “Copyright 2016 The Japan Society of Applied Physics.”

256

Gallium Oxide

Properties of Schottky diodes fabricated on a heteroepitaxial α-Ga2O3 film, grown on a sapphire substrate, and separated as a freestanding film

Table 11.2

n/n+ nm/μm

Mobility cm2/V s

Turnon V

RON mΩ cm2

Breakdown V

Jr A/cm2

430/3-4

300

1.5–1.6

0.1

530

2

2850/3-4

–

–

0.4

880

0.5

Contact metal Pt (Pt/Ti/ Au)

and 3  1019/cm3 by doping using Sn halides introduced with the Ga source during MIST EPITAXY® CVD growth of the films. They used Pt/Ti/Au as the Schottky contact and Ti/Au as the back contact. The small thickness of the structure, 4–6 μm, yielded a very low on-resistance. Relatively high breakdown voltages were also reported (see Table 11.2). High reverse-bias leakage currents were attributed to the use of an unpassivated test structure. Nevertheless, the reported results indicate the potential of α-Ga2O3 as a potential material for high-power devices. A β-Ga2O3 trench MOS-type Schottky barrier diode (MOSSBD), as shown in Fig. 11.11A, was demonstrated by Sasaki et al. [52]. The device had an n/n+ structure with a HVPE grown 7-μm-thick epitaxial film with an effective doping concentration of 6  1016/cm3 on a 350-μm-thick (001)-oriented substrate with an effective doping concentration of 2.5  1018/cm3. A periodic structure comprised of 1.2 μm wide mesas and 4.8-μm wide  2.5-μm deep trenches were formed by ICP RIE and photolithography. A 50-nm thick HfO2 film was deposited over the entire wafer by ALD. The HfO2 film from the top of the mesas was removed by CMP, followed by

Fig. 11.11 (A) Device schematic and (B) Reverse characteristics of β-Ga2O3 MOS Schottky barrier diode. © [2017] IEEE. Reprinted, with permission, from Sasaki, K., et al. Depletion-mode vertical Ga2O3 trench MOSFETs fabricated using Ga2O3 homoepitaxial films grown by halide vapor phase epitaxy. Appl. Phys. Express 10(12) (2017) 783.

Schottky contacts to β-Ga2O3

257

deposition of the contact metallization Cu/Au/Ni (200 nm/3000 nm/50 nm). A barrier height of 1.07 eV for the MOS Schottky diode was determined from the saturation current. A SBD fabricated on the same wafer had a barrier height of 0.98 eV. Ideality factors of 1.1 and 1.05 were observed for the MOS Schottky diode and the Schottky diode, respectively. On-resistances of 2.9 mΩ cm2 for the MOS Schottky diode and 2.3 mΩ cm2 for the Schottky diode were obtained. The observed higher barrier for the MOS Schottky diode was considered to be due to the difference in the potential barrier in the mesa generated by the MOS structure. The device showed a substantially lower reverse leakage current in comparison to the standard Schottky diode, shown in Fig. 11.11B. A breakdown of 240 V occurred at the edge of the contacts, indicating that a higher breakdown should be obtainable by the introduction of a field-plated structure.

11.7.2 Metal-semiconductor field-effect transistors Higashiwaki et al. [30] demonstrated the first Ga2O3 three terminal device: a MESFET, which employs a Schottky contact as the gate element (Fig. 11.12). The depletion-mode MESFET was fabricated on a MBE-grown, Sn-doped, 300-nm-thick homoepitaxial film on an Mg-doped semi-insulating (010) β-Ga2O3 substrate. A circular pattern without any surface passivation was used to fabricate source/drain ohmic contacts that were formed on the surface that had undergone RIE using BCl3 followed by metallization with Ti/Au (20/230 nm). Schottky gates were formed by depositing Pt/Ti/Au (15/5/250 nm). A device with a gate length of 4 μm, source/drain separation of 20 μm and a 200-μm-diameter drain contact showed pinch-off and conducted a maximum drain current of 26 mA/mm1 at a drain voltage of 40 V and a gate voltage of +2 V. The off-state leakage current was less than 4 μA down to 40 V, whereas the maximum transconductance was 1.4 mS, and the off-state breakdown occurred at 250 V. The device suffered from high source/drain contact resistance and a low ID on/off ratio. Since MESFETs are normally on devices, they are not typically the preferred choice in circuit applications. Pt/Ti/Au Ti/Au Source

Gate

Ti/Au Drain

Sn-doped n-Ga2O3 300 nm

Mg-doped semi insulating b-Ga2O3 (010) substrate

Fig. 11.12 Schematic representation of a β-Ga2O3 MESFET [53]. Reproduced with permission from Higashiwaki, M., et al., Semicond. Sci. Technol. 31 (2016) 034001.

258

Gallium Oxide

Recent emphasis has been on the development of MOS-type devices using atomic layer deposited (ALD) Al2O3 or SiO2 films as the gate dielectric [53]. However, most Ga2O3 MOS devices have been normally on, depletion-mode devices. An enhancement-mode, fin-array field effect transistor (FINFET) fabricated on a Sn-doped homoepitaxial film grown on a (100) Mg-doped, semi-insulating substrate using ALD Al2O3 as the gate dielectric was demonstrated by Chabak et al. [54]. Although the drain current was limited by high the on-resistance of the device, an off-state breakdown voltage of 600 V was achieved without any passivation. For additional details on Ga2O3 MOSFETs the reader is referred to Chapter 17.

11.8

Summary

Schottky contacts have been reported on β-Ga2O3 bulk single-crystal substrates grown by EFG, CZ, and float-zone methods; different crystal orientations, including
201 , (100), and (010) were employed. A minority of Schottky contacts structures were reported on β-Ga2O3 epitaxial layers. Because all of the β-Ga2O3 crystals and epilayers are n-type (or semi-insulating), high work function metals are typically used as Schottky contacts, with Ni and Pt the most commonly reported. In general, device development on β-Ga2O3 is in its very early stages, but progress is steadily advancing: β-Ga2O3 SBDs with Vbr > 1000 V and RON Ed), where Ed is the lattice displacement energy [14, 18].

14.2

Radiation damage in wide bandgap semiconductors

The strong bonding in wide bandgap semiconductors gives them an intrinsically highradiation resistance [16, 19, 22–24]. The fluence of ionizing radiation at which materials and devices such as transistors and light-emitting diodes made from SiC, GaN, and related materials start to show degradation is about two orders of magnitude higher than in their GaAs equivalents [15, 22–24]. This difference is at least partially attributed to the stronger bonding of these materials [16, 22–24]. A measure of this bond strength is the energy required to displace an atom from its lattice position or simply the atomic displacement energy, denoted by Ed. This parameter has been measured in several semiconductors and empirically determined to be inversely proportional to the volume of the unit cell as shown in Fig. 14.2 [15]. This also generally scales with bandgap, so that these wide bandgap materials have intrinsically higher radiation resistance than Si. From the known size of the Ga2O3 unit lattice, it has been suggested that it should be quite radiation hard. The displacement energy threshold plays an especially crucial role in determining the induced defect concentration for incident electrons having energies 2 order of magnitude increase in on-state resistance at the highest fluence. There was a reduction in reverse current, which scaled with electron fluence. The on/off ratio at 10 V reverse bias voltage was severely degraded by electron irradiation, decreasing from approximately 107 in the unirradiated reference diodes to approximately 2  104 for the highest fluence of 1.43  1016 cm2. The reverse recovery characteristics showed little change even at this highest fluence, with values in the range 21–25 ns for all rectifiers. The changes in device characteristics were accompanied by a decrease in electron diffusion length from 325 to 240 μm at 300 K, as shown at the bottom of Fig. 14.3. Fig. 14.4 shows compilation of carrier removal rates in Ga2O3 for different types and energy of radiation on the same scale as that of GaN-based layers and devices [32]. In the latter case, the standard GaN transistor is the high-electron-mobility transistor (HEMT), consisting of a thin layer of AlGaN or InAlN on a thicker layer of GaN. This structure has a much thinner active than a rectifier of the type used in Ga2O3, but the sheet doping concentration is also higher, so it should be more radiation resistant than a rectifier. The data reported to date shows that the carrier removal rates in Ga2O3 are basically comparable to those in GaN.

Current density (mA/cm2)

Radiation damage in Ga2O3

319

Fig. 14.3 (A) I-Vs from Ga2O3 diodes before and after 1.5 MeV electron irradiation to different doses (B) diffusion length of electrons as a function of temperature after different electron irradiation doses.

Reference Dose = 1.79 × 1015 cm–2 Dose = 1.43 × 1016 cm–2

105 103 101 10–1 10–3 –60

–45

0.0 0.4 Voltage (V)

(A)

0.8

Diffusion length, Lp (mm)

360 Nonirradiated 1.79 × 1015 (C)

340

1.43 × 1016 (C)

320 300 280 260 240 220 280

300

(B)

320

340

360

380

106

Protons Neutrons p-GaN

Carrier removal rate (cm–1)

400

Temperature (K)

InAIN MOSHEMT

104 n-GaN

Stressed AIGaN HEMT n-GaN

102

100

p-GaN

- This work - OSU [26] - UF [25]

Electrons g-Rays InAIN HEMT Ga2O3

AIGaN MOSHEMT AIGaN HEMT Ga2O3

n-GaN

Ga2O3 n-GaN

n-GaN

10–2

0.1

1 Energy (MeV)

10

Fig. 14.4 Carrier removal rate in single-layer GaN or HEMT structures as a function of energy for different types of radiation.

320

Gallium Oxide

Fig. 14.5 Optical image of singleGa2O3 nanobelt transistor (top) and the experimental configuration for studying effect of proton damage on the characteristics of these devices (bottom).

b-Ga2O3 nanobelt

Ti/Au (S)

Ti/Au (D) 10 mm

10 MeV proton beam (low vacuum)

Ti/Au (S)

Ti/Au (D) b-Ga2O3

SiO2/Si Ti/Au (BG)

(iii) Proton damage in Ga2O3 nanobelt transistors

Proton damage in back-gated field-effect transistors (FETs) fabricated on exfoliated quasi-two-dimensional β-Ga2O3 nanobelts was studied for devices exposed to fluences of 10-MeV protons [40]. Fig. 14.5 (top) shows an SEM image of the nanobelt transistor structure, in which the exfoliated nanobelt is positioned between Ti/Au source and drain contacts and SiO2/Si metallized with Ti/Au is used as the backgate. These devices were then exposed to 10 MeV proton beams, as shown at the bottom of Fig. 14.5. The radiation damaged FETs showed a decrease of 73% in the field-effect mobility and a positive shift of threshold voltage after proton irradiation at a fluence of 2  1015 cm2, which corresponds to approximately 105 times the intensity of a solar proton event. The on/off ratio of the exfoliated β-Ga2O3 FETs was maintained even after proton doses of up to 2  1015 cm2. The data are summarized in the drain-source characteristics of Fig. 14.6, which show the effect of proton dose [40]. Note that doses of 1–2  1015 cm2 both lead to significant suppression of the drain current. The radiation-induced damage in β-Ga2O3-based FETs was significantly recovered after rapid thermal annealing at 500°C [40]. This annealing temperature is also similar to that needed for removal of plasma-induced dry etch damage in Ga2O3 [41]. It will be interesting to compare the results from nanobelt transistors with more conventional devices fabricated on bulk or epi Ga2O3.

Radiation damage in Ga2O3

as-fabricated VGS +60 V +40 V +20 V 0 –20 V –40 V –60 V

9 6 3

9 6 3

0

0 0

VDS (V)

20

30

10 Mev, 2x1015 cm–2 VGS 60 V 40 V 20 V 0 –20 V –40 V –60 V

12 IDS (mA/mm)

10

9 6

12

3 0

9

10

VDS (V)

20

30

at VDS = 30 V as fabricated 10 MeV, 1x1015 cm–2 10 MeV, 2x1015 cm–2

6 3 0

0

(C)

0

(B)

IDS (mA/mm)

(A)

10 Mev, 1x1015 cm–2 VGS +60 V +40 V +20 V 0 –20 V –40 V –60 V

12 IDS (mA/mm)

12 IDS (mA/mm)

321

10

20

30

VDS (V)

–60

(D)

–30

0 VGS (V)

30

60

Fig. 14.6 Output characteristics (IDS vs VDS) of β-Ga2O3 nanobelt FET before and after 10-MeV proton irradiation at different doses: (A) as-fabricated, (B) 1  1015 cm2, and (C) 2  1015 cm2, (D) transfer characteristics (IDS vs VGS) of β-Ga2O3 nanobelt FET at VDS ¼ 30 V before and after 10-MeV proton irradiation at different doses. From G. Yang, S. Jang, F. Ren, S.J. Pearton, J. Kim, Influence of high-energy proton irradiation on β-Ga2O3 nanobelt field-effect transistors, ACS Appl. Mater. Interfaces 9 (2017) 40471–40476.

(iv) Proton damage in Ga2O3 photodetectors

As mentioned previously, the cut-off wavelength of β-Ga2O3 is 250 nm, which means it is intrinsically solar-blind and DUV photodetectors made from this material do not require any supplementary filters. To test their response to high-energy protons, planar thin film β-Ga2O3 photodetectors were irradiated with 5 MeV protons at doses from 1013 to 1015 cm2 and the resulting effects on photocurrent, responsivity, quantum efficiency, and photo-to-dark current ratio (PDCR) at 254 nm wavelength were measured at both 25°C and 150°C [37]. The devices were subjected to doses of 5 MeV protons between 1013 and 1015 cm2 at 25°C. These doses are equivalent to many decades of exposure in low Earth orbit. Protons of this energy completely traverse the Ga2O3 film and pass into the sapphire substrate. The energy loss and damage profile were calculated from the SRIM code. This Monte Carlo code (transport of ions in matter)/SRIM (stopping and range of ions in matter) is widely used to obtain information about vacancy production rates. SRIM can also be employed to calculate

322

Gallium Oxide

NIEL. The SRIM output gives the vacancy production rate as a function of position as the incident proton slows down in the target material. Combining these data with the total energy loss data, the vacancy production rate as a function of proton energy can be found [37]. Fig. 14.7 shows (A) SRIM simulation of the vacancy profile created by proton irradiation into the β-Ga2O3 photodetector structure, (B) simulation of shallow region near the surface equivalent to the entire thickness of the Ga2O3, and (C) average vacancies of gallium and oxygen within the β-Ga2O3 region as a function of different proton irradiation doses. The total vacancy concentration in the Ga2O3 scales with proton dose. It should be noted that these densities represent an upper limit because some fraction of the defects can recombine at room temperature. However, they do provide a useful indicator of the density of point defects created. The dark current increased in proportion with the implant dose, leading to a decrease in the ratio of photocurrent to dark current. The increase in photocurrent relative to the dark current measured under exposure to this light can be explained by the presence of defect levels within the bandgap. Fig. 14.8 shows the resulting photo to dark current ratio for (A) 254 nm and (B) 365 nm illumination as a function of different proton irradiation dose at room temperature and 150°C. The PDCR decreased from 60 in the control samples to 9 after proton doses of 1015 cm2 for illumination with 254 nm photons and correspondingly lower numbers for 365 nm illumination. Ga2O3 photodetectors were subject to fluences of 5 MeV protons from 1013 to 1015 cm2. The nonionizing energy loss of the protons as they traverse the Ga2O3 layers creates states in the gap that increase the photocurrent but decrease the PDCR as the proton fluence increases. (v) Defects created by proton implantation into Ga2O3 There is particular interest in the properties of hydrogen in wide bandgap semiconductors in general and in Ga2O3 in particular because of the predictions from density functional theory and total energy calculations that it should be a shallow donor in this material [45, 46]. The generally observed n-type conductivity, therefore, may at least in part be explained by the presence of residual hydrogen from the growth ambient, rather than to native defects such as Ga interstitials or O vacancies, which are suggested to be deep donors. There is some experimental support for the fact that hydrogen may be a shallow donor in Ga2O3 from experiments on its muonium counterpart and from electron paramagnetic resonance of single-crystal samples [47, 48]. To understand the defects created in Ga2O3, H, or D was introduced by ion implantation at room temperature with multiple doses and energies (up to 180 keV) to obtain H or D concentrations of approximately 1  1020 cm3 in a layer 1200 nm in depth. All of the crystals used had (2 0 1) surface orientation, with [010] and [102] edges. Fourier transform infrared spectroscopy (FTIR) experiments were carried out to examine temperature- and polarization-dependent effects as well as relative H- and D-concentrations. The results of analysis of these data, coupled with detailed theoretical calculations, show no evidence of interstitial atomic hydrogen (Hi); instead, the defects observed appear to be in a family that involves H trapped at a Ga vacancy, the primary member involving a particular two-H configuration. This configuration is shown in Fig. 14.9 [49]. For samples in which hydrogen or deuterium was inserted

Radiation damage in Ga2O3

1×1015 cm–2

20

Vacancies (vacancies/cm3)

323

10

7×1013 cm–2 1×1013 cm–2

1019 1018 1017 1016 0

20

(A)

40

80

60

100

120

Depth (mm) 1020

Vacancies (vacancies/cm3)

1×1015 cm–2 7×1013 cm–2 1×1013 cm–2

1018

1016

1014

0

200

(B)

400

600

800

1000

Depth (nm)

Ga Vacancies O Vacancies Total

Vacancies (cm–3)

1017

1016

1015 1013

(C)

1014

1015 –2

Dose (cm )

Fig. 14.7 (A) SRIM simulation for proton irradiation into β-Ga2O3 photodiode, (B) zoomed in simulation of shallow region, and (C) average vacancies of Gallium and oxygen within Si implanted β-Ga2O3 region as a function of different proton irradiation doses. Reprinted with permission from S. Ahn, Y.-H. Lin, F. Ren, S. Oh, Y. Jung, G. Yang, J. Kim, M.A. Mastro, J.K. Hite, C.R. Eddy, Jr., S.J. Pearton, Effect of 5 MeV proton irradiation damage on performance of β-Ga2O3 photodetectors, J. Vac. Sci. Technol. B34 (2016) 041213-1-6, Copyright 2016, American Vacuum Society.

324

60

PDCR

Fig. 14.8 Photo to dark current ratio of (A) 254 nm and (B) 354 nm as a function of different proton irradiation doses at room temperature and 150°C. Reprinted with permission from S. Ahn, Y.-H. Lin, F. Ren, S. Oh, Y. Jung, G. Yang, J. Kim, M.A. Mastro, J.K. Hite, C.R. Eddy, Jr., S.J. Pearton, Effect of 5 MeV proton irradiation damage on performance of β-Ga2O3 photodetectors, J. Vac. Sci. Technol. B34 (2016) 041213-1-6, Copyright 2016, American Vacuum Society.

Gallium Oxide

Room 150°C 254 nm

40

20

0 Ref.

1013

(A)

1014

1015

Dose (cm–2)

60

PDCR

Room 150°C 365 nm

40

20

0 Ref.

(B)

1013

1014

1015

–2

Dose (cm )

Fig. 14.9 Relaxed configuration of Ga(1)21 plus two H which satisfies the polarization and weak coupling constraints. Modes with zero and nonzero transition moments are illustrated.

Radiation damage in Ga2O3

325

by annealing in those ambients of H2 or D2 gas, we observe strong absorption lines at 3437 and 2545 cm1. If the samples were annealed in H2 and D2 simultaneously, these OH and OD lines become split into two lines. This requires defects that contain two equivalent H atoms, not one. This, and the fact that the lines are completely polarized, leads to the model where two H atoms are bonded to a Ga vacancy. When the samples are implanted with hydrogen, additional absorption peaks are observed beyond the usual 3437 and 2545 cm1 lines. There are three additional lines in the implanted samples not seen in samples treated in H2 gas. As they are annealed, these defects become converted into the 3437 and 2545 cm1 lines at 400°C. These lines are stable up to 700°C where they are then converted into other new lines. All of these lines have the same polarization properties which suggest that they have related structures.

14.5

Summary and conclusion

To summarize, the initial data on proton, electron, neutron, and gamma irradiation of photodetectors and transistors show fairly similar radiation resistance to GaN devices under the same conditions. The carrier removal rates in irradiated rectifier structures range from 5 for 1.5 MeV electrons to 300 for 5 MeV protons. The dominant defect formed in Ga2O3 by annealing in an H2 ambient or by the implantation of protons is a specific relaxed VGa-2H structure for the 3437 cm1 line that dominates the IR absorption spectra.

Acknowledgments The work at UF is partially supported by HDTRA1-17-1-0011 (Jacob Calkins, monitor). The project or effort depicted is sponsored by the Department of the Defense, Defense Threat Reduction Agency. The content of the information does not necessarily reflect the position or the policy of the federal government, and no official endorsement should be inferred. The work at NRL is partially supported by Defense Threat Reduction Agency grant HDTRA1-17-1-0011 and the Office of Naval Research. The work at Korea University was supported by the Korea Institute of Energy Technology Evaluation and Planning (KETEP) and the Ministry of Trade, Industry & Energy (MOTIE) of Korea (No. 20172010104830) and Space Core Technology Development Program (2017M1A3A3A02015033) through the National Research Foundation of Korea funded by the Ministry of Science, ICT, and Future Planning of Korea.

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Ga2O3 nanobelt devices Janghyuk Kim, Sooyeoun Oh, Suhyun Kim, Jihyun Kim Department of Chemical and Biological Engineering, Korea University, Seoul, South Korea

15

Chapter Outline 15.1 Introduction

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15.1.1 Bottom-up methods 332 15.1.2 Top-down methods 333

15.2 β-Ga2O3-based optoelectronic nanodevice

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15.2.1 Introduction of β-Ga2O3-based optoelectronic nanodevice 337 15.2.2 Development of β-Ga2O3-based photodetectors 340

15.3 β-Ga2O3-based nanoelectronic devices

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15.3.1 Introduction of β-Ga2O3 nanobelt-based transistors 350 15.3.2 Development of β-Ga2O3 nanobelt-based transistors 353

15.4 Summary 363 References 363

15.1

Introduction

β-Gallium oxide (β-Ga2O3) is attractive as a novel material for (opto)electronics, especially high-power electronics and solar-blind photodetectors (PDs). It has an ultrawide direct bandgap of around 4.8–4.9 eV at room temperature and high thermal and chemical stabilities [1, 2]. The theoretical electrical breakdown field (Ebr) of β-Ga2O3 is known to be
8 MV/cm, and 3.8 MV/cm of Ebr has been experimentally demonstrated in a recent report, recording a higher value than those of GaN and SiC. Baliga’s figure of merit (BFOM) of β-Ga2O3 is also superior among some of the other popular wide-bandgap semiconductors, such as 4H-SiC and GaN [3–6]. These outstanding properties have led to a large number of reports on various electrical devices based on β-Ga2O3 including metal-oxide-semiconductor field-effect transistors (MOSFETs), metal-semiconductor field-effect transistors (MESFETs), and Schottky barrier diodes [7–10]. Furthermore, the wide bandgap of β-Ga2O3 provides intrinsic solar blindness that allows fabrication of solar-blind PDs without the need for additional optical filters that block light in the range of long wavelength [11]. Singlecrystal β-Ga2O3 is commercially available as a various of growth methods exist; especially the edge-defined film-fed growth (EFG) method that can be used to grow bulk β-Ga2O3 substrates with high crystal quality [12, 13]. However, the low thermal conductivity of β-Ga2O3 has to be considered when fabricating high-power electrical Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00015-4 © 2019 Elsevier Inc. All rights reserved.

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devices [14]. Instead of using native β-Ga2O3 as the substrate, introducing substrates with higher thermal conductivity can be a solution. Many researchers have made attempts to use β-Ga2O3 in the form of film or nanostructures on different substrates other than β-Ga2O3. In particular, nanobelts (nanosheets, nanomembranes, or nanoribbons) have been highlighted for their one-dimensional (1D) or twodimensional (2D) properties when prepared by bottom-up or top-down methods. Diverse nanostructures of metal oxides have attracted great attention for their large surface-to-volume ratio, and unique optical and electrical properties, further encouraging a wide range of applications in nanoelectronics or photonics [15, 16]. The first section of this chapter focuses on the preparation of β-Ga2O3 nanobelts and their morphological, structural, and optical properties. The fabrication of devices, including PDs and transistors, based on β-Ga2O3 nanobelts is dealt with later in this chapter along with details on their characteristics.

15.1.1 Bottom-up methods β-Ga2O3 nanostructures have been synthesized using a variety of bottom-up methods including physical evaporation, arc discharge, laser ablation, and microwave-plasma chemical vapor deposition (MPCVD). Among various nanostructures, nanobelts have rectangular cross sections with belt-like or ribbon-like morphology. Pan et al. successfully prepared nanobelt structures of Ga2O3 and other semiconducting oxides, such as ZnO, SnO2, In2O3, CdO, and PbO2 through thermal evaporation of oxide powders at high temperature without any catalysts. The synthesized nanobelts had well-defined geometry and good crystallinity. Gundiah et al. [17] heated a mixture of β-Ga2O3 powder and multiwalled carbon nanotubes at 1000°C to obtain Ga2O3 nanosheets and nanobelts. Scanning electron microscope (SEM) images revealed that they have a morphology of belts and sheets [Fig. 15.1A]. The nanobelts had widths of 150–200 nm and reflections corresponding to (104), (211), and (202) as shown in the lowmagnification transmission electron microscope (TEM) image and selected area electron diffraction (SAED) pattern in Fig. 15.1B. Xiang et al. [18] synthesized β-Ga2O3 nanosheets and nanobelts by evaporating a gallium metal source; the report included X-ray powder diffraction (XRD) pattern [Fig. 15.1C] and energy-dispersive X-ray spectrum (EDX) [Fig. 15.1D] to prove that the nanostructures are pure single-phase β-Ga2O3 and consisted of Ga and O elements. Zhang et al. [19] fabricated Ga2O3 nanobelts in large scale through physical evaporation of Ga and Ga2O3 powder in oxygen atmosphere. The SEM images indicated that the lengths and the widths of the nanobelts are in the range of several tens to several hundreds of micrometers and 20–400 nm, respectively. As shown in Fig. 15.2A, photoluminescence (PL) spectra of the Ga2O3 nanobelts gave three blue-light emission peaks between 410 and 460 nm and one ultraviolet (UV) emission peak at 310 nm; the emission of blue light was stronger compared to that of Ga2O3 powder, suggesting the suitability of Ga2O3 nanobelts for optoelectronic devices. Kuo et al. [20] used the vapor-phase transport method to synthesize β-Ga2O3 nanobelts and other nanostructures on silicon substrates. Diffuse reflectance spectroscopy indicated that the nanostructures have a bandgap of 4.56 eV. Li et al. grew Ga2O3 nanobelts using gallium and oxygen as precursors in a hot-wall chemical vapor deposition (CVD)

Ga2O3 nanobelt devices

333

Fig. 15.1 (A) SEM image and (B) TEM image of synthesized β-Ga2O3 nanosheets and nanobelts. The inset is the SAED pattern of a nanobelt. (C) XRD pattern and (D) EDX spectrum of β-Ga2O3 nanostructures. (A and B) From G. Gundiah, A. Govindaraj, C.N.R. Rao, Nanowires, nanobelts and related nanostructures of Ga2O3, Chem. Phys. Lett. 351 (2002) 189–194. (C and D) From X. Xiang, C.-B. Cao, Y.-J. Guo, H.-S. Zhu, A simple method to synthesize gallium oxide nanosheets and nanobelts, Chem. Phys. Lett. 378 (2003) 660–664.

equipment [21]. Field-emission SEM images revealed synthesized nanobelts with uniform lengths, widths, and thicknesses of several tens of micrometers, 1–3 μm and 50–200 nm, respectively. A TEM image [Fig. 15.2B] was obtained for a singlepristine nanobelt and the corresponding high-resolution TEM (HRTEM) image [Fig. 15.2C] confirmed the monoclinic structure of a single-crystal Ga2O3. The individual nanobelts were then used for the fabrication of PDs.

15.1.2 Top-down methods The monoclinic structure of β-Ga2O3 allows a simple top-down method of preparing nanobelts. From the crystal structure, it can be observed that β-Ga2O3 clearly has a ˚ , b ¼ 3.0371 A ˚ , and larger lattice constant along one direction (a ¼ 12.214 A ˚ c ¼ 5.7981 A) [7, 22, 23]. This property allows cleavage into thin flakes mostly along

334

(A) 415 nm PL intensity (a.u.)

(B)

Ga2O3 nanobelts Ga2O3 powder

438 nm

(C)

460 nm

475 nm 340 nm 350 nm

300

350

400

450

500

550

600

650

700

Wavelength (nm)

Fig. 15.2 (A) PL spectra of Ga2O3 nanobelts and powder. (B) TEM image. (C) HRTEM image of an individual Ga2O3 nanobelt (lower left inset is the corresponding SAED pattern). (A) From J. Zhang, F. Jiang, L. Zhang, Fabrication, structural characterization and optical properties of semiconducting gallium oxide nanobelts, Phys. Lett. A 322 (2004) 363–368. (B and C) From L. Li, E. Auer, M. Liao, X. Fang, T. Zhai, U.K. Gautam, A. Lugstein, Y. Koide, Y. Bando, D. Golberg, Deep-ultraviolet solar-blind photoconductivity of individual gallium oxide nanobelts, Nanoscale 3 (2011) 1120–1126. Gallium Oxide

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335

the [100] direction although there exists two cleavage planes parallel to the (100) and (001) planes [Fig. 15.3C] [23]. Therefore, despite the fact that β-Ga2O3 is not a 2D material formed by the van der Waals interaction, it can be mechanically exfoliated in the same manner as 2D materials like graphene and fabricated into transistors as depicted in Fig. 15.3A, B, and D. No other impurities besides Ga and O were observed in the EDX spectrum of the β-Ga2O3 flake [Fig. 15.4A] and the optical bandgap extracted from the absorption spectrum of β-Ga2O3 flake [Fig. 15.4B] agrees with that of obtained by first-principles calculations [23]. The exfoliated β-Ga2O3 flakes had thicknesses varying from
20 to 400 nm as observed using atomic force microscopy (AFM) [7, 24–26]. The thicknesses of the thinner flakes were comparable to those of other 2D materials. The AFM image and the corresponding height profile in Fig. 15.5A and B, respectively, displayed an exfoliated flake with a thickness of
200 nm and considerably low surface roughness [7]. SEM images of the exfoliated flakes were also obtained as shown in Fig. 15.5C and D; the smooth surface and the morphology of stacked layers at the edge of the flake were revealed, proposing the feasibility of 2D β-Ga2O3 flakes via repeated mechanical exfoliation [7]. Mechanical exfoliation is not only simple in process, but it

Fig. 15.3 (A–D) Schematic of mechanical exfoliation of β-Ga2O3 flakes using Scotch-tape method and fabrication of transistor using the exfoliated flake. From J. Kim, S. Oh, M.A. Mastro, J. Kim, Exfoliated β-Ga2O3 nano-belt field-effect transistors for air-stable high power and high temperature electronics, Phys. Chem. Chem. Phys. 18 (2016) 15760–15764.

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5

4

O

Absorption (a.u.)

Counts (a.u.)

Stoichiometry Ga2O3 within 5% variation

Ga Ga

3

2

1

Cu

EG ~ 4.77 eV

Ga 0

(A)

2

4 6 8 10 Energy (keV)

12

0 4.0

(B)

4.5 5.0 5.5 Photon energy (eV)

6.0

Fig. 15.4 (A) Normalized intensity of EDX spectrum and (B) absorption spectrum of β-Ga2O3 flake with photon energy. From W.S. Hwang, A. Verma, H. Peelaers, V. Protasenko, S. Rouvimov, H. Xing, A. Seabaugh, W. Haensch, C.V. de Walle, Z. Galazka, M. Albrecht, R. Fornari, D. Jena, Highvoltage field effect transistors with wide-bandgap β-Ga2O3 nanomembranes, Appl. Phys. Lett. 104 (2014) 203111.

also results in thin flakes with good crystal quality and smooth interface when stacked with other layers to form a heterostructure. Peaks at the same values of Raman shift appeared in the Raman spectra of mechanically exfoliated β-Ga2O3 flakes as in those of bulk β-Ga2O3 sample [Fig. 15.6A] [11]. This suggests that the crystal structure was maintained after the repeated mechanical exfoliation and the following photolithography process. TEM images also displayed high crystallinity of β-Ga2O3 flakes. The HRTEM image of the cleavage plane showed the lattice symmetry and lattice parameters of the bc plane identical to those of the bulk sample. HRTEM images were also obtained for the cross section of the flakes [Fig. 15.6B and C], confirming that the flakes were exfoliated along the [100] direction. The distance between the (100) planes and the angle between [001] and [100] are known to be around 1.22 nm and 103.51°, respectively, which are comparable to the experimental results,
1.14 nm and
104° [11, 27]. An atomically smooth interface between β-Ga2O3 and insulating layer or the substrate could be observed in the crosssectional images of fabricated devices as the layers are connected by the van der Waals interaction as in 2D materials [7, 25]. As mechanically exfoliating flakes creates difficulties in controlling the thickness, a plasma-etching technique that reduces the thickness of β-Ga2O3 flakes was recently studied. Reactive-ion etching (RIE) process was performed using SF6 gas to uniformly etch the flakes both vertically and laterally. The thickness of a flake could be reduced to
60 nm and the etch rate was calculated to be 16.4 nm/min [28]. Inductively coupled plasma (ICP) discharges of either Cl2/Ar or BCl2/Ar were also used to etch bulk β-Ga2O3 and etching processes at both high and low rates were achieved by controlling ICP conditions [29].

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Fig. 15.5 (A) AFM image of a transistor based on exfoliated β-Ga2O3 flake. (B) The height profile obtained from (A). (C) SEM image of the identical transistor, and (D) the magnified SEM image of the red dotted area in (C). From J. Kim, S. Oh, M.A. Mastro, J. Kim, Exfoliated β-Ga2O3 nano-belt field-effect transistors for air-stable high power and high temperature electronics, Phys. Chem. Chem. Phys. 18 (2016) 15760–15764.

15.2

β-Ga2O3-based optoelectronic nanodevice

15.2.1 Introduction of β-Ga2O3-based optoelectronic nanodevice Direct wide-bandgap semiconductors are suitable for UV PDs due to their insensitivity to longer wavelengths and their very low dark current even at elevated temperatures [30, 31]. UV PDs are classified as “visible-blind” detectors when their cut-off wavelength is in the range of 280–400 nm and as “solar-blind” detectors when their cut-off wavelength is shorter than 280 nm [30–32]. Promising alternative widebandgap materials include SiC, diamond, and III-nitrides, especially AlxGa1 xN, ZnO, and β-Ga2O3 [30–33]. The bandgaps of these materials are presented in Table 15.1 [30–35]. Among these, GaN and AlxGa1 xN need to be prepared throuth hetero-growth techniques because their bulk substrate does not exist. When using hetero-growth methods, defects are inevitably induced due to lattice mismatch between the substrate and the grown material. These defects exist in the film, and

338

Fig. 15.6 (A) Raman spectra of a mechanically exfoliated β-Ga2O3 flake transferred on SiO2 substrate (top) and bulk β-Ga2O3 (bottom). (B) HRTEM image, and (C) SAED pattern of an exfoliated β-Ga2O3 flake. From S. Oh, J. Kim, F. Ren, S.J. Pearton, J. Kim, Quasi-two-dimensional β-gallium oxide solar-blind photodetectors with ultrahigh responsivity, J. Mater. Chem. C 4 (2016) 9245–9250.

Gallium Oxide

Ga2O3 nanobelt devices

Table 15.1

Promising materials for UV PDs with different bandgaps

UV-C 4.43 ≤ Eg ≤ 6.20

UV-B 3.88 ≤ Eg ≤ 4.43

UV-A 3.10 ≤ Eg ≤ 3.88

Visible & IR Eg ≤ 3.10

Material

Eg (eV)

Material

Eg (eV)

Material

Eg (eV)

Material

Eg (eV)

β-Ga2O3

4.84.9

–

–

GaN

3.4

6H-SiC

3.0

Diamond

5.5

–

–

ZnO

3.3

CdS

2.5

LaAlO3

5.5

–

–

4H-SiC

3.2

CdSe

1.7

h-BN

4.0–5.8

–

–

TiO2

3.2

InN

1.9–2.05

In2Ga2O7

4.4

–

–

SnO2

3.8

GaSe

2.1

Zn2GeO4

4.5

–

–

ZnS

3.6

GaAs

1.4

AlxGa1xN (x  0.4)

4.3–5.9

–

–

Si

1.2

Zn1xMgxO (x  0.4)

4.3–5.75

–

–

MoS2

1.75

A1N

6.2

–

–

BP

0.3–2.0

339

340

Gallium Oxide

inevitably increase the dark current. GaN and AlxGa1 xN are suitable for solar-blind PDs because of their direct and wide bandgap which ranges from 3.4 to 6.2 eV as the concentration of Al increases. However, growing high-quality AlxGa1 xN with high concentration of Al is quite challenging because the defect density of AlxGa1 xN also increases with Al concentration. These defects in the material generate a deep level in energy bandgap, leading to induced defect-defect transitions or defect-band transitions, which deteriorate the performance of optical devices [31]. Among the aforementioned materials, β-Ga2O3 has been extensively studied as a promising active layer material for deep UV PDs due to its inherent solar blindness. Unlike narrow-bandgap materials such as Si and GaAs, β-Ga2O3-based PDs do not need additional filters or heavy cooling systems to eliminate light of longer wavelengths. As a result, such devices have the advantages of reduced cost and launch weight, being compactness and robustness, and a better rejection of light wavelength radiation due to their sharp UV band edge. Moreover, β-Ga2O3 is both thermally and chemically robust, which makes it less subject to aging and ideal for radiation-intense space applications [12, 36]. Due to these superior features, β-Ga2O3-based PDs have been investigated using various types of β-Ga2O3 such as bulk, thin films, and nanostructures [11, 27, 28, 37–55]. Compared with conventional thin-film-based PDs, nanoflake (or nanobelt, nanoribbon, and nanomembrane)-based PDs have a higher responsivity and photoconductive gain due to the high surface-to-volume ratio of the nanostructures [31, 33, 56, 57]. As mentioned before, β-Ga2O3 nanobelts can be grown using various techniques including CVD, PLD and solution-based techniques, such as the Czochralski process and EFG. In addition, owing to its unique crystal structure, β-Ga2O3 nanobelts can be obtained from bulk β-Ga2O3 with high crystallinity using mechanical exfoliation method [12, 23, 36].

15.2.2 Development of β-Ga2O3-based photodetectors This section discusses the basic principles of PD and advancements made in the fields of processing and fabrication technologies for β-Ga2O3-based PDs. Numerous studies on β-Ga2O3-based UV PDs have been conducted and various types of detectors have been reported; they include photoconductive, metal-semiconductor-metal (MSM), and photovoltaic (especially PN and Schottky diodes) types [11, 27, 28, 37–55]. The schematics of respective types of PDs are shown in Fig. 15.7. Semiconductor-based PDs are operated by electrical transition (or interband transition). The detector ideally works when absorbed photons have sufficient energy to excite electrons from the valance band to the conduction band. In general, photons with a higher energy than the bandgap energy of the semiconductor are absorbed in the semiconductor, and the electron-hole pairs are created simultaneously. These photogenerated carriers are separated by the electric field induced by the built-in potential or/and the applied bias voltage. The generated carriers contribute to the current which can be electrically measured. Therefore, the threshold response of a PD depends on the material used in the detector and the maximum wavelength for the PD response is calculated by the following equation.

Ga2O3 nanobelt devices

341

Photoconductive

Schottky photodiode

Metal-semiconductor-metal

Metal-intrinsic-semiconductor

p-i-n photodiode

Field effect phototransistor

Ohmic contact Schottky contact insulator

Intrinsic semiconductor p-type semiconductor n-type semiconductor

Bipolar phototransistor

Fig. 15.7 Schematics of various types of PDs.

λðμmÞ ¼

1:24 Eg ðeVÞ

This is the cut-off wavelength of the material; the detectors should be properly composed of a suitable semiconductor material depending on the target detection range. The performance of the PD is evaluated by the following parameters [58, 59]. 

Responsivity: The amount of signal coming out from the device per incident photon of a given energy and wavelength. The responsivity is calculated using the following equation:

R¼

Iphoto ληext qg ¼ hc Plight

where λ is the wavelength of the incident light, ηext is the external quantum efficiency, hc is the photon energy, q is the electron charge, g is the gain, Iphoto is the photocurrent, and Plight is the power of the incident light. The unit of responsivity is amperes per watt (A/W).   

Spectral response: The variation of the responsivity as a function of wavelength. This parameter shows the detection range of the device. Detectivity: The minimum detectable power required to achieve a signal-to-noise ratio of unity. The unit of detectivity is Jones. Detectivity is a figure of merit that is inversely proportional to the noise level; the performance of PD is better for higher figure of merit. Time response: The detector response time for a change in the signal. The PD response time is defined by the time it takes for the photocurrent to increases from 10% to 90% (or drops from 90% to 10%) of its maximum current.

Each parameter is affected by the device structure or the properties of the material used in the device. The photoresponse properties and performance of the devices will be compared for different device structures.

342

Gallium Oxide

15.2.2.1 Photoconductive detectors Photoconductive detectors are semiconductor PDs with a simple structure, consisting of two Ohmic contacts and an active layer. Photoconductive detectors feature a simple fabrication process and high responsivity, but suffer from high dark currents [58, 59]. Many groups have conducted the research on β-Ga2O3-based PDs of the photoconductive type and have reported their photoresponse properties [11, 28, 38–41, 46, 47]. Li et al. grew Ga2O3 nanobelts with the length and the thickness of
1.6 μm and
100 nm, respectively, using a CVD technique, and transferred Ga2O3 nanobelts onto a SiO2/Si substrate by drop casting. Using a conventional photolithography process, Cr/Au contacts were formed as an Ohmic metal to fabricate single-β-Ga2O3nanobelt-based PD. When its spectral responsivity was measured as a function of wavelength, the peak photoresponse was obtained at 250 nm. The fabricated PD exhibited a rejection ratio (R250nm/Rvisible) of
106, a photo-to-dark current ratio (PDCR) of
104, and a fast response time ( I0. (C) Under reverse bias, the depletion layer width increases, and Ism ≪ I0 as the barrier to electron flow from the semiconductor to the metal is very large and thus Ism is negligible.

Power MOSFETs and diodes

407

redistribution of electrons to reach an equilibrium state creates a space charge region in the semiconductor with a width given by W¼

  2ε0 εs ðVD  Va Þ 1=2 , qN D

which shows that the depletion layer width is proportional to the square root of the bias voltage [12]. The maximum electric field in the semiconductor, εm ¼  qNDW/(ε0εs), is at the metal/semiconductor interface. The Fermi potential, Vn, defines the position of the Fermi level relative to the conduction band by Vn ¼ ðkB T=qÞ ln ðNc =ND Þ ¼ ðEc  EF Þ=q: The diffusion potential, VD, given by V D ¼ ϕ B  Vn ¼ ϕ M  ϕS , creates a barrier to electron traveling from the semiconductor to the metal. Under moderate reverse bias (Fig. 17.2C), e.g., Va < 3(kBT/q), Ism approaches zero as the electron flow from the semiconductor to the metal is blocked by the barrier (VD + Va), and the current is simply the negative of the saturation current, qϕB

I0 ¼ AA∗ T 2 e kT . Hence the barrier height can be obtained from the I-V relationship under moderate reverse bias. Under moderate forward bias voltage (Fig. 17.2B), e.g., Va > 3(kBT/q), the forward current can be approximated as 

 qV a I ¼ I0 nkT : e Therefore, it is common to extract the ideality factor from the slope of a log forward current as a function of linear forward voltage. A related approach exists to extract the barrier height from capacitance-voltage relationship. Since C2 ¼ 2(Va + VD)/qNdεsA2, a plot of C2 vs V will have an intercept the voltage axis equal to the diffusion potential VD. Computational modeling by Rozhkov et al. found that Schottky diodes based on Au/Ga2O3 should have a barrier height of 1.23 eV and a breakdown voltage of approximately 2500 V. Furthermore, this theoretical analysis by Rozhkov et al. revealed that Schottky diodes based on Au/Ga2O3 should present a smaller reverse current than Au/ GaN and Ni/4H-SiC Schottky diodes [13]. Tadjer et al. found that Pt and TiN Schottky contacts to (201) β-Ga2O3 displayed near-unity ideality factor values, that is, nPt ¼ 1.07 and nTiN ¼ 1.09. In addition, the authors found Schottky barrier heights of approximately 1 eV, independent of temperature. The 65 nm ALD deposited TiN contact displayed approximately an order of

408

Gallium Oxide

magnitude lower reverse current at room temperature. In contrast to the TiN contact, at high temperature the reverse current of the Pt contact was significantly lower (0.27 mA/cm2) [14]. Yang et al. reported that vertical geometry Ni/Au/β-Ga2O3 Schottky rectifiers displayed a reverse breakdown voltage (without edge termination) that was a function of the diode diameter. Specifically, they found a reverse breakdown voltage of 920— 1016 V for a diode diameter of 105 μm and 810 V for a diode diameter of 210 μm. The layers were fabricated via hydride vapor-phase epitaxy layers on conducting bulk substrates. The authors found a Schottky barrier height of 1.1 at 25°C with an ideality factor of 1.08. An increase in temperature to 100°C led to a decrease in the Schottky barrier height to 0.94 with an ideality factor 1.28. They reported a FOM (V2BR/Ron) of 154.07 MW cm2 for a 105 μm diameter diode, with an on-state resistance of Ron ¼ 6.7 mΩ/cm2. Tests on switching from +5 V to 5 V, revealed a reverse recovery time of 26 ns [15]. Ahn et al. compared Schottky diodes based on Ni/Au contacts fabricated on a Si-doped Ga2O3 epitaxial layer on bulk Sn-doped Ga2O3 vs Pt/Au contacts on bulk β-Ga2O3. The Ni/Au Schottky diodes displayed a barrier height of 1.07 eV and the Pt/Au Schottky diodes displayed a barrier height of 1.04 eV at 25°C. The on-state resistance decreased with an increase in temperature form 25°C to 200°C, and the barrier height increased with temperature over the same 25°C to 200°C range. Defining a linear reverse breakdown relationship as VRB ¼ VRB0 + βðT  T0 Þ, Ahn et al. extracted a temperature coefficient of reverse breakdown voltage, β, as 0.1 mV/K for Pt/Au and 4 mV/K for Ni/Au. For the Ni/Au diodes, the FOM (V2B/RON) was approximately 3 MW cm2 at 25°C and approximately 1 MW cm2 at 200°C [16]. A summary of the literature for Schottky barrier height obtained from current– voltage measurements is available in Table 17.2. Examination shows that the Schottky barrier height is in the range of 0.95–1.5 eV for a variety of metals and crystal orientations.

Table 17.2 Literature values for Schottky barrier height and ideality obtained from current-voltage measurements [17] Metal

Orientation

Structure

Pd Ni Pt Au Au

(010) (010) (010) (010) (100)

EFG UID bulk EFG UID bulk EFG UID bulk EFG UID bulk FZ bulk (1017–1018)

SBH (eV) 1.29 1.5 1.53 1.2–1.1

Ideality

Ref

1.05 1.04 1.03 1.09 1.8–1.1

[18] [18] [18] [18] [19]

Power MOSFETs and diodes

Table 17.2

409

Continued

Metal

Orientation

Structure

Au

(100)

Ni

(100)

Pt

(010)

Ni

(010)

Ni

(010)

Ni Ni


201
201

Ni


201

Ni


201

Cu


201

CZ bulk (UID 0.6–8  1017) CZ bulk (UID 0.6–8  1017) FZ bulk (UID 3–5  1016) EFG bulk (UID 1.1  1017) 2-μm undoped Ga2O3 epi-layer on n-type (4.1   1018) Ga2O3:Sn EFG bulk (2.8  1017) EFG single crystal (1.7  1017) EFG single crystal (UID 1.2  1017) EFG single crystal (3.9  1018) PLD epi/ZnO/sapphire (1.6  1018) HVPE epi/sapphire (1.2  1016) Ga2O3:Si epi/bulk Ga2O3:Sn Bulk Ga2O3:Sn EFG bulk (3  1017) EFG bulk (3  1017) Mist Ga2O3:Sn (2-7e17)

Pt Ni/Au Pt/Au TiN Pt Pt/Ti/ Au

(001)
201
201
201
201 α-(0001)

SBH (eV)

Ideality

Ref

0.98

1.09

[20]

0.97

[21]

1.46–1.47

1.04–1.06

[22]

1.49

1.04

[23]

0.95–1.01

3.38–1.2

[24]

1.08 1.05

1.19

[25] [26]

1.25

1.01

[27]

0.9

1.4

[28]

0.92

1.22

[29]

1.5

1

[30]

1.07

1.3

[16]

1.04 0.98 1.01 1.7–2.0

1.28 1.09 1.07 1.1

[16] [14] [14] [31]

An interesting alternatively line of research involves the deposition of α-type Ga2O3 by a mist-CVD process. Based on this material, Oda et al. reported a Schottky barrier diode fabricated on a 0.43-μm Sn-doped α-Ga2O3 film with a Ti/Au Ohmic contact and a Pt/Ti/Au Schottky contact with an on-resistance of 0.1 mΩ cm2 and breakdown voltage of 531 V, which corresponds to a breakdown field strength of 12.35 MV/cm. They also reported a second Schottky barrier diode based on a 2.58-μm Sn-doped α-Ga2O3 film with an on-resistance of 0.4 mΩ cm2 and breakdown voltage of 855 V, which corresponds to a breakdown field strength of 3.31 MV/cm [31, 32].

410

Gallium Oxide

17.4

Power MOSFETs

The FET family of devices is desirable for high-power applications as they are high input impedance devices with simple voltage controlled operation—particularly when compared to the current control of a BJT. Still, MOSFETs operating under constant current applications can be vulnerable to thermal runaway. Power MOSFETs require minimal gate drive power owing to their insulated gate. These devices have high switching speed that is limited by the rate of charge applied on and removed from the capacitances in the MOSFETs. Power MOSFET can produce significant radiated and conducted emissions owing to the fast transient switching. The main metric for a power MOSFETs is efficiency and its defining parameters are drain-source on-state resistance, rise time, and fall time [33]. The total power loss in the power MOSFET is from the combination of conduction and switching loss, Ploss_total ¼ Pswitching_loss + Pcond_loss , where the conduction loss is the product of the drain-source on-state resistance and the loading current 2  RDS,On , Pcond_loss ¼ Iload

where Iload is the RMS current in the switch. Switching losses occurs during the charging and discharging of the MOSFET’s parasitic capacitance. The switching loss occurs due this finite switching times. Specifically, switching on VGS results in an IDS current rising time, tr, and, VDS decreases to its minimum, resulting in a power loss of Pr ¼ ðVIN  ID,min  tr Þ=6: Conversely, switching of VGS results in an IDS current falling time, tf, and VDS increases to its maximum, resulting in a power loss of Pf ¼ ðVIN  ID,max  tf Þ=6: Low-voltage MOSFETs can be configured with top-side source-and-drain contacts similar to digital MOSFET designs. In this configuration, it is common to internally connect the source and the body (back side contact) to stay at the same voltage. Referencing the gate bias voltage to the source and body, allows stable formation of the conduction channel. Moderate- to high-voltage Si-based MOSFETs employ a vertical design where the drain contact is formed on the bottom of the device. In the on-state, the alignment of the source and drain contacts enables a vertical flow of the electrons through a thick n drift layer. The thickness and doping level of the n drift layer determines breakdown voltage and on-resistance of the device. As discussed previously, the material

Power MOSFETs and diodes

411

properties of SiC, GaN, and Ga2O3 advantageously shift this breakdown-voltage/onresistance trade off. Commercial 1200 V SiC MOSFETs have been available in high-volume for many years with rated continuous drain current (Id) up to approximately 60 A at room temperature. These SiC MOSFETs allow stable switching up to approximately 50 kHz. The SiC-based MOSFETs have a specific on-resistance approximately four times better than comparable Si-based devices. The low total gate charge of SiC-based MOSFETs when compared to Si-based MOSFETs enable a further reduction in switching loss. In addition, SiC-based FETs are able to operate at elevated temperature (typically 200°C) with stable specific on-resistance over the temperature range. Despite years of progress in defect reduction, SiC-based devices suffer from low reverse blocking performance [34, 35]. Comparatively to SiC, GaN-based power FETs have lagged in high-voltage capability with 600 V parts appearing in 2011 and 1200 V parts only first appearing in 2012. Still, these GaN parts enable operation beyond 2 MHz, which enables large step down ratios in buck converters and decreases passive components sizes. Transistors based on Ga2O3 are less developed but early research has shown that Ga2O3 with a bandgap of 4.9 eV can operate a high-voltage level. An early work by Higashiwaki et al. employed molecular-beam epitaxy to grow a Sn-doped Ga2O3 layer on a semi-insulating (010) β-Ga2O3 substrate. The circular MESFET had a gate length of 4 μm and a source–drain spacing of 20 μm. The transistor displayed complete drain current pinch-off for VGS <  20 V, and the three-terminal off-state breakdown voltage of >250 V. A high on/off drain current ratio of about 10,000 A was obtained with a low drain leakage current of 3 μA at the off-state [36, 37]. Wong et al. developed field-plated Ga2O3 MOSFETs with a breakdown voltage of 755 V. Selective-area Si ion implantation doping of an undoped Ga2O3 epitaxial layer defined the device channel. The device gate length was 2 μm, source-drain spacing was 22 μm, the source-gate spacing was 5 μm, and the gate width was 15 μm. The MOSFETs exhibited an off-state breakdown field of 0.503 MV/cm and a high drain current on/off ratio >109. In addition, stable operation was observed up to a 300°C thermal stress. The maximum drain current, Idmax, was 78 mA/mm and the maximum transconductance was 3.40 mS/mm [38]. The relatively large a-plane lattice constant of 1.2 nm for β-Ga2O3 allows mechanically exfoliation of thin films from the (001) plane of a substrate [39]. Tadjer et al. used this tape method to transfer a 300 nm Ga2O3 flake onto a SiO2/Si substrate. Atomic layer deposition was used to form a high-k HfO2 gate dielectric layer. The insulated-gate (001) β-Ga2O3 MOSFET displayed a positive threshold voltage. The conductivity was due principally to the presence of oxygen vacancies in the substrate. A positive threshold voltage at low drain bias was confirmed by capacitance-voltage and current-voltage measurements [40]. Kim et al. used a similar exfoliation process to transfer a two-dimensional sheet of hexagonal boron nitride (h-BN) on a quasi-two-dimensional β-Ga2O3 layer. This heterostructure was fabricated into a metal-insulator-semiconductor field-effect transistor (MISFET) where the β-Ga2O3 acted as the channel and the h-BN acted as the gate dielectric. The h-BN dielectric was shown to be extremely flat and possess a

412

Gallium Oxide

pristine surface that provides a minimal density of charged impurities at the interface. The device demonstrated low gate leakage as well as an IDS on/off ratio of 107 [41]. Green et al. recently measured 3.8 MV/cm breakdown strength for a MOSFET fabricated via mesa isolation using a BCl3 inductively coupled plasma (ICP)/reactive-ion etch (RIE) process. An Sn-doped (100) β-Ga2O3 epitaxial layer was deposited by MOCVD onto a single-crystal Mg-doped semi-insulating (100) β-Ga2O3 substrate. The transistors had a 2-μm gate length, 3.4-μm source-drain spacing, and 0.6-μm gate-drain spacing (LGD). A gate-to-drain voltage of 230 V was maintained in the off-state. The gate-to-drain electric field corresponds to 3.8 MV/cm, which is approximately one-half the theoretical limit. It will be important to note that this breakdown voltage was measured in a sub-micron channel device [42]. Recently, Chabak et al. addressed the inherent tradeoff that high doping levels in Ga2O3 will increase the current density but will result in an undesirable negative threshold voltage [43]. Their approach is to form nonplanar fin-shaped channels, which allows carrier depletion along the sidewalls in addition to the surface. Fin-array field-effect transistor (finFET) devices were fabricated from a 300-nm Sn-doped Ga2O3 film that was grown homoepitaxially by MOVPE on Mg-doped semiinsulating (100) Ga2O3 substrate [44–46]. Ni gated Arrays of 300-nm wide fin channels with a 900-nm pitch were fabricated in the Sn-doped Ga2O3 layer. This device displayed an exceptionally high breakdown voltage of >600 V at a VGS ¼ 0 V off-state. Hu et al. argue that a fin or nanowire pillars are necessary geometries to achieve normally off operation in a wide bandgap FET, particularly for proper electrostatic gate control [47]. Moving to a vertical design is advantageous for high-voltage operation. In a recent work, Wong et al. applied the current aperture vertical electron transistor design to Ga2O3 [48, 49]. Sasaki et al. developed a β-Ga2O3 trench MOS-type Schottky barrier diodes (MOSSBDs). Halide vapor-phase epitaxy was employed to grow a 7-μm Si-doped Ga2O3 layer (6  1016 cm3) on a single-crystal Sn-doped β-Ga2O3 (001) substrate. A dry etch process was used to form a trench structure with a trench width of 4.8 μm, a mesa width of 1.2 μm, and mesa depth of 2.5 μm. A 50-nm HfO2 film was used as the dielectric film on the trench MOS structure. The trench MOSSBD displayed a specific on-resistance of 2.9 mΩ cm2 and the Schottky barrier diode component displayed a specific on-resistance of 2.3 mΩ cm2. This design (without a field plate or guard ring) displayed a breakdown voltage of approximately 240 V [50].

17.5

Application space and competing technologies

Semiconductor electronics can be broadly classified as bipolar (or minority carrier), or unipolar (or majority carrier). Majority carrier devices are usually faster and include Schottky diode, power vertical diffused MOSFET, and JFET. Minority carrier devices usually have better on-state performance and include variants of the PIN diode, IGBT, BJT, and thyristor.

Power MOSFETs and diodes

413

An important parameter in bipolar power devices, such as pin diodes, is the minority carrier lifetime. Indirect semiconductors such as Si and SiC intrinsically have a longer minority carrier lifetime compared to direct bandgap semiconductors such as GaN. There are several types of bipolar devices built in Si and SiC that display high blocking voltages. Silicon-based laterally diffused MOS (LDMOS) display relatively high-blocking voltage for a low on-resistance and is popular for RF power amplifiers in wireless base stations. Most applications for semiconductor electronic devices are classified as high speed or high power. The current and voltage handling capability as well as several factors influence the selection of a certain technology for a given application including reliability (ruggedness and thermal considerations), performance (linearity, efficiency), size, cost, and legacy. Scaling the device size in general increases the speed, i.e., frequency of the device but it is necessary to look at the particular device design as well as the underlying material parameters. The major applications for power electronics and corresponding semiconductor technology are depicted in Fig. 17.3.

Fig. 17.3 Current-voltage diagram of the application space for several semiconductor materials. At extremely high power and/or frequency, vacuum electronics are still implemented (not shown). The high-power/high-voltage market is primarily served by high-voltage vertical FETs (HVVFET), and silicon carbide-based devices such as 12 kV Schottky diodes are now commercially available. Given its exceptional properties, Ga2O3 is the leading candidate to address the ultrahigh-power market (>1 MW). The on-going question is how quickly will the Ga2O3 substrate cost decrease and size increase, thus enabling Ga2O3 devices to compete with GaN FETs, SiC Schottky and FETs, and Si-IGBTs in the medium- to high-power market.

414

Gallium Oxide

Low-voltage devices (up to 100 V) have major markets in mobile phones and personal computer power supplies as well as automotive and LED lighting. The phone and computer application space can be further delineated into three device types. The first device type is switching power supplies for AC-DC converters include half-bridge type (up to 400 W) for servers and televisions, forward-type (up to 200 W) for desktop PCs, game consoles, and printers, and flyback-type (up to 100 W) for phone charges and laptop adaptors. The second device type is DC-DC converters (up to 50 W) primarily used for on-board DC-DC voltage converters for GPU cards, motherboard, and communication equipment voltage supply points, and related voltage regulator modules. The third type is linear power supplies (up to 20 W) for electronic device power supplies. High-voltage power conditioning systems to convert between DC and 60 Hz AC are needed throughout a modern energy grid/distribution system. For example, DC-AC and AC-AC conversions and conditioning is necessary to connect the electric grid to wind-turbine farms, fuel-cell storage, thermal energy storage, and hydroelectric dams, as well as plug-in electric vehicles. A boost converter shown in Fig. 17.4 is the critical component of a switch-mode power supply. The freewheeling diode, also known as a suppressor diode, catch diode, snubber diode, or flyback diode, is used to eliminate a sudden voltage spike across an inductive load when its supply voltage is abruptly removed or reduced. These diodes are almost always found in circuitry where inductive loads are opened/closed or controlled by silicon components such as a silicon laterally diffused MOS transistor. Examples include motor drivers and relay drivers. Schottky diodes are advantageous due to their low forward voltage drop and fast reverse recovery voltage. Gallium oxide technology can play a role in all but the highest frequency variants of conversion circuits: AC-AC converters (frequency changers, cycloconverters); AC-DC converters (rectifiers, off-line converters); DC-AC converters (inverters); and DC-DC converters.

Fig. 17.4 Boost (step-up) converter where two main semiconductor components are the transistor and the rectifying diode. Current power supply technology necessitates designs (not included in the schematic) to account for the limited capability of Si-based transistors and rectifiers.

+

–

L

+

C FET –

VGS = 0, Von

+

–

Schottky barrier diode

Power MOSFETs and diodes

415

The low-voltage power device market is addressed by silicon-based MOSFET with operation up to 100 kHz. The primary approach is double-diffused MOSFET that can address from 10 to 500 V. In this vertical design, the on-resistance is determined by the thickness of the epitaxial n-type drift region. A 30 V device has an on-resistance of a few mΩ, while at 500 V the on-resistance is several Ω. To address this high-resistance issue, super junction MOSFETs were developed that operate from 500 V to 1 kV. This device reduces its on-resistance by defining a vertical conduction path with deep p-type epitaxial layers—although this step raises the cost of the device [51]. The medium-power/medium-voltage market (400 V to 10 kV) is primarily served by silicon-based devices such as insulated-gate bipolar-transistor (IGBT) and integrated gate-commuted thyristors (IGCT). These technologies deliver approximately 6.5 kV blocking voltage and require multi-level inverter to operate at 4160 V AC. The inherent low switching frequency (98.5% for output of 30%–100%. Twelve 83 kW booster circuits were combined to create a 1-MW power conditioning system. Commercial SiC MOSFET switches and Schottky diodes are available up to 12 kV at 20 kHz. Even higher voltages of 15 kV (at 5 kHz) are currently addressed with SiC IGBT switches and SiC PiN diodes [54]. Silicon Carbide Schottky barrier diodes are commonly used in power factor correction circuits and IGBT power modules. SiC power devices are used in the US Navy DDG 1000 Zumwalk Class destroyer to supply 78 MW at 4160 V. Similarly, more-electric fighter aircraft have eliminated heavier hydraulic and pneumatic system with electric switching systems [55].

17.6

Future

Gallium oxide possesses material properties that are advantageous for power electronic devices. Combined with the availability of native Ga2O3 substrates, Ga2O3based devices are projected to supplant Si-based as well as GaN- and SiC-based power electronics over a certain slice of the high-voltage, low-frequency application space. The β-type of Ga2O3 is clearly the most developed polymorph of Ga2O3. Devices based on β-Ga2O3 have not reached below the expected 8 MV/cm theoretical value for breakdown voltage. The best Schottky barrier diodes based on β-Ga2O3 have achieved a breakdown strength of only approximately 1 MV/cm. Recent results in Schottky diodes fabricated from 0.43 μm mist grown α-Ga2O3 displayed a breakdown strength of 12 MV/cm [5]. For transistors, a 3.8 MV/cm breakdown strength was obtained for a 0.65-μm channel length device. The best reported breakdown voltage of 755 V only corresponds to 0.5 MV/cm breakdown strength that was obtained for a 15-μm channel length device. As material quality improves, it is expected that the experimental breakdown voltages will approach the theoretical predictions. A clear research path is that large-area and vertical devices will mirror the designs of silicon and silicon carbide power devices particularly with an emphasis on vertical design with defined conduction paths. The relatively low thermal conductivity of Ga2O3 creates self-heating effects that must be mitigated in order to utilize Ga2O3 for high-power at even moderate switching frequencies. Lastly, the current state of p-type β-Ga2O3 is currently an issue to utilized pn junctions to define the current path in vertical devices. This issue, like many aspects of Ga2O3, is rapidly evolving.

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[32] M. Oda, J. Kikawa, A. Takatsuka, R. Tokuda, T. Sasaki, K. Kaneko, S. Fujita, T. Hitora, 73rd Annual Device Research Conference (DRC), 137, 2015. [33] E. Chen, A. Leng, Power MOSFET Switching Loss Precise Analysis, NIKO Semiconductor F020810M Report, 2010. [34] N. Mohan, T.M. Underland, W.P. Robbins, Power Electronics, second ed., John Wiley & Sons, New York, NY, 1995. [35] R.W. Erickson, D. Maksimovic, Fundamentals of Power Electronics, second ed., Kluwer Academic Publishers, Norwell, MA, 2001. [36] M. Higashiwaki, K. Sasaki, A. Kuramata, T. Masui, S. Yamakoshi, Appl. Phys. Lett. 100 (1) (2012) 013504. [37] M. Higashiwaki, K. Sasaki, T. Kamimura, M.H. Wong, D. Krishnamurthy, A. Kuramata, T. Masui, S. Yamakoshi, Appl. Phys. Lett. 103 (12) (2013)123511. [38] M.H. Wong, K. Sasaki, A. Kuramata, S. Yamakoshi, M. Higashiwaki, IEEE Electron Device Lett. 37 (2) (2016) 212. [39] J. Kim, S. Oh, M.A. Mastro, J. Kim, Phys. Chem. Chem. Phys. 18 (2016) 15760. [40] M.J. Tadjer, N.A. Mahadik, V.D. Wheeler, ECS J. Solid State Sci. Technol. 5 (9) (2016) 468. [41] J. Kim, M.A. Mastro, M.J. Tadjer, J. Kim, ACS Appl. Mater. Interfaces 9 (25) (2017) 21322. [42] A.J. Green, K.D. Chabak, E.R. Heller, R.C. Fitch, M. Baldini, A. Fiedler, K. Irmscher, G. Wagner, Z. Galazka, S.E. Tetlak, A. Crespo, K.D. Leedy, G. Jessen, IEEE Electron Device Lett. 37 (2016) 902. [43] K.D. Chabak, N. Moser, A.J. Green, D.E. Walker Jr, S.E. Tetlak, E. Heller, A. Crespo, R. Fitch, J.P. McCandless, K. Leedy, M. Baldini, G. Wagner, Z. Galazka, X. Li, G. Jessen, Appl. Phys. Lett. 109 (2016) 213501. [44] G. Wagner, M. Baldini, D. Gogova, M. Schmidbauer, R. Schewski, M. Albrecht, Z. Galazka, D. Klimm, R. Fornari, Phys. Status Solidi A 211 (1) (2014) 27. [45] M. Baldini, M. Albrecht, A. Fiedler, K. Irmscher, D. Klimm, R. Schewski, G. Wagner, J. Mater. Sci. 51 (7) (2016) 3650. [46] Z. Galazka, K. Irmscher, R. Uecker, R. Bertram, M. Pietsch, A. Kwasniewski, M. Naumann, T. Schulz, R. Schewski, D. Klimm, J. Cryst. Growth 404 (2014) 184. [47] Z. Hu, W. Li, K. Nomoto, M. Zhu, X. Gao, M. Pilla, D. Jena, H.G. Xing, Device Research Conference, (2017) pp. 1–2. [48] Z. Hu, K. Nomoto, W. Li, L.J. Zhang, J.-H. Shin, N. Tanen, T. Nakamura, D. Jena, H.G. Xing, Ext. Abstr. Device Research Conf., 2017. [49] M.H. Wong, K. Goto, A. Kuramata, S. Yamakoshi, H. Murakami, Y. Kumagai, M. Higashiwaki, Ext. Abstr. Device Research Conf., 2017. [50] K. Sasaki, D. Wakimoto, Q.T. Thieu, Y. Koishikawa, A. Kuramata, M. Higashiwaki, S. Yamakoshi, IEEE Electron Device Lett. 38 (2017) 783. [51] ROHM Semiconductor Power Devices, NE Handbook Series, Nikkei Business Publications, 2012. [52] S.H. Ryu, S. Krishnaswami, M. O’Loughlin, J. Richmond, A. Agarwal, J. Palmour, A.R. Hefner, IEEE Electron Device Lett. 25 (2004) 556. [53] Y. Furusho, K. Fuji, IEEE 9th International Conference on Integrated Power Electronics Systems, Proceedings of, 2016. [54] A. Hefner, R. Singh, J. Lai, D.W. Berning, S. Bouche, C. Chapuy, IEEE Trans. Power Electron. 3 (2001) 273. [55] J.G. Kassakian, T.M. Jahns, IEEE J. Emerg. Sel. Top. Power Electron. 1 (2) (2013) 47.

18

Ga2O3-photoassisted decomposition of insecticides

Hisao Hidaka, Tohru Tsukamoto Department of Chemistry, Meisei University, Hodokubo, Hino-shi, Tokyo, Japan

Chapter Outline 18.1 Introduction 419 18.2 Photochemical and photocatalytic degradation reactions 18.2.1 18.2.2 18.2.3 18.2.4 18.2.5 18.2.6 18.2.7 18.2.8

421

Photochemical reactions 421 Effect of oxygen on photodegradation 421 Charge separation by UVA and UVC irradiation 422 Loading effect of the metal-oxide photocatalyst 426 pH dependence 427 Temperature dependence 427 Concentration dependence 428 Effect of light intensity 428

18.3 Toxicity test of photodegraded intermediate products 428 18.4 Photocatalytic mineralization of chemical substances 428 18.4.1 18.4.2 18.4.3 18.4.4

Photocatalytic degradation of miscellaneous organic substances with gallium oxide 428 Photocatalytic degradation of insecticides with gallium oxide 430 The degradation rate for insecticide in mixed aqueous/organic media 432 The tendency of defluorination based on chemical structures 433

18.5 Concluding remarks Acknowledgments 434 References 434 Further reading 437

18.1

434

Introduction

Recent years have witnessed an ever increase in the production of many insecticides that contain fluorine and/or trifluoromethyl groups [1]. These fluorinated chemicals have excellent physicochemical characteristics such as, adsorption, hydrophobicity, thermal chemical stability, and bioactivity [2]. Accordingly, they have attracted intense investigations and have been applied in various fields [3]. Reports have appeared regarding contamination of common insecticides, health risks, and pollutant-treatment systems. However, scant research is being carried out on the decomposition of fluorinated substances. The perfluorinated chemicals such as perfluorooctanoic acid (PFOA) and the perfluorooctanesulfonate (PFOS) surfactant Gallium Oxide. https://doi.org/10.1016/B978-0-12-814521-0.00018-X © 2019 Elsevier Inc. All rights reserved.

420

Gallium Oxide

were produced by DuPont in the United States in 1947. Seven decades have elapsed since their developments [4]. PFOA and PFOS have also been used in firefighting water sprays in great amounts resulting in the worldwide contamination of aquatic environments. Because the quantities used were low in the beginning, people failed to recognize the potential hazards of these perfluorinated chemicals. Yet recently the production has increased considerably with its inevitable serious ecological and environmental results caused by the use of these fluorinated chemicals. As one specific example, eggs in the European Union have been contaminated with fipronil [5], which is suspected to cause severe health risks to the elderly and interference in the functioning of the brain, thyroid gland, pineal, kidney, and reproductive system in animals and human beings [6]. Drinking water contaminations by fluoride systems are now under reevaluation by the United States Environment Protection Agency [7–11]. Pharmaceuticals present conflicting faces of positive action [12] and negative after effects for the risk of heart attacks and strokes regarding the cholesterol ester transfer protein (CETP). Fluorinated medicinals are statistically evaluated for medical effects. Fundamentally, a fluorinated medicinal is dosed at specified bioactivities for anticancer treatments. They are excreted from the human body, but biodegradation through bacteria cannot be undertaken. Although their amount used was slight in the beginning of utilization, the original increase in the number of fluorinated drugs has gradually increased. On the whole, pharmaceuticals have analogous structure to agrochemicals. Detoxification procedures have been tried for wastewater remediation. The degradation of agrochemicals and pharmaceuticals is a challenging task. There are difficult cases even by the active sludge wastewater treatment through bacteria. To the extent that agrochemicals and pharmaceuticals have bioactive properties, toxicity to bacteria is not to be underestimated. Additionally, since halide bonds involving fluorine and chlorine present significant bond energies, their degradation present serious difficulties. The halide, by itself, is a poisonous material. The fluorine groups dCF3 and dF are introduced in the heteroaromatic skeleton because of the unique specificity of fluorinated activities toward adsorption. Hence, cleavage of the CdF bond is generally difficult even by the oxidative power presented by the valence band holes and/or related oxidative species photogenerated by UV-irradiated TiO2. Not to be outdone, many agrochemicals bearing complex skeletons have been developed: for instance, triazine skeletons, organic phosphorous, anilite, thiocarbonate, phenylurea, phenyl azide, carbamate, and sulfonylurea among others, all of which contain atoms of nitrogen, sulfur, phosphor, and others [2, 5]. Agrochemical insecticides bearing fluoride functions are widely sprayed or scattered on various plants to deter parasites and vermin. Furthermore, defluorination of fluorinated agrochemicals with gallium oxide as a potential photocatalyst presents a possible remediation/treatment of wastewaters contaminated with fluorinated substances.

Ga2O3-photoassisted decomposition of insecticides

18.2

421

Photochemical and photocatalytic degradation reactions

18.2.1 Photochemical reactions Since the wavelengths 200–800 nm nearly correspond to the bond dissociation energy of organic substances, a variety of photochemical reactions can take place. The illuminated substrate converts to the excited singlet state and then to the low-lying triplet state. The energy is emitted as either heat or light. The Jablonski scheme shows the schematic illustration on the light excitation mechanism as well as the decay process of the excited states. Light is absorbed by a substance allowing an electron in a lower orbital level to be elevated to a higher energy level. Consequently, various photoreactions can take place such as bond cleavage, isomerization, ring opening, ring cyclization, addition, substitution, oxidation, and reduction. The type of fission reactions that occur during the photodegradation is well known; it is summarized in Fig. 18.1 [13, 14]. Under photochemical nitrogen-purged and airequilibrated conditions, an insecticide can undergo various photochemical reactions via direct photolysis (viz., homolysis, heterolysis, and photoionization) [15].

18.2.2 Effect of oxygen on photodegradation Oxygen molecules cause oxidation through the excited singlet state and superoxide radical anion as demonstrated in the following equations. O*

O

hn

C

O

C

C

C6H5

C6H5

C6 H5

O C C6H5

H

Norrish type I reaction H

H O C C6H5

O

CH2

H

CH2

O

C

C

C6H5

C6H5 O C C 6 H5

Norrish type II reaction

Fig. 18.1 Radical photocleavage by Norrish type I and II reactions.

CH2

422

Gallium Oxide

Reaction of oxygen under UV light in the presence and absence of a photocatalyst. O2 + hv ! 1 O2 1

O2 + insecticide ! oxidized products

(18.1) (18.2)

O2 + e ðcatalystÞ ! O2  

(18.3)

O2   + H + ! OOH

(18.4)

Oxygen molecules are also highly reactive toward other radicals, and they play an important role in photodegradations. Since oxygen molecules have electronwithdrawing properties, they capture excited electrons generated by photocatalysis. The oxygen molecules and excited electrons react to form superoxide anions, which then interact with hydrogen ions to form hydroperoxyl radicals having stronger oxidizing power.

18.2.3 Charge separation by UVA and UVC irradiation This section clarifies the difference between gallium oxide under UVC irradiation and the conventional titanium dioxide under UVA irradiation. When a metal oxide is exposed to light energy that exceeds the bandgap, charge separation occurs resulting in the oxidation of a substrate by the valence band holes and to reduction of a substance by the conduction band electrons. The bandgap of TiO2 is 3.2 eV (387 nm) so that light irradiation at wavelengths below 387 nm will photogenerate these charge carriers [16]. A high-pressure mercury lamp is suitable for UVA light irradiation suitable to activate the titanium dioxide. On the other hand, the band gap of Ga2O3 is about 4.8 eV (absorption edge is at ca. 258 nm) requiring higher energy to achieve charge separation [16]. Therefore, a low-pressure mercury lamp that emits UVC light around 254 nm is used. Unlike the CdF bond in fluorinated organics, the CdC, CdH, and CdCl bonds can be cleaved by UVC light. However, under UVC irradiation gallium oxide is effective toward the defluorination and the mineralization of fluorinated insecticides. As a potential photocatalyst, gallium oxide has been examined for its application in wastewater treatments, particularly for the defluorination of fluorinated substances [16]. To understand the photocatalytic activity of gallium oxide, we examine now what kind of influence UVC emitted light from the low-pressure mercury lamp may have. This lamp mainly emits light at both 254 and 185 nm. However, the absorption of vacuum ultraviolet light (VUV; 185 nm) by air molecules significantly attenuates the radiation reaching the reactor. Thus, special equipment has to be used. In addition, a vessel made of heat-resist glass cannot pass light shorter than 300 nm, so that such a reactor cannot be used for photodegradation processes occurring with gallium oxide. A quartz vessel that is transparent above 150 nm is therefore utilized as the photoreactor.

Ga2O3-photoassisted decomposition of insecticides

423

Since the light wavelength of 254 nm cannot cleave water, no OH radicals are formed with the radiation provided by the low-pressure mercury lamp [17]. The dissociation energy of HdOH is ca. 471 kJ/mol [18]. Therefore, an energy of ca. 26 kJ/mol is deficient in this respect. In addition, the light from the low-pressure mercury lamp does not produce ozone from molecular oxygen in air. Concerning the bond dissociation energy of fluorinated substances, the CdF bond energy of FdCHFCl structure is 460 kJ/mol [18]. Consequently, the fission of this CdF bond is possible even with the UVC radiation from the low-pressure lamp. However, the CdF bond dissociation energy of almost all chemicals having a fluorine moiety lies in the range of 480–510 kJ/mol [18]. Therefore, no decomposition of CdF bonds in fluorinated organics occurs only by UVC irradiation. It is common knowledge that gallium oxide has several crystal types. In gallium oxide, the valence band is constituted of O-2p orbitals. The conduction band is composed of gallium atom 4s4p hybrid orbitals. Since Ga2O3 has a higher bandgap, the redox potential is higher than that of titanium dioxide. Hence, the photogenerated hole and excited electron on the surface of Ga2O3 possess higher oxidation and reduction power. Therefore, gallium oxide is expected to efficiently degrade organic substances attached to the catalyst surface; the elimination efficiency is higher in comparison with the TiO2. The bandgap of each crystal type of gallium oxide is different (4.5–5.0 eV). The Ga2O3 crystal has α-type, β-type, γ-type, and others of which the photocatalytic activities have been previously evaluated for the photocatalytic degradation of organic substances as summarized in Table 24.1, which shows the specific hydrocarbon removal rate and CO2 production rate in a dry stream UV irradiation through the use of α-, β-, γ-gallium oxide, and titanium dioxide. It can be seen from this table that Ga2O3 exhibits a higher photoactivity than TiO2 [19–21]. The β-type Ga2O3 crystal exhibits the highest photoactivity among the α-, β-, and γ-type crystallites. The α-type crystallite has the second highest activity. It follows from this that β-gallium oxide is a better photocatalyst for the decomposition of organic substances. It is a well-known fact that the bandgap of β-Ga2O3 is approximately 4.8 eV. The valence band potential of β-Ga2O3 is 7.75 eVAVS (AVS: absolute vacuum scale), l

Comparison of photocatalytic activities for aromatic substances by different crystals of gallium oxide and titanium dioxide (P-25)

Table 24.1

α-Ga2O3

β-Ga2O3

γ-Ga2O3

TiO2

Hydrocarbon removal rate [μmol h1 m2]

Benzene

0.32

0.42

0.21

0.08

Toluene

0.46

0.52

0.26

0.15

Ethylbenzene

0.31

0.36

0.20

0.18

CO2 production rate [μmol h1 m2]

Benzene

1.7

2.4

0.95

0.17

Toluene

1.3

1.8

0.85

0.13

Ethylbenzene

1.2

1.6

0.73

0.15

424

Gallium Oxide

Fig. 18.2 Speculation models of the simplified events occurring in the photooxidative and photoreductive degradation on the gallium oxide surface illuminated with UVC light.

whereas that of TiO2 is 7.41 eVAVS. On the other hand, the conduction band potential of β-Ga2O3 is at 2.95 eVAVS and that of TiO2 is at 4.21 eVAVS [22]. Since the conduction band potential of gallium oxide is much higher than that of titanium dioxide, the reaction between substances and excited electrons is more important for the photodegradation processes. As described earlier, oxygen acts as a scavenger of the excited electrons. For this reason, it is necessary to delineate gallium oxide photocatalytic reactions into those that occur in the presence of oxygen and those that are not. Fig. 18.2 depicts a graphical representation of the reactions that take place on the surface of gallium oxide in both the presence and absence of oxygen. These various reactions are classified into those that occur in the presence of oxygen (reactions 18.5–18.10) and those that occur in oxygen-free (reactions 18.11–18.14) conditions. [I] Photoassisted Ga2O3 reaction in the presence of oxygen Insecticide + h + ðVBÞ ! oxidized intermediates

(18.5)

ðO2 Þabs: + e ðCB ! ðO2  Þabs:

(18.6)

H2 O + h + ðVBÞ ! OH + H +

(18.7)

ðO2  Þabs: + H + ! OOH

(18.8)

Insecticide +  OH ! Hydroxylated intermediates

(18.9)

Insecticide +  OOH ! Peroxidized intermediates

(18.10)

[II] Photoassisted Ga2O3 reaction in the absence of oxygen Insecticide + e ðCBÞ ! reduced intermediates

(18.11)

Ga2O3-photoassisted decomposition of insecticides

425

Insecticide + h + ðVBÞ ! oxidized intermediates

(18.12)

H2 O + h + ðVBÞ ! OH + H +

(18.13)

Insecticide +  OH ! Hydroxylated intermediates

(18.14)

As described earlier, a large numbers of oxygen molecules are adsorbed onto the surface of gallium oxide under an oxidative atmosphere. When charge separation occurs on the surface of a photocatalyst, oxygen molecules are reduced by the excited electrons in the conduction band. Consequently superoxide radical anions are generated (18.6) and most of the excited electrons are consumed by the large number of oxygen molecules under an oxidative atmosphere. The hydroperoxyl radical, OOH, was generated from the reaction of the superoxide anions with protons (18.8). When insecticides are oxidized by the hydroperoxyl radical, peroxo intermediates are generated (18.10). On the other hand, water molecules are oxidized by the valence band holes (18.7), yielding OH radicals and protons. These hydroxyl radicals can react with the insecticides and produce hydroxylated intermediates (18.9). In addition to the oxidizing radicals, oxidation of the insecticides directly by the valence band holes is not precluded (18.5). There is a significant difference of the reaction on the conduction band between an oxygen-free atmosphere and an oxidative atmosphere. Reduced by the excited electrons (18.11), the insecticide releases an elimination group (E), for example a fluoride ion, followed by hydrogenation of the residual radical or alternatively the possible formation of a double bond. Previous research has suggested that as a catalyst gallium oxide can transform propane into propene through a dehydrogenation process [23]. The valence band holes oxidize water under oxygen-free conditions (18.13); the OH radical degrades the insecticide (18.14). The subsequent sequential reactions to (18.12) under oxygen-free conditions are the same as those that occur under oxygenated conditions [compare reactions (18.13) and (18.14) versus reactions (18.7) and (18.9)]. The Langmuir-Hinshelwood (L-H) expression is often erroneously used [24] to suggest that if the data could be fitted to the L-H model [see Eq. (18.15)], then the photoreaction likely takes place on the surface of the metal oxide. Nonetheless, application of the L-H expression (18.15) to the photoassisted defluorination of two fluorinated substrates, namely fluoxetine and fluvoxamine, in the presence of Ga2O3 could provide a qualitative comparison of the processes between the two substrates (see Fig. 18.3): 1 1 1 1 ¼  + ri kKLH ci k

(18.15)

where ri is the initial rate of photodegradation in the heterogeneous Ga2O3 dispersion, k is an apparent rate constant, ci is the initial concentration, and KLH is the apparent L-H adsorption constant.

426

Gallium Oxide

1/rate (µM/min)–1

2.0 1.6

FLX

1.2 0.8

FOM

0.4 0.0 0.00

0.02

0.04

0.08

0.06

0.10

1/Ci (µM)–1

Fig. 18.3 Langmuir-Hinshelwood (L-H) plots of the photoassisted defluorination of fluoxetine and fluvoxamine maleate solution (100 mL) at various initial concentrations in the presence of Ga2O3 particles (loading 50 mg) under UV irradiation in an inert nitrogen atmosphere. H2N H N

O

O

N CF3

CF3

O

(I) Fluoxetine (II) Fluvoxamine

The apparent defluorination rate constant (k) and the L-H apparent adsorption constant (KLH) for FLX were k ¼ 3.01 μM min1 and KLH ¼ 0.026 μM1, whereas for FOM, k ¼ 3.37 μM min1 and KLH ¼ 0.14 μM1. Evidently, it would appear that the defluorination of fluvoxamine (II) was, within experimental error, rather similar to that of fluoxetine (I); fluvoxamine appears to be adsorbed more strongly onto the surface of the metal oxide [25].

18.2.4 Loading effect of the metal-oxide photocatalyst The optimal added amount is influenced by various experimental parameters such as particle size and surface charge. However, most of the parameters have a common effect on these experiments. Therefore, the general tendency is described next. When the catalyst amount is considerably low, the degradation yield tends to increase with an increase of the catalyst loading (weight). The degradation attains a maximum efficiency at a certain loading, and then the efficiency tends to plateau with further increase in loading, probably because a larger quantity of the photocatalyst inhibits light transmission through the heterogeneous dispersion; most of the catalyst is not exposed to UVC irradiation [26–28]. Typically, an increase in surface area of the catalyst leads to an increase in the degradation efficiency. Various procedures to increase the surface area have been

Ga2O3-photoassisted decomposition of insecticides

427

published [29]. However, to the extent that photocatalyst particles in aqueous solution tend to aggregate causes the active surface area to decrease. Therefore, the catalyst with the higher surface area is not necessarily the more active one. This calls attention to carry out preliminary experiments so as to optimize catalyst.

18.2.5 pH dependence The pH of the dispersion has an influence on the efficiency of the photocatalytic degradation of organic pollutants. Varying the pH lead to the changes of the photocatalyst’s surface such as the surface charge, and the aggregation state among others. It has been reported extensively [2–17] that the pH of most solutions of organic substances lie in the acidic region (from 4.5 to 7). In this pH range, the surface potential of the metal-oxide photocatalyst is near zero. The isoelectric point is the point of zero charge (PZC). The contact between an organic compound and a photocatalyst is minimized at this PZC. When the proton concentration is below the PZC in the solution, the charge potential of the catalytic surface is positive. Therefore, negatively charged compounds are apt to adsorb strongly on the photocatalyst in this low pH range [30, 31]. On the other hand, when the pH is higher than PZC, a positively charged insecticide can readily make contact with the negatively charged surface. As a result, the photocatalytic degradation is likely enhanced. Since the photocatalytic degradation of pollutants likely occurs on the catalytic surface, the degradation efficiency is then modulated by the pH. In addition, the solution to be degraded commonly becomes more acidic with the passing of irradiation time [32]. High pH conditions have been shown to be effective at the photodegradation. At high pH, a large number of hydroxide ions adsorb on the catalytic surface of the metal oxide [33], which could result in a larger number of hydroxyl radicals to be generated simultaneously. Oxygen captures the electrons at the photocatalyst surface to generate superoxide radical anions, through which various substrates are hydroxylated and/or oxidized. The excited electrons can also reduce organic substrates. In this manner, the degradation is effective at high pH. However, the inhibitory effect of impurities like counter ions can in some cases be documented [30].

18.2.6 Temperature dependence It is thought that the temperature effect on Ga2O3 catalyst is similar to that of TiO2 catalyst [34]. At high temperature, recombination of valence band holes and excited electrons on the catalyst surface is enhanced, thereby hindering charge separation [26]. Additionally, adsorption of a substance onto the photocatalyst also decreases at high temperature. Another point to consider is the desorption of intermediate products formed with the photocatalyst, which is the rate-determining step at high temperature. Adsorption occurs between 20°C and 80°C. Below 0°C, the increase in the activation energy is confirmed. Therefore, the range from 20 °C to 80 °C is generally preferred for a photocatalytic reaction [35].

428

Gallium Oxide

18.2.7 Concentration dependence The photocatalytic degradation efficiency is dependent upon the concentration of contaminants. The complete mineralization of a concentrated solution requires a longer reaction time. Since the excess organic substances adsorb on the surface of the catalyst, the degradation efficiency is generally lower at a high concentration of insecticide solutions [36].

18.2.8 Effect of light intensity The light intensity also affects the efficiency in the photodegradation of contaminants. However, there are several reports on the photocatalytic activity which have shown that process is independent of light intensity because a small number of photons is sufficient to lead to charge separation for the surface reaction [37].

18.3

Toxicity test of photodegraded intermediate products

The need for verification of intermediates which are toxic assumes greater importance for the photodegradation system. To our knowledge, there have been a few reports about identifying whether or not the intermediates formed are toxic. Furthermore, some naturally degraded intermediates are no doubt more toxic than the original agrochemicals. Our earlier studies have demonstrated that photodegraded intermediates generated in a gallium oxide/UVC system were not toxic. The procedure employed to check carcinogenicity is the Ames screening test [38–40]. The toxicity tests for the degraded intermediates of fipronil and flubendiamide were examined using this test [41, 42]. If the result of the Ames test is positive, the possibility of cancer-causing materials is estimated to be high. The tested method is for the photodegraded samples using the mixed organic aqueous solvent. Since the common insecticide is insoluble in water, a mixed solvent is employed. The organic solvent used in the Ames assay is acetone or dimethyl sulfoxide, which have no influence on the Salmonella bioactivity. Consequently, the intermediate products were no more toxic or mutagenic than the starting insecticides.

18.4

Photocatalytic mineralization of chemical substances

18.4.1 Photocatalytic degradation of miscellaneous organic substances with gallium oxide Although gallium oxide as a photocatalyst has drawn intense research because of its potential application in refractory wastewater treatment, little reports have been published on the degradation of insecticides. Model compounds such as ethylenediaminetetraacetic acid (EDTA), phenol, PFOA, PFOS, rhodamine-B, methyl orange,

Ga2O3-photoassisted decomposition of insecticides

429

methylene blue, 4-chlorophenol, benzene, toluene, and ethylene were investigated using gallium oxide as the photocatalyst. The overview of these experimental conditions and results are described below. Results of the photocatalytic degradation of EDTA have been described by Seshadri et al. [22]. In these experiments, the photocatalytic activity of decomposition of EDTA by β- and γ-gallium oxide and titanium dioxide was compared. β- and γ-gallium oxide was synthesized from gallium metal. The prepared photocatalysts were characterized by X-ray diffraction, differential thermal analysis, and Raman spectral measurements. A cylindrical photoreactor was surrounded with four lowpressure mercury lamps (16 W). Experimental conditions were maintained at ambient temperatures of 25–30°C with a cooling fan. One liter of reaction solution that included the photocatalyst (3 mg), 0.5 mL of 30% H2O2 and 0.1% EDTA. Most EDTA could be degraded within 60 min. On the other hand, in a catalyst-free/H2O2 system, degradation efficiency after 180 min was only 50%. Their catalytic activities decreased in the following order: β-gallium oxide/H2O2 > γ-gallium oxide/ H2O2 >titanium dioxide/H2O2 > H2O2. The degradation efficiency of EDTA was higher at pH 10 than it was under neutral conditions (pH 7). A comparison of the photocatalytic activity of γ- and β-gallium oxide and titanium dioxide showed that γ-gallium oxide is effective for the decomposition of 1,4-dioxane [43]. Over 90% of the 1,4-dioxiane (1 g/L) was decomposed by γ-gallium oxide (10 mg/L)/H2O2 (0.5 mL) at the end of 180 min of irradiation. According to a report on the photodegradation of 4-chlorophenol with α-gallium oxide, the degradation efficiency was approximately 98% after 2 h at pH 7.8 [19]. The 4-chlorophenol solution (20 mg/L, 200 mL) was dosed with α-gallium oxide (0.4 g/L); then the mixture was exposed to UVC irradiation provided by a 15-W low-pressure mercury UVC lamp placed in the center of the reactor. Under the same conditions, the extent of degradation of the phenol after 2 h of irradiation by gallium oxide was 98% whereas only 81% was observed with titanium dioxide. There is one reference on the degradation of model compounds by gallium oxide microspheres [44]. In this chapter, gallium oxide microspheres were prepared by the surfactant-assisted hydrothermal method. The gallium oxide microspheres were characterized by scanning electron microscopy, X-ray diffraction and transmission electron microscopy, Fourier-transform infrared spectroscopy, UV-vis spectroscopy, and surface area analysis. The photocatalytic degradation activity of gallium oxide microspheres of rhodamine-B and methylene blue in aqueous solution was investigated. For rhodamine-B (50 mL, 0.02 mM) or methylene blue (50 mL, 0.02 mM) solutions were dosed with the photocatalyst (50 mg). These solutions were irradiated with a 150-W xenon lamp. Decomposition of methylene blue and the mineralization of rhodamine-B occurred within 90 min and 180 min, respectively. Tri-n-butyl phosphate was degraded after 40 min using nanostructured β-gallium oxide [21]. The concentration of a tri-n-butyl phosphate solution was 400 ppm (1000 mL), onto which were added 10 mg of nanostructured β-gallium oxide and 0.5 mL of 30% H2O2. The reactor was irradiated with UVC light from four lowpressure Hg lamps. In the case of titanium dioxide, the decomposition of tri-n-butyl phosphate needed 70 min.

430

Gallium Oxide

Water treatment by gallium oxide has been widely studied to degrade perfluorinated substrates, in particular PFOA [4, 45, 46]. In one study, the degradation ratio of PFOA (con. 40 mg/L) by synthesized β-gallium oxide (0.5 g/L) was 38% after 4 h [45]. This experimental solution was irradiated with UVC light from a lowpressure mercury lamp under a nitrogen atmosphere. Degradation yield under different atmospheres drew a parallel between N2 (ca. 38%), air (ca. 31%), and O2 (700°C), oxygen defects in the Ga2O3 are in equilibrium with the surrounding ambient, and oxygen vacancies quickly diffuse out from the interior to the surface and vice versa [53]. Once the oxygen concentration is reduced at high temperatures, oxygen vacancies are created in the Ga2O3 film and electrons are released owing to the diffusion of oxygen out of the film, resulting in a decrease in resistance. Because other reducing gases can also involve in this equilibrium reaction at low temperature, an operating temperature above 900°C is required for the complete selectivity of oxygen sensing.

19.2.3 Ga2O3-based CO sensors 19.2.3.1 Pt and La0.8Sr0.2FeO3 (LSFO) for CO sensing Carbon monoxide is one of the most common hazardous gases produced by combustion processes in the automotive and stationary energy industries. Accurate monitoring and precise controlling of CO during the reaction is essential to reduce its emission and to obtain high efficiency in energy conversion [2, 3, 54]. Wide-bandgap β-Ga2O3 is a semiconductor metal oxide suitable for high-temperature applications in harsh environments, owing to its thermal and chemical stabilities [2, 4, 55, 56]. Generally, Pt on Ga2O3 is used as a functionalizing material for the catalytic reduction of CO. However, the replacement of high-cost noble metal is researched to reduce or eliminate the amount of Pt required for gas sensor fabrication. Recently, Lin et al. reported sensors based on the perovskite oxide, LSFO decorated Ga2O3 nanorods for hightemperature CO detection [2]. Ga2O3 nanorods were grown by a hydrothermal method. Vertically aligned GaOOH nanorods were grown on SnO2 seed layer with Ga(NO3)3 solutions at 150°C for 12 h. After subsequent annealing at 1000°C for 4 h, 2-μm long Ga2O3 nanorods with diameters of 100–300 nm were obtained. LSFO or Pt nanoparticles were decorated on Ga2O3 nanorods for CO sensing at 500°C. The sensitivity of the devices for CO gas exposure was measured, which is defined as the value of resistance under pure N2 over that under a mixture of CO and N2. For 100 ppm CO, the sensitivities were 8, 70, and 95 for pristine, LSFO, and Pt functionalized devices, respectively. On a bare Ga2O3 surface at 500°C, CO molecules react with pre-adsorbed O– ions, produce electrons, and decrease the resistance of the Ga2O3 nanorod array [2]. For a Pt and Ga2O3 heterostructure, the spillover process occurs [2, 5, 57]. CO molecules are adsorbed on the Pt surface, and react with pre-adsorbed

Ga2O3-based gas sensors

455

O– ions with reduced activation energy through a catalytic path resulting in enhanced reduction in the resistance of Ga2O3. For the LSFO and Ga2O3, the spillover process was confirmed by broadening of the O 1 s peak from XPS analysis after CO gas exposure [2]. In addition, LSFO is a p-type semiconductor material. Hence, the depletion region of Ga2O3 by the built-in potential at the interface with LSFO is greater than for the pristine Ga2O3 depleted by just O– ions. After reduction process of CO with O– ions, the depletion region might be reduced more than the pristine Ga2O3 surface that has fewer adsorbed O– ions, and thus the electron channel of Ga2O3 with LSFO is further recovered, resulting in large resistance change. Additionally, the perovskite oxide is known to have a considerable oxygen storage capability in the lattice [2, 58, 59].

19.2.3.2 UV-enhanced CO sensor The sensitivity of G2O3-based gas sensors can be enhanced by illumination of UV light on the device while detecting the target gas. Lin et al. investigated the effect of UV radiation on CO sensing of bare, LSFO, or Pt-decorated Ga2O3 resistive sensors [1]. LSFO perovskite oxide or Pt nanoparticles were deposited on β-Ga2O3 nanorod arrays by rf magnetron sputtering. A sensor with 1  0.5 cm2 exposed sensing area was installed in the middle of a tube furnace, and the device was heated to 500°C with a ramp rate of 20°C/min. The resistance change upon 100 ppm CO exposure was monitored at the bias of 1 V with a constant flow rate of 1.51 L/min. UV light of wavelength 254 nm, which has higher energy than the bandgap of β-G2O3 (4.9 eV), was illuminated for the UV-enhanced sensing. The sensitivity was defined as the ratio of the resistance under N2 ambient to under 100 ppm CO. Under the UV illumination, the sensitivities were increased for all the devices compared to the non-UV condition. Improvements of 30%, 20%, and 50% in the sensitivity were achieved for the bare, LSFO, and Pt decorated sensors, respectively. Additionally, the response times were reduced for all cases under the UV illumination. CO molecules react with adsorbed oxygen ions on the β-Ga2O3 surface, and produce electrons, thereby decreasing the resistance of the sensor. Upon UV illumination, electron-hole pairs are generated, and this photocurrent increases the density of adsorbed oxygen ions reacting with CO molecules. Therefore, the illumination of UV light with higher energy than the bandgap of β-Ga2O3 improves the sensitivity of the β-Ga2O3 based CO sensors. For the LSFO and Pt-decorated sensors, the sensitivities would be further enhanced by photosensitizing effect and spillover effect [1, 2]. UV illumination is one of the most effective ways to increase the sensitivity of Ga2O3-based gas sensors, but side reactions such as the decomposition of the target gas need to be considered under UV illumination condition, especially at high temperatures.

19.2.4 Ga2O3-based H2 sensors Semiconductor metal oxides are regarded as highly promising for hydrogen sensor applications, especially at high temperatures in harsh environments. Gu et al. made a comprehensive review of hydrogen gas sensors using metal oxide thin films and nanostructures [60]. The hydrogen sensing of metal oxide is believed to result from

456

Gallium Oxide

the decrease in resistance via the reactions of chemisorbed oxygen species with hydrogen. Fleisher et al. reported the potential of Ga2O3 thin films for sensing reducing gases at high temperatures [51, 61, 62]. In particular, a hydrogen sensor using a Pt Schottky diode on Ga2O3 thin film was first demonstrated by Trinch et al. [63, 64]. The authors showed that the Schottky barrier height of Pt/Ga2O3 diode sensor was effectively modulated with increasing hydrogen concentrations. Nakaomi et al. reported a field-effect hydrogen sensor using Ga2O3 thin film and Schottky diodes based on β-Ga2O3 single crystals, showing enhanced response to hydrogen and stable operation at elevated temperatures above 400°C [20, 65, 66]. The hydrogen sensing
characteristics of Pt Schottky diodes have been investigated using 201 and (010) β-Ga2O3 single crystals, as discussed in this section. For the device fabrication, Ohmic contacts were formed on β-Ga2O3 by Ti/Al, which was deposited by e-beam evaporation and patterned by a lift-off process, followed by annealing under N2 ambient in a rapid thermal annealing system. A SiNx passivation layer with 200 nm thickness was deposited for diode isolation by plasma-enhanced chemical vapor deposition. The windows for the active sensing area were opened by buffered oxide etching. The Schottky contact area was evaporated with 10-nm-thick Pt film, followed by Ti/Au contact pads for probing and wire
bonding. The I-V characteristics of the Pt Schottky diode sensors on 201 and (010) β-Ga2O3 were measured under the flammable limit of hydrogen (4%) balanced with nitrogen in a gas test chamber using an Agilent 4155C semiconductor parameter analyzer. Fig. 19.9 shows a schematic illustration and microscope image of the basic Schottky diode on Ga2O3, consisting of an Ohmic contact, Schottky metal, and a Si3N4 passivation layer. When the Schottky diode sensors are exposed to hydrogen, catalytic decomposition of the hydrogen molecules occurs on the Pt layer. The dissociated hydrogen atoms adsorbed on oxygen atoms on Ga2O3, thus forming dipole layers at the Pt/Ga2O3 interface. The Schottky barrier height is gradually reduced to ΦB, H2, resulting in the increased forward current, as illustrated in Fig. 19.10.

Contact pads

Pt

Ohmic window

Silicon nitride

57 Ga2O3 crystal

(A)

(B)

Fig. 19.9 (A) The illustrated schematic of Ga2O3-based Schottky diode, (B) optical microscope image of fabricated diode sensor.

Ga2O3-based gas sensors

457

Before exposure to hydrogen Upon exposure to hydrogen

ΔF B F B,N2

EC

F B,H2

EF

Pt

Ga2O3

EV

Fig. 19.10 Schematic band diagram of Ga2O3-based Pt Schottky diode for hydrogen exposure. 102

Current (mA)

100

Fig. 19.11 Current-voltage (I-V) characteristics of Pt Schottky diodes on (A) (010) and (B) (201) Ga2O3 single crystals before and after exposure to 500 ppm H2 in N2 at 25°C.

Before exposure After 500 ppm H2 exposure

10–2 10–4 10–6 10–8 10–10 –3

–2

(A)

0

1

2

3

Voltage (V) 102 0

10

Current (mA)

–1

Before exposure After 500 ppm H2 exposure

10–2 10–4 10–6 10–8 10–10 10–12 –3

(B)

–2

–1

0

1

2

Voltage (V)

Fig. 19.11 shows the I-V characteristics of the Pt Schottky diodes on (a) (010) and
(b) 201 Ga2O3 crystals measured under 500 ppm H2 in N2 at room temperature. Clearly, the forward current of both diodes increased in the sweeping bias range when exposed to hydrogen. The turn-on voltage was reduced in the forward bias region from 0.9 to 0.3 V for (010) Ga2O3 and from 0.7 to 0.2 V for (201) Ga2O3. The increased current and voltage shift occur owing to a decrease in the series resistance and the Schottky barrier height (SBH) due to the formation of hydrogen dipoles [67, 68].

458

Gallium Oxide


Fig. 19.12 shows the change and the recovery of SBH in both (010) and 201 Ga2O3 diodes after exposure to 500 ppm H2 in N2 and switching back to normal condition. The SBH changes (ΔΦB) were measured to be 0.09 and 0.27 eV for (010) and

201 Ga2O3 diodes, respectively. The higher SBH change of 201 Ga2O3 diode may
result from more hydrogen adsorption sites available on 201 surface owing to the higher density of oxygen dangling bonds [69]. As discussed in the previous chapter,
the oxygen atomic density on the 201 Ga2O3 surface is 1.34  1015 cm2, which is 1.5 times higher than that on (010) Ga2O3. A larger number of hydrogen dipoles can be formed on the Pt/Ga2O3 interface, leading to lowering the SBH. This indicates that the surface atomic configuration and atomic density for adsorption sites have a key role in determining the hydrogen sensing properties. The sensitivity, the relative current change as a percentage, is defined as Sð % Þ ¼

ðIH2  IN2 Þ  100% IN2

Schottky barrier height (eV)

where IH2 and IN2 denote diode currents measured in H2 and N2 ambient, respectively. The sensitivity increased as the forward bias increased to 0.7–0.8 V, and fell at higher
voltage than 1 V. The maximum value of the sensitivity for the 201 Ga2O3 diode 0.70 0.65 0.60

0.50

Schottky barrier height (eV)

(A)

(B)

ΔF = 0.09 eV

0.55

(010) Ga2O3

0

20

40

Time (min) 0.7 0.6 0.5 0.4

ΔF = 0.27 eV

0.3 0.2

(–201) Ga2O3

0

20

40

Time (min)

Fig. 19.12 Time dependence of the Schottky barrier height of Pt Schottky diodes on (A) (010) and (B) (201) Ga2O3 single crystals after exposure to 500 ppm H2 in N2 and switching back to N2 ambient at 25°C.

Ga2O3-based gas sensors

(010) Ga2O3

400

Current (nA)

459

H2 exposure (2 s) Applied voltage=0.8 V

300

40,000 ppm 30,000 ppm

Fig. 19.13 Time-dependent current changes of Pt Schottky diodes on (010) Ga2O3 single crystals as a function of H2 concentrations.

20,000 ppm

200

10,000 ppm 5000 ppm 3000 ppm 2000 ppm

100 0 0

2500

5000

7500

10,000

12,500

Time (s)

was 7.86  107%, which is slightly higher than that for (010) Ga2O3. Even though the
SBH reduction of 201 Ga2O3 is larger than that of (010) Ga2O3, the maximum sensitivities for both Ga2O3 crystals are similar, indicating that there exists another factor resulting in the oxide resistance change. The hydrogen atoms absorbed within Ga2O3 react with Ga2O3 and then release the electrons, thus reducing the oxide resistance. The slope of the forward I-V curve for the (010) Ga2O3 Schottky diode increased a
factor of two whereas that for 201 diode remained the same (not shown here). Fig. 19.13 shows the time dependence of current changes of the (010) Ga2O3 diode sensor with different hydrogen concentrations. A fixed voltage of 0.8 V was applied and the current changes were measured in response to repeated exposure to hydrogen for 2 s with increasing hydrogen concentrations. It shows that the current cycled with stable repeatability and decayed exponentially back to its initial value when hydrogen was switched off. Fig. 19.14 shows the sensitivity change of Pt Schottky diodes on (010) Ga2O3 single crystals as a function of H2 concentration at the forward bias of 0.7 V. The sensitivity increased with a reasonable degree of linearity as H2 concentration increase from 0.01 to 40,000 ppm. Fig. 19.15 shows the time-dependent current changes of Pt Schottky diodes on (010) Ga2O3 single crystals with various gas species, including N2, CO, CO2, O2, CH4, NO2, and NH3. No clear response to gas species other than hydrogen was observed. Clearly, Pt Schottky diodes on Ga2O3 single crystals hold great promise as sensitive hydrogen gas sensors for use at high temperatures in harsh environments.

19.2.5 Capacitance-based gas sensor Most Ga2O3 sensors are based on changes in resistance. Thus, heating source units are required to provide high-temperature operating conditions for the activation of target gas on Ga2O3 surface. In contrast, capacitance-based gas sensors, particularly with

460

Gallium Oxide

105 Sensitivity (%)

(010) Ga2O3 10

4

103 102 101

0.01

0.1 1 H2 concentration (ppm)

10

Fig. 19.14 The variation of the sensitivity (%) of Pt Schottky diodes on (010) Ga2O3 single crystals as a function of H2 concentrations. 80

60

40

Current (nA)

80

Current (nA)

Fig. 19.15 The time-dependent current change of Pt Schottky diodes on (010) Ga2O3 single crystals with various gas species, including N2, CO, CO2, O2, CH4, NO2, and NH3.

60 40

CO O2 NO2 CH4 20 N NH3 CO2 2 0 0

20

500 ppm H2 500 ppm H2 level

1000

Time (s)

2000

0 0

2000 Time (s)

4000

Ga2O3 nanowires, operated at room temperature with low power consumption [7, 70]. The response and recovery times of capacitance-type sensors are short, and auxiliary heating energy is not necessary for rapid response and recovery of the device. Arnold et al. reported an interdigital capacitor gas sensor with β-Ga2O3 nanowires as dielectric [70]. Responses to acetone, methanol, and some hydrocarbons including toluene were investigated. On a Si substrate with a 250-nm-thick SiO2 layer, 250-nm-thick interdigitated Pt electrodes were deposited by standard photolithography and a liftoff process. The sample was covered with a 5-nm Au layer by e-beam evaporation for the formation of seeding islands for the following Ga2O3 nanowire growth. The nanowires were grown on the device by a vapor-liquid-solid growth method at 900°C. The capacitance of the sensor was measured by a balanced ac bridge circuit. The capacitance changes increased with the exposed acetone concentrations, following the Freundlich isotherm. The response time was