Technology of Gallium Nitride Crystal Growth (Springer Series in Materials Science)

Springer Series in

materials science

133

Springer Series in

materials science Editors: R. Hull C. Jagadish R.M. Osgood, Jr. J. Parisi Z. Wang H. Warlimont The Springer Series in Materials Science covers the complete spectrum of materials physics, including fundamental principles, physical properties, materials theory and design. Recognizing the increasing importance of materials science in future device technologies, the book titles in this series ref lect the state-of-the-art in understanding and controlling the structure and properties of all important classes of materials.

Please view available titles in Springer Series in Materials Science on series homepage http://www.springer.com/series/856

Dirk Ehrentraut Elke Meissner Michal Bockowski Editors

Technology of Gallium Nitride Crystal Growth With 190 Figures

123

Editors

Professor Dirk Ehrentraut

Dr. Elke Meissner

Tohoku University World Premier International Research Center - Advanced Institute for Materials Research (WPI - AIMR) Katahira 2-1-1, Aoba-ku 980-8577 Sendai, Japan E-mail: [email protected]

Fraunhofer-Institut Integrierte Systeme und Bauelementetechnologie (IISB) Abteilung Kristallz¨uchtung Schottkystr. 10, 91058 Erlangen, Germany E-mail: [email protected]

Dr. Michal Bockowski Polish Academy of Sciences, Institute of High-Pressure Physics Sokolowska 29/37, 01-142 Warszaw, Poland E-mail: [email protected]

Series Editors:

Professor Robert Hull

Professor J¨urgen Parisi

University of Virginia Dept. of Materials Science and Engineering Thornton Hall Charlottesville, VA 22903-2442, USA

Universit¨at Oldenburg, Fachbereich Physik Abt. Energie- und Halbleiterforschung Carl-von-Ossietzky-Straße 9–11 26129 Oldenburg, Germany

Professor Chennupati Jagadish

Dr. Zhiming Wang

Australian National University Research School of Physics and Engineering J4-22, Carver Building Canberra ACT 0200, Australia

University of Arkansas Department of Physics 835 W. Dicknson St. Fayetteville, AR 72701, USA

Professor R. M. Osgood, Jr.

Professor Hans Warlimont

Microelectronics Science Laboratory Department of Electrical Engineering Columbia University Seeley W. Mudd Building New York, NY 10027, USA

DSL Dresden Material-Innovation GmbH Pirnaer Landstr. 176 01257 Dresden, Germany

Springer Series in Materials Science ISSN 0933-033X ISBN 978-3-642-04828-9 e-ISBN 978-3-642-04830-2 DOI 10.1007/978-3-642-04830-2 Springer Heidelberg Dordrecht London New York Library of Congress Control Number: 2010920285 © Springer-Verlag Berlin Heidelberg 2010 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specif ically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microf ilm or in any other way, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer-Verlag. Violations are liable to prosecution under the German Copyright Law. The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specif ic statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Cover design: eStudio Calamar Steinen Printed on acid-free paper Springer is part of Springer Science+Business Media (www.springer.com)

Foreword

Semiconductor materials have been studied intensively since the birth of silicon technology more than 50 years ago. The ability to physically and chemically tailor their properties with precision is the key factor responsible for the electronic revolution in our society over the past few decades. Semiconductor material systems (like silicon and GaAs-related materials) have now matured and found well established applications in electronics, optoelectronics, and several other fields. Other materials such as III-Nitrides were developed later, in response to needs that the above mentioned semiconductors were unable to fulfill. The properties of IIInitrides (AlN, GaN InN, and related alloy systems) make them an excellent choice for efficient light emitters in the visible as well as the UV region, UV detectors, and for a variety of electronic device such as high frequency unipolar power devices. There was a major upsurge in the research of the GaN material system around 1970. Its prime focus was to grow bulk materials. The failure to produce p-type materials at the time discouraged most research groups, and their activities faded after a few years. Several important developments in growth procedures in the mid to late 1980s led to a resurgence and revival of the research interest in GaN and related materials. A method to grow smooth single crystalline GaN layers on a foreign substrate (i.e., heteroepitaxy on sapphire) using a thin (about 30 nm) low temperature grown AlN buffer layer was suggested by Akasaki et al. They also discovered a method to produce low Ohmic p-type GaN and demonstrated a pn-junction based light emission device. This resulted in the rapid development and commercialization of III-Nitride light emitters (LEDs and lasers). Using heteroepitaxy over sapphire (later SiC), Khan et al. also demonstrated the two-dimensional electron gas at the interface of high quality AlGaN-GaN layers and high electron mobility transistors, in the late 1980s. This led to a rapid development and improvement of electronic microwave power devices with impressive power handling capacity. These are now being employed in commercial applications. To date, a major part of the GaN research has used heteroepitaxy because IIIN bulk substrates are not yet sufficiently developed. Large size, high quality bulk GaN substrates are still much more expensive than sapphire or SiC, which are a preferred choice for GaN heteroepitaxy. The future need for bulk GaN substrates is currently under debate. These substrates provide a clear advantage of increased lifetime in continuous operation of lasers because of reduced dislocations (threading v

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dislocation density <106 cm3 /. For light emitting diodes, however, the advantage is not so obvious. Nevertheless, it is probably true that all GaN based devices would benefit in some respect from the reduced defect density possible from growth on high quality bulk GaN substrates. Hence, the development of suitable growth techniques to produce bulk GaN substrates of high quality, at an affordable cost for commercial devices, has become a very important research area during the last decade. This book is aimed at reflecting the present status of this development. The scope is limited to the GaN material, which is by far the most developed in terms of the growth procedures and material quality. Similar efforts are ongoing for AlN as well, but they are not covered in this book. At present, as discussed in detail by the different authors in this book, several techniques are used to produce bulk GaN. It is clear from the outset that the growth techniques have to be different from the ones used to prepare bulk material of Si or the lower bandgap III–V materials. For these “conventional” semiconductor materials, growth of a boule directly from the melt employing a seed crystal is the most common technique, needing some moderate pressure in the case of some compounds. Purification of the material can also be done using melting and recrystallization, e.g., “zone melting.” For the III-nitrides (like many other wide bandgap semiconductors, such as SiC), this is not a convenient option. The phase diagram for the nitride materials is such that at moderate pressures, the material does not melt upon raising the temperature; it decomposes into liquid metal (Ga for GaN) and N2 gas. Melting of GaN has been calculated to occur at about 2,500ı C, but only if the pressure is >4.5 GPa. Experimentally, if GaN powder is compressed at a pressure of 6 GPa and heated to 2,400ı C maintaining that pressure, small GaN single crystals about 100 m in size have been produced upon cooling. This is the closest analog of growth from a melt such as for silicon. The need for these extremely high pressures and temperatures (and low growth rates as well) makes the melt growth an impractical technique for realizing large single crystals of GaN. Growth techniques for bulk GaN at or near atmospheric pressure, restricting the temperatures to the range <1,200ı C, are more interesting. The so called Halide (or Hydride) Vapor Phase Epitaxy (HVPE) method is the one that has been most popular so far. The technique uses GaCl and NH3 as gaseous sources, and the growth temperature is typically 1,000–1,100ı C, at atmospheric pressure. Most often, sapphire substrates are used, with various seed layers allowing tailoring of the near substrate region. The growth rate can be as high as 0.5 mm/h, and a boule thickness of up to >1 cm has been reported on 2" substrates. Si-doping during growth produces n-type material; p-type can be obtained by Mg doping. By elaboration with different patterning of the substrate or on a buffer layer deposited on the substrate, a pattern of cavities can be produced near the substrate, acting as a weak mechanical link after (or during) cool down. This can be optimized as a self separation technique where the substrate is automatically removed during cool down from growth. Wafers are otherwise cut from the boule by standard sawing, and polished. Obviously, such a polished bulk wafer can be used as a high quality substrate in subsequent bulk growth runs.

Foreword

vii

Another vapor phase growth technique is the Metal Organic Vapor Phase Epitaxy (MOVPE) technique, very important for growth of thin epitaxial layers in device structures. Metalorganic precursors (such as triethyl or trimethyl-gallium) are used together with NH3 in the growth chemistry. For growth of thin epi-layers the growth rate is typically adjusted to about 1–2 m/h, to get high quality layers. For the purpose of bulk growth, the growth rate can be optimized to much higher values, up to about 50 m/h. Solution growth of bulk crystals at moderate temperatures (<1,000ı C) and high pressures is another technique. For decades, this technique has been commercially used for the growth of large quartz crystals. In the so called ammonothermal method for growth of GaN, a solution of GaN polycrystals in mineralizer-containing supercritical NH3 is employed, and kept under a high pressure, typically in the 100– 300 MPa range. Crystals grow on a GaN seed in a temperature gradient inside the apparatus, at a temperature of typically 400–600ı C. So far, the obtained growth rate is limited to about 1–4 m/h. There is however a potential that scaling up the size of the growth system and optimizing growth conditions will allow growth of very large single crystal boules in the future. So far, 2" diameter boules have been reported. Another approach to grow bulk GaN is the flux growth method, where a metal flux is used to dissolve GaN. At high temperature and pressure, a certain solubility of GaN into Ga can be obtained, so that bulk crystals can be obtained by solution growth. Experimentally, at a pressure of 1.5 GPa, GaN thin crystal platelets of cm size have been grown from Ga solution in N2 atmosphere by lowering the temperature in the range of 1,600ı C. Scaling up to large crystals and seeding are still a problem with this method. A vibration of the flux growth method uses Na as the flux metal. Growth takes place on a seed in a temperature gradient at moderate temperatures (700–800ı C) and pressures (of the order of 50 bar). Recently, this technique has also shown rapid progress; growth rates have been increased to about 30 m/h, and boules of a diameter of 2" and thickness of several mm have been demonstrated. Many other flux techniques operating at atmospheric pressure are also under development. The quality observed for the bulk GaN crystals produced by the above techniques is generally excellent, far better than the material produced by epitaxy on foreign substrates. The emphasis has generally not been on producing very high purity material, the best purity so far in terms of residual impurity density is in the range 1016 –1017 cm3 . For commercial purposes, the main demand is on n-doped and semi-insulating wafers. Both the insulating and the conducting crystals can generally be produced by the above techniques. However, the approaches used usually vary slightly on the basis of the technique type. The XRD rocking curve width is typically of the order 20 arcsec, and threading dislocation density is reported to be in the range 104 –106 cm2 , which is regarded as proper for even the most demanding applications. Over the last few years, impressive progress has been made in the growth of bulk GaN materials. The growth rate and the size of the grown boules have increased significantly. Several academic and industrial laboratories now readily produce GaN substrates. This will allow a proper examination of the advantages of using high

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Foreword

quality bulk GaN substrates for the epitaxial growth of various optical and electronic device structures. It seems clear that for high power light emitting structures (such as those used for solid state lighting) it may be necessary to use non-polar or semipolar surfaces for the growth, because of the limitations in radiative efficiency imposed by intrinsic Auger recombination in the quantum well active region. Large boules or crystals will be needed to cut wafers in the desired crystal direction for this development. This demand for substrates for non-polar or semi-polar growth might further stimulate the development and the use of bulk GaN substrates, as heteroepitaxial growth of device structures in non-polar or semi-polar directions with a low defect density has so far proven to be quite difficult. Ultimately, the use of bulk GaN substrates will primarily be decided on the basis of the economic factors. Like their counterparts used in heteroepitaxy, they will also have to be high quality and low-cost. The challenge will thus be to scale up the growth systems to efficiently produce large quantities of high quality material. At present, many of the techniques discussed here are in a phase of transfer to industry, where such a process scale up will be carried out. Columbia, USA Linköping, Sweden March 2010

M. Asif Khan B. Monemar

Preface

Gallium nitride (GaN) has emerged as one of the most important semiconductors in modern technology. Its future shines even brighter as we see the advances towards solid state lighting and high-power electronics. What was mainly pushing, actually creating this entirely new sector of GaN-based device technology, was the success in achieving reliable p-type doping and consequently, the ability for fabrication of light emitter devices (LEDs and LDs). A pioneer in the field, Shuji Nakamura, has summarized this work in S. Nakamura, G. Fasol, The Blue Laser Diode, SpringerVerlag, 1st edition, 1997. Much has been done, since then, in the development of better and more efficient GaN-based devices, already creating a multibillion dollar market. This is even more astounding in the light of the relatively underdeveloped technology for lattice and thermally-matched substrates for GaN-based devices. Only since around the year 2000, has the crystal growth technology of GaN been developed to a now widely recognized field in academia and industry. This book is designed to bring to the readership for the first time, a comprehensive overview of the state-of-the-art GaN crystal growth technology, reflecting the tremendous progress made particularly over the last decade, drawing the possible path we still have to cover to realize our common goal: large-size, dislocation-free GaN crystals to fabricate non-polar, semi-polar, and polar GaN wafers in sufficient quantity and at reasonable price. A good number of recognized leaders from industry and academia have contributed to this book for which the editors are very grateful. We sincerely hope that this book will be a unique resource for engineers, researchers, and students dealing with the crystal growth of GaN, processing, and device fabrication, both in industry and academia. Last but not the least, we are grateful to everyone who was involved in preparing this book. In particular, D.E. is indebted to Prof. T. Fukuda for the collaboration over the last few years; he also thanks his wife, Yumiko, for her patience and support while he was working on this book.

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Finally, Dr. C. Ascheron and Ms. A. Duhm from the Springer-Verlag at Heidelberg are acknowledged for the cooperation extended by them toward the realization of the book. Sendai, Erlangen, Warsaw, March 2010

Dirk Ehrentraut Elke Meissner Michał Bo´ckowski

Contents

Part I Market for Bulk GaN Crystals 1

Development of the Bulk GaN Substrate Market . . . . . . . . . . . . . .. . . . . . . . . . . Andrew D. Hanser and Keith R. Evans 1.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.2 III-N Device Market Drivers and Forecast .. . . . . . . . . . . . . . . .. . . . . . . . . . . 1.2.1 Light Generation and Solid State Lighting in the III-Ns . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.2.2 Electrical Systems and Power Electronics in the III-Ns . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.2.3 Positioning GaN Substrates for SSL and Power Electronics Markets . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.2.4 Key Drivers for Bulk GaN Substrate Commercialization Success . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.3 Benefits and Importance of Bulk GaN Substrates . . . . . . . . .. . . . . . . . . . . 1.3.1 Device Performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.3.2 Thermal Conductivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.3.3 Thermally Activated Device Failure .. . . . . . . . . . . . .. . . . . . . . . . . 1.3.4 Device Cost . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.4 GaN Device Trends for Bulk GaN Substrates . . . . . . . . . . . . .. . . . . . . . . . . 1.4.1 Lasers and LEDs. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.4.2 Power Switches .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.4.3 High-Frequency, High-Power HEMTs . . . . . . . . . . .. . . . . . . . . . . 1.5 Bulk GaN Substrate Trends .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.5.1 Hydride Vapor Phase Epitaxy .. . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.5.2 Ammonothermal Growth .. . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.5.3 Solution Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.5.4 Combined Growth Techniques .. . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 1.6 Summary.. .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .

3 3 4 4 5 7 9 11 12 16 17 18 19 19 20 21 21 22 22 23 23 24 25

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Part II Vapor Phase Growth Technology 2

3

Hydride Vapor Phase Epitaxy of GaN . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . Akinori Koukitu and Yoshinao Kumagai 2.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.2 Thermodynamic Analysis on HVPE Growth of GaN . . . . .. . . . . . . . . . . 2.2.1 Calculation Procedure . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.2.2 Equilibrium Partial Pressure and Driving Force for the GaN Deposition .. . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.3 Cubic GaN Epitaxial Growth on (100) GaAs Substrate .. .. . . . . . . . . . . 2.3.1 Experimental .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.3.2 Cubic GaN Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.4 Comparison of GaN Growth on (111)A and (111)B GaAs Substrates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.4.1 Experimental .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.4.2 Comparison of GaN Growth on (111)A and (111)B GaAs Surfaces .. . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.5 Ab Initio Calculations of GaN Initial Growth Processes on (111)A and (111)B GaAs Surfaces . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.5.1 Calculation Procedure . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.5.2 GaN Initial Growth Processes on (111)A and (111)B GaAs Surfaces .. . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.6 Thick GaN Growth on (111)A GaAs Substrate . . . . . . . . . . . .. . . . . . . . . . . 2.6.1 Experimental .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.6.2 Thick GaN Growth on (111)A GaAs Surface .. . . . . . . . . . . . . . 2.7 Preparation of Fe-Doped Semi-insulating GaN Substrates.. . . . . . . . . . 2.7.1 Experimental .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 2.7.2 Fe-Doped GaN Layer Grown on Sapphire and GaAs. . . . . . . References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . Growth of Bulk GaN Crystals by HVPE on Single Crystalline GaN Seeds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . B. Łucznik, B. Pastuszka, G. Kamler, I. Grzegory, and S. Porowski 3.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 3.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 3.2.1 Seed Crystals. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 3.2.2 HVPE Reactor and Growth Conditions . . . . . . . . . .. . . . . . . . . . . 3.2.3 Characterization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 3.3 Experimental Results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 3.3.1 Crystals Grown on Small Near Dislocation Free GaN Platelet-Like Seeds .. . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 3.3.2 Crystals Grown on Small Near Dislocation Free GaN Needle-Like Seeds . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .

31 31 32 32 34 37 38 39 42 42 43 45 46 46 49 49 50 53 55 56 59

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3.3.3

Crystals Grown on Large (0001) Oriented GaN Substrates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 73 3.4 Conclusions .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 76 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 77 4

5

Freestanding GaN Wafers by Hydride Vapor Phase Epitaxy Using Void-Assisted Separation Technology .. . . . . . . . .. . . . . . . . . . . Y. Oshima, T. Yoshida, T. Eri, K. Watanabe, M. Shibata, and T. Mishima 4.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.2 Outline of the HVPE-VAS Technology . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.2.1 Concept of the HVPE-VAS Technology .. . . . . . . . .. . . . . . . . . . . 4.2.2 Overview of the Process .. . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.3 Preparation of a GaN Template with a Porous TiN Film . .. . . . . . . . . . . 4.3.1 Experimental .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.3.2 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.3.3 Mechanisms for the Formation of the Porous Structure . . . . 4.4 HVPE Growth on GaN Templates with a Porous TiN Film . . . . . . . . . . 4.4.1 Process of HVPE Growth and Separation .. . . . . . .. . . . . . . . . . . 4.4.2 Mechanisms of Growth and Separation . . . . . . . . . .. . . . . . . . . . . 4.5 Properties of GaN Wafers Fabricated by HVPE–VAS Technology .. 4.5.1 Structural Properties .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.5.2 Electrical Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.5.3 Thermal Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 4.6 Summary.. .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .

79

79 80 80 81 81 81 81 82 83 84 85 88 88 90 91 94 95

Nonpolar and Semipolar GaN Growth by HVPE . . . . . . . . . . . . . .. . . . . . . . . . . 97 Paul T. Fini and Benjamin A. Haskell 5.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 97 5.2 Heteroepitaxial Films, Including Substrate Selection . . . . .. . . . . . . . . . . 99 5.2.1 Planar a-Plane GaN Films . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . 99 5.2.2 Planar m-Plane GaN Films . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .102 5.2.3 Planar Semipolar GaN Films. . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .104 5.3 Lateral Epitaxial Overgrowth of Nonpolar, Semipolar GaN . . . . . . . . .109 5.3.1 LEO of a-Plane GaN . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .109 5.3.2 LEO of m-Plane GaN . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .112 5.3.3 LEO of Semipolar GaN . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .115 5.4 Conclusions and Future Development.. . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .116 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .117

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6

Contents

High Growth Rate MOVPE . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .119 K. Matsumoto, H. Tokunaga, A. Ubukata, K. Ikenaga, Y. Fukuda, Y. Yano, T. Tabuchi, Y. Kitamura, S. Koseki, A. Yamaguchi, and K. Uematsu 6.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .119 6.2 Growth Characteristics of AlGaN and GaN by Conventional MOVPE .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .120 6.3 Quantum Chemical Study of Vapor-Phase Reaction . . . . . .. . . . . . . . . . .123 6.4 Result of High-Growth-Rate GaN by Using a High-Flow-speed Reactor .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .127 6.5 Discussion and Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .130 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .132

Part III

Solution Growth Technology

7

Ammonothermal Growth of GaN Under Ammono-Basic Conditions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .137 R. Doradzi´nski, R. Dwili´nski, J. Garczy´nski, L.P. Sierzputowski, and Y. Kanbara 7.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .137 7.2 The Growth Method .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .138 7.2.1 Physico-Chemical Basics . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .139 7.2.2 Solubility Measurements .. . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .141 7.2.3 Equipment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .142 7.2.4 Seeded Recrystallization . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .143 7.2.5 Doping . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .144 7.2.6 Crystal Machining . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .144 7.3 Crystal Characterization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .146 7.3.1 Structural Properties .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .146 7.3.2 Optical Properties.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .151 7.3.3 Electrical Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .152 7.4 Homoepitaxy on Ammonothermal GaN . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .152 7.5 Conclusions .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .157 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .158

8

A Pathway Toward Bulk Growth of GaN by the Ammonothermal Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .161 Tadao Hashimoto and Shuji Nakamura 8.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .161 8.2 Impact of Mineralizer on Ammonothermal Synthesis of GaN . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .162 8.3 Solubility of GaN in Ammonobasic Solutions . . . . . . . . . . . . .. . . . . . . . . . .165 8.4 Seeded Growth of GaN with Metallic Ga Nutrient . . . . . . . .. . . . . . . . . . .170 8.5 Seeded Growth of GaN with Polycrystalline GaN Nutrient.. . . . . . . . .174 8.6 Growth of Bulk GaN Crystals and Sliced Wafers . . . . . . . . .. . . . . . . . . . .177 8.7 Summary.. .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .180 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .181

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Acidic Ammonothermal Growth Technology for GaN . . . . . . . .. . . . . . . . . . .183 Dirk Ehrentraut and Yuji Kagamitani 9.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .183 9.2 Brief History of the Ammonothermal Growth Technique of GaN .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .185 9.3 Growth Technology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .186 9.4 Chemistry of the Solution and Growth Mechanism . . . . . . .. . . . . . . . . . .188 9.4.1 Solubility .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .188 9.4.2 Growth Rate and Chemistry of the Solution . . . . .. . . . . . . . . . .190 9.4.3 Effect of Acidity on Formation of GaN . . . . . . . . . .. . . . . . . . . . .193 9.5 Properties of Ammonothermal GaN. . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .195 9.6 Prospects and Future developments for Ammonothermal GaN. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .200 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .202

Part IV

Flux Growth Technology

10 High Pressure Solution Growth of Gallium Nitride . . . . . . . . . . .. . . . . . . . . . .207 Michal Bo´ckowski, Pawel Strak, ˛ Izabella Grzegory, and Sylwester Porowski 10.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .207 10.2 Growth Method .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .209 10.2.1 Thermodynamic and Kinetic Aspects of HPS Growth.. . . . .210 10.2.2 Experimental .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .211 10.3 Spontaneous Crystallization by HPS Growth Method. . . . .. . . . . . . . . . .214 10.3.1 Habit and Morphology of the Crystals . . . . . . . . . . .. . . . . . . . . . .214 10.3.2 Physical Properties of the Crystals. . . . . . . . . . . . . . . .. . . . . . . . . . .215 10.4 Seeded Growth by HPS Method .. . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .216 10.4.1 Liquid Phase Epitaxy in the c-Direction on Various Substrates .. . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .217 10.4.2 Modeling of the Convective Transport in Gallium for LPE Growth (Time Independent Solutions) 219 10.4.3 Seeded Growth with Convective Flow of Gallium Under Control . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .222 10.4.4 Growth on HVPE Seeds in the c-Directions . . . . .. . . . . . . . . . .224 10.4.5 Growth on HVPE Seeds in Nonpolar Directions . . . . . . . . . . .227 10.4.6 Modeling of the Convective Transport in Gallium (Time Dependent Solutions) . . . . . . . . . . . .. . . . . . . . . . .227 10.5 Applications of Pressure Grown GaN Substrates: Blue Laser Diodes in TopGaN Ltd . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .230 10.6 Summary and Perspectives of HPS Growth Method . . . . . .. . . . . . . . . . .232 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .233

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Contents

11 A Brief Review on the Na-Flux Method Toward Growth of Large-Size GaN Crystal.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .235 Dirk Ehrentraut and Elke Meissner 11.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .235 11.2 Historical Development in Brief .. . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .235 11.3 Experimental Conditions for the LPE Growth of GaN by the Na-Flux Method . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .238 11.4 Growth Mechanism and Dislocations . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .239 11.4.1 Effect of Flux Composition on Growth Stability and Crystal Morphology . . . . . . . . . . . . . . . .. . . . . . . . . . .239 11.4.2 Growth Mechanism and Effect on Dislocation Population .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .240 11.4.3 Solubility and Growth Rate . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .241 11.5 Properties of GaN . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .242 11.6 Industrialization Potential for the Na-Flux Method . . . . . . .. . . . . . . . . . .243 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .243 12 Low Pressure Solution Growth of Gallium Nitride . . . . . . . . . . . .. . . . . . . . . . .245 E. Meissner, S. Hussy, and J. Friedrich 12.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .245 12.2 Technology of Solution Growth Under Ambient Pressure, the LPSG Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .247 12.2.1 The Formation Reaction of GaN Under Ammonia Atmosphere . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .248 12.2.2 Solubility of Nitrogen in Gallium-Metal Solutions . . . . . . . . .251 12.2.3 Growth Setup, Process, and Basic Challenges . . .. . . . . . . . . . .254 12.2.4 The Influence of Process Parameters on the Epitaxial and Parasitic Growth of GaN . . .. . . . . . . . . . .256 12.3 Evolution of Structure and Morphology of the GaN Layers . . . . . . . . .261 12.3.1 Importance of the Initial Growth Stage . . . . . . . . . . .. . . . . . . . . . .261 12.3.2 Macroscopic Defects in LPSG GaN . . . . . . . . . . . . . .. . . . . . . . . . .263 12.4 Properties of the LPSG GaN Material.. . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .265 12.4.1 Structural Properties and Dislocation Density in LPSG GaN Material . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .265 12.4.2 Electrical Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .270 12.4.3 Impurities . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .270 12.5 Summary and Prospect .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .271 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .272 Part V

Characterization of GaN Crystals

13 Optical Properties of GaN Substrates .. . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .277 Shigefusa F. Chichibu 13.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .277

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13.2 Optical Properties of GaN Substrates Grown by Metalorganic Vapor Phase Epitaxy and Halide Vapor Phase Epitaxy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .279 13.2.1 Photoreflectance Spectra of Excitonic Polaritons in MOVPE-LEO GaN Substrate . . . . . .. . . . . . . . . . .279 13.2.2 TRPL of MOVPE-LEO GaN Substrate . . . . . . . . . .. . . . . . . . . . .283 13.2.3 Low Temperature PL Spectra of HVPE-LEO GaN Substrate . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .285 13.3 Effects of Growth Polar Direction on the Optical Properties of Seeded GaN Substrates Grown by Ammonothermal Method .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .286 13.4 Effects of Dislocation Bending on the Optical Properties of Seeded GaN Substrates Grown by Ammonothermal Method .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .288 13.5 Summary.. .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .290 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .291 14 Point Defects and Impurities in Bulk GaN Studied by Positron Annihilation Spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .295 Filip Tuomisto 14.1 Introduction .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .295 14.2 Positron Annihilation Spectroscopy .. . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .296 14.2.1 Positrons in Solids . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .296 14.2.2 Positrons at Defects.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .298 14.2.3 Experimental Techniques . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .301 14.3 In-Grown Defects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .304 14.3.1 Defect Formation: Growth Methods and Doping . . . . . . . . . . .304 14.3.2 Defects and Growth Polarity .. . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .307 14.4 Defect Engineering .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .309 14.4.1 High Pressure Thermal Annealing .. . . . . . . . . . . . . . .. . . . . . . . . . .310 14.4.2 Electron Irradiation Experiments . . . . . . . . . . . . . . . . .. . . . . . . . . . .311 14.5 Summary.. .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .314 References .. . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .314 Index . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . .317

−

Contributors

Michal Bo´ckowski Institute of High Pressure Physics, Polish Academy of Sciences, 01-142 Warsaw, Sokołowska 29/37, Poland, [email protected] Shigefusa F. Chichibu Center for Advanced Nitride Technology (CANTech), Institute of Multidisciplinary Research for Advanced Materials (IMRAM), Tohoku University, 2-1-1 Katahira, Aoba, Sendai 980-8577, Japan, [email protected] ´ R. Doradzinski AMMONO Sp. z.o.o., ul. Czerwonego Krzyz˙ a 2 /31, 00-377 Warsaw, Poland, [email protected] ´ R. Dwilinski AMMONO Sp. z.o.o., ul. Czerwonego Krzyz˙ a 2 /31, 00-377 Warsaw, Poland, [email protected] Dirk Ehrentraut WPI-AIMR, Tohoku University, 2-1-1 Katahira, Aoba-ku, 980-8577 Sendai, Japan, [email protected] T. Eri Takasago Works, Compound Semiconductor Division, Hitachi-cable Ltd., 880 Isagozawa, Hitachi, Ibaraki 319-1418, Japan, [email protected] Keith R. Evans Kyma Technologies, Inc., 8829 Midway West Road, Raleigh, NC 27617-4606 USA, [email protected] Paul T. Fini Inlustra Technologies, LLC, Santa Barbara, CA 93111, USA, [email protected] Jochen Friedrich Fraunhofer Institute of Integrated Systems and Device Technology, Department of Crystal Growth, Schottkystr. 10, 91085 Erlangen, Germany Y. Fukuda Taiyo Nippon Sanso, Tsukuba Laboratories, 10 Okubo, Tsukuba, 300-2611 Japan, [email protected] ´ J. Garczynski AMMONO Sp. z.o.o., ul. Czerwonego Krzyz˙ a 2 /31, 00-377 Warsaw, Poland, [email protected] Izabella Grzegory Institute of High Pressure Physics, Polish Academy of Sciences 01-142 Warsaw, Sokołowska 29/37, Poland, [email protected]

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xx

Contributors

Andrew D. Hanser SRI International, 7935 114th Avenue N, Largo, FL 33773, USA, [email protected] T. Hashimoto SixPoint Materials, Inc., Buellton, CA 93427, USA, [email protected] Benjamin A. Haskell Inlustra Technologies, LLC, Santa Barbara, CA 93111, USA Stephan Hussy Fraunhofer Institute of Integrated Systems and Device Technology, Department of Crystal Growth, Schottkystr. 10, 91085 Erlangen, Germany K. Ikenaga Taiyo Nippon Sanso, Tsukuba Laboratories, 10 Okubo, Tsukuba, 300-2611 Japan, [email protected] Yuji Kagamitani Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, 2-1-1 Katahira, Aoba-ku, 980-8577 Sendai, Japan, [email protected] G. Kamler Institute of High Pressure Physics, Polish Academy of Sciences, ul. Sokolowska 29/37 01-142 Warsaw, POLAND, [email protected] Y. Kanbara Nichia Corp., Oka, Kaminaka-cho, 491-100 Anan-shi, Tokushima, Japan, [email protected] Y. Kitamura Taiyo Nippon Sanso, Engineering and Development Div., 1-3-26 Koyama, Shinagawa, Tokyo 142-8558 Japan, [email protected] S. Koseki Taiyo Nippon Sanso, Compound Semiconductor Div., 6-2 Kojimacho, Kawasaki, Kanagawa 210-0861 Japan, [email protected] Akinori Koukitu Department of Applied Chemistry, Tokyo University of Agriculture and Technology, Koganei, Tokyo 184-8588, Japan, [email protected] Yoshinao Kumagai Department of Applied Chemistry, Tokyo University of Agriculture and Technology, Koganei, Tokyo 184-8588, Japan, 4470kuma@cc. tuat.ac.jp B. Łucznik Institute of High Pressure Physics, Polish Academy of Sciences, ul. Sokolowska 29/37 01-142 Warsaw, POLAND, [email protected] K. Matsumoto TN EMC Ltd., 2008-2 Wada, Tama-shi, Tokyo 206-0001 Japan, [email protected] Elke Meissner Department of Crystal Growth, Fraunhofer Institute of Integrated Systems and Device Technology, Schottkystrasse 10, 91058 Erlangen, Germany, [email protected] T. Mishima Advanced Electronic Materials Research Dept. Research & Development Laboratory, Hitachi Cable Ltd., 3550 Kidamari, Tsuchiura, Ibaraki 300-0026, Japan, [email protected]

Contributors

xxi

S. Nakamura Materials Department, University of California, Santa Barbara, CA 93106, USA Y. Oshima Advanced Electronic Materials Research Dept. Research & Development Laboratory, Hitachi Cable Ltd., 3550 Kidamari, Tsuchiura, Ibaraki 300-0026, Japan, [email protected] B. Pastuszka Institute of High Pressure Physics, Polish Academy of Sciences, ul. Sokolowska 29/37 01-142 Warsaw, POLAND Sylwester Porowski Institute of High Pressure Physics, Polish Academy of Sciences, ul. Sokolowska 29/37, 01-142 Warsaw, POLAND, [email protected] M. Shibata Takasago works, Compound Semiconductor Division, Hitachi-Cable Ltd., 880 Isagozawa, Hitachi, Ibaraki 319-1418, Japan, [email protected] L.P. Sierzputowski AMMONO Sp. z.o.o., ul. Czerwonego Krzyz˙ a 2 /31, 00-377 Warsaw, Poland, [email protected] Pawel Stra¸k Institute of High Pressure Physics, Polish Academy of Sciences, ul. Sokolowska 29/37 01-142 Warsaw, POLAND, [email protected] T. Tabuchi Taiyo Nippon Sanso, Tsukuba Laboratories, 10 Okubo, Tsukuba, 300-2611 Japan, [email protected] H. Tokunaga Taiyo Nippon Sanso, Tsukuba Laboratories, 10 Okubo, Tsukuba, 300-2611 Japan, [email protected] Filip Tuomisto Department of Applied Physics, Helsinki University of Technology, P.O. Box 1100, FI-02015, TKK Espoo, Finland, [email protected] A. Ubukata Taiyo Nippon Sanso, Tsukuba Laboratories, 10 Okubo, Tsukuba, 300-2611 Japan, [email protected] K. Uematsu Taiyo Nippon Sanso, Compound Semiconductor Div., 6-2 Kojimacho, Kawasaki, Kanagawa 210-0861 Japan, [email protected] K. Watanabe Takasago Works, Compound Semiconductor Division, Hitachicable Ltd., 880 Isagozawa, Hitachi, Ibaraki 319-1418, Japan, [email protected] A. Yamaguchi Taiyo Nippon Sanso, Compound Semiconductor Div., 6-2 Kojimacho, Kawasaki, Kanagawa 210-0861 Japan, [email protected] Y. Yano Taiyo Nippon Sanso, Tsukuba Laboratories, 10 Okubo, Tsukuba, 300-2611 Japan, [email protected] T. Yoshida Advanced Electronic Materials Research Department, Research and Development Laboratory, Hitachi Cable Ltd., 3550 Kidamari, Tsuchiura, Ibaraki 300-0026, Japan, [email protected]

Part I

Market for Bulk GaN Crystals

Chapter 1

Development of the Bulk GaN Substrate Market Andrew D. Hanser and Keith R. Evans

Abstract Near- and long-term market applications for bulk GaN substrates are examined, along with motivations and challenges for adopting the substrate technology for specific device applications. The near-term demand for bulk GaN is driven primarily by laser diodes, while solid state lighting and power electronics will drive the long-term demand. Challenges to achieving broad market penetration include increased volume production and reduced manufacturing cost, which are needed to penetrate incumbent GaN device technologies based on foreign substrates.

1.1 Introduction Technology is pervasive throughout most of today’s society. Around the world, many of us are mass consumers of technology and vast industries have been created to meet our growing technological needs. Even in the third world, technology is increasingly being introduced in various forms including wireless communications, solar power, and modern pharmaceuticals. Advanced technologies have benefitted our lives in countless ways, but have also extracted a price in terms of resource and energy consumption and waste generation. A growing challenge is to employ technologies that advance our standard of living while addressing an increasingly important need to do more with less energy in order to reduce the negative impact on our environment. In the ongoing development and application of new technologies, every so often there appears a fundamental technology that can shift the way the world operates. The development of silicon semiconductor materials, which enabled transistors, integrated circuits, microprocessors, the computer, and the information age, has influenced virtually every aspect of modern life. In a similar manner, the III-Nitride (III-N) semiconductor materials are poised in such a way to fundamentally change our lives. These materials, which include aluminum nitride (AlN), gallium nitride (GaN), and indium nitride (InN), will enable semiconductor devices with new capabilities and will make possible the reinvention of existing technologies. While there are many possible and even likely applications for these materials, their impact over 3

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the next decade will focus on two main applications: light generation and the control of electrical power. These applications can anticipate large new commercial markets and this possibility is already stimulating aggressive development of supporting materials and device technologies. The great promise arising from these innovations is the ability to do more (more light, more power) with less (higher energy efficiency, less electricity, less waste heat). In this chapter, we discuss the opportunities and the challenges GaN-based materials face in addressing these markets and applications, and, in particular, we examine how GaN substrates will be used in new market applications. We present how GaN substrates will be introduced into near-term market opportunities that will create the driving force for manufacturing improvements and will allow GaN substrates to effectively address future market opportunities.

1.2 III-N Device Market Drivers and Forecast 1.2.1 Light Generation and Solid State Lighting in the III-Ns One of the unique properties of the III-Ns is that they can enable efficient light emission from deep ultraviolet to infrared wavelengths. Accessing these wavelengths makes possible the development of highly efficient semiconductor-based white light sources. Solid-state lighting (SSL) based on light-emitting diodes (LEDs) allows the consumer to replace less efficient and less reliable lighting technologies, such as incandescent, fluorescent, and metal halide light sources. A tremendous amount of energy goes into generating light for our everyday lives. For example, in the United States, about 22% of the nation’s electrical energy is used for lighting [1]. This equals the output of about 100 large power plants, but more than three times this amount is needed to produce the electricity due to electrical inefficiencies in lighting systems. The cost of this electricity is about $55 billion and growing, as energy costs increase, populations grow, and technology penetration throughout the world rises. Today, the need for lighting translates into a global market for lamps (light bulbs), ballasts, lighting fixtures, and lighting controls that is valued at about $40 billion annually [1]. Highly efficient lighting systems based on SSL are being developed to replace many of these low efficiency lighting technologies. As one considers the penetration of SSL into lighting applications and the energy savings resulting from these changes, it is easily seen how this technology can bring a significant benefit to our lives: improved energy efficiency and reduced energy consumption, carbon emission reduction, and lower total cost to the consumer. These benefits create the impetus for the development of the new GaN-based technologies for SSL. According to the U.S. Department of Energy (US DOE), the long-term research and development goals for SSL call for a luminous efficiency of 160 lm/W in cost-effective, market-ready systems in 2025 [2]. This efficiency level of

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Table 1.1 Potential impact of conversion to solid state lighting on U.S. electrical energy consumption General illumination lighting Performance estimates SSL Product 60 W incandescent 23 W compact light bulb fluorescent lamp Light output (lm) 1000 1000 1200 Power (W) 6.67 60 23 Lumens/W (system) 150 16.7 52 Annual energy consumption 19.5 175.2 67.2 (8 h/day, 365 days) (kWh) Factor higher than LED 1 9.0 3.4 Annual energy cost per lamp $1.81 $16.29 $6.25 (9.3 cents/kWh) Estimated annual energy savings with LED lighting: 2020 estimated baseline energy consumption for lighting: 7.5 quads % U.S. lighting conversion Quads saved $ Saved (Billions) to SSL Assumption: equal 1% 0.05 0.33 replacement of 10% 0.49 3.23 incandescent and 25% 1.21 7.99 fluorescent 50% 2.43 16.04 lighting Notes: “quad” is one quadrillion BTU; approximately $6.6 billion per quad of electrical energy

performance represents a 60% improvement over the current high-efficiency sources such as fluorescent and metal halide lighting and a much larger >800% improvement over incandescent lighting, by far the most prevalent form of lighting in the world today. As can be seen from the summary of luminous efficacies for different light sources shown in Table 1.1, significant advances over the current state-of-the-art technology are required to meet the US DOE’s requirements. In 2008, the GaN LED market was estimated to be around $7 billion and is estimated to grow to approximately $15 billion by 2012 [4]. In spite of the already significant market for GaN LEDs, their use in SSL today is quite limited to a small but growing number of applications, restricted both by the cost and the overall technology readiness level of GaN-based LED lighting solutions compared to the existing approaches. As improved device technologies are developed, the market for LEDs for SSL is forecast to grow to $1.2 billion in 2012 [4] and will reach $20B or more over the long term (Table 1.2).

1.2.2 Electrical Systems and Power Electronics in the III-Ns Improving energy efficiency in electrical systems reduces the total energy requirements for many applications. For example, standby electrical power for consumer electronics can account for 5–10% of annual energy consumption in the home, adding up to more than $3 billion in annual energy costs [5, 6]. According to the

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Table 1.2 Luminous efficacy ranges for different light sources [3] Light Typical luminous efficacy source range (lm/W) Incandescent 10–18 Halogen 15–20 Compact fluorescent (incl. ballast) 35–60 Linear fluorescent (incl. ballast) 50–100 Metal halide (incl. ballast) 50–90 Current SSL products for general illumination Up to 62 SSL system luminous efficacy target (2025) 160

US DOE, this standby electricity consumption was 64.5 million megawatt hours in 2004 [7]. This amount of wasted energy is equivalent to the output of 18 typical power stations. This standby power consumption is a result of the use of low-cost yet inefficient power supplies that consume a significant amount of power when the electronic device is not in use. Additionally, the implementation of SSL will require high-efficiency power suppliers and controllers to maintain the overall highefficiency of the lighting system. Electrical systems are also finding increased use in automotive applications, where hybrid electric vehicles (HEVs) are growing in importance as a way to improve efficiency in transport systems. Using an electric motor in combination with an internal combustion engine boosts vehicle energy efficiency and enables reduction in transportation related oil demand and associated atmospheric emissions. One particular study shows that increased HEV use offers the possibility of reducing the total oil demand for light-duty vehicles by 2030 relative to current levels of consumption, in spite of an increase in both the number of vehicles and the average annual vehicle mileage [8]. Using an electric motor alone or in combination with an internal combustion engine boosts vehicle energy efficiency and enables a reduction in CO2 emissions, along with a reduction in transportation related oil demand. Finally, the electrical grid that includes the traditional generation and transmission of electrical power from power plants to users is growing to include new sources of power generation, including distributed and renewable power networks, such as wind, solar, biomass, and others. Current generation systems can benefit from more efficient transformation and switching of power, and smarter systems improve the efficiency of renewable energy sources through monitoring and management of the varying cycles of energy generation and load requirements. Improvements in the transformation of voltages and enhanced dynamic control of the voltage, impedance, and phase angle of high-voltage ac transmission lines can be achieved through solid-state power control devices [9]. Energy efficiency in power conversion for industrial and consumer electronics, electronics for fully electric and hybrid electric vehicles, and power generation and distribution for the electric grid, can all benefit from higher efficiency electrical components. Improving efficiency in these systems starts with the high power electronic (HPE) components that make up the power inverters and converters in power supplies. For modern systems these components are semiconductor devices made

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primarily from silicon. Due to intrinsic materials limitations, silicon devices are limited in power handling capabilities and temperature of operation, which negatively impacts the overall electrical system efficiency. Gallium nitride has several properties that enable improved power performance with higher efficiency when compared to silicon. Gallium nitride has a wide bandgap (3.4 eV), a large critical breakdown electric field .3 MV=cm/, and a high electron mobility (1500 cm2=V s at RT). These properties result in low on-state resistance, low switching losses, high-temperature performance capabilities, and high power conversion efficiencies for electronic devices such as Schottky diodes and field effect transistors (FETs). Employing these devices would improve the efficiency of electrical systems and could have a significant impact in reducing overall electricity consumption worldwide, impacting virtually every aspect of electrical usage with potential savings in excess of $30 billion per year [10]. In 2008, the total available market for GaN power electronics was greater than $350 million, possibly growing to a large portion of the $22 billion power electronics market over the long term [11].

1.2.3 Positioning GaN Substrates for SSL and Power Electronics Markets Considerable technology development is needed to enable penetration of GaN devices into the SSL and power electronics markets. Accessing these markets requires the maturation of the processing technologies for fabricating the materials and devices that are used in the lighting and power components systems. Highly manufacturable, low-cost processes for fabricating high-performance devices are needed. Two significant elements will contribute to the successful implementation of GaN-based devices in these applications. First, the performance of the GaN devices must improve. For example, luminous efficacy needs to increase by more than a factor of three in order to meet and exceed the performance of competing approaches. GaN-based LEDs have steadily improved their efficacy over a period of time through improvements in several areas such as the device design; however, it remains to be seen if the current technology approach will be able to continue on its current performance trajectory. Most of the standard LEDs today are grown on foreign (non-nitride) substrates, which results in an extensive number of materials defects due to the imperfect match of materials properties between the substrate and epitaxial layers. Such an approach results in an imperfect match of materials properties between the substrate and epitaxial layers, which generates defects and degrades device performance. Researchers are investigating the effect of these defects on the performance of LEDs and whether device performance remains fundamentally limited due to material quality. This is of particular importance for high current density devices that are used for high total light output. As with lighting applications, device performance and demonstrated reliability remain as significant barriers to realizing the full potential of GaN-based power electronics devices.

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The second element is cost reduction of GaN-based device technology, where substrate cost can be a significant component of the total cost of the device. Substrate cost reduction is accomplished through increasing production volume. Navigating the cost–volume curve is one of the main challenges in semiconductor manufacturing and market demand drives the need for increased volume production. Demand is strong for GaN-based devices, but for devices based on GaN substrates, current markets such as those for LEDs are only beginning to use GaN substrates and only for high end applications, due to incumbent device technologies with low price points. Devices on foreign substrates have adequate performance for current LED applications, and the price point for these GaN epiwafers and devices is low due in part to the high volumes and low cost of substrates. As a result, emerging LED technologies utilizing bulk GaN substrates cannot take advantage of these markets yet, due to the high costs of the substrates that result from low production volumes and low manufacturing yields. Thus, new markets need to be identified that can accommodate these near-term limitations for bulk GaN substrates. Over the long term, GaN substrate manufacturing costs should approach that of GaAs substrates since the cost of the growth equipment, the starting materials, and the growth rates are similar; however, it must also be recognized that substrate cost is only one ingredient in determining the manufacturing cost of a given device for a given application’s requirement in terms of device performance, reliability, and form factor (Fig. 1.1).

GaN

Substrate Cost ($ / cm2)

$100.00

$10.00 SiC

InP

Sapphire $1.00 GaAs

$0.10 Silicon $0.01 1.0E+05

1.0E+06

1.0E+07

1.0E+08

1.0E+09

1.0E+10

1.0E+11

Estimated Annual Wafer Production (cm2)

Fig. 1.1 Estimated substrate cost per square centimeter vs. annual wafer production for semiconductor substrates

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Table 1.3 Applications for GaN-based LED, power, and laser devices LED applications Power applications Laser diode applications Backlighting (mobile RF power amplifiers for Optical data storage devices, LCD displays) communications and radar Full-color, large-format Switch mode power Projectors video screens supplies Interior and exterior Motor controls Displays automotive lighting Signals and signage Satellite power systems Commercial printing; Lighting for general High-temperature Testing and measurement illumination electronics (spectroscopy, sensing)

Bulk GaN substrate manufacturers have recognized that intermediate and niche opportunities exist and can be leveraged to help meet the needs performance and cost requirements of the larger SSL and power electronics markets. It is these near and intermediate term markets that present the chance to establish manufacturing capacities, grow commercial volumes, increase manufacturing yield, reduce cost, and thus enable access to new and larger markets through growth and continuous improvements (see Table 1.3 for applications for GaN-based devices). Initial target market niches are those that will command a premium for higher performance devices, or benefit from some other aspect of the device, such as higher electrical efficiency, smaller size, and/or less temperature-sensitive performance. These initial markets will rapidly adopt devices that utilize a bulk GaN substrate as they address a need unfilled or poorly met by competitive approaches. We will now examine some of the factors that will contribute to bulk GaN substrates meeting market requirements and how these drivers will facilitate their entry into additional markets.

1.2.4 Key Drivers for Bulk GaN Substrate Commercialization Success One of the areas where GaN substrates provide a clear and immediate benefit is in GaN laser diodes (LDs). The applications for GaN-based LDs are primarily for next-generation, high-definition DVD players and recorders, although other applications include high definition video projectors and displays, commercial printing, and testing and measurement applications, such as spectroscopy and bio-sensing. The high-definition DVD systems utilize a 405 nm laser diode to read and write data on an optical disc, and GaN is uniquely suited to address this wavelength for semiconductor lasers. The format war between Blu-ray and HD DVD in the early 2000s is thought to have potentially hindered the development of the market for these next generation players [12], but in early 2008 the industry settled on the Bluray format, paving the way for increased volumes of laser diodes and optical drives. Sales of Blu-ray DVD players and recorders were expected to reach 5 million units

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in 2008 and were predicted to more than double to 12 million units in 2009 [13], while the GaN-based laser diode market is expected to grow to $1.2 billion by 2011 [14]. Read-only LDs for Blu-ray disc players typically operate at 20 mW [15] and have been developed with device lifetimes exceeding 10,000 h under continuous wave (CW) operation, but developing higher power lasers will enable faster write speeds in DVD recorders. Read/write powers for laser diodes vary from 125 mW for 2 standard write speeds to 170 mW (pulsed) for 4 standard write speeds [16] and up to 250 mW for 8 write speeds [17]. For LDs with higher output power the amount of heat generated during operation becomes significant. Extending the reliable performance of LDs to higher output powers requires improvements in material quality [18]. Manufacturing of GaN-based lasers remains focused in Japan. In 2008, Nichia Corporation, the leading producer of GaN LDs, was reported to hold an 80% market share of the GaN-based LD market [19]. Nichia is reported to use both a sapphirebased process and a bulk GaN substrate process [20], and has worked with Sony in developing and licensing GaN LD technology. Sony, along with Sharp and Sanyo, are developing LDs using a GaN substrate-based manufacturing process [21]. Significant challenges in the manufacturing of GaN-based LDs slowed the ramp to volume production and delayed the delivery of the Sony Playstation 3 and Blu-ray players. The main problem in manufacturing GaN LDs is getting a high production yield, which is reportedly due in part to the limited availability of a high-quality bulk GaN substrate [22]. Manufacturing approaches for GaN LDs based on sapphire have higher dislocation densities and require careful placement and alignment of the devices on the substrate. While virtually no data has been publicly released regarding the manufacturing yields for GaN-based LDs on sapphire or bulk GaN substrates, it is expected that the use of a high quality bulk GaN substrate will improve the yield of an optimized epitaxial growth process due to lower defects, improved uniformity, and simplified epitaxial growth processes. Empirical evidence of this is seen in the move from sapphire to GaN substrates by all LD manufacturers, particularly for high-power LDs. Markets for the GaN LDs capture the several key drivers for success in bulk GaN substrate commercialization:  Strong and growing commercial demand for advanced III-N semiconductor

devices for a broad range of applications  Improvements in bulk GaN materials properties that enable the device structural

quality and thermal properties not attainable through other approaches  Demonstration of high-performance devices with an advantage derived from the

properties of the bulk GaN substrate GaN substrate manufacturers will use these market drivers to justify the deployment of resources for increasing the manufacturing capacity that results in the increased availability of high-quality bulk GaN substrates. This supply of substrates enables growth in the commercial device market, which motivates further substrate capacity expansion and makes possible manufacturing yield improvements. As this GaN substrate manufacturing sequence progresses, other new device applications can benefit

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700 000 # of 4" GaN wafers to process per year for power electronics applications in units

600 000

4" wafers (Units)

500 000 400 000 300 000 200 000 100 000 0 2007

2008

2009

2010

2011

2012

2013

2014

2015

Fig. 1.2 400 GaN wafer forecast for power electronics applications Source: Yole Développement

from the increased availability of bulk GaN substrates through the reduction of R&D costs and a stable supply of substrates. When new devices are developed and demonstrated and new areas of commercial demand are generated, this can provide further motivation for substrate manufacturers to increase manufacturing capacities. For example, Fig. 1.2 shows the forecast for 400 GaN wafer volume for power electronics applications, with unit volumes at nearly 600,000 in 2015. This volume forecast represents a tremendous opportunity for GaN substrates, but will require significant cost reduction in order to meet expected cost requirements.

1.3 Benefits and Importance of Bulk GaN Substrates Within the different families of semiconductor materials, we observe that essentially all silicon, SiC, GaAs, and InP devices are fabricated on their associated bulk substrate, in spite of many attempts to develop SiC and especially GaAs and InP devices on cheaper substrates such as silicon and sapphire. The situation for nitride devices is different: most commercial nitride semiconductor devices are currently fabricated on foreign substrates such as Si, SiC, and sapphire, which reflects the slow rate of progress in the worldwide effort to develop a source of costeffective, high-quality bulk GaN substrates, and the resulting limited availability of these substrates. The selection of foreign substrates for the development and commercialization of GaN devices represents a compromise in the suitability of the substrate as a starting material for GaN epitaxial growth. In spite of some advantages of foreign substrates, which include readily available substrates of large size,

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1µm

0.25µm

Fig. 1.3 Atomic force microscopy images of epitaxial GaN films on sapphire (left) and on a bulk GaN substrate (right). The dislocation density of the film grown on sapphire is >109 cm2 and the steps on the surface are pinned by threading dislocations, resulting in nonparallel growth steps. The dislocation density of the film grown on bulk GaN is approximately 106 cm2 and the monolayer steps are nearly parallel and not pinned Source: A. Allerman, Sandia National Laboratories, and R. Dupuis, Georgia Institute of Technology

high quality, and low cost, there are a number of well-known disadvantages to this heteroepitaxial approach: lattice mismatch, thermal expansion mismatch, and chemical incompatibility between the substrate and the epilayers. Nevertheless, due to years of development with these foreign substrate approaches, epitaxial growth processes have been developed that have allowed the commercialization of GaN-based LEDs for many applications. This development path has also led to a number of novel epitaxial growth approaches and substrate configurations that reduce somewhat the dislocation density of the heteroepitaxially grown GaN layers. Examples of these techniques include epitaxial lateral overgrowth (ELOG) and pendeo- or cantilever-epitaxy, but each of these adds to the cost and complexity of the epitaxial process. Many LED manufacturers are transitioning to patterned sapphire substrates (PSS) which improves performance with only a modest cost increase (Fig. 1.3). The main benefits of the bulk GaN substrate specifically address the weaknesses of the foreign substrate approaches. Indeed, the driving force for bringing GaN substrates to the market is strong because of the inherent benefits that most device applications will realize through the application of a bulk substrate. How a bulk GaN substrate influences the cost of making a device is more complex and we discuss that in some detail below.

1.3.1 Device Performance For device performance optimization, there is a simple rule: the best devices come from the highest quality epitaxial device layers. High quality in this case means several things, but most notably that the epitaxial layer structure matches very well to the original design, and that structural imperfections are not present, at least

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significantly. High-quality device epitaxy, in turn, depends on the combination of substrate choice plus the detailed and optimized epitaxial growth approach. As discussed above, a high-quality device epiwafer means high structural quality: accurate thicknesses and compositions of the various device layers, flat heterointerfaces, abrupt doping profiles, and low defect densities. While the primary benefit of a bulk GaN substrate for essentially all GaN devices could simply be stated as that of device epilayer structural quality, it is also sometimes difficult to understand how improved structural quality might impact the performance or reliability of a specific device and application. The answer lies to a large extent in what are the most important properties of the epitaxial device structure and then how the bulk GaN substrate affects those properties. For example, for blue/near-UV laser diodes for optical storage, device output power and reliability must meet certain threshold values for commercial acceptance. Systematic studies have shown that laser diode reliability depends strongly on defect density in the device active regions for a given output power and is accordingly benefited by use of a bulk GaN substrate. For many other devices there exists a limited amount of evidence that a bulk GaN substrate is beneficial. Here, we present in Table 1.4 a summary of demonstrated and predicted benefits of bulk GaN substrates for the more common device applications today, and we examine the benefits of bulk GaN substrates for several specific applications.

Laser Diodes Significant progress has been made in GaN-based LDs since the first demonstration by Nakamura et al. [23]. Much of this work was developed and demonstrated through crystal growth using sapphire as the growth substrate. As a consequence of this heteroepitaxial development approach, the high dislocation density and biaxial strain in the GaN layers has hindered the improvement in the performance of these nitride-based laser diodes. Dislocations can cause deterioration in the operation of quantum well based LDs mainly by three mechanisms [24]: (a) by serving as nonradiative recombination centers for electrons and holes leading to heat generation instead of optical emission; (b) by introducing fast diffusion along the dislocation lines, smearing out quantum wells and shorting p–n junctions; and (c) by disturbing the epitaxial growth front, so that atomically flat structures cannot be obtained. Dislocations also cause an increase in the threshold current density of the laser diode and ultimately limit the lifetime of the devices [25, 26]. It is well known that the device lifetime is increased when LDs are fabricated on the lower dislocation regions realized by the epitaxial lateral overgrowth (ELOG) technique, which can produce GaN with the dislocation density down to the mid-106 cm2 level [27, 28]. Improved results for LDs fabricated on hydride vapor phase epitaxy (HVPE)-grown bulk GaN substrates with dislocation density of less than 105 cm2 have been demonstrated [29]. Also, as discussed previously, for high power LDs, the thermal conductivity of the device affects performance and reliability. The failure of LDs is closely related to the heat generated by joule heating or nonradiative recombination [28]. Thermal issues in LDs are discussed further below.

14 Table 1.4 Summary of bulk GaN substrate demonstrated applications Device application Demonstrated benefits Laser Increased lifetime; reduced threshold current; higher output power; reduced leakage current; improved thermal conductivity; reduced dislocation enhanced diffusion LED Reduction of nonradiative recombination centers; reduced leakage current; improved thermal conductivity; improved efficiency; decreased In segregation associated with dislocations; reduced aging effects caused by electromigration of metals along dislocations RF FET Greater channel mobility, greater channel charge, greater channel mobility  charge product; reduced buffer leakage; reduced surface states Power switches and High breakdown voltage; faster diodes switching speed; higher efficiency; reduced reverse recovery time Photodetectors, solar Low dark currents in cells photodetectors and solar cells; Geiger mode photodetector operation

A.D. Hanser and K.R. Evans and predicted benefits for device Predicted benefits Increased power capabilities and power density

Further improvements in IQE; reduced efficiency droop at longer wavelengths; increased output powers

Higher operating frequencies; higher operating powers; greater reliability

Higher reliability; lower on-resistance; greater power carrying capability; higher channel mobility Improved detectivity; improved photocurrent; improved efficiency

UV and Visible LEDs GaN-based LEDs cover a wide range of wavelengths and it has been shown that device performance (total lumen or power output, and device efficiency) depends to a large extent on the output wavelength and the current density of the device. Blue LEDs, which are combined with phosphors to produce the majority of white light LEDs produced today, perform at the highest efficiency level of GaN-based LEDs. Devices operating at shorter (UV) or longer (green) wavelengths have lower efficiencies. There is still extensive discussion on what mechanisms lead to the changing internal quantum efficiency (IQE) with different wavelengths and operating conditions, but these are the main issues that hinder LED performance. While there are several proposed mechanisms to explain the decrease in LED efficiency with increasing current, including carrier injection and confinement effects [30, 31] piezoelectric polarization effects [32], and Auger loss [33], there is also evidence that the material quality of the device may play a large role in governing the LED

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performance. For visible wavelengths, the emission mechanisms in the InGaN active regions of the LED determine the IQE of the device. At blue wavelengths, the InGaN quantum wells require 15–20% indium content, while the indium content is in the range of 25–35% for green wavelengths. These compositional changes and associated material quality differences are believed to contribute to the efficiency drop seen at longer wavelengths [34]. InGaN is a notoriously difficult material to grow due to its low decomposition temperature, a reported miscibility gap [35], and increasing lattice mismatch with GaN as the indium content is increased. Material quality tends to decrease with increasing indium content, and compositional uniformity in quantum wells tends to decrease as well. Indium-rich regions form in quantum wells and result in carrier localization, which enhances efficiency at blue wavelengths but at longer wavelengths the effective localization is reduced, allowing higher indium content LEDs to be more susceptible to nonradiative recombination at dislocations in InGaN [36]. V-defect formation in the growth of InGaN has also been shown to be associated with threading dislocations [37], and it is believed that reducing the density of such defects will improve the performance of green LEDs [38]. Growth of InGaN via MBE on bulk GaN substrates has been shown to have a layer-by-layer growth mode when no dislocations are present [39]. Reducing the dislocation density for InGaN growth can improve its material quality and should improve LED performance. Nonradiative recombination is also an issue for UV LEDs, where the IQE is lower than visible wavelength LEDs due to the lack of carrier localization [34] and the typically high number of dislocations. UV devices operating near the GaN bandedge benefit significantly from a GaN substrate, while deep UV LEDs .200 nm <  < 300 nm/ may benefit from an AlN substrate [40].

RF FETs RF transistors built of GaN on SiC show reliability issues that have been shown to be related to trapping of carriers in the devices [41], but presently there is only limited proof that a device fabricated on bulk GaN substrates has improved reliability. A notable study involving bulk GaN substrates showed improved GaN FET reliability when growing on bulk GaN vs. SiC substrates [42]. Recent studies by the U.S. Naval Research Laboratory researchers on semi-insulating GaN showed good X-band RF FET behavior and also showed that improved sheet resistance and higher carrier mobilities were generally observed in comparison to similar structures grown on SiC substrates [43]. Ongoing research at Kyma Technologies and elsewhere may provide additional evidence as more bulk semi-insulating GaN substrates are grown on and processed into transistors for high-speed, high-power RF applications.

High-Power Switches High-power switching devices based on GaN substrates such as GaN Schottky diodes show defect density dependent performance and therefore should benefit

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from the lowest possible defect density. Pure screw dislocations have been identified as a path for reverse-bias leakage current in GaN [44, 45]. Growing on a GaN substrate will reduce the total dislocation density and, therefore/hence, the screw dislocation density at the surface, reducing the gate leakage current. Schottky barrier diodes have been fabricated that have shown some benefits of using bulk GaN over sapphire as the substrate material, where improvements in the reverse recover time, on-resistance, reverse breakdown, and reverse leakage current were reported [46,47]. Recent activities by several groups in the power electronics industry, including Velox and International Rectifier, show strong motivation to build GaN power switching devices on Si and sapphire, yet bulk GaN may be required to significantly further advance the performance levels. Several papers by Japanese researchers over the past two years suggest an intense move to develop bulk GaN based power switching devices [48–50].

New Device Applications There is an ample amount of R&D into new device applications for GaN materials including terahertz applications and quantum well intersubband transition based devices such as the quantum cascade laser (QCL) initially developed at Bell Labs [51]. There are plenty of reasons to believe that bulk GaN will be important in these new applications. Indeed, past failed attempts to build a good GaN-based QCL were ascribed to the lack of a good GaN substrate [52]. Other applications include photovoltaics, where researchers have developed UV-blind avalanche photodetectors [53] and InGaN-based solar cells that offer the potential to develop ultra-high-efficiency devices [54]. Reduced dislocation densities enabled by GaN substrates are important for reduced dark currents and improved solar cell photoresponse.

1.3.2 Thermal Conductivity Many of the targeted high-end applications using GaN-based devices require high output power and high current density. Luminaires for SSL call for hundreds and thousands of lumens of light output, which require high current levels in highbrightness LEDs, while laser diodes operate at high current densities .kA=cm2 / for high power outputs. RF FETs in power amplifiers for wireless communications and radar target power densities of 5 W/mm and higher, while switches and diodes for power electronics will operate at high currents and current densities (up to hundreds of amps and several hundred A=cm2 ). All these operating conditions result in high temperatures in the devices through self-heating, and the performance of virtually all these devices is significantly affected by elevated temperatures. Changes in the device temperature can change the device performance, including dominant emission wavelength, intensity, and efficiency for optical devices, and on-resistance, gain, and efficiency in electronic devices.

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High-power GaN devices cannot be operated continuously without proper thermal management and an integral part of managing the heat in a device structure is to determine how the heat propagates from the device action region through the epitaxial device layers. Exploring these issues, researchers have investigated the effect of the thermal conductivity of the substrate and the epitaxial layers on the device junction or channel temperature. Published experimental measurements show that the room temperature thermal conductivity (RT-TC) of GaN can be as low as 1:0 W=cm K for highly defective .1010 cm2 / GaN grown on sapphire or SiC; for high-quality .106 cm2 / bulk GaN substrates the RT-TC has been measured up to 2.3–2.5 W/cm K [55, 56]. As a result, the benefit of a highly thermally conductive substrate such as SiC may be reduced due to the high defect density of the buffer layer and device active region. The impact of the thermal conductivity of defective epitaxial AlGaN/GaN layers grown on SiC for RF applications have been investigated using micro-Raman thermography. The high spatial resolution of the micro-Raman technique (in the micron to submicron range) allows one to see temperatures that are significantly higher in the GaN epi than in the underlying SiC substrate, due to a high thermal impedance in the GaN layer. Additionally, a thermal boundary resistance is observed between the GaN and the substrate in AlGaN/GaN electronic devices [57]. These factors contribute to a high channel temperature in AlGaN/GaN FETs for RF applications, which may have an impact on device performance and reliability. Substrate thermal effect comparisons have been tested in AlGaN/GaN RF FET modeling studies for GaN and SiC substrates with experimental tests conducted in parallel [58]. Thermal models show that the introduction of defects in the epitaxial GaN and AlN layers grown on SiC influences the flow of heat to the substrate and increases the peak channel temperature in the device, and microRaman measurements show similar channel temperatures for the two substrate cases. For RF applications, bulk GaN substrates offer the possibility of similar thermal performance to SiC, while improving the material quality and the potential for improving device reliability. For GaN LDs, improved junction temperatures have been demonstrated using bulk GaN substrates relative to sapphire substrates, which are attributed to the improved thermal conductivity of the GaN substrate [28].

1.3.3 Thermally Activated Device Failure Failure in semiconductor devices due to a thermally induced mechanism with a first-order dependence on a particular defect concentration follows the rule:  1 D AŒD exp.Ea =kT /; where A is the preexponential factor, Ea is the activation energy associated with the failure mechanism, ŒD is the concentration of the defect associated with the rate limiting step in the failure mechanism, and T is the temperature in the region

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of failure. Such behavior has been observed in GaAs-based HBTs [59], solar cells [60], and GaAs-based laser diodes [61]. For GaN devices, researchers at Sony have demonstrated the clear first order dependence of GaN laser diode lifetime on threading defect density in the device active region [27]. For devices that fail due to thermal activation of a mechanism that involves a threading defect, bulk GaN offers a great advantage over foreign substrate based approaches. Other benefits may exist as well, based on increased efficiency and improved thermal conductivity. Reliability studies are underway for several other device types, particularly RF FETs [62], and associated informative publications should result over the next couple of years.

1.3.4 Device Cost The cost of manufacturing a device that has specific performance, reliability, and form factor requirements can be a complex function of the costs of starting materials, epitaxial growth equipment, device fabrication steps and device yield. (Packaging costs and yields and performance and reliability testing costs and yields can also be important but are not discussed here.) These cost factors need to be considered when evaluating the benefit of a bulk GaN substrate. For example, a low-performance LED can be made at low cost on sapphire substrates due to a thin epitaxial layer structure and relatively high yields. On the other hand, a high-performance LED structure generally requires a thicker epitaxial layer structure and may have lower yields in epitaxial growth and processing due to tighter specifications, such as lumen output or wavelength variation. Some increases in certain cost factors result in overall cost reductions: some LED manufacturers have moved to slightly more costly patterned sapphire substrates, and this has resulted in improved LED performance, effectively reducing the overall cost of a higher performance LED. Beyond materials cost considerations, process development costs need to be considered for new growth processes. In the development of GaN device technology, III-N epitaxial growth capabilities preceded the availability of bulk GaN substrates. Because of this, GaN device developers used foreign substrates exclusively, working primarily with sapphire, SiC, and silicon. For other semiconductor materials, such as GaAs, SiC, and InP, a native substrate was available during the device development stage. With these materials, a move away from the native substrate may offer the potential for cost savings but can also present challenges in terms of development cost and yield of epitaxy and device processing. In the case of GaN, device manufacturing processes have been optimized on foreign substrates. As GaN substrates become available, device manufacturers must contend with not only an initially more expensive starting substrate, but the additional cost of reoptimizing growth processes for the new substrate. The use of GaN substrates should simplify the growth process when compared to foreign substrates and may improve device yields, but the optimization process will depend on many factors, such as substrate surface preparation, off-cut angle and orientation, doping, etc. Device

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manufacturers will make the switch to bulk GaN substrates when the benefits of the device performance outweigh the costs of the materials and process development. Improved device yield is potentially the biggest cost benefit for GaN substrates but also the least discussed and the most difficult to predict. GaN LD developers quickly discovered that defects in the epilayers dramatically affected yield through poor device performance and reliability, which provided strong impetus to accelerate the development of bulk GaN substrates [22, 27]. Current developers and manufacturers of LEDs, RF FETs, and power switching electronics utilize foreign substrates primarily, which indicates that device performance requirements are currently met with the lower cost manufacturing approaches. As GaN substrates are further developed and as the GaN substrate cost is reduced, we believe that essentially all GaN devices, except perhaps for the lowest performance LEDs, will eventually transition to GaN substrates, where the rate of changeover will depend on the specific device and the cost, quality, and diameter availability of bulk GaN. In some cases, additional cost is incurred when employing a foreign substrate. In LED manufacturing, there is a trend toward removing the substrate to improve light extraction and possibly to improve thermal conductivity [63]. RF FETs fabricated from GaN on Si have thermally conductive epoxies used to contact the backside of the epi nominally to improve heat removal [64]. As discussed earlier, bulk GaN enables good thermal conductivity in both the buffer layer and the device active region, which is to be contrasted with poor thermal conductivity in GaN buffers and device layers grown on foreign substrates, including SiC. History shows that all devices get pushed harder in terms of operating conditions as their respective markets mature, and thus all devices become thermally limited. As a result, while GaN on Si RF FETs operating at a moderate level of performance can be made cheaper today than on bulk GaN, that will change as GaN substrate costs come down and diameter availability increases. In some cases the number of RF FETs used can be reduced by increasing the power per FET, which will eventually be achievable by transitioning to bulk GaN substrates.

1.4 GaN Device Trends for Bulk GaN Substrates 1.4.1 Lasers and LEDs The markets for GaN-based optoelectronic devices are large and will continue to generate new applications for LEDs and LDs. For example, many systems would greatly benefit from the availability of LDs with wavelengths >500 nm, including laser-based video projectors and analytical systems. Extrapolations of performance and material quality requirements toward longer wavelengths based on observed trends in blue/violet optical storage laser development are such that most experts agree that a low-defect-density free-standing GaN substrate is a necessary building block of such a laser diode. However, conventional c-plane quantum well structures

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suffer from the undesirable quantum-confined Stark effect (QCSE), due to the existence of strong piezoelectric and spontaneous polarizations. The strong built-in electric fields along the c-direction cause spatial separation of electron and holes that in turn give rise to lower carrier recombination efficiency, reduced oscillator strength, and red-shifted emission. The QCSE becomes more pronounced with increasing indium composition in InGaN QWs needed for longer wavelength LEDs. Avoiding the QCSE is important for high efficiency LDs and LEDs and has led to intense investigation of nonpolar and semipolar GaN orientations. Nonpolar orientations of GaN eliminate the spontaneous polarization fields, the absence of which has been confirmed experimentally for both violet laser diodes [65, 66] and blue LEDs [67] and suggest nonpolar substrate orientations may benefit longer wavelength GaN-based LDs. Semipolar orientations mostly eliminate the spontaneous polarization effect and may be easier to achieve high-quality epitaxy [68]. It is believed that a high-quality nonpolar and semipolar GaN substrates represent a potential enabler of GaN-based LDs with wavelengths >500 nm. Based on recent results, power LEDs for SSL will continue to be developed on sapphire substrates for the foreseeable future. Companies such as Nichia, Cree, and Osram continue to report improvements in LED efficacy at increasing output power [69], and in total lumen output [70]. Surface emission LED designs improve light extraction efficiency, and removal of the GaN epitaxial layers from the underlying substrate is simplified through the use of laser lift-off techniques with a transparent sapphire substrate [71]. While many groups are performing R&D on LEDs on GaN substrates, fundamental limits in LED performance due to material quality have yet to be definitively identified, and so the demand for a GaN substrate for LEDs in a production application has not yet materialized. Panasonic has introduced white LEDs based on blue LEDs grown on GaN substrates, which the company says have very high thermal and electrical conductivity that substantially improved LED performance at high current levels [72]. Also, the brightest white light LEDs will continue to be those based on blue LEDs with phosphors until the efficiency of green LEDs is improved to allow RGB white light sources to perform at comparable levels.

1.4.2 Power Switches Normally-off switches for HPE are of growing interest and many market applications will benefit from FET device designs that enable a zero bias off-state and high breakdown voltage and current handling capabilities. Vertical power Field effect transistors (FETs) device designs enable smaller chip sizes at higher voltages than lateral designs and are the preferred design approach in most cases. Vertical devices make a conductive substrate necessary, while low leakage currents and high breakdown voltage require low defect densities. As a result, high-quality conductive bulk GaN substrates are desirable for these devices. Several groups are investigating the use of an AlGaN/GaN high electron mobility transistor (HEMT) structure in power

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FETs to take advantage of the structure’s high carrier mobility and low sheet resistance, although the HEMT structure is a normally-on device [50]. GaN MOSFET designs are also being investigated, although there are many challenges in developing a viable gate dielectric for GaN. Standard SiO2 gate dielectrics have generated devices that exhibit very good channel mobilities [73], and other dielectrics are being investigated for improved performance [74]. A fully integrated approach of high-quality GaN materials enabled by GaN substrates, along with gate dielectrics developed and optimized for the GaN materials system, would make possible highefficiency, high-performance, normally-off FET devices that meet standard design requirements for power electronic circuits.

1.4.3 High-Frequency, High-Power HEMTs GaN-based RF devices for wireless communications systems outperform Si-based devices in terms of operating voltage and power density, but cannot yet match Si in cost. As a result, Si remains very competitive in wireless base station applications in the 2.4–5.8 GHz frequency range. With cost being a large driver for the commercial applications, lower cost manufacturing approaches for GaN RF devices dominate, which include GaN devices fabricated on Si and SiC substrates. These device architectures have delivered the performance and reliability needed to enter the market and are expected to continue to improve with the ongoing development. Thus, the demand for GaN substrates in high-frequency applications is relegated to frequencies starting at 10 GHz. Since the exact relation between RF device performance and material quality is not clearly known, additional development is needed to identify benefits associated with bulk GaN substrates. These higher frequency applications are smaller in market size compared to the LED, LD, and HPE markets, and likely present only niche opportunities for bulk substrate manufacturers [75]. A key challenge for GaN RF FET developers is realization of a normally off device; nonpolar GaN offers clear potential for such an “enhancement mode” or emode FET to be developed. As nonpolar GaN becomes more available, emode GaN FET development will become important.

1.5 Bulk GaN Substrate Trends The bulk GaN substrate industry is just at the start of establishing and expanding techniques for manufacturing substrates for commercial applications. Researchers continue to explore methods of improving and optimizing existing crystal growth techniques for fabricating bulk GaN substrates, while exploring new methods for creating larger, high-quality crystals with the desired structural, optical and electrical properties at an even lower cost. Three main techniques are currently employed for the development and manufacturing of bulk GaN crystals for substrates. We will

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briefly examine these three crystal growth techniques and the trends for their use in producing substrates for GaN devices.

1.5.1 Hydride Vapor Phase Epitaxy The HVPE growth process has several benefits for manufacturing GaN substrates, including high growth rates (up to 300 m=h have been demonstrated for manufacturing), high purity, and doping control, which make bulk materials a possibility and have led to successful substrate commercialization. The HVPE manufacturing method currently dominates in the delivery of GaN substrates to the market in terms of volume. There are a number of challenges, however, that the HVPE process faces in achieving high quality bulk substrates. GaN substrates tend to have either an engineered variation of the quality and material structure of the substrate, or a domain-like structure with a number of subgrains and low angle grain boundaries within the crystal. Due to these defects in the GaN, the crystal quality is nonuniform and the dislocation density can vary across the substrate and from top to bottom. As a result, the materials uniformity of HVPE material has yet to approach that of higher structural quality, albeit smaller, crystals grown by other processes, such as the ammonothermal or solution growth techniques. Additionally, crystal lattice tilt, wafer bow, and precise crystalline orientation control over larger area substrates remain as challenges. Steady improvements to the GaN HVPE process has led recently to the demonstration of long length boules, on the order of 1 cm in the [0001] direction, enabling the development of nonpolar substrates to be sliced from c-plane oriented crystals. These substrates have greatly improved material quality relative to high defect density heteroepitaxial material grown in nonpolar directions, although the size of these nonpolar substrates is small (up to 10 mm10 mm [56]). This remains the primary limitation for this technology at this time. Increasing the size of bulk nonpolar substrates will occur through several developments, including extended c-direction growth or growth on nonpolar GaN seeds. Significant issues remain, however, in maturing both of these processes to achieve longer and larger crystals. Once optimized though, GaN seeded boule growth by HVPE is expected to produce excellent quality material at an attractive cost. More detail on the HVPE growth process for bulk GaN substrates is found in Sect. II of this book.

1.5.2 Ammonothermal Growth The ammonothermal (or solvothermal) growth process is similar to the hydrothermal growth process used for the growth of quartz crystals in supercritical water, where elevated heat and pressure is used to enhance the solubility of the solute. In the ammonothermal growth process, supercritical ammonia in conjunction with

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special mineralizers is used as the solvent to dissolve and transport GaN to seed crystals where growth occurs through recrystallization. Solvothermal processes have been demonstrated on very large scales and with very high crystal quality for other materials, including quartz and zinc oxide (ZnO). The scale of the manufacturing process offers great potential for cost reduction, which is one of the strengths of the approach. Additionally, very high-quality GaN crystals with low dislocation densities and low lattice tilt have been demonstrated using the ammonothermal growth process [76]. Drawbacks to the growth technique include a slow growth rate .100 m=day/, the extended development time for developing the growth chemistry, and scaling the process to obtain the economies of scale needed to reduce the manufacturing cost. Details of the ammonothermal process are examined in Sect. III of this book.

1.5.3 Solution Growth Solution (or flux) growth techniques are liquid phase processes that involve solutions that dissolve the growth feedstock and provide transport of the growth species for recrystallization and growth on seed crystals. In high-pressure solution growth using a gallium melt as the solvent, nitrogen is dissolved at high pressures (1 GPa or 10 kbar) and temperatures .1500 ıC/ to form crystals that are either thin hexagonal platelets or elongated hexagonal needles [77]. In low-pressure solution growth, a solvent containing gallium is used to dissolve nitrogen and transport it to seed crystals where GaN crystallizes and grows. These techniques are typically conducted at low to moderate pressures .50 bar/ and temperatures .750 ıC/ [78], with several variations on the process. The driving force for growth from the solution is a thermal gradient within the growth system and growth occurs through precipitation and recrystallization due to a change in solubility resulting from the change in temperature. While the high-pressure process appears to have significant barriers to mass production, the low-pressure processes show promise. Strengths of the solution growth process are the near-equilibrium growth conditions under which the crystal grows, a higher growth rate than the ammonothermal process (up to 240 m=day [79]), and the low pressure and temperature are beneficial. There is also the potential for scaling the process, as solution growth processes have been demonstrated in other materials systems for high-quality, low-cost, large crystals. Challenges in the development of a GaN solution growth process include a relatively low solubility of nitrogen, complex solvent chemistries, and impurity control. An overview of flux growth processes is presented in Sect. IV of this book.

1.5.4 Combined Growth Techniques While MOCVD templates on foreign substrates have been used extensively as seeds for crystal growth development, researchers are beginning to recognize the possible

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benefits of combining features from different growth techniques to enhance the growth of bulk GaN crystals for substrates. For example, the use of high-quality seeds grown by either the solution growth process or the ammonothermal process can be used to supply seeds for the HVPE process, where higher growth rates and larger existing manufacturing capacities can produce high-quality substrates more rapidly. Additionally, the HVPE process can generate high-purity GaN as a source material for the solution and ammonothermal processes. These methods for crystal growth are just now beginning to be explored.

1.6 Summary The III-N materials enable new semiconductor devices with previously unobtainable performance capabilities in terms of light output, power handling, and efficiency, and these attributes will make possible the reinvention of existing technologies in ways that benefit many facets of our lives. While there are many potential applications for these materials, the biggest applications appear to be light generation and the control of electrical power. Bulk GaN substrates are expected to be used for many devices addressing these market opportunities, but there are numerous challenges facing the adoption of bulk GaN substrate technology. Considerable technology development is needed to enable penetration of GaN devices into the SSL and power electronics markets. The performance of the GaN devices must improve in many areas: LED luminous efficacy at high lumen output; FET and diode breakdown voltage and on-resistance; and power switching reliability, among others. Substrate cost must also be significantly reduced in order to become economically practical in the device applications, which is particularly challenging due to incumbent GaN device technology based on foreign substrates. Bulk GaN substrates provide clear and immediate benefits for GaN-based LD applications over sapphire-based approaches, particularly at higher operating current density, and GaN substrate manufacturers are currently focused on addressing this market. Other benefits for LEDs, RF FETs, power devices, and photodetectors have been shown, but commercialization of these devices requires high volumes of low cost substrates. These improvements will occur through the continued development of bulk growth techniques, such as HVPE, ammonothermal, solution growth, and combinations thereof. As these improvements are made, the potential for III-N devices to do more (more light, more power) with less (higher energy efficiency, less electricity, less waste heat) will be enabled. Acknowledgements The authors would like to thank their colleagues Tanya Paskova and Edward Preble at Kyma Technologies for their valuable contributions to the research and writing of this chapter. We would also like to thank Philippe Roussel of Yole Développement for the power electronics market data and Andrew Allerman of Sandia National Laboratories and Russell Dupuis of the Georgia Institute of Technology for the AFM data.

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33. Y.C. Shen, G.O. Mueller, S. Watanabe, N.F. Gardner, A. Munkholm, M.R. Krames, Appl. Phys. Lett. 91, 141101 (2007) 34. Y. Yang, X.A. Cao, C. Yan, IEEE Trans. Electron. Dev. 55, 1771 (2008) 35. I. Ho, G.B. Stringfellow, Appl. Phys. Lett. 69, 3701 (1996) 36. Y.-H. Cho, S.K. Lee, H.S. Kwack, J.Y. Kim, K.S. Lim, H.M. Kim, T.W. Kang, S.N. Lee, M.S. Seon, O.H. Nam, Y.J. Park, Appl. Phys. Lett. 83, 2578 (2003) 37. Y. Chen, T. Takeuch, H. Amano, I. Akasaki, N. Yamada, Y. Kaneko, S.Y. Wang, Appl. Phys. Lett. 72, 710 (1998) 38. C. Wetzel, T. Salagaj, T. Detchprohm, P. Li, J.S. Nelson, Appl. Phys. Lett. 85, 866 (2004) 39. Unipress InGaN MBE 40. Z. Ren, Q. Sun, S.-Y. Kwon, J. Han, K. Davitt, Y.K. Song, A.V. Nurmikko, W. Liu, J. Smart, L. Schowalter, Phys. Stat. Solidi (c) 4, 2482 (2007) 41. M. Faqir, G. Verzellesi, G. Meneghesso, E. Zanoni, F. Fantini, IEEE Trans. Electron Dev. 55, 1592 (2008). 42. K.K. Chu, P.C. Chao, M.T. Pizzella, R. Actis, D.E. Meharry, K.B. Nichols, R.P. Vaudo, X. Xu, J.S. Flynn, J. Dion, G.R. Brandes, IEEE Electron Dev. Lett. 25, 596 (2004) 43. D.F. Storm, D.S. Katzer, J.A. Roussos, J.A. Mittereder, R. Bass, S.C. Binari, D. Hanser, E.A. Preble, K.R. Evans, J. Cryst. Growth, 301–302, 429 (2007) 44. J.W. Hsu, M.J. Manfra, R.J. Molnar, B. Heying, J.S. Speck, Appl. Phys. Lett. 81, 79 (2002) 45. B.B. Simpkins, E.T. Yu, P. Waltereit, J.S. Speck, J. Appl. Phys. 94, 1448 (2003) 46. K.H. Baik, Y. Irokawa, J. Kim, J.R. LaRoche, F. Ren, S.S. Park, Y.J. Park, S.J. Pearton, Appl. Phys. Lett. 83, 3192 (2003) 47. Y. Zhou, M. Li, D. Wang, C. Ahyi, C.-C. Tin, J. Williams, M. Park, N.M. Williams, A. Hanser, Appl. Phys. Lett. 88, 113509 (2006) 48. H. Lu, R. Zhang, X. Xiu, Z. Xie, Y. Zheng, Z. Li, Appl. Phys. Lett. 91, 172113 (2007) 49. T. Uesugi, R&D Review of Toyota CRDL 35 1 (2000.6) 50. H. Ueda, M. Sugimoto, T. Uesugi, and T. Kachi, Wide-bandgap semiconductors for automobile applications, In CS MANTECH Conference, April 24–27, 2006, Vancouver, British Columbia, Canada 51. C. Gmachl, H.M. Ng, S.N.G. Chu, A.Y. Cho, Appl. Phys. Lett. 77, 3722 (2000) 52. G. Sun, R.A. Soref, J.B. Khurgin, Superlattices Microstruct. 37, 107 (2005) 53. J.B. Limb et al. Appl. Phys. Lett. 89, 011112 (2006) 54. O. Jani, H. Yu, E. Trybus, B. Jampana, I. Ferguson, A. Doolittle, C. Honsburg, Effect of phase separation on performance of III–V nitride solar cells, In 22nd European Photovoltaic Solar Energy Conference, September 3–7, 2007, Milan, Italy 55. C. Mion, J.F. Muth, E.A. Preble, D. Hanser, Appl. Phys. Lett. 89, 092123 (2006). 56. S. Kubo, K. Fujito, H. Nagaoka, T. Mochizuki, H. Namita, S. Nagao, Bulk GaN crystals grown by HVPE, In Second International Symposium on Growth of III-Nitrides, July 6–9, 2008, Laforet Shuzenji, Izu, Japan 57. A. Sarua, H. Ji, M. Kuball, M.J. Uren, T. Martin, K.P. Hilton, R.S. Balmer, IEEE Trans. Electron Dev. 53, 2438 (2006) 58. A. Sarua, H. Ji, K.P. Hilton, D.J. Wallis, M.J. Uren, T. Martin, M. Kuball, IEEE Trans. Electron Dev. 54, 3152 (2007) 59. A.D. Hanser, E.A. Preble, L. Liu, K. R. Evans, B.K. Wagner, H.M. Harris, Thermal benefits of native GaN substrates, In Advanced Technology Workshop on Advanced Substrates and Next-Generation Semiconductors, April 30–May 1, 2008, Linthicum Heights, Maryland, USA 60. M. Harris, B. Wagner, S. Halpern, M. Dobbs, C. Pagel, B. Stuffle, J. Henderson, K. Johnson, Full two-dimensional electroluminescent (EL) analysis of GaAs/AlGaAs HBTs, In Proc. IEEEIRPS, vol. 121 (1999) 61. C. G. Michel, G. J. Vendura, H. S. Marek, J. Elect. Mat. 16, 295 (1987) 62. G. Mueller, P. Berwian, E. Buhrig, and B. Weinert, GaAs substrate for high power laser diodes, in High Power Diode Lasers, ed by R. Diehl, Topics Appl. Phys. 78, 121 (2000) 63. M.J. Rosker, in IEEE Compound Semiconductor Integrated Circuit Symposium 2007, pp 1–4 64. I.P. Smirnova, L.K. Markov, D.A. Zakheim, E.M. Arakcheeva, M.R. Rymalis, Semicond. 40, 1363 (2006)

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65. http://compoundsemiconductor.net/cws/article/fab/36330. (October 2008) 66. K. Okamoto, H. Ohta, S.F. Chichibu, J. Ichihara, J. Takasu, Jpn. J. Appl. Phys. 46, L190 (2007) 67. M.C. Schmidt et al. Jpn. J. Appl. Phys. 46, L190 (2007) 68. J.P. Liu et al. Appl. Phys. Lett. 92, 011123 (2008) 69. H. Zhong, A. Tyagi, N.N. Fellows, F. Wu, R.B. Chung, M. Saito, K. Fujito, J.S. Speck, S.P. DenBaars, S. Nakamura, Appl. Phys. Lett. 90, 233504 (2007) 70. http://www.ledsmagazine.com/news/5/7/22. (October 2008) 71. http://www.ledsmagazine.com/news/4/9/9. (October 2008) 72. O.B. Shchekin, J.E. Epler, T.A. Trottier, T. Margalith, D.A. Steigerwald, M.O. Holcomb, P.S. Martin, M.R. Krames, Appl. Phys. Lett. 89, 071109 (2006) 73. http://www.ledsmagazine.com/news/4/3/7. (October 2008) 74. W. Huang, T.P. Chow, and T. Khan, Phys. Stat. Sol (a) 204, 2064 (2007) 75. J.S. Jur, V.D. Wheeler, M.T. Veety, D.J. Lichtenwalner, D.W. Barlage, and M.A.L. Johnson, Epitaxial rare earth oxide growth on GaN for enhancement-mode MOSFETs, in CS MANTECH Conference, April 14–17, 2008, Wheeling, Illinois, USA 76. Yole Développement, 400 epi-wafer needs for GaN-based RF devices market 2005–2012, 2008 77. R. Dwili´nski, R. Doradzi´nski, J. Garczy´nski, L.P. Sierzputowski, A. Puchalski, Y. Kanbara, K. Yagi, H. Minakuchi, H. Hayashi, J. Cryst. Growth 310, 3911 (2008) 78. I. Grzegory, M. Bo´ckowski, S. Porowski, in Bulk Crystal Growth of Electronic, Optical, and Optoelectronic Materials, ed. by P. Capper (Wiley, UK, 2005), Chap. 6, pp. 173–207 79. F. Kawamura, M. Morishita, M. Tanpo, M. Imade, M. Yoshimura, Y. Kitaoka, Y. Mori, T. Sasaki, J. Cryst. Growth 310, 3946 (2008)

Part II

Vapor Phase Growth Technology

Chapter 2

Hydride Vapor Phase Epitaxy of GaN Akinori Koukitu and Yoshinao Kumagai

Abstract Hydride vapor phase epitaxy (HVPE) is the growth method known for a long time, where a halide vapor precursor, such as GaCl, and a hydride, such as NH3 , are used as the group III and the group V precursors, respectively. HVPE growth can achieve a high growth rate and a high crystal quality due to thermal stability of a halide source, the use of high purity starting materials without containing carbon, and a high surface migration of halide molecules. Therefore, HVPE growth attracts attention as a thick film growth method.

2.1 Introduction The first hydride vapor phase epitaxy (HVPE) of GaN was reported by Maruska et al. in 1969, using a sapphire substrate [1]. The growth temperature of GaN was 825 ıC, which is extremely lower than that used for a high quality GaN HVPE growth recently. The GaN epitaxial layer was enough for measuring the band-edge absorption and lattice constant the first time, although epitaxial layer should not be a high quality GaN from the viewpoint of growth temperature. In the 1990s, HVPE growth of GaN attracted attention for the preparation of freestanding GaN wafers. The first thick GaN (approximately 400 m) with a smooth surface was reported by Detchprohm et al. in 1992, using a ZnO buffer layer on sapphire [2]. Although the epitaxial layer of GaN on the ZnO buffer layer was transparent and smooth, the ZnO buffer layer was so thin that the ZnO layer could not be etched off by a solution. From the late 1990s, most of the attention has been focused on the heteroepitaxy of GaN on foreign substrates such as sapphire, Si, NdGaO3 , and GaAs to realize freestanding GaN wafers [3–7]. Usui et al. [6] achieved preparation of a freestanding GaN wafer of 2-in. (approximately 5 cm) diameter size by the use of a new technique, epitaxial lateral overgrowth (ELO), to reduce significantly the density of dislocations, and by a laser irradiation from behind to remove the GaN thick layer from the sapphire substrate. Unfortunately, the laser irradiation method did not seem to satisfy the need for high reproducibility in mass production.

31

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A. Koukitu and Y. Kumagai

An important characteristic of GaN is that depending on the atomic configuration of the substrate, it crystallizes in either a wurtzite structure having hexagonal symmetry or a zincblende structure with cubic symmetry. From the point of view of atomic configuration, the use of GaAs as the substrate has some advantages, i.e., (a) the atomic configuration on the GaAs (100) surface is equivalent to the cubic GaN structure, (b) the atomic configuration on the GaAs (111) surface is equivalent to the wurtzite GaN structure, (c) the difference of the thermal expansion coefficient between GaAs and GaN is about a quarter of that between sapphire and GaN, and (d) the GaAs substrate is easy to lap and etch off to separate the GaN layer from the GaAs substrate. In this chapter, we describe HVPE of GaN using the GaAs substrate. First, in order to get the information on suitable growth conditions and understand a growth mechanism, a thermodynamic analysis of HVPE is described for GaN. Then, we show the HVPE growth of cubic GaN on GaAs (100) substrates, and the comparison of GaN growth on (111)A and (111)B GaAs substrates. The orientation dependence of GaN buffer layer growth on GaAs is also investigated using density-function theory (DFT) to better understand GaN growth on GaAs. Next, it is shown that the growth of thick GaN layers on GaAs (111)A surface at 1;000 ı C or above is promising for the preparation of a freestanding GaN substrate. Finally, characteristics of undoped and Fe-doped freestanding GaN substrates obtained by removing GaAs (111)A substrates are presented.

2.2 Thermodynamic Analysis on HVPE Growth of GaN Knowledge of the mechanisms that determine the growth rate is important for the growth of compound semiconductors. It provides us with information on suitable growth conditions for the preparation of compound semiconductors, as well as insight into the crystal growth mechanism. We have shown that an equilibrium model is useful for predicting the growth rate and the element incorporation of III–V compound semiconductors grown by vapor phase epitaxy (VPE) [8–17], including metalorganic vapor phase epitaxy (MOVPE), HVPE and molecular beam epitaxy (MBE). In this chapter, thermodynamic analysis on the HVPE growth of GaN is described for understanding the equilibrium partial pressures of vapor species and the driving force for the GaN deposition by HVPE.

2.2.1 Calculation Procedure The following six species were chosen as the necessary vapor species in analyzing the HVPE growth of GaN: GaCl, GaCl3 , NH3 , HCl, H2 and inert gas (IG) such as nitrogen or helium. The chemical reactions which connect the species at

2

Hydride Vapor Phase Epitaxy of GaN

33

the deposition zone are: GaCl.g/ C NH3 .g/ D GaN.s/ C HCl.g/ C H2 .g/ GaCl.g/ C 2HCl.g/ D GaCl3 .g/ C H2 .g/

(2.1) (2.2)

In HVPE, GaCl as the major source material of Ga is generally formed by the reaction between Ga metal and HCl at the source zone heated up about 800–900 ıC. The growth equipment of HVPE becomes complex compared with those of MOVPE and MBE, because two zones called source zone and deposition zone are required in the HVPE growth reactor. On the other hand, metalorganic hydrogen chloride vapor phase epitaxy (MOHVPE) [5, 18] was developed with a view to simplifying the growth reactor. In this system, alkyl gallium (trimethyl or triethyl gallium, TMG or TEG) and ammonia .NH3 / are used as Ga and nitrogen sources, respectively. Alkyl Ga is supplied with H2 , and mixed with HCl in the upstream region called mixing zone of the reactor heated at 750 ıC. The reaction between alkyl Ga and HCl forms gaseous GaCl, which is transported to the downstream region called deposition zone. When GaCl and NH3 are mixed, the final reaction for the GaN growth is the same as that in conventional HVPE growth, as shown in reaction (2.1). The equilibrium equations for reactions (2.1) and (2.2) are K1 D

PHCl PH2 ; PGaCl PNH3

(2.3)

K2 D

PGaCl3 PH2 : 2 PGaCl PHCl

(2.4)

and

The total pressure of the system and the stoichiometric relationship for III-nitride deposition in the vapor phase can be written as X

Pi D PGaCl C PGaCl3 C PNH3 C PHCl C PH2 C PIG ;

(2.5)

and o o PGaCl  PGaCl D PNH  PNH3 ; 3

(2.6)

where P ı i ’s indicates the input partial pressure of the source materials. In addition, the parameter A and F are introduced. AD

1=2PGaCl C 3=2PGaCl3 C 1=2PHCl ; PH2 C 3=2PNH3 C 1=2PHCl C PIG

(2.7)

F D

PH2 C 3=2PNH3 C 1=2PHCl : PH2 C 3=2PNH3 C 1=2PHCl C PIG

(2.8)

34

A. Koukitu and Y. Kumagai

A is the ratio of the number of chloride atoms to the number of hydrogen and IG atoms in the system. F is the mole fraction of hydrogen relative to IG atoms. These parameters are kept invariant under a given growth condition. Thermodynamically, almost all the NH3 is decomposed into N2 and H2 at temperature higher than 300 ı C. However, it is well known that the decomposition rate of NH3 under typical growth conditions is very slow without a catalyst, and the extent of the decomposition strongly depends on the growth conditions and equipment. Therefore, a parameter ’ is introduced into the calculation as the mole fraction of decomposed NH3 . NH3 .g/ ! .1  ’/NH3 .g/ C ’=2N2 .g/ C 3’=2H2 .g/:

(2.9)

In the following calculation, the change of ’ affects the change of H2 ; N2 , and NH3 partial pressures in the growth system. A similar result is produced by changing the parameter F . Consequently, we performed the following calculation by changing the parameter F, and fixed ’ as 0.03, according to Ban’s experiments in hot-wall quartz reactor at 950 ı C [19]. The equilibrium partial pressures can be calculated using (2.3)–(2.8).

2.2.2 Equilibrium Partial Pressure and Driving Force for the GaN Deposition Variation of the equilibrium partial pressures for each vapor species is shown in Figs. 2.1 and 2.2 as functions of growth temperature and input partial pressure of GaCl. The general features of the equilibrium partial pressures are similar to those of the hydride systems of arsenides and phosphides of Ga and In [20,21]. It is seen that NH3 , GaCl, N2 , and HCl except the hydrogen carrier gas are major vapor species, and their pressures increase with an increase of the input partial pressure of GaCl. Since Ga is transported as GaCl from the mixing zone and GaN is deposited to form HCl, the reaction governing the growth at the substrate may be described as GaCl C NH3 D GaN C HCl C H2 :

(2.10)

Figure 2.3 shows the relationship between the partial pressures versus parameter F, including those of H2 and IG. The partial pressure of IG in the figure involves N2 produced by the decomposition of NH3 . In the present calculation, IG brings about the same effect irrespective of the IG species. The partial pressure of HCl increases at the expense of GaCl while decreasing the mole fraction of H2 in carrier gas. This means that GaN tends to deposit into the solid phase with an increase of the IG mole fraction. Consequently, hydrogen plays an important role in the deposition of GaN. In Fig. 2.4, the driving force for the GaN deposition, PGa ŒP ı GaCl  .PGaCl C PGaCl3 / is shown as a function of temperature. The driving force for the deposition decreases with an increase of F . Therefore, it is expected that a higher growth rate is obtained in the inert carrier gas system than in the H2 system. This fact agrees

2

Hydride Vapor Phase Epitaxy of GaN

35

SPi:1.0 atm PGaCl :5.0×10–3 atm V/III:50 F:1.0 a:0.03

1

H2 NH3

10–2

GaCl

Partial Pressure (atm)

N2 HCl

10–4

10–6

10–8 GaCl3

10–10

10–12 500 600 700 800 900 1000 1100 1200

Temperature (°C) Fig. 2.1 Equilibrium partial pressures over GaN as a function of growth temperature. Total pressure: 1.0 atm, input partial pressure of GaCl: 5  103 atm, input V/III ratio: 50, F : 1.0 and ’: 0.03

1

SPi:1.0 atm

Temp.:1000°C F:1.0 a:0.03 H2

10–2

NH3

GaCl N2

Partial Pressure (atm)

V/III:50

HCl

10–4 10–6 10–8

GaCl3

10–10 10–12 10–14 0.001

0.005

0.01

0.015

Input Partial Pressure of GaCl (atm) Fig. 2.2 Equilibrium partial pressures over GaN as a function of input partial pressure of GaCl at 1;000 ı C. Total pressure: 1.0 atm, input V/III ratio: 50, F : 1.0 and ’: 0.03

36

A. Koukitu and Y. Kumagai

1

Temp.:1000°C SPi:1.0 atm PGaCl :5.0×10–3 atm V/III:50 a:0.03 NH3

H2

IG

Partial Pressure (atm)

10–2 GaCl HCl

10–4

10–6

10–8 GaCl3

10–10 0.0

0.2

0.4

0.6

0.8

1.0

F Fig. 2.3 Equilibrium partial pressures over GaN as a function of hydrogen partial pressure in the carrier gas, parameter F at 1;000 ı C. Total pressure: 1.0 atm, input partial pressure of GaCl: 5  103 atm, input V/III ratio: 50 and ’: 0.03. The partial pressure of IG involves N2 , which is produced by the decomposition of NH3

DPGa (atm)

5.0×10

–3

SPi:1.0 atm F:0.0 – 1.0

1.0×10–3

1.0×10–4 500

PGaCl :5.0×10–3 atm a:0.03

V/III:50

F:0.0 0.01 0.05 0.1 0.2 0.4 0.6 0.8 1.0

600

700

800

900 1000 1100 1200

Temperature (°C) Fig. 2.4 Driving force for the GaN deposition as a function of growth temperature with various parameters F . Total pressure: 1.0 atm, input partial pressure of GaCl: 5  103 atm, input V/III ratio: 50 and ’: 0.03

2

Hydride Vapor Phase Epitaxy of GaN

Growth Rate (mm/h)

100

37

experimental calculated

80 60 40

Temp.: 1000°C

20

F :1.0 a :0.03

PNH3 :0.26 atm

2.0×10–3

4.0×10–3

6.0×10–3

8.0×10–3

1.0×10–2

1.2×10–2

Input Partial Pressure of GaCl (atm) Fig. 2.5 Comparison between calculated growth rates and experimental data

well with Ban’s experiments [19]. In [19], it has been shown that the extent of GaCl consumption, i.e., the deposition of GaN, was significantly greater for He as the carrier gas than for H2 . The growth rate, r, under the mass transport type II or diffusion limit at constant pressure is expressed as (2.11) r D KgPGa ; where Kg is the mass transport coefficient [22]. Figure 2.5 shows a comparison between the calculated growth rate using (2.11) and the experimental data reported by Usui et al [6]. In the calculation of the growth rate, the value of Kg was determined to be 1:18  105 m=.h atm/. The agreement between the calculated and experimental growth rate is quite good, and shows that the growth of GaN using HVPE is thermodynamically controlled.

2.3 Cubic GaN Epitaxial Growth on (100) GaAs Substrate [5, 18] In the 1990s, epitaxial growth of cubic GaN was investigated on various substrates [23–28] such as GaAs, Si, SiC, and MgO. The growth of cubic GaN on GaAs (100) substrate had been reported using gas source MBE (GSMBE) and MOVPE [27,28]. Tsuchiya and Hasegawa et al. [29] demonstrated the homoepitaxial growth of cubic GaN by HVPE on cubic GaN/GaAs (100) substrates prepared by GSMBE. They obtained cubic GaN films with a growth rate as high as 1:6 m=h. At that time, it was generally supposed that a low-temperature buffer layer of GaN could not be grown on a foreign substrate using HVPE system since only nonequilibrium growth system realized a GaN buffer layer.

38 Fig. 2.6 Schematic diagram of the metalorganic hydrogen chloride vapor phase epitaxy (MOHVPE) system

A. Koukitu and Y. Kumagai TMG + H2 HCl + H2 NH3 + H2 Furnace Substrate

Exhaust

However, from the present knowledge, the cubic GaN layers grown, at the time of writing of those reports should be a mixture of some wurtzite type structure and zinc blende structure. In this section, the HVPE growth on GaAs (100) substrate is described to grow cubic GaN. It will be seen in the following chapters that the predeposited GaN buffer layer plays an essential role in the subsequent growth of high-quality cubic GaN films on GaAs (100) substrate by HVPE.

2.3.1 Experimental A schematic illustration of the reactor used in this work is shown in Fig. 2.6. The quartz reactor was vertically placed and the growth of buffer layers and epilayers was carried out at atmospheric pressure. The substrate was GaAs (100) which was etched in a solution of 3 W 1 W 1 D H2 SO4 W H2 O2 W H2 O for about 1 min. The MOHVPE system [5, 18] used is a simple-structured HVPE using TMG  HCl  NH3  H2 system as described above. TMG was coinjected with HCl into a hot-wall reactor to form gaseous GaCl and CH4 . GaN was then deposited on GaAs substrates through the reaction with NH3 . Prior to the growth of cubic GaN, a GaN buffer layer was grown at 500 ı C. Then the substrate temperature was raised to the growth temperature of 850 ı C in the flow of H2 and NH3 . Subsequently, TMG and HCl were supplied simultaneously to grow the GaN layer. Typical input partial pressures of TMG, HCl and NH3 over the substrate were 8  104 atm; 8  104 atm and 1:6  101 atm, respectively.

39

30

c-GaN (200)

GaAs (200)

Hydride Vapor Phase Epitaxy of GaN

Intensity (arb. units)

2

40

50

2q (deg) Fig. 2.7 XRD profile of the GaN film grown on (100) GaAs

2.3.2 Cubic GaN Growth The growth rate of the GaN films obtained under these conditions was about 3:0 m=h. Figure 2.7 shows a typical X-ray diffraction (XRD) profile for the obtained GaN films. The strong diffraction peak at about 40:0ı corresponding to (200) of cubic GaN can be seen for the grown layer. However, a quality of cubic GaN cannot be evaluated only by a typical XRD, ™  2™ method. In addition, the full-width at half-maximum (FWHM) of XRD rocking curve of the (200) cubic GaN was about 37.7 min. When the growth was carried out without the GaN buffer layer at high temperature, the GaAs substrate broke into pieces due to the reactions with HCl and NH3 . These findings indicate that the buffer layer prevents the reaction between the GaAs substrate and HCl or NH3 , which leads to high-quality GaN films. Figure 2.8 shows the FWHM values of the XRD rocking curve of the (200) cubic GaN with the epitaxial layer thickness of about 3 m as a function of the thickness of the GaN buffer layer on GaAs (100). The FWHM value initially decreases with increasing thickness of the GaN buffer layer up to 30 nm and then gradually increases. These results indicate that the optimum thickness of the GaN buffer is around 40 nm. GaN epitaxial layers with about 30 nm thick GaN buffer layer were observed by high resolution TEM and electron diffraction patterns to evaluate the quality of the GaN epitaxial layers. Three samples, A, B and C were evaluated. Sample A was taken out from the growth system just after the growth of a 30 nm thick GaN buffer layer, sample B was taken out after annealing at 850 ı C for 10 min and sample C was taken out after growth of the GaN epitaxial layer. Figure 2.9a shows a high-resolution TEM image observed from the [110] direction for sample A. Grain structures consisting of many small polycrystals with various orientations are observed in the GaN buffer layer (arrow 1 in Fig. 2.9a). The interface between the GaN buffer layer and the GaAs substrate is almost flat, but dark regions caused

40

A. Koukitu and Y. Kumagai PTMG :8.0×10–4 atm PHCI :8.0×10–4 atm

140

V/III:200

120

FWHM (min)

100 80 60 40 20 0

0 20 40 60 80 100 120 140 Thickness of GaN buffer layer (nm)

Fig. 2.8 The FWHM values of the X-ray rocking curve of cubic GaN (200) as a function of the GaN buffer layer thickness

a

GaN buffer layer GaAs substrate

b

5 nm

Fig. 2.9 (a) High-resolution TEM image and (b) electron diffraction pattern of sample A

2

Hydride Vapor Phase Epitaxy of GaN

41

by stress are observed in part of the GaAs near the GaN buffer layer (arrow 2 in Fig. 2.9a). The stress may be due to the difference in the lattice parameters between GaAs and GaN. Figure 2.9b shows an electron diffraction pattern taken at the interface with the incident beam along the [110] zone axis of GaAs in sample A. The pattern consists of strong diffraction spots (arrow 1 in Fig. 2.9b) for the GaAs substrate and a ring pattern (arrow 2 in Fig. 2.9b) for the GaN buffer layer. These findings show that the GaN buffer layer consists mainly of polycrystals after the growth at 500 ıC. Figure 2.10a shows a TEM image of sample B. A TEM image and an electron diffraction pattern of the interface for sample C are very similar to those of sample B (not shown here). The dark regions observed in the interface in Fig. 2.9a have disappeared due to stress relaxation by thermal annealing. No (111) facet can be observed near the interface region of the GaAs substrate, and the GaAs interface is fairly flat. The period of the dark lines (arrow 1 in Fig. 2.10a) is 4 in terms of the dot spacing in the GaAs image, and 5 in terms of that of the GaN image [30]. The dark lines (arrow 1 in Fig. 2.10a) indicate misfit dislocations due to the difference in the lattice parameters. Therefore, the lattices of GaN are almost strain-free and form single cubic crystals with some dislocations in the <111> direction (arrow 2 in Fig. 2.10a). Also, dark and bright regions are observed in the GaN buffer layer near the GaAs substrate (arrow 3 in Fig. 2.10a), implying that the GaN buffer layer near the GaAs substrate consists of a set of crystals having slightly different orientations. Figure 2.10b shows an electron diffraction pattern observed from the [110] direction for sample B. It is seen that the ring pattern observed in sample A changed to a spot pattern (arrow 1 in Fig. 2.10b). This finding indicates that the GaN buffer layer becomes a single crystal upon thermal annealing. Line patterns caused by stacking faults in the <111> direction are also observed (arrow 2 in Fig. 2.10b). These

a

GaN buffer layer GaAs substrate

b

Fig. 2.10 (a) High-resolution TEM image and (b) electron diffraction pattern of sample B

5 nm

42

A. Koukitu and Y. Kumagai

results show that thermal annealing after the growth of a GaN buffer layer plays an important role in changing the polycrystalline GaN to a single crystal.

2.4 Comparison of GaN Growth on (111)A and (111)B GaAs Substrates [7, 31] There have been several attempts to prepare a bulk GaN crystal to avoid defect generation. Porowski et al. reported the synthesis of GaN bulk crystal under high pressure and high temperature [32]. Kurai et al. employed the sublimation method [33]. However, GaN crystals obtained by these methods are too small for use as substrate at this stage. Another potential method for a GaN substrate preparation is to use a high growth rate of HVPE [6, 34, 35]. Large freestanding GaN substrates have been prepared by separating thick HVPE GaN layers from sapphire substrates [35]. However, it is very difficult to separate the GaN layer from the sapphire substrate because sapphire is very hard and stable. Growth of GaN on GaAs (111) substrate is very attractive, since the GaAs starting substrate can be easily removed by aqua regia after the GaN growth. In addition, the polarity of wurtzite GaN (0001) grown on GaAs (111) is expected to be controlled by the polarity of GaAs used, i.e., GaN (0001) (Ga-polarity) on GaAs (111)A N (N-polarity) on GaAs (111)B. However, there is the problem of and GaN .0001/ how to prevent deterioration of GaAs due to the unstable nature of GaAs at high temperature around 1;000 ıC (widely used temperature for the epitaxial growth of GaN). Hasegawa et al. introduced a 3 m thick GaN intermediate layer grown at 850 ıC on a 50 nm thick GaN buffer layer to prevent deterioration of a GaAs substrate at high temperature, and obtained a hexagonal GaN layer with a fairly smooth surface at 1;000 ıC on GaAs (111)B [4]. However, there is no report of the GaN growth on GaAs (111)A surfaces at temperatures as high as 1;000 ı C. In this chapter, comparison of the GaN buffer layers and the GaN epitaxial layers grown on GaAs (111)A and (111)B surfaces is described.

2.4.1 Experimental The MOHVPE system used in this study employs a vertical quartz reactor as shown in Fig. 2.6. A GaAs (111) substrate with both the sides polished was used and placed parallel to the gas stream so that GaN could be grown on both GaAs (111)A and (111)B surfaces simultaneously. The GaAs substrate was first cleaned in a solution of HCl W H2 O D 1 W 19 for 1 min and then set into the reactor. TMG and NH3 were used as Ga and N sources. The mixing zone of TMG and HCl was maintained at 750 ıC throughout the entire growth procedure to form gaseous GaCl by the reaction between TMG and HCl. GaN was deposited on the both surfaces of GaAs substrate in the downstream region, where GaCl and NH3 were mixed.

2

Hydride Vapor Phase Epitaxy of GaN

43

First, a thin GaN buffer layer of about 50 nm thick was grown on both sides at 550 ıC. The growth was performed under atmospheric pressure and the input partial pressures of TMG, HCl and NH3 were 8  104 atm; 8  104 atm and 1:6  101 atm, respectively. Then the substrate temperature was raised to 1;000 ıC in an NH3 ambient. Subsequently, TMG and HCl were supplied again and the growth of GaN was performed for 1 h. The input partial pressures of TMG, HCl and NH3 at this stage were 5:8  103 atm; 5:8  103 atm and 2:9  101 atm, respectively, where the growth rate of GaN was about 25 m=h.

2.4.2 Comparison of GaN Growth on (111)A and (111)B GaAs Surfaces Figure 2.11 shows the thickness of the low temperature GaN buffer layers grown on GaAs (111)A and (111)B surfaces as a function of growth time. Although the growth rate of buffer layers on GaAs (111)A and (111)B surfaces is almost the same after about 3 min of growth, the thickness on GaAs (111)A surface is about 1.3 times larger than that on GaAs (111)B surface due to the difference in growth rate in the early stages of the growth. The reason for this result may be related to the nucleation manner of the buffer layer. From this point of view, GaN initial growth processes on GaAs (111)A and B surfaces will be described in the next chapter. Figure 2.12 shows AFM images of as-grown buffer layers on GaAs (111)A and (111)B surfaces after 15 min of growth. The difference in surface morphology is clearly seen. The buffer layer grown on GaAs (111)A surface shows a smooth sur-

180 160

Thickness (nm)

140 120

on GaAs(111)A

100 80

on GaAs(111)B

60 40 20 0

0

5

10

15

20

25

30

Growth time (min) Fig. 2.11 Thickness of the GaN buffer layers grown by MOHVPE as a function of growth time: on GaAs (111)A surface (circles) and on GaAs (111)B surface (triangles)

44

A. Koukitu and Y. Kumagai

a

b

nm 25

nm 25

0 0.5

0 0.5

0.5

0.25

0.25

μm

0.5

0.25

0.25

μm

μm

0 0

μm

0 0

Fig. 2.12 AFM images of as-grown GaN buffer layers after 15 min of growth: (a) on GaAs (111)A and (b) on GaAs (111)B surface

a

b

5m

m

Fig. 2.13 Photographs of the GaAs substrate covered with a 50-nm-thick GaN buffer layer grown at 550 ı C and heated to 1;000 ı C in NH3 ambient. (a) on GaAs (111)A and (b) on GaAs (111)B. The word “GaN” is projected onto the surface

face and covers the GaAs surface homogeneously. Peak-to-valley height of this surface is within 5% of the grown layer. On the other hand, the GaN buffer layer on GaAs (111)B surface shows a very rough surface with many pits. Peak-to-valley height of this surface is about 30% of the grown layer. Also, there are some pinholes reaching the GaAs substrate on the surface. Next, we investigated the surface feature of 50 nm thick GaN buffer layers grown on GaAs (111)A and (111)B surfaces after heating in NH3 ambient up to 1;000 ı C. Photographs of the substrate after heating are shown in Fig. 2.13. In this figure, the word “GaN” has been projected onto the surface in order to reveal the surface features clearly. The difference in surface morphology between the two surfaces is very clear. The surface on the GaAs (111)A still appears to be mirror-like even after 1;000 ıC heating (Fig. 2.13a). In addition, there is no change in the GaAs substrate beneath the GaN buffer layer. On the contrary, numerous pinholes are seen in the GaN buffer layer grown on the GaAs (111)B surface (Fig. 2.13b). Furthermore, serious deterioration of the GaAs substrate beneath the GaN buffer layer is observed. This may be caused by arsenic (As) desorption from the GaAs substrate through the pinholes. The above results suggest two important points. One is the difference between the surface polarities of the GaN buffer layers grown on GaAs (111)A and (111)B surfaces, and the other is the possibility of high-quality GaN layer growth on GaAs (111)A surface. It was found that the uniform GaN layer with a mirror-like surface can be grown at 1;000 ıC on the 50 nm thick GaN buffer layer on GaAs (111)A surface.

2

Hydride Vapor Phase Epitaxy of GaN

45

Fig. 2.14 SEM image of the top and cleaved surface of a GaN layer grown at 1;000 ı C for 1 h on a 50-nm-thick GaN buffer layer grown on a GaAs (111)A

Figure 2.14 shows a scanning electron microscopy (SEM) image of the top and cleaved surface of GaN layer grown at 1;000 ı C for 1 h on the 50 nm thick GaN buffer layer grown on the GaAs (111)A surface. It is clearly seen that the surface of the grown layer is mirror-like, although there are a few hexagonal pits. As not shown here, the surface behind the substrate (GaAs (111)B surface) is seriously damaged due to the deterioration of the 50 nm thick GaN buffer layer at 1;000 ı C. Another interesting point we can see from Fig. 2.14 is that the porous shape of the GaAs substrate proceeds during GaN growth at 1;000 ıC from the surface behind the substrate, i.e., from the GaAs (111)B surface.

2.5 Ab Initio Calculations of GaN Initial Growth Processes on (111)A and (111)B GaAs Surfaces [36] As described in the preceding section, the GaAs (111)A substrate covered with a 50 nm thick GaN buffer layer did not deteriorate even after subsequent heating up to 1;000 ı C, whereas numerous holes were formed in the GaN buffer layer that was grown on GaAs (111)B surface. However, little is known why the GaN buffer layer on the GaAs (111)A does not deteriorate at 1;000 ı C. Therefore, the growth mechanisms, especially between GaN buffer layer and GaAs (111) surfaces, should be investigated to better understand GaN growth on GaAs. In this chapter, we show the orientation dependence of GaN buffer layer growth on GaAs by employing density-function theory (DFT). At the initial stage, we

46

A. Koukitu and Y. Kumagai

assume that As atoms on a GaAs surface react with hydrogen in the carrier gas, and then As atoms desorb from GaAs surface through AsHn molecules. In the next stage, Ga atoms on the surface react with NHn (NH3 , NH2 , or NH) to form GaN on the GaAs surface.

2.5.1 Calculation Procedure The total energy calculations and geometry optimization were performed employing DFT [37–39] within the generalized gradient approximation (GGA) [40–42] and the first-principles pseudopotential approach [43]. The energy cutoff value, which determines the number of plane-waves in the basis set, was 15 Ry for GaAs. The calculations were carried out using the package fhi98md [44]. The models used for the ab initio calculations represented the As-terminated GaAs (111)A and (111)B surfaces by supercell geometries with four GaAs atomic layers containing eight Ga atoms and eight As atoms per layer, and 20 Å vacuum region. The Ga atoms in the bottom layer are terminated by fictitious H atoms with 1.25 elementary charges in order to form perfect covalent bonds at the bottom surface of GaAs [45]. The initial structures of the models were obtained from ab initio calculations, where only the positions of fictitious hydrogen atoms terminated on the surface were varied. The positions of Ga and As atoms were fixed to keep the crystal structure of bulk GaAs. Then, in order to perform energy minimization, the position of Ga and As atoms in the top three layers were varied, and the positions of As precursors over the GaAs surface were varied along the a and b axes. The positions of Ga atoms in the fourth layer and the optimized fictitious H atoms were fixed in their bulk GaAs positions.

2.5.2 GaN Initial Growth Processes on (111)A and (111)B GaAs Surfaces The surface on the GaAs (111)A still appeared to be mirror-like even after 1;000 ıC heating as shown in Fig. 2.13a, whereas numerous pinholes were seen in the GaN buffer layer grown on the GaAs (111)B surface as shown in Fig. 2.13b. Therefore, we assumed that the As atoms, which remain on the GaAs surface, prevent N precursors such as NH3 ; NH2 , and NH from reacting with Ga atoms on the surface, and we calculated the As desorption energies from the GaAs (111)A and (111)B surfaces by ab initio calculations. First, the GaAs (111)A and (111)B surfaces before As desorption were optimized. In VPE, the uppermost atoms on the GaAs surface might be adsorbed by hydrogen atoms because hydrogen is generally used as a carrier gas. Therefore, we selected four models as the initial surfaces; Model 1: the bare surface of GaAs (111) As plane, and the GaAs (111) As surfaces terminated with monohydride (Model 2), dihydride (Model 3), and trihydride (Model 4). Next, we simulated the As

2

Hydride Vapor Phase Epitaxy of GaN

47

Potential energy change(eV)

–2147

Non-hydride (-As)

–2148 –2149

461 kJ/mol

Monohydride (-AsH)

As

–2150

400 kJ/mol

Dihydride (-AsH2)

(111)A –2151

164 kJ/mol

Trihydride (-AsH3)

–2152

71.9 kJ/mol

–2153 –2154

0

1

2

3

4

5

6

Distance from surface (Å) Fig. 2.15 Potential energy changes for the desorption processes of AsHn molecules (n D 0, 1, 2, 3) from nonhydride .As/, monohydride .AsH/, dihydride .AsH2 / and trihydride .AsH3 / surfaces of GaAs (111)A

desorption process from these initial surfaces, where it was thought that As atoms on the four surfaces were removed as As, and AsHn with n D 1–3. In the calculations, As or AsHn were moved from the initial surface perpendicularly to the surface in steps of about 1.0 Å, and then the total energies and the optimized structures were obtained for each height of the molecule from the surface. Figure 2.15 shows the potential energy change in the As desorption process on the GaAs (111)A surface with respect to the distance between the GaAs (111)A surface and the desorbed AsHn (n D 0, 1, 2, 3) molecules. The total energy of the initial monohydride surface .AsH/ is the smallest, and then comes trihydride .AsH3 /, nonhydride .As/, and dihydride .AsH2 / in that order. However, the differences of the potential energies in the initial surface are small. As shown in the figure, the activation energies on the GaAs (111)A surface are 71.9 kJ/mol for desorption of AsH3 molecules from the AsH3 surface, 164 kJ/mol for that of AsH2 from the AsH2 surface, 400 kJ/mol for that of AsH from the AsH surface, and 461 kJ/mol for As from the As surface. These results directly suggest that As atoms on GaAs (111)A desorb from the GaAs surface in AsH3 molecules, and it is likely that As atoms completely desorb from the surface at about 550 ı C before the growth of GaN buffer layer because the activation energy of the reaction is extremely small. The potential energy change in the As desorption process from GaAs (111)B is shown in Fig. 2.16. The initial surfaces are AsH2 ; AsH; As and AsH3 surfaces in increasing order of the value of total energy. The activation energy for desorption of AsH3 from the GaAs (111)B surface is shown with a negative value. However, the AsH3 desorption directly from the AsH3 surface does not occur, because the formation of the AsH3 surface is impossible due to the large potential

48

A. Koukitu and Y. Kumagai –1954 Non-hydride (-As)

Potential energy change(eV)

–1955 –1956 As

–1957 –1958

397 kJ/mol

Monohydride (-AsH)

364 kJ/mol

Dihydride (-AsH2)

(111)B

235 kJ/mol

Trihydride (-AsH3)

–1959 –1960 –1961

0

1

2

3

4

5

6

Distance from surface (Å) Fig. 2.16 Potential energy changes for the desorption processes of AsHn molecules (n D 0, 1, 2, 3) from nonhydride .As/, monohydride .AsH/, dihydride .AsH2 / and trihydride .AsH3 / surfaces of GaAs (111)B

energy of the AsH3 surface as shown in the figure. From these calculations, we can speculate on the As desorption process from the GaAs (111)B surface by considering the stability of AsH3 molecules as follows. First, hydrogen molecules from the carrier gas attach to the surface, and then each of them react with As atoms on the surface to form an AsH2 surface. Next, the AsH2 molecules desorb from the surface with activation energy of 235 kJ/mol, and then the AsH2 molecules react with hydrogen in the vapor phase to form AsH3 . Figure 2.17 shows the arithmetical mean deviation of GaN buffer surface as a function of the activation energy for As desorption process from the surface. The activation energy from the GaAs (001) surface shown in this figure (144 kJ/mol) was obtained by the in situ gravimetric monitoring method [46, 47]. From the figure, the small activation energy on the GaAs (111)A suggests that all As atoms can desorb as AsH3 from the surface thus leaving a completely Ga-terminated surface for the subsequent GaN buffer growth. On the other hand, the large activation energy on GaAs (111)B prevents the complete desorption of all As atoms before the GaN buffer growth. Accordingly, the As atoms remaining on the surface probably prevent N precursors from reacting with the Ga atoms on the surface during GaN buffer growth. The above results can explain the experimental finding that the GaN buffer layers on GaAs (111)B surface had many holes. Therefore, the complete removal of As atoms from the surface is required for the growth of high quality GaN buffer layers.

Hydride Vapor Phase Epitaxy of GaN

Arithmetical mean deviation (nm)

2

49

calculated experimental

As

2.0 As

(111)B

(001)

As 1.0

(111)A

50

100

150

200

250

Activation energy (kJ/mol) Fig. 2.17 Arithmetical mean deviation of the surface as a function of the activation energy for As desorption from GaAs (111)A, (001) and (111)B surfaces

2.6 Thick GaN Growth on (111)A GaAs Substrate [48, 49] As described in the above section, it was found that the uniform GaN layer with a mirror-like surface could be grown at 1;000 ıC on (111)A GaAs surface. Growth of a thick GaN layer on a GaAs substrate is worth investigating because GaAs is easily removed by lapping or etching. In this chapter, the growth temperature dependence of the surface morphology, optical property and crystalline quality of the GaN epitaxial layers are described. In addition, it is shown that the growth of thick GaN layers on (111)A GaAs surface at 1;000 ı C or above is promising for the preparation of a freestanding GaN substrate.

2.6.1 Experimental The MOHVPE system used in this study is the same as that described above. In this system, a vertical quartz reactor is employed and the growth of GaN is carried out under atmospheric pressure using H2 carrier gas. TMG and NH3 were used as Ga and N sources, respectively. The mixing zone of TMG and HCl was kept at 750 ıC throughout the entire growth procedure to form gaseous GaCl by the reaction between TMG and HCl. The substrate used was (111)A GaAs. First, a thin GaN buffer layer of about 50 nm in thickness was grown on the GaAs (111)A surface at 550 ı C. The input partial pressures of TMG, HCl and NH3 were 8  104 atm; 8  104 atm and 1:6  101 atm, respectively. Then the substrate temperature was raised to the growth temperature in an NH3 ambient. Subsequently,

50

A. Koukitu and Y. Kumagai

TMG and HCl were supplied again and the growth of GaN was performed at various temperatures ranging from 920 ıC to 1;000 ıC. The input partial pressures of TMG, HCl and NH3 at this stage were 5:8  103 atm; 5:8  103 atm and 2:9  101 atm, respectively.

2.6.2 Thick GaN Growth on (111)A GaAs Surface First, we investigated the growth rate of GaN on the GaAs (111)A surface in the temperature range from 920 ıC to 1;000 ı C. It was found that the growth rate of GaN was almost constant at around 25 m=h in this temperature range and increased linearly with/while increasing the input partial pressure of TMG and HCl. Therefore, the growth process is limited by mass transportation in the temperature range from 920 ıC to 1;000 ıC. On the other hand, the surface morphology of the grown layer strongly depends on the growth temperature. Figure 2.18 shows surface morphologies of 25 m thick GaN layers grown at various temperatures ranging from 920 ıC to 1;000 ı C. Improvement of the surface morphology with the increasing growth temperature is

Fig. 2.18 Surface morphologies of the 25-m-thick GaN layers grown on the GaAs (111)A surfaces at various temperatures observed by SEM: (a) 920 ı C, (b) 940 ı C, (c) 960 ı C, (d) 980 ı C and (e) 1;000 ı C

2

Hydride Vapor Phase Epitaxy of GaN

51

quite clear. The GaN layer grown at 920 ı C shows a rugged surface, and no-mirrorlike regions can be seen in the surface. However, the area of the mirror-like regions significantly increases with the increase of the growth temperature. As shown in the figure, the GaN layer grown at 1;000 ıC shows a mirror-like surface without hexagonal pits. These results indicate that the growth temperature of 1;000 ı C or above is essential to obtain a GaN layer with a mirror-like surface. The photoluminescence (PL) spectra measured at room temperature using a He  Cd laser excitation are shown in Fig. 2.19. One important feature in this figure is the increase of the near band edge emission intensity at 363.1 nm with the increase of growth temperature. Although the band edge emission cannot be seen in the layer grown at 920 ıC, an intense emission with a narrower FWHM of 96 meV can be seen in the layer grown at 1;000 ıC. Another important feature is the decrease of the broad emission peaking at 568.3 nm with the increase of growth temperature. The broad emission, or so-called yellow emission, is related to the deep-level traps due to lattice imperfections. PL analysis also suggests that the growth temperature at 1;000 ıC or above is essential for the growth of high quality GaN layers on GaAs (111)A surface. The crystalline structure and/or quality of the GaN layers grown on GaAs (111)A surfaces were examined by means of XRD and double crystal XRD (DCXRD). First, the 25-m-thick GaN layers grown at various temperatures shown in Fig. 2.18

568.3

363.1

920 °C

×0.5

940 °C

PL intensity (arb. units)

×0.5

×0.5

960 °C

×0.5

980 °C

×0.5

400

500 600 Wavelength (nm)

1000 °C 700

800

Fig. 2.19 Room temperature PL spectra of the GaN layers grown on GaAs (111)A surface at various temperatures shown in Fig. 2.18

52

A. Koukitu and Y. Kumagai 60

FWHM (min)

50

25-μm-thick MOHVPE GaN

40 30 20 10 0 900 920 940 960 980 1000 1020

Temperature(°C) Fig. 2.20 FWHM values of the ¨-mode DCXRD profiles of the (0002) GaN as a function of growth temperature. Note that the thickness of the GaN layer is 25 m at each temperature

were analyzed by XRD in ™–2™ mode. In the obtained XRD profiles, only the hexagonal GaN (0002) and (0004) peaks were seen, irrespective of the growth temperature, which indicates that the hexagonal GaN grows epitaxially on the GaAs (111)A surface to satisfy the relation of GaN (0001) jj GaAs(111) in the temperature range from 920 ı C to 1;000 ıC. Therefore, we examined the FWHM value of the ¨–scan for the (0002) GaN plane by DCXRD. The dependence of the FWHM value on the growth temperature is shown in Fig. 2.20. It is seen in this figure that the FWHM value decreases with increasing growth temperature up to 960 ı C, which is consistent with the results obtained in Figs. 2.18 and 2.19 which show, that both surface morphology and optical qualities are improved with an increase of the growth temperature. On the other hand, at growth temperatures above 960 ı C, the FWHM value increases with increasing growth temperature. This opposite tendency is thought to be due to the bending of the thin grown layer .25 m/ with improvement of crystallinity. Fortunately, however, an answer to the problem was also found. In Fig. 2.21, the FWHM values in the ¨ mode scan for the GaN (0002) plane by DCXRD are shown as a function of grown layer thickness. The growth was performed at 1;000 ı C in this case. While the FWHM value is large when the grown thickness is around 13 m (41.8 min), the FWHM value decreases with the increasing grown layer thickness. By growing a 100 m thick GaN layer with a mirror-like surface, the FWHM value of 4.6 min can be obtained. From these results, we believe that the growth of a thick GaN layer on a (111)A GaAs surface at 1;000 ı C or above is a promising method for the preparation of freestanding GaN substrates. In actual fact, the HVPE growth of a freestanding GaN substrate over 2 in. in size was reported according to the above mentioned findings [48]. In that case, a 0.1m-thick SiO2 layer with 2-m-diameter round openings is first formed directly on a (111)A GaAs substrate, and a 60-nm-thick GaN buffer layer was selectively grown in the openings at 500 ıC. Then, a thick GaN layer was grown at 1;030 ı C. After growing an about 500-m-thick GaN layer, the (111)A GaAs substrate was

2

Hydride Vapor Phase Epitaxy of GaN

53

60 MOHVPE GaN Temp.: 1000°C

FWHM (min)

50 40 30 20 10 0

0

20

40

60

80

100

Thickness (mm) Fig. 2.21 FWHM values of the ¨-mode DCXRD profiles of the GaN (0002) as a function of grown layer thickness. The GaN layer was grown at 1;000 ı C

dissolved by aqua regia. Surface SEM micrographs in the GaN growth process at 1;030 ıC is shown in Fig. 2.22. It is clearly seen that the growth of GaN begins selectively in the openings in the SiO2 mask (Fig. 2.22a), and hexagonal pyramid of GaN N surrounded by f1101g appear (Fig. 2.22b). As the growth continues, overgrowth of GaN on the SiO2 mask occurs, and coalescence of the GaN begins (Fig. 2.22c). After removing the (111)A GaAs substrate, a freestanding GaN substrate of about 500 m thick was prepared by lapping and polishing. The SiO2 layer on backside was also removed during this process. Figure 2.23 shows a photograph of the freestanding GaN substrate obtained. DCXRD profile of the (0002) GaN showed a narrow FWHM of 106 arcsec, and plan-view TEM observation at the top surface revealed the dislocation density of the GaN substrate was as low as 2  105 cm2 . Hall measurement showed that the GaN substrate had n-type conductivity with a carrier concentration and electron mobility of 5  1018 cm3 and 170 cm2 V1 s1 , respectively, indicating that the obtained freestanding GaN substrate is suitable for use in fabrication of violet laser diodes on it.

2.7 Preparation of Fe-Doped Semi-insulating GaN Substrates [50, 51] Recent progress in the preparation of 2-in. -diameter freestanding GaN substrates has enabled mass-production of high-power and long-lifetime violet laser diodes. However, nominally undoped GaN layers grown by any technique, including HVPE, generally show n-type conductivity due to unintentionally incorporated background donor impurities such as Si and O [52]. A next target of production of freestanding GaN substrates is preparation of high-quality semi-insulating (SI) GaN substrates

54

A. Koukitu and Y. Kumagai

Fig. 2.22 SEM micrographs in the process of GaN growth at 1;030 ı C on a (111)A GaAs substrate with a SiO2 mask: (a) 0.5 min, (b) 4 min and (c) 10 min of GaN growth

2

Hydride Vapor Phase Epitaxy of GaN

55

Fig. 2.23 Photograph of a freestanding GaN substrate prepared by HVPE using a (111)A GaAs as an initial substrate

for high-frequency applications. One approach for the growth of SI GaN is to introduce deep acceptor states to compensate for the background donors. It has been reported that the growth of SI GaAs and InP is possible by Fe doping [53, 54]. Fe has been also known to act as a deep acceptor in GaN [55]. Therefore, the method for efficient generation of Fe source and its transportation is essential for growing SI GaN substrate by HVPE. In this chapter, the growth of Fe-doped GaN layers on both (0001) sapphire and (111)A GaAs substrates is shown. Then, preparation of a SI GaN substrate using (111)A GaAs substrate and properties of the SI GaN substrate is described.

2.7.1 Experimental A MOHVPE system with a vertical hot-wall quartz reactor was used. Ga source was GaCl formed by the reaction between TMG and HCl, and Fe source was FeCl2 formed by the reaction between ferrocene ŒCp2 Fe W .C5 H5 /2 Fe and HCl. Both GaCl and FeCl2 were formed independently in the upstream region of the reactor maintained at 750 ıC. The substrates used were (0001) sapphire and (111)A GaAs. First, a 50-nm-thick Fe-doped GaN buffer layer was grown with input partial pressures of TMG, HCl to form GaCl, and NH3 of 8:0  104 atm; 8:0  104 atm and 1:6  101 atm, respectively. The growth was carried out at 500 ı C and 550 ıC for the sapphire and GaAs substrates, respectively. Then, Fe-doped thick GaN layers were grown at 1;000 ıC for 1 h with input partial pressures of TMG, HCl to form GaCl, and NH3 of 3:0  103 atm; 3:0  103 atm and 1:5  101 atm, respectively. During the growth, CP2 Fe and HCl to form FeCl2 were additionally introduced into the reactor. Input partial pressure of the HCl to form FeCl2 was twice as that of Cp2 Fe, and FeCl2 supply was varied by using the input partial pressure ratio of Cp2 Fe to TMG as an index.

56

A. Koukitu and Y. Kumagai

2.7.2 Fe-Doped GaN Layer Grown on Sapphire and GaAs First, GaN layers were grown on sapphire and GaAs substrates with input partial pressure ratios, Cp2 Fe=TMG, of 0 (undoped) and 0.34 (Fe-doped). The growth rate of GaN was 11–12 m=h, regardless of the substrate type and the supply of FeCl2 . After the growth, resistivity was determined by measuring the direct current through the GaN layer at a fixed bias. It was found that the undoped GaN layers grown on sapphire and GaAs substrates show low resistivities about 2:5  103  cm. On the other hand, the resistivities of Fe-doped GaN layers were 2:5  104  cm on sapphire substrate and 3:3  102  cm on GaAs substrate. These results indicate that resistivity of the Fe-doped GaN layer grown on GaAs substrate remained low. To clarify the mechanism of the Fe doping, secondary ion mass spectrometry (SIMS) analysis of GaN layers was performed. It was found that undoped GaN layers grown on sapphire and GaAs substrates contain large quantities .>1015 cm3 / of O (donor), C (acceptor) and Si (donor) impurities, among which O had the highest concentration of about 5  1019 cm3 . This is the reason why undoped GaN layers show n-type conductivity. Figure 2.24 shows the SIMS profiles of GaN layers grown at a CP2 Fe=TMG ratio of 0.34 on sapphire (left) and GaAs (right) substrates. It was found that the Fe-doped GaN layer on a sapphire substrate has impurity concentrations comparable to those of the undoped layer, while a high concentration of Fe in excess of O is incorporated. Thus, it is clear that a high doping level of Fe above 1019 cm3 is possible and that the increase in resistivity by Fe doping on sapphire substrate is attributed to the compensation of O donors by Fe. On the other hand, the Fe-doped GaN layer grown on GaAs substrate also contains O, C and Si impurities with concentrations at roughly the same level as those on (111)A GaAs

on (0001) sapphire GaN

Concentration (cm–3)

1021 1020

GaN

O

Fe

O

1019

C

C 1018

Fe As

1017

Si

Si

1016 0

4

8

Depth (mm)

12

16

0

4

8

12

Depth (mm)

Fig. 2.24 SIMS profiles of Fe-doped GaN layers grown on (0001) sapphire (a) and (111)A GaAs substrate (b). Growth was performed at 1;000 ı C with an input ratio of FeCl2 /TMG of 0.34

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Hydride Vapor Phase Epitaxy of GaN

57

in Fe-doped GaN layer grown on sapphire substrate. However, it is found that the Fe concentration drops by two orders of magnitude with respect to that on sapphire. In addition, a high concentration of As (neutral) impurity is incorporated on the order of 1017 cm3 . Thus, the low-resistivity of the Fe-doped GaN layer grown on GaAs is due to incomplete compensation of O donors by the decrease of Fe incorporation, even though the same growth conditions were used as on the sapphire substrate. The trade-off between the incorporation of Fe and As is surprising since Fe atoms occupy Ga sites [56] while As atoms is thought to occupy N sites in GaN lattice. As shown later, it was found that the use of a GaAs substrate coated on its back by NiTi layer could increase the concentration of Fe. Thus, it is speculated that the increase in As vapor-species in the growth reactor due to the degradation of the back surface of the GaAs substrate hinders FeCl2 from reacting with NH3 . In order to grow a SI GaN substrate using GaAs substrate, protection of the back surface of the GaAs substrate is essential. Utilizing the above findings, preparation of a SI freestanding GaN substrate was successfully achieved using (111)A GaAs substrate. In this case, GaCl was formed by the reaction between Ga metal and HCl gas for a long period of growth. Before placing a (111)A GaAs substrate in the reactor, the back of the GaAs substrate was covered by a 0:1 m-thick NiTi protective layer. First, a 60-nm-thick GaN buffer layer was grown at 550 ıC. Then, a 400-m-thick Fe-doped GaN layer was grown. Input partial pressures of GaCl, FeCl2 and NH3 at this stage were

Fig. 2.25 Photograph of a 400-m-thick Fe-doped SI GaN substrate after removal of GaAs substrate

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Fe concentration (cm–3)

1021 1020 Fe-doped GaN 1019 1018 1017 undoped GaN 1016 1015 0

2

4

6

8

Depth from surface (μm) Fig. 2.26 SIMS depth profiles of Fe in undoped (dashed line) and Fe-doped (solid line) freestanding GaN substrates

GaN (1010)

GaN (0002) undoped GaN

undoped GaN FWHM = 353 arcsec

Fe-doped GaN

Fe-doped GaN

FWHM = 410 arcsec

FWHM = 360 arcsec

Intensity (normalized)

FWHM = 443 arcsec

–1000

0

ω (arcsec)

1000

–1000

0

1000

φ (arcsec)

Fig. 2.27 XRD rocking curves of (0002) and .101N 0/ GaN measured by ¨- and ¥-mode scan, respectively: undoped (dashed lines) and Fe-doped (solid lines) freestanding GaN substrates

8:0  103 atm; 8:0  104 atm and 2:4  101 atm, respectively, which yielded a GaN growth of about 75 m=h. A photograph of the Fe-doped freestanding GaN substrate taken after removal of the GaAs substrate is shown in Fig. 2.25. The shape of the GaN substrate is the

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59

same as that of the GaAs substrate, and no polishing process was employed. It is seen that a transparent GaN substrate with a smooth surface is obtained although the GaN substrate is slight grayish. Figure 2.26 shows SIMS depth profiles of Fe for the undoped and Fe-doped GaN substrates. It is found that the Fe-doped GaN substrate has a Fe concentration of 1:5  1019 cm3 , while the undoped GaN substrate has a Fe concentration almost equal to the detection limit .5  1015 cm3 /. Thus, a high Fe doping level above 1019 cm3 is achieved even on a GaAs substrate. The resistivity of the Fe-doped freestanding GaN substrate measured at room temperature by three-electrode guard method was 8:8  1012  cm. Thus, a SI freestanding GaN substrate can be successfully prepared. Crystalline quality of the Fe-doped SI freestanding GaN substrate was examined by XRD. Figure 2.27 shows the ¨- and ¥-mode XRD rocking curves of the (0002) N of the undoped (dashed lines) and Fe-doped (solid lines) freestanding and .1010/ N of the FeGaN substrates. It is found that the FWHMs of the (0002) and .1010/ doped SI freestanding GaN substrate are fairly small and are 410 and 360 arcsec, respectively. It is also found that the FWHMs do not vary by the Fe doping, indicating that the crystalline quality of the Fe-doped GaN substrate is comparable to that of the undoped GaN substrate.

References 1. H.P. Maruska, J.J. Tietjen, Appl. Phys. Lett. 15, 327 (1969) 2. T. Detchprohm, K. Hiramatsu, H. Amano, I. Akasaki, Appl. Phys. Lett. 61, 2688 (1992) 3. A. Wakahara, T. Yamamoto, K. Ishio, A. Yoshida, Y. Seki, K. Kainosho, O. Oda, Jpn. J. Appl. Phys. 39, 2399 (2000) 4. F. Hasegawa, M. Minami, K. Sunaba, T. Suemasu, Jpn. J. Appl. Phys. 38, L700 (1999) 5. Y. Miura, N. Takahashi, A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 34, L401 (1995) 6. A. Usui, H. Sunakawa, A. Sakai, A.A. Yamaguchi, Jpn. J. Appl. Phys. 36, L899 (1997) 7. Y. Kumagai, A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 39, L149 (2000) 8. H. Seki, A. Koukitu, J. Cryst. Growth 98, 118 (1989) 9. A. Koukitu, N. Takahashi, T. Taki, H. Seki, Jpn. J. Appl. Phys. 35, L673 (1996) 10. A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 36, L750 (1997) 11. A. Koukitu, N. Takahashi, T. Taki, H. Seki, J. Cryst. Growth 170, 306 (1997) 12. A. Koukitu, N. Takahashi, H. Seki, Jpn. J. Appl. Phys. 36, L1136 (1997) 13. A. Koukitu, S. Hama, T. Taki, H. Seki, Jpn. J. Appl. Phys. 37, 762 (1998) 14. A. Koukitu, Y. Kumagai, N. Kubota, H. Seki, Phys. Stat. Sol. B 216, 707 (1999) 15. A. Koukitu, Y. Kumagai, H. Seki, Phys. Stat. Sol. A 180, 115 (2000) 16. A. Koukitu, Y. Kumagai, H. Seki, J. Cryst. Growth 221, 743 (2000) 17. Y. Kumagai, K. Takemoto, T. Hasegawa, A. Koukitu, H. Seki, J. Cryst. Growth 231, 57 (2001) 18. Y. Miura, N. Takahashi, A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 35, 546 (1996) 19. V.S. Ban, J. Electrochem. Soc. 119, 761 (1972) 20. H. Seki, S. Minagawa, Jpn. J. Appl. Phys. 11, 850 (1972) 21. A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 16, 1967 (1977) 22. D.W. Shaw, in Crystal Growth, vol. 1, ed. by C.H.L. Goodman (Plenum, New York, 1978), p.1 23. M. Mizuta, S. Fujieda, Y. Matsumoto, T. Kawamura, Jpn. J. Appl. Phys. 25, L945 (1986) 24. T. Lei, M. Fanciulli, R.J. Molnar, T.D. Moustakas, R.J. Graham, J. Scanlon, Appl. Phys. Lett. 59, 944 (1991)

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25. M.J. Paisley, Z. Sitar, J.B. Posthill, R.F. Davis, J. Vac. Sci. Technol. A 7, 701 (1989) 26. R.C. Powell, G.A. Tomasoch, Y.W. Kim, J.A. Thornton, J.E. Greene, Mater. Res. Soc. Symp. Proc. 162, 525 (1990) 27. H. Okumura, S. Yoshida, S. Misawa, E. Sakuma, J. Cryst. Growth 120, 114 (1992) 28. S. Miyoshi, K. Onabe, N. Ohkouchi, H. Yaguchi, R. Ito, S. Fukatsu, Y. Shiraki, J. Cryst. Growth 124, 439 (1992) 29. H. Tsuchiya, T. Okahisa, F. Hasegawa, H. Okumura, S. Yoshida, Jpn. J. Appl. Phys. 33, 1747 (1994) 30. N. Kuwano, Y. Nagatomo, K. Kobayashi, K. Oki, S. Miyoshi, H. Yaguchi, K. Onabe, Y. Shiraki, Jpn. J. Appl. Phys. 33, 18 (1994) 31. Y. Kumagai, H. Murakami, H. Seki, A. Koukitu, Phys. Stat. Sol. A 188, 549 (2001) 32. S. Porowski, J. M. Baranowski, M. Leszczynski, J. Jun, M. Bockowski, I. Grzegory, S. Krukowski, M. Wroblewski, B. Lucznik, G. Nowak, K. Pakula, A. Wysmolek, K.P. Korona, R. Stepniewski, Proc. Int. Symp. Blue Laser and Light Emitting Diodes, Chiba, 1996 (Ohmsha, Tokyo, 1996) p.38 33. S. Kurai, Y. Naoi, T. Abe. S. Ohmi, S. Sakai, Jpn. J. Appl. Phys. 35, L77 (1996) 34. S.T. Kim, Y.J. Lee, D.C. Moon, C.H. Hong, T.K. Yoo, J. Cryst. Growth 194, 37 (1998) 35. M.K. Kelly, R.P. Vaudo, V.M. Phanse, L. Görgens, O. Ambacher, M. Stutzmann, Jpn. J. Appl. Phys. 38, L217 (1999) 36. Y. Matsuo, Y. Kumagai, T. Irisawa, A. Koukitu, Phys. Stat. Sol. A 188, 553 (2001) 37. P. Hohenberg, W. Kohn, Phys. Rev. B 136, 864 (1964) 38. R.M. Dreuzler, E.K.U. Gross, Density Functional Theory (Springer, Berlin 1990) 39. W. Kohn, A.D. Bedcke, R.G. Parr, J. Phys. 100, 12974 (1996) 40. A.D. Becke, Phys. Rev. A 38, 3098 (1988) 41. J.P. Perdew, Phys. Rev. B 33, 8822 (1986) 42. J.P. Perdew, J.A. Chevary, S.H. Vosko, K.A. Jackson, M.R. Pederson, D.J. Singh, C. Fiolhais, Phys. Rev. B 46, 6671 (1992) 43. D.R. Hamann, Phys. Rev. B 40, 2980 (1982) 44. M. Bockstedte, A. Kley, J. Neugebauer, M. Scheffler, Comput. Phys. Commun. 107, 187 (1997) 45. K. Shiraishi, J. Phys. Soc. Jpn. 59, 3455 (1990) 46. A. Koukitu, N. Takahashi, Y. Miura, H. Seki, J. Cryst. Growth 146, 239 (1995) 47. A. Koukitu, T. Taki, K. Narita, H. Seki, J. Cryst. Growth 198/199, 1111 (1999) 48. Y. Kumagai, H. Murakami, A. Koukitu, K. Takemoto, H. Seki, Jpn. J. Appl. Phys. 39, L703 (2000) 49. K. Motoki, T. Okahisa, N. Matsumoto, M. Matsushima, H. Kimura, H. Kasai, K. Takemoto, K. Uematsu, T. Hirano, M. Nakayama, S. Nakahata, M. Ueno, D. Hara, Y. Kumagai, A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 40, L140 (2001) 50. Y. Kumagai, K. Takemoto, H. Murakami, A. Koukitu, Jpn. J. Appl. Phys. 44, L1072 (2005) 51. Y. Kumagai, F. Satoh, R. Togashi, H. Murakami, K. Takemoto, J. Iihara, K. Yamaguchi, A. Koukitu, J. Cryst. Growth 296, 11 (2006) 52. W.J. Moore, J.A. Freitas, Jr., G.C.B. Braga, R.J. Molnar, S.K. Lee, K.Y. Lee, I.J. Song, Appl. Phys. Lett. 79, 2570 (2001) 53. P.L. Hoyt, R.W. Haisty, J. Electrochem. Soc. 113, 296 (1966) 54. K. Tanaka, K. Nakai, O. Aoki, M. Sugawara, K. Wakao, S. Yamakoshi, J. Appl. Phys. 61, 4698 (1987) 55. B. Monemar, O. Lagerstedt, J. Appl. Phys. 50, 6480 (1979) 56. R. Togashi, F. Satoh, H. Murakami, J. Iihara, K. Yamaguchi, Y. Kumagai, A. Koukitu, Phys. Stat. Sol. B 244, 1862 (2007)

Chapter 3

Growth of Bulk GaN Crystals by HVPE on Single Crystalline GaN Seeds B. Łucznik, B. Pastuszka, G. Kamler, I. Grzegory, and S. Porowski

Abstract In this chapter, the problems related to bulk crystallization of GaN by HVPE are discussed. High quality small single crystalline GaN seeds grown under high pressure and large, flat, free standing GaN substrates grown by HVPE were used in long duration HVPE experiments. The characterization results of the obtained bulk crystals suggest that strong dependence of physical properties (mostly oxygen content) on the orientation of the crystallization front is the main reason for strain and structural defects generation in the new grown material.

3.1 Introduction The hydride vapor phase epitaxy (HVPE) method is, at present, the only one allowing mass production of free standing GaN substrates with quality sufficient for fabrication of laser diodes. However, the technology of making the GaN substrates by HVPE is again very specific (the GaN-based technologies are usually “unusual”) because it is heteroepitaxy (on GaAs [1] or sapphire [i.e., 2]) followed by subsequent substrate lift-off giving just one GaN substrate in one HVPE run (if the single wafer HVPE reactor is used). This is in contrast with the typical approach in semiconductor technology where the substrates (for example of Si or GaAs) are obtained by slicing and polishing of bulk high quality single crystals typically grown from stoichiometric liquids by Czochralski or Bridgman methods. For GaN, it is difficult because of its extreme conditions of melting (TM D 2;220 ıC and pM D 6:0 GPa) [3]. It is possible to grow GaN by HVPE with relatively high rates of the order of 100 m=h in various crystallographic directions, which suggests that bulk crystallization by this method should be attainable in principle. There are a few reports [i.e., 4, 5] on the long duration .>10 h/ growth of GaN by HVPE on a large area, flat sapphire or GaN seeds (substrates) giving 2 in. diameter crystals with thickness approaching 10 mm. However, the progress (in terms of crystal size and quality) in bulk crystallization of GaN by HVPE is not very fast, which is related mostly with technical limitations of the method like parasitic deposition of GaN out of the seed, 61

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deposition of solid reaction products (mostly NH4 Cl) or exhausting of the liquid gallium source. It seems that the sensitivity of physical properties of GaN on the orientation of the growth surface can also cause serious problems in the crystallization of large crystals by HVPE. It is well established [i.e., 6] that for GaN growth with intentional variations of growth direction (i.e., in Epitaxial Lateral Overgrowth – ELOG technique) by MOVPE or HVPE methods, significant differences in optical and electrical properties are observed in different sectors of the crystals. It can be particularly important for growth starting from small seed crystals because the extension of the crystal size proceeds in three dimensions and therefore in different crystallographic directions. In this chapter, the HVPE growth on both small and large (flat) seed crystals is discussed. The discussion is based mostly on the results obtained in the Institute of High Pressure Physics PAS in Warsaw, Poland.

3.2 Experimental 3.2.1 Seed Crystals The crystals used in this study for seeding of bulk crystallization of GaN by HVPE were grown from solutions in liquid gallium at nitrogen gas pressure of 1.0 GPa [7]. This method results in crystals of two dominating morphological forms (Fig. 3.1): thin hexagonal platelets and hexagonal needles elongated into c-direction of wurzite structure. The needles grown usually at high supersaturations are morphologically unstable N faces being partially hollow or skeletal but they often have well developed f1010g which makes these crystals good seeds for growth by HVPE. Crystals of both forms are limited in size to about 1 cm due to still insufficient control of supersaturation in gallium solutions in large volume high pressure reactors. The important advantage of the High Pressure Solution method is that the resulting crystals of GaN are almost free of structural defects. Defect selective etching (DSE) in molten bases [8, 9] of both (0001) polar and nonpolar faces of the crystals

Fig. 3.1 Two morphological forms of GaN crystals grown under high N2 pressure: (a) hexagonal platelets, (b) hexagonal needles. Grid: 1 mm

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Fig. 3.2 Free standing 1–2 mm thick GaN crystals separated spontaneously from GaN/sapphire substrates after 10–12 h HVPE growth: (a) susceptor with large hexagonal crystal just after the HVPE process, (b) the same crystal, (c) and (d) – similar crystals. Grid: 1 mm

of both forms has shown that dislocation densities are negligible and usually do not exceed 100 per cm2 . On the other hand, the crystals contain a high number of point defects due to the high activity of oxygen always present as nonintentional impurity in the high pressure growth system. Due to this, the GaN crystals grown without intentional doping are highly n-type with free electron concentrations exceeding 1019 cm3 . This leads to the slight expansion of the crystal lattice, corresponding to a mismatch to pure GaN of the order of 0.01%. The free electrons can be removed by compensation of the oxygen donor by magnesium added to the growth solution. Then the lattice constants become the same as for undoped material. Both n-type GaN and highly resistive GaN:Mg platelets have been used as seeds for HVPE growth. The (0001) polar surfaces of the seed crystals were prepared for the growth by mechanical and chemical polishing. The n-type GaN needles were used for seeding in their “as grown” state. The large (0001) oriented free standing GaN substrates were grown on GaN– sapphire templates as thick (1–2 mm) GaN crystals which usually separated from sapphire in large pieces if their thickness was sufficiently uniform [10]. Figure 3.2 shows some examples of the thick GaN crystals with characteristic hexagonal shape often observed as a result of spontaneous separation of the nitride from sapphire in the HVPE system used in this study. Such crystals were polished mechanically and etched by reactive ions to be used as substrates for further crystallization. Dislocation density in these substrates was in the range 1–5  106 cm2 depending on the “as grown” crystal thickness.

3.2.2 HVPE Reactor and Growth Conditions The HVPE system used in this work was the horizontal home-built quartz reactor with the rotating quartz susceptor designed for 2 in. diameter substrates. GaCl was supplied vertically, just over the surface of the susceptor with the use of shower head type quartz line. A growth temperature of about 1;050 ı C, temperature of GaCl synthesis of 870 ıC, HCl flow in the range of 5–30 ml=min diluted in 500 ml/min of

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N2 ; NH3 flow of 1,200 ml/min and 3,000 ml/min of N2 or H2 as carrier gases were applied for runs of 4–15 h. The growth rates observed for this geometry and the set of conditions varied, from 20 to 200 m=h and were dependent mainly on the flow of HCl. The main factor limiting the duration of the experiments was deposition of GaN microcrystals at the outlet of GaCl line thus changing the effective size of the openings during growth. All the crystals discussed in this study were grown without an intentional doping therefore their electrical properties were due to either intrinsic point defects or impurities incorporated nonintentionally during growth.

3.2.3 Characterization The bulk GaN crystals grown by HVPE were characterized by various methods allowing determination of structural quality as well as specification and distribution of point defects in the samples. The structural properties were evaluated by X-ray diffraction (XRD) DSE in molten KOH–NaOH eutectics [8, 9] of different surfaces of the crystals and also by transmission electron microscopy (TEM) [11]. For mapping of electron concentration, photo-electro-chemical etching (PEC) in basic solutions under UV illumination was used [12, 13]. The PEC etching rate is dependent on the free carrier concentration in the crystal so the crystal sectors with different electrical properties can be visualized. The PEC method was calibrated with the use of micro-Raman scattering technique [12] allowing mapping of local electron concentration (in certain range) with resolution of about 1 m.

3.3 Experimental Results 3.3.1 Crystals Grown on Small Near Dislocation Free GaN Platelet-Like Seeds The growth of GaN by HVPE on (0001) Ga-polar surface of the pressure grown platelets (HP platelets)is usually stable that is the crystallization front remains flat and the material is continuous. Some typical crystals are shown in Figs. 3.3a–d. The nonpolar cross-sections of the HVPE-grown crystals are shown in Fig. 3.3e, f. It can be seen that the crystal is macroscopically continuous (no inclusions or voids) and that the growth proceeds in both normal and lateral directions. Beside {0001} N and semipolar (mostly polar faces, the low index side walls of both nonpolar {1010} N {1011}) orientations are present. On polar (0001) surfaces of the crystals, macroscopic defects in the form of inverted polyhedral pits are observed. They can be due to defects related to the nonperfect surface preparation or noncoherent inclusions of GaN microcrystals

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Fig. 3.3 Morphology of GaN crystals grown by HVPE on HP platelets: (a–d) General view of the crystals, (e) (101N 0) nonpolar cross-section including HP substrate (dark stripe) and the HVPE GaN, (f) nonpolar cross-sections non including the HP substrate: upper edges correspond to semipolar {101N 1} and nonpolar {101N 0} side walls of the crystal, bottom edges are cut., Grid 1 mm on all images

especially at high supersaturations where nucleation and growth is relatively fast. The latter can be particularly relevant for shower head type geometry of GaCl outlet in the HVPE reactor. N nonpolar surfaces of such crystals revealed DSE of both (0001) polar and (1010) that despite a very high structural quality of the substrates (in terms of dislocation density) a significant number of lattice defects such as dislocations (at a density of 104 –106 cm2 ) and low angle grain boundaries is observed in the HVPE grown thick GaN [14, 15]. Typical DSE result is shown in Fig. 3.4. For growth on the n-type HP substrates, in (0001) direction with a rate of 100 m=h or higher and in N2 atmosphere, the defects are generated in HVPE crystals with thickness exceeding 30–50 m. Such critical thickness can be explained by a small lattice mismatch between strongly n-type substrate and HVPE grown GaN showing typical free electron concentration of less than 1017 cm3 for HVPE growth in Ga-polar (0001) direction. This problem was analyzed by Kry´sko et al [16] who determined curvature of HPGaN/HVPEGaN system in dependence on the HP substrate thickness. The substrates were gradually removed by reactive ion etching (RIE) in order to follow changes of deformation of the crystal. It was found that for “subcritical” HVPE layers, the system behaved in quite a good agreement with Stoney equation modified for the case of thick layers [17]. For layers thicker than 36 m, the crystal shape was constant, which indicated plastic deformation of the system. It seems however that the simple lattice mismatch between the substrate and the new grown crystal is not the only explanation of strain and defect generation in HVPE GaN crystals grown on small seeds because the phenomenon is observed also for HP substrates doped with magnesium for which the free electrons are fully compensated by the acceptor and no lattice mismatch between these crystals and

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dislocations

Low angle grain marker 50 µm

200 µm

Fig. 3.4 GaN crystal grown by HVPE on the n-type HP platelet after DSE: defects visualized as etch pits on both (0001) and (101N 0) surfaces of the crystal

Fig. 3.5 X-ray rocking curves for (0002) reflection for: (a) Mg-doped 110 m thick GaN substrate, (b) the same substrate with 100 m thick HVPE layer, (c) the HVPE layer after removing of the substrate. Both intensity and angle scales are the same for all diagrams

undoped GaN has been found. For the Mg-doped substrates and similar growth conditions (100 m=h; N2 atmosphere) the “critical” thickness was about 100 m and the behavior of the system was quite similar to the one with the n-type substrates. Such a behavior is illustrated in Fig. 3.5 where the X-ray rocking curves (XRC) are shown for the Mg-doped substrate, for the substrate with 100 m HVPE

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layer deposited on (0001) surface and for the HVPE layer after removing of the substrate by mechanical polishing. Deposition of the HVPE layer induced a strong bowing of the system reflected by widening of the rocking curve with corresponding decrease of its intensity. The initial width of the XRC can be partially recovered by the removal of the substrate. Such a behavior shows presence of strain in this virtually lattice and thermally matched system. Difficulties in reproduction of excellent quality of the small pressure grown GaN crystals by bulk crystallization by HVPE can be better understood after analysis of physical properties of HVPE grown GaN in dependence of orientation of the crystallization front. The application of characterization methods sensitive to the concentration of free carriers to the bulk GaN crystals grown by HVPE has shown [12, 15, 18, 19] that physical properties of GaN grown by HVPE are very much dependent on the orientation of the growth surface. This is illustrated in Figs. 3.6a and b, where some typical results of PEC of HVPE grown GaN crystals, are shown. The results are schematically summarized in Fig. 3.6c. For the Ga-polar (0001) growth surface, typical free electron concentration is usually lower than 1017 cm3 N one as well as for semipolar (1011) N surfaces the conwhereas for the opposite (0001) centration is higher by more than two orders of magnitude and exceeds 1019 cm3 . N surface (see next section) the concentration is For growth on the non polar (1010) also quite high and usually achieves 5–8  1018 cm3 . In the considered experimental system, the main source of free electrons in GaN is the oxygen donor. We have shown by secondary ion mass spectroscopy (SIMS) that GaN grown by HVPE contains much more oxygen at places where the crystallization front is rough with predominant semipolar orientation in contrast to the one corresponding to the flat (0001) face of the crystal. Figure 3.7 shows one such result where oxygen concentration has been measured in cross-sectional sample showing two types of growth front shapes: rough with a large fraction of semipolar orientation and a smooth (0001) oriented one. The shape of the crystallization front was recovered again by PEC.

<0001>

High n – no etching

100 µm

High n – no etching

> 1019

Low n – etching

< 1017

5 × 1018

Low n

> 1019 5 × 1018

> 1019

a

100 µm

b

C

Fig. 3.6 Anisotropy in physical properties of the GaN crystals revealed by photo-electro-chemical etching: (a) (101N 0) cross-section of GaN crystal with flat crystallization front grown by HVPE without an intentional doping, on n-type HP substrate: the substrate and “semipolar” part of the crystal shows no etching features due to very high electron concentration, the (0001) polar part of the crystal is etched, (b) similar crystal with inverted pyramid pits: both substrate and “semipolar” GaN inside the pits are not etched, (c) schematic illustration of electron concentration distribution for main low index faces of GaN during growth by HVPE without an intentional doping

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Fig. 3.7 Oxygen concentration in GaN crystal grown by HVPE in dependence on the shape of the crystallization front: (a) SIMS spectrum showing high oxygen content in part 1 of the crystal shown on the right, (b) cross-section of GaN crystal grown by HVPE after PEC showing changing shape of the crystallization front, (c) SIMS spectrum showing low oxygen content in part 2 (flat growth front) of the crystal

The observed dependences are very important and should be taken into account if bulk crystallization of GaN by HVPE is designed. It is particularly important for fast growth where the local changes of the growth surface orientation can take place. It can happen at the edges of three dimensional islands, macroscopic growth steps or at the walls of the pits. The effect of bowing of the substrate-HVPE layer system as well as structural properties of the final bulk crystals (never as good as the quality of the substrates) observed for both n-type and even the Mg-doped substrates support that growth mechanisms play a crucial role in the generation of strain and defects in GaN crystals grown by HVPE. One can easily imagine that the laterally grown part of the crystal (with high electron concentration) surrounding the internal low concentration sector of the crystal grown on (0001) face of the seed, can lead to deformation (bowing) of the substrate-layer sandwich. An additional (positive) support for this interpretation is the recent observation of a strong influence of the carrier gas on behavior of the investigated system. It seems that it is much easier to grow low defect density crystal on the HP plate-like seeds using hydrogen as a carrier gas in the HVPE experiment. It was shown that near dislocation free, over 1 mm thick crystals can be grown on (0001) surface of the n-type GaN platelet at a rate of about 100 m=h in H2 atmosphere. The only structural defects revealed by DSE of the (0001) surface of the resulting thick crystal were the ones corresponding to defects in the substrates (Fig. 3.8). This effect can be due to relative decrease of the rate of growth in lateral directions (observed also in [20]). The nonpolar GaN quasi wafers were made by slicing the GaN crystals grown by HVPE on the pressure grown GaN substrates Fig. 3.9a. Up to now this approach has been tried for GaN crystals grown on the high pressure platelets with relatively high rates, in N2 atmosphere. Therefore, the nonpolar wafers contained some disN oriented slices locations and grain boundaries (Fig. 3.4). The 300 m thick, (11 20) were prepared for epitaxy by mechanical and chemical polishing. Then, 10-fold

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b

Defects in the substrate

a Defects in HVPE-GaN

Defects in the substrate

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c

500 μm

Fig. 3.8 Ga-polar surface of 1 mm thick GaN crystal grown by HVPE on the n-type HP GaN substrate in H2 atmosphere: (a) (middle) – Nomarski contrast (DIC) image focused on (0001) surface of the HVPE grown GaN after DSE: dark etch pits indicate dislocations, large part of the surface is almost free of dislocations, (b,c) DIC images focused on (0001) surface of the substrate showing etch pits corresponding defected area in HVPE grown crystal

Fig. 3.9 GaN/AlGaN MQW deposited by PA MBE on GaN bulk substrates: (a) the substrates, (b) TEM image of the structure grown in nonpolar <11 2N 0> direction (courtesy of Julita Smalc), (c) excitonic spectra of the nonpolar GaN–AlGaN MQWs with different thicknesses

GaN/AlGaN sequences of 2, 3, and 4 nm thick GaN quantum wells separated by 7 nm Al0:11 Ga0:89 N barriers were deposited by plasma assisted MBE [15, 21] on both the nonpolar and the “usual” polar GaN substrates (for comparison of the optical properties). An advantage of the PA MBE method was that both thicknesses and compositions of the epitaxial layers composing the MQWs were the same for both polar and nonpolar growth directions. The structural quality of the “nonpolar” MQW was checked by XRD measurements and TEM. The TEM image of the nonpolar quantum structure is shown in Fig. 3.9b whereas excitonic photoluminescence (PL) spectra for the nonpolar MQWs of different thicknesses in Fig. 3.9c. The well resolved excitons in the PL spectra are good evidence of an excellent quality of the “nonpolar” quantum walls.

3.3.2 Crystals Grown on Small Near Dislocation Free GaN Needle-Like Seeds In case of crystallization on the needle-like HP seeds, the HVPE growth occurs N mostly in the nonpolar <10 10> directions, since the seed crystals are elongated in the direction <0001>. For conditions corresponding to the fast growth on the

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Fig. 3.10 GaN bulk crystals grown by HVPE on the needle-shaped high pressure seeds: (a) schematic illustration of the process (b) example crystals, (c) polar cross-section slice of the crystal grown in threefold 10 h. HVPE processes, interfaces between the HVPE growth runs are visible in optical microscope. Grid: 1 mm

Fig. 3.11 Images of (0001) cross-section of four times HVPE grown GaN crystal on HP GaN seed after DSE: (a) photograph of the cut platelet, HP seed crystal (area of dark contrast) and interfaces between subsequent HVPE growths are indicated. (b) optical microscopy image of the same platelet as in image (a) after etching in molten bases. Feature 1 shows the seed crystal, features 2 and 4 are areas with etch pits spreading from morphological imperfections of the seed, features 3 and 5 show low defect density parts of the HVPE bulk GaN, features 6 and 7 are parasitic, strongly defected grains. (c) scanning electron microscopy image of the boxed area on the image (b)

platelets in N2 atmosphere, the growth on the needles results in prismatic crystals with surprisingly stable morphology. Some examples are shown in Fig. 3.10b. Stability of the growth front can be perturbed by parasitic nucleation due to strain induced by macroscopic imperfections in the seed crystals (see further explanations in this section). The crystals can be reused for further crystallization which allows increasing their size (Figs. 3.10c and 3.11a). The crystals grown by HVPE on the needleshaped seeds were of n-type electrical conductivity with free electron concentration of about 5  1018 cm3 . However, as it was shown by micro-Raman scattering measurements [22], at the interfaces between subsequent HVPE runs the concentration of electrons was more than one order of magnitude higher than in the bulk. Due to this difference in point defect concentration the interfaces between subsequent HVPE runs are visible in optical microscope as shown in Figs. 3.10c and 3.11a. Structural defects in the new grown crystals have been studied by DSE in molten KOH/NaOH eutectics of both polar and nonpolar cross-section samples sliced from

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the bulk material as well as by TEM [11]. Current results of the DSE studies can be summarized as follows (also Fig 3.11): 1. DSE of [0001] surfaces indicate that its significant part (about 50%) is of very low defect density .<104 cm2 / and the seed crystal is almost dislocation free – Fig. 3.11b,c. 2. There is no deterioration in quality (in terms of dislocation density) between materials grown in subsequent HVPE runs. 3. There are areas on the [0001] surfaces with nonuniformly distributed hexagonal etch pits spreading from the morphological imperfections in the seed – Fig. 3.11. In some cases the array of dislocations converts into a misoriented parasitic grain. N surfaces (example in Fig. 3.12c) especially at regions 4. DSE of nonpolar [11 20] close to the parasitic grains, reveals lines perpendicular to the c-axis of the crystal being basal stacking faults as was recently confirmed by TEM measurements [11]. Both specification and distribution of structural defects deduced from the results of TEM and DSE studies of nonpolar cross-sections of the crystal grown by double HVPE process on the needle shaped seed is shown in Fig. 3.12. The TEM analysis has shown that despite high structural quality of the seed crystal, a number of lattice defects is present in the HVPE-grown part of the sample. There are both basal and prismatic stacking faults as well as dislocation loops at concentrations depending on the area of the crystal. One can distinguish a few possible reasons for generation and

Fig. 3.12 Structural properties of bulk GaN crystal grown by double HVPE process on HP needle shaped seed: (a) optical image of (112N 0) nonpolar cross-section of the crystal: x; y; z – areas of the sample analyzed by TEM, (b) schematic illustration of structural defects in the crystal based on TEM studies [11], (c) SEM image of nonpolar surface after DSE: BSFs and dislocation loops are revealed by the etching, (d) TEM image of the area “y” of Fig. 13.12a – basal stacking faults closer (left a)) and further (right b)) from the “parasitic grain”, inset c): HRTEM image of a stacking fault

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relaxation of strain leading to the formation of structural defects in the investigated system. As was mentioned earlier, the needle-like seeds grown under pressure are strongly n-type crystals with free electron concentrations approaching 1020 cm3 due to high oxygen content incorporated during the high pressure growth. On the other hand the electron concentration in GaN obtained by the HVPE method depends strongly on the orientation of the crystallization front during the HVPE growth. As was indicated in the previous section, for growth in the Ga-polar [0001] direction, typical electron concentration is about 1017 cm3 or lower, whereas for the growth in the N direction it can be more than two orders of magnitude opposite N-polar [0001] higher, and usually exceeds 1019 cm3 [7]. This very high electron concentration N faces of GaN. So for was also observed for growth on semipolar (inclined) {1011} the investigated system, in the case of bulk growth one can expect strain generation at the interface between the seed crystal and the HVPE-grown GaN, as well as in the areas where the crystal sectors growing in different directions thus having different electron concentrations, meet. The latter situation can take place at the apex of the considered needle-shaped seed where the growth proceeds in various directions and where a significant number of structural defects has been found. Another possible reason for the deterioration of the crystal quality is related with morphological imperfections of the seed. If a side face of the seed is not flat, but conN tains surfaces of orientations different from {1010}, the start of growth can again be accompanied by formation of domains with large relative variations in free electron concentration, which causes strain in the crystal. This situation is illustrated in Fig. 3.11 where a polar (0001) slice of the crystal similar to the one investigated by TEM is shown. The (0001) Ga-polar surface of this slice was subjected to a DSE procedure [9] in molten KOH–NaOH eutectics to reveal structural defects. The crystal was overgrown in four separate HVPE growth cycles. The DSE result indicates that the seed crystal is almost defect-free (no etch pits), as seen in area 1 in Fig. 3.11b. The etch pits spread mostly from places of strong deviation from the N orientation of the side faces of the seed, as seen in areas 2 and 4 in Fig. 3.11b. (1010) and also in Fig. 3.11c. The strain induced at the seed was so high that its relaxation required the formation of parasitic, strongly defected grains, indicated in areas 6 and 7 of Fig. 3.11b. The highest quality material containing BSFs at density lower than 105 cm1 and no dislocations was found in the areas for which the HVPE growth started from N face of the seed. It seems therefore that for crystallization of bulk GaN flat (1010) by HVPE method seeded with GaN single crystals elongated in <0001> directions the crucial is like in the previous case of the plate-like seeds, to avoid abrupt changes of free electron concentration in the growth system. It is difficult for small seeds in general if an increase of the crystal size in three dimensions corresponding to the growth in different directions is expected. For the needle-like prismatic N seeds, the growth in the nonpolar <1010> directions leading to an increase of diameter of the needle is most suitable for obtaining good quality crystal with uniform physical properties. The growth leading to an elongation of the needle (in the polar and semipolar directions) should be suppressed to avoid the strains caused

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by the different concentrations of free electrons in the neighboring sectors. Obviously the multiple growth should be replaced by a single, long duration run to avoid inhomogeneities related to the beginning of the HVPE crystallization process. For obtaining low defect densities, larger bulk crystals by HVPE seeded with high pressure needles, careful selection and preparation of seeds are necessary to avoid parasitic 3D nucleation on seed imperfections. It is shown that the HVPE growth on the needle-shaped seeds is a very stable process. Therefore, it seems that there are no physical obstacles for growing much larger crystals in this approach.

3.3.3 Crystals Grown on Large (0001) Oriented GaN Substrates As it follows from the previous considerations, the best way to avoid non uniform properties of GaN crystals grown by HVPE is the growth on large area, flat substrates with an orientation corresponding to the one of low index crystallographic planes. The best and obvious solution is to use large area free standing GaN substrate with low defect density. Since such substrates are still not commonly available, the (0001) oriented GaN–sapphire templates are often applied at the initial stage of this approach. In this study the free standing GaN substrates separated from sapphire were used (see Sect. 3.2.1) for HVPE growth of thick GaN crystals. Since the horizontal reactor allowed processes up to app. 10–12 h duration time, the crystals thicker than 2–2.5 mm were deposited in a few subsequent growth runs similar to the crystals grown on the needle shaped seeds (previous section). The (0001) surfaces of the thick crystals were polished mechanically and ion etched before being used for further growth whereas the side faces were not subjected to any special treatment. An example of the 5 mm thick GaN crystal grown in two runs on the substrate presented already in Fig. 3.2, is shown in Fig. 3.13c. The flat shape of (0001) Ga-polar surfaces is usually perturbed by the presence of deep inverted pyramid pits. The origin of these pits has been investigated by DSE of (0001) surfaces of slices obtained by cutting the thick crystals with diamond saw along the (0001) plane.

Fig. 3.13 GaN crystals grown by HVPE: (a) free standing GaN crystal grown by HVPE used as substrate for further growth, (b) GaN crystal grown by HVPE on the substrate “a” in 12 h process: total thickness 2.5 mm, (c) the same crystal after next 12 h. process: total thickness 5 mm. Grid 1 mm

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Inversion domains

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Fig. 3.14 Defects in polar (0001) slices of bulk GaN crystals grown by HVPE: (a) optical image: dark spots and dark border part of the upper slice indicate heavily oxygen doped GaN, (b) part of the slice after DSE, (c) Nomarski contrast image showing small etch pits and large dark feature at the center of the pit area, (d) central part of the pit area: etching result typical for N-polarity (0001N ) surface. DSE – courtesy of J. Weyher

The examples of the polar slices are shown in Fig. 3.14a. Despite geometrical flatness of the hexagonal surfaces, the pits are still visible as dark spots. Also the border area is darker than the dominating part of the crystal. This difference in color of different areas is due to dependence of point defect concentration on the growth direction which was already discussed in the previous sections. The DSE in molten bases reveals dislocations (etch pits in Figs. 3.14c and 3.16c) but also the domains of inverted polarity at the centers of the dark spots corresponding to the macroscopic inverted pyramid pits. This observation allows explanation why the macroscopic pits are so difficult to overgrow once they appear. The formation of the pits (at least some of them) is due to difference in growth rate of GaN with different polarity. Since the pits are composed of semipolar surfaces incorporating much more oxygen than the Ga-polar one, it is possible to recover their formation and growth history by optical microscopy and photochemical etching sensitive to the fluctuations in concentration of point defects. For example, the dark spots on the slice in Fig. 3.14 are traces of pits moving along growth direction thus leaving highly oxygen doped trace underneath. The question is if the inversion domains are present already in the GaN–sapphire template used for fabrication of the substrate are they generated during the HVPE growth. The behavior of the pits during subsequent HVPE growth processes can be followed by observation of nonpolar slices of the thick GaN crystals. The nonpolar (10–10) slices of the thick GaN crystal grown in four HVPE runs are shown in Fig. 3.15. The slices were obtained by cutting of the thick GaN crystals with a diamond blade. The result of the cutting is shown in Fig. 3.15a where the slices fixed on the holder are imaged. Figure 3.15b presents the nonpolar slices just after the sawing. Two areas “c” and “d” containing typical defects are indicated and

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d

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b 1

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4th HVPE growth 3rd HVPE growth 2nd HVPE growth 1st HVPE growth

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Fig. 3.15 Nonpolar (101N 0) slices of thick GaN crystal grown in 4 HVPE runs: (a) – the crystal on the diamond saw holder after slicing: side view on the left, top view on the right, (b) the nonpolar slices just after sawing, the areas labeled “c” and “d” are shown in bigger magnification in Figures 3.14c and 3.14d respectively, (c) the nonpolar slice after two side polishing: the inverted pyramid pit no 1 starts from the bottom and recovers in every subsequent growth run, (d) another nonpolar slice: the inverted pyramid pit no 2 behaves like the no 1 one, the one no 3 starts from the interface between first and second HVPE process

magnified in Fig. 3.15c, d respectively. The crystals in these figures were two side polished before being photographed. The inverted pyramid pit defects are well visible in these nonpolar cross-sections as dark triangles corresponding to heavily doped material growing in semipolar directions inside of the pits. When seen from the surface, the pits are usually partially overgrown (like the one no 3). Most of the pits start from the bottom of the crystal however some of them (like the one no 2) start at interfaces corresponding to the beginning of next HVPE growth. This clearly suggests that surface preparation of the thick crystals for subsequent HVPE processes has to be improved. A quite obvious solution of the problem is to grow thick crystals in a single long duration process which is in preparation. It is well visible from Fig. 3.15c, d, that once a pit is formed it is recovered in subsequent HVPE growth process despite the surface prepared for this process has been initially flat. Therefore chains of triangles reflecting “pit formation – polishing – new pit formation exactly over the previous one” cycle are observed. A good explanation for this phenomenon is the inversion domain at the center of the pit. Such a defect can propagate along the growth direction even if the growth is interrupted and the growth surface subjected to polishing and etching procedures.

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a

b

50 mm

c

Fig. 3.16 Defect selective etching of (0001) surface of 3.5 mm thick GaN crystal grown by HVPE: (a) GaN substrate used for the HVPE growth, (b) – 3.5 mm thick GaN crystal (“as grown”), (c) (0001) surface after DSE, etch pits density – 4  105 cm2

Besides the macroscopic defects just described, that have to be eliminated, the thick GaN crystals grown by HVPE are of good quality in terms of threading dislocation density measured by DSE of both polar and nonpolar surfaces. Fig. 3.16 shows an example of such a selective etching result. The (0001) “as grown” surface of the 3.5 mm thick crystal (Fig. 3.16b) grown on the substrate of Fig. 3.16a, after the DSE treatment is shown in Fig. 3.16c. The etch pit density is 4  105 cm2 . The results obtained in this study, have confirmed that high quality bulk crystals of GaN can be grown by HVPE in long duration crystallization processes, on large, flat seeds of (0001) orientation. The macroscopic defects in the form of large pits observed in our crystals are most probably due to imperfect surface preparation, leading to the appearance of inversion domains that locally inhibit the growth that induces formation of the pits. The limitation for obtaining larger bulk GaN crystals by HVPE method seems to be rather of a technical than physical nature.

3.4 Conclusions The experimental results of growth of bulk GaN crystals by HVPE method were presented. Both low dislocation density relatively small pressure grown crystals of various shapes and large flat (0001) oriented free standing GaN crystals were used as seeds for bulk crystallization. The most important problem in obtaining large uniform GaN crystals by HVPE encountered during our study were big differences in physical properties of GaN grown in different crystallographic directions. The differences in free electron concentration caused by differences in oxygen incorporation efficiency on particular crystal faces can exceed two orders of magnitude leading to strain and defect generation in the bulk crystals even if the seed is almost

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free of structural defects. The effect is particularly detrimental for growth on small seeds when an increase of the crystal size proceeds in three dimensions. The application of large flat seeds is a good way for obtaining big three dimensional high quality GaN crystals with uniform physical properties by HVPE because the growth proceeds in one crystallographic direction. The local changes of orientation of the crystallization front (macroscopic inverted pyramid pits) are most probably due to technical imperfections in both seed preparation and construction of HVPE reactor used in our study. Acknowledgements This work was supported by the Polish Ministry of Science and Higher Education in frame of Development Project Nr R02 036 01. The authors would like to thank Dr Rafał Jakieła from the Institute of Physics PAS (Warsaw, Poland) for SIMS measurements.

References 1. K. Motoki, T. Okahisa, R. Sirota, S. Nakahata, K. Uemats, N. Matsumoto, J. Cryst. Growth 305(2), 377–383 (2007) 2. Y. Oshima, T. Eri, M. Shibata, H. Sunakawa, K. Kobayashi, T. Ichihashi, A. Usui, Jpn. J. Appl. Phys. 42, L1–L3 (2003) 3. W. Utsumi, H. Saitoh, H. Kaneko, T. Watanuki, K. Aoki, O. Shimomura, Nat. Mater. 2, 735– 738 (2003) 4. T. Paskova, R. Kroeger, S. Figge, D. Hommel, V. Darakchieva, B. Monemar, E. Preble, A. Hanser, N.W. Williams, M. Tutor, Appl. Phys. Lett. 89, 051914 (2006) 5. A. Tyagi, H. Zhong, R.B. Chung, D.F. Feezel, M. Saito, K. Fujito, J.S. Speck, S.P. DenBaars, S. Nakamura, Jpn. J. Appl. Phys. 46, L444–L445 (2007) 6. V. Wagner, O. Parillaud, H.J. Buhlmann, M. Ilegems, S. Gradecak, P. Stadelman, T. Riemann, J. Christen, J. Appl. Phys. 92(3), 1307–1316 (2002) 7. I. Grzegory, M. Bo´ckowski, S. Porowski, in Bulk Crystal Growth of Electronic, Optical and Optoelectronic Materials, Chap. 6, ed. by P. Capper (UK, 2005), pp. 173–207 8. J.L. Weyher, P.D. Brown, J.L. Rouviere, T. Wosinski, A.R.A. Zauner, I. Grzegory, J. Cryst. Growth 210, 151 (2000) 9. G. Kamler, J.L. Weyher, I. Grzegory, E. Jezierska, T. Wosi´nski, J. Cryst. Growth 246, 21 (2002) 10. B. Łucznik, B. Pastuszka, I. Grzegory, M. Bo´ckowski, G. Kamler, J. Domagała, G. Nowak, P. Prystawko, S. Krukowski, S. Porowski, Phys. Stat. Sol. 3(6), 1453–1456 (2006) 11. J. Smalc-Koziorowska, G. Kamler, B. Łucznik, I. Grzegory submitted to J. Cryst. Growth 12. J.L. Wejher, R. Lewandowska, L. Macht, B. Łucznik, I. Grzegory, Mater. Sci. Semicond. Process. 9, 175 (2006) 13. G. Kamler, B. Łucznik, B. Pastuszka, I. Grzegory, S. Porowski, J. Cryst. Growth 310(15), 3478–3481 (2008) 14. B. Łucznik, B. Pastuszka, I. Grzegory, M. Bo´ckowski, G. Kamler, E. Litwin – Staszewska, S. Porowski, J. Cryst. Growth 281, 38–46 (2005) 15. I. Grzegory, H. Teisseyre, B. Łucznik, B. Pastuszka, M. Bo´ckowski, S. Porowski, in Nitrides With Nonpolar Surfaces, Chap. 3, ed. by T. Paskova (Wiley, New York, 2007), pp 53–71 16. M. Kry´sko, M. Sarzy´nski, J. Domagała, I. Grzegory, B. Łucznik, G. Kamler, S. Porowski, M. Leszczy´nski, J. Alloys Compounds 401, 261–264 (2005) 17. L. B. Freund, J. A. Floro, E. Chason, Appl. Phys. Lett. 74, 1987 (1999) 18. R. Lewandowska, J. L. Weyher, J.J. Kelly, L. Ko´nczewicz, B. Łucznik, J. Cyst. Growth 307, 298–301 (2007) 19. F. Tuomisto, K. Saarinen, B. Łucznik, I. Grzegory, H. Teisseyre, T. Suski, S. Porowski, J. Likonen, Apel. Phys. Lett. 86, 031915 (2005)

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20. H. Miyake, S. Bohyama, M. Fukus, K. Hiramatsu, Y. Iyechika, T. Maeda, J. Cryst. Growth 237–239, 1055–1059 (2002) 21. H. Teisseyre, C. Skierbiszewski, B. Łucznik, G. Kamler, A. Feduniewicz, M. Siekacz, T. Suski, P. Perlin, I. Grzegory, S. Porowski, Appl. Phys. Lett. 86, 162112 (2005) 22. R. Lewandowska, Unpublished result

Chapter 4

Freestanding GaN Wafers by Hydride Vapor Phase Epitaxy Using Void-Assisted Separation Technology Y. Oshima, T. Yoshida, T. Eri, K. Watanabe, M. Shibata, and T. Mishima

Abstract An outline is presented of the fabrication technique of freestanding GaN wafers by hydride vapor phase epitaxy using the void-assisted separation method and the properties of resulting crystals. A thick GaN layer of large area can be separated with excellent reproducibility from a base substrate by the application of thermal stress. This process is assisted by numerous voids formed near the interface between the thick GaN layer and the base substrate. By using this method, high-quality GaN wafers of large area with diameters of over 3 in. have been prepared.

4.1 Introduction Currently, hydride vapor-phase epitaxy (HVPE) provides the highest growth rate of GaN single crystals, reaching 1 mm/h or more. In addition, because the growth is usually performed at atmospheric pressure, large-scale HVPE reactors are available. For these reasons, HVPE has considerable advantages for the production of large-area GaN crystals for freestanding wafers. Freestanding GaN crystals can be obtained by growing a thick layer of GaN on a base substrate (e.g., sapphire, SiC, GaAs, etc.) and then removing the base substrate when growth of the GaN layer is complete (Fig. 4.1). However, the process can suffer from serious problems of cracking of the thick GaN layer due to the large mismatch in the thermal expansion coefficients between the GaN layer and the base substrate. Despite strenuous attempts to overcome this problem [1–3], difficulties still exist in reproducibly manufacturing GaN crystals that have a sufficiently large area for practical use. Reducing the density of dislocations is another important issue, because the HVPE method still involves hetero-epitaxial growth on a foreign substrate with a large lattice mismatch. As a solution to these problems, we have successfully established an original void-assisted separation (VAS) technology that permits the fabrication of large-area, high-quality GaN wafers with excellent reproducibility [4–6]. In this chapter, we

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Fig. 4.1 Schematic representation of the fabrication of a freestanding GaN wafer by HVPE

will describe the features of this HVPE-VAS technology and the properties of the GaN wafers that are produced by this method.

4.2 Outline of the HVPE-VAS Technology In this section, we discuss the basic concept of the HVPE-VAS technology and present a brief overview of the process.

4.2.1 Concept of the HVPE-VAS Technology In our method, we form a fragile layer that contains numerous small voids at the interface between a thick GaN layer and a base substrate. The thick GaN layer can be separated from the base substrate by breakage of the fragile layer due to the thermal stresses generated during the cooling process after the HVPE growth. The mechanical strength of the fragile layer can be controlled by changing the density of voids within the layer. It is therefore possible to cause separation by the application of a very small thermal stress. This enables highly reproducible large-area separation without producing any cracks. The reduction of dislocations is another important issue. Epitaxial lateral overgrowth (ELO) technology is one of the most successful methods for the reduction of dislocations in hetero-epitaxial GaN crystals [7–9]. The ELO-GaN layers make a considerable contribution in improving the performance and reliability of devices such as laser diodes [10]. However, the process, which involves photolithographic steps, is rather complicated for mass production of GaN wafers. In addition, the distribution of dislocations tends to be nonuniform. In the HVPE-VAS technology, we have developed an original method using thin porous interlayer with nanoscale holes; this is self-assembled by thermal agglomeration of a polycrystalline film. This technology permits the fabrication of the uniformly high-quality GaN crystals without the use of any complicated processes.

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4.2.2 Overview of the Process The VAS process consists of two basic steps: the preparation of the base substrate and HVPE growth on the base substrate. The base substrate comprises a GaN template layer, with many small voids on a sapphire substrate covered by a thin film of porous metal nitride. The HVPE growth is carried out on the base substrate. A proportion of the voids remain that are not completely refilled by HVPE-GaN. In addition, new voids are formed just above the metal nitride film. These voids assist in the process of self-separation by thermal stress. At room temperature, a considerable proportion of these voids break to release thermal stresses, so that the void-containing layer becomes very fragile. The thick GaN layer can then be readily separated from the base substrate, for example, by using small tweezers. The detail of each process and the related mechanisms are described below.

4.3 Preparation of a GaN Template with a Porous TiN Film In this section, we present details of the preparation of the base substrates and discuss the related mechanisms.

4.3.1 Experimental A GaN template with a thickness around 300 nm was first prepared on a (0001) sapphire substrate. A 20-nm-thick layer of Ti was then deposited on the GaN template by vacuum evaporation, and the Ti-deposited GaN template was annealed in a gaseous mixture of H2 and NH3 at 1;060 ıC for 30 min. The structural properties of the resulting annealed template were examined by scanning electron microscopy (SEM) and X-ray diffraction (XRD).

4.3.2 Results Figure 4.2 shows a cross-sectional SEM image of the GaN template after annealing (bird’s-eye view). The surface of the template is covered with a thin film having a net-like structure with many nanoscale holes. The typical size of these holes is in the range 20–30 nm. We therefore named this structure the “nano-net.” XRD measurements revealed that the nano-net consists of TiN oriented along the (111) direction of the cubic system. We also found that the GaN template was heavily attacked during annealing, resulting in a pyramidal void structure under the TiN nano-net. Formation of these voids was enhanced by increasing the partial pressure of H2 , increasing the temperature during the annealing process, or increasing the thickness of the Ti layer.

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Fig. 4.2 SEM image of the GaN template after annealing

Base layer

Annealing hole

Base layer

Fig. 4.3 Schematic representation of the thermal agglomeration of a polycrystalline film

4.3.3 Mechanisms for the Formation of the Porous Structure 1. Formation of the nano-net structure by thermal agglomeration The nano-net structure automatically develops as a result of thermal agglomeration of the polycrystalline film. Thin films are thermodynamically quasi-stable, and they agglomerate into a net-like structure during thermal treatment [11]. Figure 4.3 is a schematic showing the thermal agglomeration of a polycrystalline film. During heating, deformation starts at grain boundaries to minimize the total energy of grain surfaces, grain boundaries, and interfaces between the film and the base layer. As a result, numerous small holes form if the initial thickness of the film is sufficiently small. The Ti film obtained by our process should have many grain boundaries as a result of the lattice mismatch between the Ti film and the underlying GaN. This leads to the formation of the nano-net structure. The conservation of the number of Ti atoms during the annealing process (i.e., no significant etching of Ti occurs) can be confirmed by comparing the volume of the film before and after annealing,

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0.5 m

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Fig. 4.4 Cross-sectional SEM images of Ti/GaN templates after annealing at various partial pressures of H2

taking into account the increase in the volume as a result of nitridation. Theoretically, the increase should be approximately 8.1%. The volume of the nano-net can be estimated from surface and cross-sectional SEM images. The volume of the nano-net was found to be almost identical to the theoretical value. Thus, the thermal agglomeration mechanism was corroborated. 2. Void formation in the GaN template layer The voids in the template GaN layer are formed by the etching of GaN by H2 gas. Figure 4.4 shows cross-sectional SEM images of GaN templates covered by Ti films after annealing under three sets of conditions with different H2 partial pressures P .H2 /. The development of the voids in the GaN layer was enhanced by increasing P .H2 /. In the case of P .H2 / D 100 kPa, the GaN layer virtually disappeared. A similar dependence on the temperature during the annealing was also observed. No Ga droplets or GaN-related residuals were found on the surface of the template or inside the voids under any of the conditions. This indicates that Ga generated by the decomposition of GaN was removed through the formation of gaseous gallium hydride [12, 13].

4.4 HVPE Growth on GaN Templates with a Porous TiN Film In this section, we describe the HVPE growth of GaN on the porous GaN template covered by a TiN nano-net. First, we describe the growth process of HVPE-GaN on the template, and secondly we discuss the mechanism of the separation process.

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4.4.1 Process of HVPE Growth and Separation 1. HVPE reactor and source materials In our study, we used conventional, hot-wall-type, horizontal HVPE equipment with a quartz reactor tube. GaCl and NH3 were used as sources of gallium and nitrogen, respectively. GaCl was formed in situ upstream in the reactor by the reaction between liquid Ga and HCl gas at 850 ı C. The temperature of the growth region was kept constant at 1;060 ıC during the growth. N2 ; H2 , and their mixtures were used as carrier gases. 2. Process of HVPE growth on GaN template with nano-net and substrate separation SEM images at each stage of growth are shown in Fig. 4.5a–e: (a) Before growth .t D 0 min/. (b) t D 0:5 min. The surface image shows no significant change. However, the cross-sectional image shows that some voids in the underlying GaN layer are about to be refilled. This indicates that the GaN layer plays a role as a seeding layer.

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(c) t D 1 min. Small islands of GaN with diameters of several micrometers are randomly formed N side facets. Note that on the surface. Each island is surrounded by inclined .1011/ the spacing between these islands is of the order of micrometers, which is much larger than the spacing of holes in the nano-net (several tens of nanometers). This implies that only a limited number of holes in the nano-net can contribute to island formation. (d) t D 5 min. GaN islands grow vertically and laterally to coalesce with each other. The planar (0001) surface starts to develop. On the other hand, voids are found near the nanonet. Note that voids remain not only under the nano-net (lower voids) but also on the nano-net (upper voids). The upper voids are newly formed during the growth. (e) t D 15 min. The flat (0001) surface is significantly developed. Incomplete coalescence fronts can be seen as pits. Further growth results in a fully coalesced pit-free surface.

4.4.2 Mechanisms of Growth and Separation 1. Separation of the base substrate After HVPE growth, the thick GaN layer is separated from the base substrate without cracking. Practically, the separation can readily be accomplished by inserting a sharp knife-edge into the interface between the thick GaN layer and the base substrate at room temperature. Figure 4.6 shows a photograph of a freestanding GaN wafer and the base substrate just after separation. The diameter of the GaN wafer is 3.2 in.; this is the largest single crystalline GaN ever reported [14, 15]. Its surface is specular. Ga and TiN adhering to the backside of the GaN wafer can be seen through the wafer because the crystal is quite transparent. Figures 4.7a, b show the SEM images of the backside of the GaN wafer and the upper surface of the base substrate, respectively. Almost all the nano-net remained on the base substrate: this shows that separation occurred at the interface of the thick GaN layer and the nano-mask, suggesting that the voids formed on the nano-net (upper voids) make an essential contribution to the separation process. 2. Mechanisms of the formation of the upper voids To investigate the mechanisms of formation of the upper voids, we carefully examined by SEM the interface near the nano-net in the early stages of HVPE growth (Fig. 4.8). First, we found that the upper voids do not have openings at the periphery of the GaN islands. This implies that the upper voids are not grown-in structures, but that they develop after the formation of the GaN islands. Secondly, we found that the

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GaN

Base substrate

Fig. 4.6 Photograph of an as-separated freestanding GaN and the base substrate

Fig. 4.7 SEM images of (a) the substrate-facing side of the GaN wafer and (b) the surface of the base substrate

5 m Fig. 4.8 Cross-sectional SEM image of a region near the nano-net at an early stage of HVPE growth

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Fig. 4.9 Schematic representation of the process for formation of an upper void

upper voids are always accompanied by lower voids just below them. Therefore, the lower voids probably make an essential contribution to the formation of the upper voids. Based on these considerations, the mechanism of the upper void formation is described as follows (see Fig. 4.9 for a schematic model). (a) Nonuniform island formation GaN islands do not form uniformly on the nano-net because of the nonuniformity of the sizes of the lower voids and the nano-net openings. Therefore, island formation progresses faster in some regions and is retarded in other regions. These regions are intermixed adjacently. (b) Lateral overgrowth of previously formed GaN islands and confinement of lower voids Previously formed GaN islands grow laterally over the delayed-growth region to confine the lower voids. (c) Formation and development of upper void by the decomposition of the N-polar face of the GaN islands Once the lower voids are confined, the supply of source materials for GaN, i.e., GaCl and NH3 , is virtually stopped. Decomposition of GaN in the confined space then begins as a result of the high growth temperature, which is in excess of 1;000 ı C. The decomposition is significant on the N-face of the island, and upper voids form as a result.

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4.5 Properties of GaN Wafers Fabricated by HVPE–VAS Technology 4.5.1 Structural Properties N Figure 4.10 shows X-ray diffraction curves from the (0002) plane and the .1010/ plane of the VAS-GaN wafer. The full widths at half-maximum (FWHMs) of the N peaks, i.e., the tilt and twist angles, were narrow, namely, GaN (0002) and .1010/ 35 and 52 arcsec, respectively, whereas the tilt and twist angles of the GaN template layer were approximately 600 and 1,400 arcsec, respectively. The dislocation density of the VAS-GaN, as determined by cathode luminescence measurements, was not more than 106 cm2 , whereas the dislocation density of the GaN template layer was around 1010 cm2 . Usui et al. reported that tilt/twist angles become narrow as the dislocation density decreases [16]. The relationship between the dislocation density and tilt/twist angle of the VAS-GaN wafer described above agrees well with this tendency. To clarify the mechanism of the dramatic reduction in the dislocation density, we used transmission electron microscopy (TEM) to investigate the behavior of dislocations in GaN islands formed in the early stage of the HVPE growth. Figure 4.11 shows a cross-sectional TEM image of a GaN island. The cleaved plane intersects the area within 200 nm of the island center. A considerable number of dislocations can be seen in the GaN template layer. Note that there is a region with a lower dislocation density just below the center of the island. Bending of dislocations can be observed in this area. The dislocation density in the island is much lower than that in the underlying GaN layer. Dislocations propagating from the underlying GaN layer were observed in the central bottom region of the island: these bent and formed loops. On the other hand, few dislocations were seen in the upper part of the island. In the laterally grown area of the island, we sometimes observed bending of dislocations and their horizontal propagation. Figure 4.12 shows a cross-sectional TEM image of the area near the coalescence front of two islands. No significant generation

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Fig. 4.11 Cross-sectional TEM image of a GaN island

Fig. 4.12 Cross-sectional TEM image of a region near the coalescence front of two islands

of dislocations is seen at the coalescence front. This is in contrast to the case of conventional ELO. On the basis of these TEM results, the mechanism of dislocation reduction can be described as follows. 1. Dislocation reduction of GaN template layer by refilling of voids As discussed in Sect. 4.4.1, a proportion of the voids in GaN template layer are refilled during the early stages of HVPE growth. In this process, bending of dislocations could occur through refilling of the pyramidal voids. As a result, propagation of dislocations beyond the nano-net could be suppressed. 2. Dislocation bending by growth of faceted GaN islands In the second stage of HVPE growth, GaN islands of submicrometer size form through the openings of the nano-mask. A proportion of dislocations in the GaN

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template layer may extend into the islands. However, the propagation of these dislocations toward [0001] is suppressed by bending of the dislocations toward the horizontal direction because of growth of the faceted islands. Some of the dislocations may be annihilated by meeting together and forming dislocation loops. Basically, this process of dislocation reduction is similar to that of facet-initiated ELO (FIELO) [7]. 3. Further dislocation bending by large inclined facets develops as a result of island growth and coalescence On continuing the growth, the small GaN islands grow laterally and vertically to coalesce with each other. As a result, larger islands with larger inclined facets develop so that the dislocation bending is further enhanced. This process occurs repeatedly until complete coverage is accomplished. Consequently, a considerable reduction in the number of dislocations is achieved.

4.5.2 Electrical Properties [17, 18]

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The electrical properties (carrier concentration, resistivity, and electron mobility) of the VAS-GaN wafers were investigated by the van der Pauw method at room temperature. Figure 4.13 shows the donor concentration dependences of carrier concentration and electrical resistivity. The carrier concentration is almost identical to the donor concentration and increases linearly with increasing donor concentration up to 6:7  1018 cm3 , although a slight discrepancy is observed at 1:5  1019 cm3 . Electrical resistivity decreases with increasing donor concentration, reaching a minimum of 2:5 m cm at n D 1:24  1019 cm3 . The dependence of electron mobility on the carrier concentration is shown in Fig. 4.14. Data reported in the literature [19] are also shown for comparison. Our values were significantly larger than the reported values.

0 20

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Fig. 4.13 The dependence of the carrier concentration and of the electrical resistivity on the donor concentration

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Fig. 4.14 The dependence of the electron mobility on the carrier concentration

Let us consider the origin of this difference. The major difference between the two sets of crystals is in the dislocation density, which is of the order of 106 cm2 for our crystal, and of the order of 109 cm2 for the films grown through conventional MOVPE (Metalorganic Vapor Phase Eiptaxy) process[19], as estimated by using the relationship between dislocation density and FWHM of the X-ray rocking curve deduced by Gay et al. [20]. It is known that dislocations having an edge component introduce acceptor centers along the dislocation line, which capture electrons from the conduction band in n-type semiconductors. This results in the dislocation lines becoming negatively charged and a space charge forming around these dislocation lines; this scatters electrons traveling across the dislocations, thus reducing the mobility [21]. However, note that a dislocation density of 1010 cm2 corresponds to a concentration of only 2  1017 cm3 of dangling bonds. Thus, the dislocation effect is significant for crystals having a high dislocation density and a low carrier concentration. For crystals with a carrier concentration of more than 1  1018 cm3 , the dislocation effect for densities below 109 cm2 is negligible. Therefore, it is likely that the difference in mobility arises from the difference in the density of point defects (vacancies and/or their complexes), which could be affected by the strain field of dislocations.

4.5.3 Thermal Properties 1. Thermal conductivity [17, 18, 29] Thermal conductivity is one of the most important material parameters for heat dissipation in high-power devices. We present the results of thermal conductivity measurement of VAS-GaN wafers with various carrier concentrations. The thermal conductivity › was calculated by using the relationship D A cp ;

(4.1)

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Thermal conductivity [W cm–1 K–1]

where A is the thermal diffusivity, is the mass density, and cp is the specific heat under constant pressure. A was measured by the laser-flash method [22]. This is a standard method for measuring the thermal diffusivity of bulk materials above room temperature. The surface of a GaN platelet is heated uniformly by using a pulsed laser, and the temperature of the rear surface increases. This is a one-dimensional thermal diffusion phenomenon. Thermal diffusivity is simply estimated from the time dependence of the temperature increase at the rear surface in conjunction with the sample thickness. was measured by the Archimedes method at 296 K. cp was measured by differential scanning calorimetry (DSC). The accuracy of these results was estimated to be better than ˙5%. Figure 4.15 shows the carrier concentration dependence of . The data for the HVPE-GaN/sapphire samples of Florescu et al. [23] and for the samples produced by the high-pressure, high-temperature method by Jezowski et al. [24] are also shown for comparison. The result obtained by Florescu et al., D 1:95 W cm1 K1 at n D 6:9  1016 cm3 , is one of the largest values ever reported. However, the thermal conductivity decreased linearly with log n by about a factor of two for every decade increase in n. As a result, the thermal conductivity decreased to 0:76–1:12 W cm1 K1 at approximately n D 2  1018 cm3 , which is the carrier concentration usually needed to establish a reliable ohmic contact on the surface. On the other hand, we found that the thermal conductivity of our crystal exhibited a much weaker dependence, and maintained high values even in the region of high carrier concentration, i.e., D 2:0 W cm1 K1 for n D 1:0  1018 cm3 , and D 1:87 W cm1 K1 for n D 1:24  1019 cm3 . From a practical point of view, it is very important that a high thermal conductivity is confirmed to be compatible with high carrier concentrations. In GaN crystals, the dominant contribution to the room-temperature thermal conductivity comes from acoustic phonons. The contribution of electrons to the thermal conductivity is three orders of magnitude smaller than the lattice contribution in the measured range of carrier concentrations [23]. Therefore, the observed difference is probably due to a difference in the lattice contribution. According to Zou

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et al., the contribution of dislocations at a density below 1010 cm2 is negligible [25]. Although the dislocation density is not presented in the literature, it does not usually exceed a value of 1010 cm2 by a significant amount. Therefore, the predominant source of the difference in thermal conductivity is, presumably, not the difference in the dislocation densities themselves. Zou et al. attributed the decrease in thermal conductivity reported by Florescu et al. [23] to the contribution of point defects [25]. However, Jezowski et al. reported that bulk materials with ultralow dislocation densities of approximately 102 cm2 or less, obtained by the highpressure, high-temperature method, showed large values of the thermal conductivity, i.e., D 1:60–2:26 W cm1 K1 , despite the inclusion of various impurities at high concentrations (e.g., ŒMg D 1  1018 cm3 ; ŒC D 1  1019 cm3 ; ŒH D 1  1017 cm3 ; ŒSi D 1  1017 cm3 , and ŒO D 1  1020 cm3 / [24]. These results imply that the decrease in thermal conductivity with increasing concentration of point defects, such as impurities and vacancies, could be suppressed when the quality of the matrix crystal is very high. Therefore, the weak dependence of the thermal conductivity of our crystals on the impurity concentration might originate from their excellent crystal quality. Further investigations should be conducted to understand the combined effect of point defects and dislocations on thermal conductivity. 2. Thermal expansion coefficient [18, 26] We report the results of our measurements of the thermal expansion coefficient of VAS-GaN crystals. The thermal expansion coefficient ’ is especially important for the design of hetero-epitaxial structures. Several studies on thermal expansion coefficients have been reported [27–29]. For example, in 1969, Maruska and Tietjen determined the value of ˛ perpendicular to the C -axis .<0001>/ to be 5:59  106 K1 (300–900 K) by using material grown by HVPE [27]. In 1995, Grzegory et al. obtained values of 3:1  106 K1 (294 K) and 6:2  106 K1 (700 K) for a material synthesized by a high-pressure, high-temperature method [28]. These values were obtained by powder XRD measurements. However, actual changes in specimen dimensions are also significant from the practical point of view. Therefore, it is of great interest to examine the thermal expansion coefficients of GaN by measuring the dimensional change in high-quality single crystals. However, there are only a limited number of reports on the determination of ˛ by such direct measurement techniques [26], and there are no reports on the measurement of the dependence of ˛ on impurity concentrations. The temperature dependence of thermal expansion .L=L0 / along the N N A <1120> and M <1010> axes was measured to an accuracy of ˙0:2  106 K1 for VAS-GaN crystals containing various concentrations of impurity (silicon) by means of thermal mechanical analysis (TMA) at 298–573 K. Values of ˛ along each direction (˛A and ˛M ) were then calculated from the resulting curves. Impurity concentration dependences of the thermal expansion coefficients are shown in Fig. 4.16. L=L0 (not shown) increases approximately proportionately to the temperature within the measured temperature range. ˛A and ˛M were approximately the same and they are confirmed to be nearly constant, i.e., 5:1  106 (averaged value from 298 to 573 K), in the measured impurity concentration range.

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Fig. 4.16 The dependence of the thermal expansion coefficient on impurity concentrations

Fig. 4.17 Photograph of a 2-in. and a 3-in. VAS–GaN wafer

4.6 Summary An outline is presented of the HVPE-VAS technology and the properties of VASGaN crystals. By means of the HVPE-VAS method, a thick layer of GaN can be separated with excellent reproducibility from a base substrate by the application of thermal stress. This process is assisted by numerous voids formed near a TiN nanonet. By using the HVPE-VAS method, high-quality GaN wafers with diameters of over 3 in. have been prepared. Figure 4.17 shows a photograph of a 2-in. and a 3-in.

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Table 4.1 Properties of a VAS-GaN crystal Property Value Tilt angle 35 arcsec Twist angle 52 arcsec Dislocation 106 cm2 density Carrier n D 1  1018 cm3 concentration Mobility  D 441 cm2 V1 s1 Thermal ˛A .298–573 K/ D .5:1 ˙ 0:2/  106 K1 expansion coefficient ˛M .298–573 K/ D .5:1 ˙ 0:2/  106 K1 Mass density

D 6:076 ˙ 0:006 g cm3 Specific heat cp D .4:22 ˙ 0:04/  101 Jg1 K1 Thermal AC== D .7:72 ˙ 0:39/  101 cm2 s1 diffusivity AC? D .7:77 ˙ 0:39/  101 cm2 s1 Thermal C== D 2:0 ˙ 0:1 W cm1 K1 conductivity C? D 2:0 ˙ 0:1 W cm1 K1 Refractive index 2.58 (at 358 nm)

Optical absorption coefficient

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Archimedes method DSC Laser flash method

i D Ai ¡cp

Transmittance and reflectance measurement

<1:4 cm1 .<358 nm/

VAS-GaN wafer. The principal properties of the VAS–GaN crystals are summarized in Table 4.1.

References 1. Yu.V. Melnik, K.V. Vassilevski, I.P. Nikitina, A.I. Babanin, V. Yu. Davydov, V.A. Dmitriev, MRS Internet J. Nitride Semicond. Res. 2, 39 (1997) 2. S. Nakamura, M. Senoh, S. Nagahama, N. Iwasa, T. Yamada, T. Matsushita, H. Kiyoku, Y. Sugimoto, T. Kozaki, H. Umemoto, M. Sano, K. Chocho, Appl. Phys. Lett. 72, 2014 (1998) 3. M.K. Kelly, O. Ambacher, R. Dimitrov, R. Handschuh, M. Stutzmann, Phys. Status Solidi (A) 159, R3 (1997) 4. Y. Oshima, T. Eri, M. Shibata, H. Sunakawa, A. Usui, Phys. Stat. Sol. (A) 194(2), 554 (2002) 5. A. Usui, T. Ichihashi, K. Kobayashi, H. Sunakawa, Y. Oshima, T. Eri, M. Shibata, Phys. Stat. Sol. (A) 194(2), 572 (2002) 6. Y. Oshima, T. Eri, M. Shibata, H. Sunakawa, K. Kobayashi, T. Ichihashi, A. Usui, Jpn. J. Appl. Phys. 42, L1 (2003) 7. A. Usui, H. Sunakawa, A. Sakai, A.A. Yamaguchi, Jpn. J. Appl. Phys. 36, L899 (1997) 8. H. Hiramatsu, K. Nishiyama, A. Motogaito, H. Miyake, Y. Iyechika, T. Maeda, Phys. Stat. Sol. (A) 176, 535 (1999)

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9. T.S. Zheleva, S.A. Smith, D.B. Thomson, T. Gehrke, K.J. Linthicum, P. Rajagopal, E. Carlson, W.M. Ashmawi, R.F. Davis, Mater. Res. Soc. Symp. Proc. 537, G3.38 (1999) 10. S. Nakamura, M. Senoh, S. Nagahama, N. Iwasa, T. Yamada, T. Matsushita, H. Kiyoku, Y. Sugimoto, T. Kozaki, H. Umemoto, M. Sano, K. Chocho, Jpn. J. Appl. Phys. 36, L1568 (1997) 11. D.J. Srolovitz, M.G. Goldiner, J. Miner. Met. Mater. Soc., Mar. 31 (1995) 12. R. Groh, G. Gerey, L. Bartha, J. I. Pankove, Phys. Stat. Sol. (A) 26, 353 (1974) 13. A. Koukitu, M. Mayumi, Y. Kumagai, J. Cryst. Growth 246, 230 (2002) 14. T. Yoshida, Y. Oshima, T. Eri, K. Ikeda, S. Yamamoto, K. Watanabe, M. Shibata, T. Mishima, J. Cryst. Growth 310, 5 (2008) 15. T. Yoshida, Y. Oshima, T. Eri, K. Watanabe, M. Shibata, T. Mishima, Phys. Stat. Sol. (A) 205(5), 1053 (2008) 16. A. Usui, H. Sunakawa, K. Kobayashi, H. Watanabe, M. Mizuta, Mater. Res. Soc. Symp. Proc. 482, G5.6.1 (2001) 17. Y. Oshima, T. Yoshida, T. Eri, M. Shibata, T. Mishima, Jpn. J. Appl. Phys. 45, 7685 (2006) 18. Y. Oshima, T. Yoshida, T. Eri, M. Shibata, T. Mishima, Phys. Stat. Sol. (C) 4(7), 2215 (2007) 19. S. Nakamura, T. Mukai, M. Senoh, Jpn. J. Appl. Phys. 31, 2883 (1992) 20. P. Gay, P.B. Hirsch, A. Kelly, Acta Metall. 1, 315 (1953) 21. N.G. Weimann, L.F. Eastman, D. Doppalapudi, H.M. Ng, T.D. Moustakas, J. Appl. Phys. 83, 3656 (1998) 22. W.J. Parker, R.J. Jenkins, C.P. Butler, G.L. Abbott, J. Appl. Phys. 32, 1679 (1961) 23. D.I. Florescu, V.M. Asnin, F.H. Pollak, J. Appl. Phys. 88, 3295 (2000) 24. A. Jezowski, B.A. Danilchenko, M. Bockowski, I. Grzegory, S. Krukowski, T. Suski, T. Paszkiewicz, Solid State Commun. 128, 69 (2003) 25. J. Zou, D. Kotchetkov, A.A. Balandin, D.I. Florescu, F.H. Pollak, J. Appl. Phys. 92, 2354 (2002) 26. Y. Oshima, T. Suzuki, T. Eri, Y. Kawaguchi, K. Watanabe, M. Shibata, T. Mishima, J. Appl. Phys. 98, 103509 (2005) 27. H.P. Maruska, J.J. Tietjen, Appl. Phys. Lett. 15, 327 (1969) 28. I. Grzegory, J. Jun, M. Bockowski, M. Wroblewski, B. Lucznik, S. Porowski, J. Phys. Chem. Solids 56, 639 (1995) 29. E. Ejder, Phys. Stat. Sol. (A) 23, K87 (1974)

Chapter 5

Nonpolar and Semipolar GaN Growth by HVPE Paul T. Fini and Benjamin A. Haskell

Abstract This chapter describes the structural and morphological characteristics of planar nonpolar and semipolar GaN films grown by hydride vapor phase epitaxy (HVPE). While smooth enough to allow device regrowth and fabrication, these films often contain 1010 cm2 threading dislocations and 105 cm1 basal plane stacking faults. Lateral epitaxial overgrowth (LEO) has been developed to largely eliminate dislocations and stacking faults in the overgrown material, resulting in significant improvements in surface morphology and luminescence characteristics of these films and optoelectronic devices grown upon them.

5.1 Introduction Gallium nitride is a wide bandgap semiconductor that is most stable in the wurtzite structure, depicted in Fig. 5.1. This structure is described by a hexagonal coordinate system having three equivalent basal plane lattice vectors, a1 ; a2 , and a3 , and a unique c translation vector. In three-index notation (a1 ; a2 , c), gallium atoms are located at (0, 0, 0) and (1/3, 2/3, 1/2), while the nitrogen atoms sit on an equivalent sublattice at (0, 0, u) and (2/3, 1/3, 1=2 C u), where u D 0:377. Referring to Fig. 5.1, N a-planes, f1010g N mlow-index planes in the wurtzite structure include f1120g planes, as well as the .0001/ C c plane (also referred to the as the “Ga-face”) and N  c plane (also referred to as the “N-face”). Other low-index planes the .0001/ N N and f1013g, N which include the so-called semipolar planes, such as f1011g; f1122g, are inclined with respect to both the nonpolar planes and the c-plane, as shown in Fig. 5.2. The potential utility of semipolar planes is discussed below. Unfortunately, attempts over 30 years of GaN epitaxy research repeatedly suggested that nonpolar and semipolar planes were “unstable” [1–8] compared to the conventionally grown c-plane. One study after another yielded nonpolar films with surfaces too rough and/or faceted for device growth. In 2000, Waltereit et al. first demonstrated planar m-plane GaN growth via molecular beam epitaxy (MBE) [9]. This demonstration was followed by Craven et al.’s metal organic chemical vapor deposition (MOCVD) growth of planar a-plane GaN films in 2002 [10]. 97

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Fig. 5.1 Scale representation of the GaN crystal structure in hexagonal coordinates. Ga atoms are shaded grey, with N atoms shaded blue

Fig. 5.2 Example schematics of low-index semi-polar GaN crystal planes

While considerable progress was subsequently made in thin-film growth of nonpolar GaN, thick-film or bulk growth of nonpolar (and later, semipolar) orientations continued to be elusive until recently. The performance of the nonpolar GaN-based devices would therefore be limited by the lack of low-defect-density film and substrate options. This chapter describes the progress achieved in thick-film nonpolar and semipolar GaN growth via hydride vapor phase epitaxy (HVPE), toward the ultimate goal of producing low-defect-density nonpolar and semipolar GaN thick films and substrates. As a well-established vapor-phase growth technique, GaN HVPE currently offers the most attractive combination of low cost, high growth rate, and scalability for commercial production of GaN films and substrates.

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5.2 Heteroepitaxial Films, Including Substrate Selection 5.2.1 Planar a-Plane GaN Films While heteroepitaxial growth of planar a-plane GaN films presents considerable challenges, substrate selection fortunately is not among them. Previous efforts, while failing to achieve smooth film surface morphology, had demonstrated that r-plane sapphire substrates provide an acceptable lattice match and yield uniformly oriented a-plane GaN films. The sapphire r-plane contains a rectangular grid of oxygen ions with an additional O2 ion within each rectangle. One edge of the grid N direction, while the other is the Œ1101 N is the Œ1120 direction, which lies parallel to the projection of the sapphire c-axis on the r-plane surface. The lattice mismatch between r-plane Al2 O3 and a-plane GaN is 1% along the GaN c-direction and 16% along the GaN m-direction [11]. The planar a-plane GaN films described below were grown by HVPE directly on r-plane sapphire substrates without the use of any low-temperature buffer or nucleation layers. Typical V/III ratios ranged from 9 to 40, with nitrogen and hydrogen mixtures serving as carrier gases. Typical growth rates ranged from 15 to 95 m=h at substrate temperatures between 1,040 and 1;070 ıC [12]. Figure 5.3a shows a Nomarski optical contrast micrograph of a representative HVPE-grown 50 m thick a-plane GaN film. Its surface was characterized by long-range “flow patterns” that had peak-to-valley heights of 100–500 nm over 75–500 m lateral extents, as measured by profilometry. These low-angle surface features scattered light minimally; the films were specular and optically transparent. Other long-range surface features included faint ridges or “scales” of similar heights but with varying directions over the wafer surface. No correlation was found between these features and the underlying crystal structure of the films. In contrast to the previously reported a-plane GaN surfaces having facets inclined 30–50ı from the surface normal [13–15], the angular variation of these surface features was 0:2–0:8ı and could not be attributed to faceting. Subsurface cracks oriented nearly perpendicular to the GaN c-axis were observed in such films, one of which is detailed in the cross-sectional SEM image in Fig. 5.3b. These internal cracks, similar to those observed in c-plane GaN films [16], generally “healed” within 5–10 m and generally did not reach the free surface during growth. We believe these cracks result from plastic relief of tensile strain that may be a consequence of grain–grain coalescence, such as been proposed for the MOCVD c-plane GaN films [17]. Figure 5.4 shows a representative atomic force microscopy (AFM) scan of an HVPE-grown a-plane GaN film. Local root mean-square (RMS) roughness over 5  5 m sampling areas was typically 0.5–0.8 nm. The RMS roughness over 100–400 m2 areas typically remained below 2 nm. The surface was dominated by a high density of small (3–7 nm deep) pits with densities ranging from 2  109 to 9  109 cm2 . As in the case of c-plane GaN [18] and also correlated by plan-view transmission electron microscopy measurements, these pits decorate intersections

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Fig. 5.3 (a) Plan-view Nomarski optical contrast micrograph of an a-plane GaN film. (b) Crosssectional SEM image revealing an internal crack in an a-plane GaN film which healed upon subsequent growth

Fig. 5.4 AFM image .5  5 m/ of a characteristic a-plane GaN surface

of threading dislocations with the film surface. Additionally, 1-nm-high lines with a density of 7  104 cm1 are apparent in the image. In contrast to the long-range surface features, these steps were oriented roughly perpendicular to the GaN c-axis regardless of HVPE gas flow conditions. These lines have been correlated with the presence of {0001} basal plane stacking faults in the films. Microstructural characteristics of the planar a-plane GaN films were evaluated by X-ray diffraction (XRD) and plan-view transmission electron microscopy (TEM). XRD was performed using Cu K’ radiation in a four-circle X-ray diffractometer

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Fig. 5.5 (a) On-axis XRD ! rocking curves for a planar a-plane GaN film. (b) Off-axis skewgeometry rocking curves

operating in receiving slit mode. !  2 scans of the a-GaN films exhibited peaks N N N reflections, that were indexed solely to the r-plane sapphire 1102; 2204, and 3306 N GaN reflection. No GaN 0002 reflection was observed, demonas well as the 1120 strating that within the detection limits of this technique the films were uniformly a-plane-oriented. To measure crystalline mosaic, ! rocking curves, shown in Fig. 5.5a, were N reflection for geometries in which the GaN measured on the on-axis GaN 1120 N Œ1100 and [0001] directions were in a coplanar configuration. Typical full width N reflection in these geometries at half-maximum (FWHM) values for the 1120 were generally comparable at 1,040–1,250 arcsec, though a FWHM as small as 660 arcsec has been observed. The on-axis peak widths were comparable to those observed for planar MOCVD-grown a-plane GaN films [19]. However, in contrast to MOCVD-grown a-plane GaN films [10], the peak width of the on-axis reflection measured with the c-axis within the scattering plane . D 90ı / was comparable to the peak width with the c-axis normal to the scattering plane, whereas a substantial asymmetry is commonly observed in MOCVD-grown a-plane films. The asymmetric mosaic of the MOCVD-grown films was correlated with growth condition variation; no such correlation has been identified for the HVPE-grown a-plane films. XRD rocking curves were measured for several off-axis reflections, as shown N reflections were measured by tilting the samples in Fig. 5.5b. The off-axis 1010 ı 30 relative to the scattering plane (skew geometry), yielding an FWHM ranging from 2,900 to 3,700 arcsec from sample to sample. Typical FWHM for the 33ı offN N ranged from 1,150 to 1,950 arcsec and 1,330 to axis 2021 and 54ı off-axis 1012 2,660 arcsec, respectively. Minimal variation in the off-axis peak widths with the in-plane rotation of the crystal was observed. Figure 5.6 shows plan-view transmission electron micrographs of a thick aplane GaN film. Figure 5.6a was imaged under the g D 0002 diffraction condition, revealing threading dislocations having a Burgers vector component parallel

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Fig. 5.6 Plan-view TEM images as observed with (a) g D 0002 to reveal dislocations having a nonzero [0001] component, and (b) g D 11N 00 to observe stacking faults parallel to the (0001) plane. After [12]

to the GaN [0001] direction. Thus, these are edge-component dislocations. The ccomponent dislocation density ranged from 9  109 cm2 to 2  1010 cm2 in these N diffraction consamples. The TEM image in Fig. 5.6b, taken under the g D 1100 5 1 dition, shows a stacking fault density (SFD) of 4  10 cm . These basal-plane N stacking faults are likely related to the presence of exposed nitrogen-face .0001/ surfaces during the early stages of growth, prior to island–island coalescence. AddiN tional imaging with varying sample tilt in the g D 1100 diffraction condition 9 2 N revealed 7 10 cm partial dislocations having Burgers vectors b D 13 <1010>, which bound stacking faults. Subsequent TEM studies of a-plane GaN films grown on a-plane SiC substrates showed similar results, with I1 and I2 stacking faults predominating [20].

5.2.2 Planar m-Plane GaN Films The growth of planar m-plane GaN thin films predated their a-plane counterparts by several years [9]. Waltereit et al.’s MBE-grown films first exhibited sufficiently low surface roughness for the growth of quantum wells and other fine heterostructures. While significant progress has been made in the fabrication of a-plane GaN quantum structures and devices by vapor phase growth techniques, m-plane GaN film growth and devices has since developed at a slower pace. This can be largely attributed to two key issues: the lack of suitable substrates and difficulties in growing high-quality, low-defect density planar m-plane GaN films. In this section, the morphological and structural characteristics of planar HVPE m-plane GaN films grown N on (100) -LiAlO2 and .1100/ m-plane 6H-SiC are discussed. We note that recently it was reported that m-plane GaN thin films can be grown by MOCVD on m-plane sapphire substrates using optimized buffer layers [21,22]. Until this was discovered, m-plane sapphire was a substrate that only yielded semipolar GaN films. If a similar approach can be utilized in HVPE m-plane growth, the development of thick films and eventually freestanding wafers could be accelerated.

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The growth of high-quality m-plane GaN films has historically been hindered by the lack of suitable substrates offering a combination of adequate lattice matching, chemical stability, and thermal stability. Lithium aluminate . -LiAlO2 /, though satisfies the first conditions, is otherwise a poor substrate for use in GaN HVPE growth conditions. It is attacked by hydrogen gas and hydrogen chloride at typical GaN HVPE growth temperatures, releasing LiC and O2 into the growth environment. LiAlO2 also exhibits poor mechanical stability, frequently fracturing upon cooldown from GaN growth temperatures. Despite these drawbacks, we and other groups have demonstrated thick, planar m-plane GaN growth on LiAlO2 substrates. Growth rates ranged from 30 to 240 m=h at substrate temperatures of 860–890 ıC and V/III ratios of 10–60, yielding freestanding films up to 250 m thickness. Figure 5.7 shows a representative 5  5 m AFM image of a planar m-plane GaN surface. This striated or “slate” morphology is dominated by ridges oriented along N the GaN <1120> direction. The linear density of these ridges along the <0001> direction averaged 105 cm1 . This morphology is also commonly observed for MBE-grown m-plane GaN films [23]. The surface exhibited an RMS roughness of 0.894 nm, which was smoother than is typically observed for MBE-grown films. The microstructural characteristics of the planar m-plane GaN films were evaluated using XRD and TEM. XRD 2 ! scans of m-plane GaN films grown on (100) N N LiAlO2 substrates revealed three peaks, which were indexed to the GaN 1100; 2200, N and 3300 reflections, demonstrating that the films were uniformly m-plane-oriented within the sensitivity of the measurement. XRD rocking curves were also measured on such films. In contrast to the a-plane GaN films described above, these m-plane GaN films exhibited in-plane asymmetry in their on-axis rocking curve peak widths. This asymmetry was, however, consistent with a-plane GaN films grown by MOCVD [10]. The FWHM of the on-axis reflection of the m-plane GaN films when the c axis was normal to the scattering plane . D 0ı / was 920 arcsec, whereas the FWHM when the GaN c axis was within the scattering plane . D 90ı /

Fig. 5.7 Atomic force micrograph of a thick m-plane GaN film grown by HVPE

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Fig. 5.8 (a) Plan-view and (b) cross-sectional transmission electron micrographs of an m-plane GaN thick film

was roughly 6,300 arcsec. The asymmetry in the on-axis reflection FWHMs is a consequence of asymmetrical tilt mosaic in the film. The peak widths of the offN N and 1012 N reflections, all measured in a skew geometry, were 880, axis 2021; 1120, 3,250, and 1,750 arcsec, respectively. Figure 5.8 shows representative plan-view and cross-sectional TEM images of an m-plane GaN film [24]. Figure 5.8a was imaged under the g D 0002 diffraction condition, revealing threading dislocations having a Burgers vector component parallel to the GaN [0001] direction. The threading dislocation density (TDD) observed in this image was 4109 cm2 , which is roughly an order of magnitude greater than was previously reported for HVPE-grown c-plane GaN films of comparable thickness [25]. The cross-sectional TEM image shown in Fig. 5.8b was imaged under N diffraction condition, thus revealing basal plane stacking faults. The the g D 1100 SFD was 1  105 cm1 , again consistent with previous reports for both m-plane and a-plane GaN thin films [12, 23]. This SFD was comparable to the ridge density observed in the AFM images in Fig. 5.7. Since the intersection of the basal plane N and the m-plane is a line along the <1120> direction, the close correlation in fault and ridge linear densities suggests that the surface ridges observed in the “slate” morphology described above are caused by displacement of the m-plane surface by basal plane stacking faults.

5.2.3 Planar Semipolar GaN Films Nonpolar GaN planes offer the most direct means of eliminating the built-in polarization fields that hinder carrier recombination in the quantum wells at the heart of GaN-based optoelectronic devices. So-called semipolar GaN orientations are also attractive in that they may minimize or also remove these fields via wellchosen combinations of crystallographic orientation and quantum well composition

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and thickness. This is possible by virtue of the fact that net polarization field strength does not vary monotonically with orientation from the c-plane to the nonpolar planes, and in fact it is strongly dependent on composition and strain state [26]. Thus “tailoring” of quantum well structure for each semipolar orientation (or vice versa) for minimization of net polarization is feasible. In addition, the range of growth parameters for “optimized” device layer growth may be favorable for semipolar films compared to their nonpolar counterparts; this of course depends on the characteristics (e.g., emission wavelength) being optimized. Semipolar planes have often been observed within stable “V defects” (pits) in c-plane GaN [27] and as the inclined sidewalls on c-plane HVPE and MOCVD LEO GaN stripes [28], suggesting their stability under a variety of conditions. HVPE of smooth, large-area epitaxial semipolar GaN films largely developed out of the search for favorable nonpolar GaN substrates. It was found that m-plane sapphire and MgAl2 O4 (spinel) yielded smooth semipolar films of a few different orientations, as described below. Both substrate materials are fortunately stable in typical GaN HVPE growth chemistry and temperature, and are commercially available in reasonable sizes and prices. HVPE growth on two orientations of single-crystal spinel was pursued: the (100) and (110) planes. Initial HVPE studies on (100) spinel substrates using conditions N GaN (i.e., the quite similar to a-plane GaN described above resulted in f101N 1g “pyramidal” plane), with a growth rate of 60 m=h. Presumably due to the fourfold symmetry of the (100) spinel plane, two domains of 90ı relative orientation formed, as evident in Fig. 5.9 and confirmed by XRD. It was found that a substrate miscut of 0:5–1ı toward the Œ011spinel direction was necessary to yield a singleN GaN [30], with an in-plane relationship of Œ101N 2 N GaN ==Œ011spinel and domain f101N 1g N N Œ1210GaN==Œ011spinel . Note that the GaN planar indices here convey “N-polarity” (i.e., the [0001] direction, which was confirmed by convergent-beam TEM. Interestingly, XRD measurements indicated that these films were tilted by 5:6ı relative to

Spinel [0 00

[011]

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[1

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Fig. 5.9 Left: Two rotational domains visible in a f101N 1N g GaN film grown on (100) spinel of nominally zero miscut. Right: miscut toward the Œ011spinel direction yielding smooth, single-domain GaN (After [29])

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N angle, which is likely the result of strain accommodation. XRD the nominal 1011 N peaks in two in-plane rocking curve scans were measured for the on-axis 1011 azimuths. The FWHM toward the Œ0002GaN direction was 0:7ı , while towards the N Œ1210 direction it was 1:3ı . This mosaic anisotropy is similar to that observed for some a-plane and m-plane films. N films revealed striated surfaces similar to those on m-plane GaN AFM of f101N 1g films, which has been attributed to stacking fault terminations [31]. As shown in N films as observed in TEM [29] were Fig. 5.10, the microstructures of the f101N 1g found to also contain a similar dislocation density .2  109 cm2 / and basal plane SFD .2  105 cm1 / as nonpolar GaN films grown by HVPE. The dislocations, N however, lay within the basal plane and had line directions in <1010>, such that they were inclined with respect to the growth direction.

Fig. 5.10 TEM micrographs of f101N 1N g GaN films, with diffraction conditions (noted) highlighting (a,b) threading dislocations and (c) stacking faults. Dislocations are depicted schematically in (d). After [29]

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Spinel substrates of (110) orientation were also utilized for semipolar GaN N GaN films resulted [29]. These substrates growth, and in this case planar f101N 3g have twofold symmetry and could be expected to encourage “twinning” in the overN lying GaN films (analogous to the rotational domains described above for f101N 1g films). This was indeed occasionally observed on nominally on-axis .˙0:2ı / substrates, yielding 10% of a twinned phase by volume, as measured by XRD. Little or no GaN twinning was seen for substrate surface miscuts greater than 0:1ı , so similar domain suppression could be achieved as had been obtained on miscut (100) spinel. The in-plane epitaxial relationship between the GaN and spinel as established N 32 N GaN ==Œ001spinel and Œ1210 N GaN ==Œ110 N spinel . by XRD was Œ30 N N AFM of these f1013g films again showed a striated surface morphology due to the presence of stacking faults, but the RMS roughness (3.5 nm) was slightly lower N GaN films. Likewise, the XRD rocking curve widths for than comparable f101N 1g N N the f1013g films were lower: the on-axis 101N 3N peak was 0:21ı toward the GaN N Œ1210 and 0:25ı toward the [0001] direction, while the off-axis 0002 peak was 0:24ı (in skew symmetric geometry). Despite the lower mosaic that these measurements implied, TEM micrographs (not shown) showed very similar TDD and SFD values N films above. This origin of this discrepancy is not clear, but could be as the f101N 1g due to differences in bowing-related broadening of XRD rocking curves between the two types of films. Although spinel proved to be robust, high-quality substrate for semipolar GaN growth, sapphire substrates were also investigated because of their superior quality, large area, and low cost. The sapphire m-plane, initially considered for m-plane GaN growth (which, as mentioned above, was observed later in MOCVD), actually yielded two semipolar HVPE GaN orientations, depending on pregrowth treatment [29]. A “high-temperature” nitridation step in NH3 (10 min prior to growth) N GaN films, whereas a “continuous” nitridation (ramp was used to obtain f101N 3g N in NH3 from room temperature up to growth temperature) resulted in f1122g films. Although quite reproducible, the exact mechanism for this dependence is not understood; it is likely due to the formation of an (oxy)nitride interlayer on the substrate whose composition, thickness, or other properties depend on temperature and duration of nitridation. N GaN films that resulted after a high-T nitridation had a surface morThe f101N 3g phology somewhat analogous to the films on spinel discussed previously, in that striations due to the presence of stacking faults were visible. However, in Fig. 5.11 shows that they also exhibited a marked “arrowhead” morphology (pointed toward N 32 N direction) readily visible in an optical microscope. TEM revealed that the Œ30 the microstructure of these films was otherwise quite similar to their counterparts grown on spinel, with a slightly lower TDD of 9  108 cm2 and a comparable SFD N of 2  105 cm1 . There were some TDs with line directions other than <1010>, but these were few. It should also be noted that minor “twinning” of the GaN was occasionally observed, likely due to the symmetry of the underlying sapphire substrate. Substrate miscut was not systematically measured, but it is likely that a consistent N would be sufficient to suppress it. miscut of 0:5–1ı toward the sapphire Œ1210

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Fig. 5.11 (a) Nomarski contrast optical micrograph and (b) AFM micrograph of a f101N 3N g GaN film on m-plane sapphire. After [29]

Fig. 5.12 Nomarski contrast optical micrographs of a f112N 2g GaN film on m-plane sapphire. After [31]

It was found that unlike some of the other semipolar film orientations disN GaN films on m-plane sapphire consistently grew in a cussed above, the f1122g single rotational domain without the need for substrate miscut. The in-plane epiN GaN ==Œ0001sapphire and taxial relationship as revealed by XRD and TEM was Œ1N 123 N N N Œ1100 ==Œ 12 10 . GaN sapphire TEM measurements indicated an elevated TDD of 2  1010 cm2 (with TDs of the same line direction as above) and an SFD of 2  105 cm1 comparable to other semipolar films. Interestingly, TEM also revealed the presence of inversion domains, i.e., the presence of both “Ga-polar” and “N-polar” regions. These inversion domains were later found to be directly correlated with surface morphology in the form of teardrop-shaped hillocks plainly visible in an optical microscope, as evident in Fig. 5.12. Their density varied significantly over a 2-in. wafer, for reasons that remain unclear. However, there were indications that microparticulate from

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the gas phase may have been at least partially responsible, due to the tendency of hillocks to form when the growth chamber was in a “dirtier” state than normal.

5.3 Lateral Epitaxial Overgrowth of Nonpolar, Semipolar GaN 5.3.1 LEO of a-Plane GaN While the surfaces of the planar a-plane GaN films described above were smooth, TEM and XRD measurements indicated that their microstructural quality needed improvement to be considered useful for optoelectronic device regrowth and fabrication. Threading dislocations in c-plane GaN have been found to act as nonradiative recombination centers, carrier scattering centers, and possibly fast diffusion pathways, limiting device efficiency, speed, and lifetime. Based on cathodoluminescence measurements, the same can be said for threading dislocations in nonpolar and semipolar films. At present, less is known about the influence of stacking faults on device performance. In order to improve the microstructural quality of a-plane GaN films, we developed HVPE-based lateral epitaxial overgrowth (LEO) for nonpolar GaN films [32]. LEO a-plane GaN films were grown either directly on r-plane Al2 O3 substrates or ona-plane GaN template layers first grown by MOCVD or HVPE. The patterned masks for the LEO process were prepared by applying conventional photolithographic processing and wet etching to 130 nm-thick SiO2 layers deposited by plasma-enhanced chemical vapor deposition. A variety of mask designs were investigated, including nonparallel stripes arranged in a “wagon-wheel” pattern, arrays of circular apertures, parallel stripes oriented along the <0001> direction, paralN lel stripes oriented along the <1100> direction, and parallel stripes oriented along higher index directions. LEO regrowth was carried out under similar growth conditions to those used for the a-plane GaN films described above. It was found that N masks consisting of parallel stripes oriented along the <1100> direction exhibited reproducibly vertical stripe sidewalls, and thus were most effective at lowering dislocation and stacking fault densities in the overgrown regions (“wings”). N Figure 5.13a shows a schematic representation of the <1100> stripe geometry that was used for the samples discussed in this section. Stripes oriented along this direction exhibited vertical {0001} sidewalls, one of which was Ga-terminated and the other N-terminated. Interrupted growths have shown that the (0001) Ga-face N N-face wing. This ratio wing advances roughly six times as rapidly as the .0001/ N indicates that the relative growth rate of the .0001/ wing is measurably greater than in MOCVD LEO of a-plane GaN, in which the ratio of Ga- to N-face growth is roughly 10 [33]. One benefit of the large difference in lateral growth rates between the {0001} faces is that the coalescence front was offset toward the N-face side of the window region, yielding broad wing regions uninterrupted by defective coalescence fronts. These stripes are shown in greater detail in the cross-sectional SEM

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Fig. 5.13 (a) Schematics of HVPE LEO a-plane growth using mask stripes/openings oriented along the GaN <11N 00> direction. (1). Cross-section of uncoalesced stripes. (2) Cross-section of coalesced stripes showing offset coalescence fronts relative to “windows.” (3) Plan view of coalesced stripes identifying the large low-defect wing region. (b) Inclined cross-sectional SEM micrograph of uncoalesced LEO stripes. (c) SEM cross-sectional image of coalesced a-plane LEO GaN film

Fig. 5.14 AFM micrograph .10  10/ m of coalesced <11N 00>-oriented a-plane LEO stripes

images of Fig. 5.13b, c. The inclined cross-section in Fig. 5.13b of nearly coalesced N stripes demonstrates the vertical {0001} sidewalls that are prevalent on <1100> oriented stripes throughout lateral growth and immediately preceding coalescence. Figure 5.13c shows a cross-section view of coalesced stripes. Only contrast variation at the film–template interface due to SEM charging effects allows the window and wing regions to be distinguished. AFM was performed to compare the surface quality in the window and wing regions of the a-plane LEO films. Figure 5.14b shows a 10  10 m AFM micrograph of two coalesced stripes. The window region appeared as the darker band of pitted material, with the coalescence front roughly 1 m to the left of the window. The Ga-face wing, apparent on the left side of the image, had superior surface quality and was nearly featureless. The wing regions exhibited average pit densities

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of less than 3  106 cm2 , compared to 109 cm2 in the window regions. The RMS roughness of the wing regions was less than 0.9 nm (11.9 nm peak-to-valley height), compared to 1.3 nm (20.8 nm peak-to-valley) in the window regions. The N faint lines parallel to the <1100> direction in the window region were not present in the Ga-face wing region. The structural quality of the a-plane LEO films was characterized by XRD N GaN reflection taken perpendicular to the and TEM. Rocking curves of the 1120 N <1100> LEO stripe direction of coalesced films were single-peaked, indicating a lack of wing tilt within the structural mosaic width of the coalesced films. In isolated cases, wing tilt of less than 0:2ı was observed for interrupted growths with uncoalesced stripes. Narrowing of both on- and off-axis reflections was observed in the LEO films compared to planar a-plane GaN films grown directly on r-plane N and 1010 N reflections were 750 and sapphire [12]. Typical FWHMs for the 1120 1,250 arcsec, respectively. The non-LEO a-plane GaN films described above exhibited comparable on-axis rocking curve widths for all orientations of the c-axis relative to the X-ray scattering plane. In contrast, the FWHM of LEO a-plane GaN films when the GaN c-axis was within the scattering plane . D 90ı / was up to five times larger than when the c-axis was perpendicular to the scattering plane . D 0/. This peak width variation was independent of stripe orientation and was rather a consequence of the mosaic of the crystal. Figure 5.15 shows plan-view and cross-sectional TEM images of a LEO film N and 0110, N respectively. In agreement with observaimaged with g vectors of 1100 tions from AFM, the window regions exhibited high threading dislocation .9  109 cm2 / and stacking fault .4  105 cm1 / densities. In contrast, the Ga-face wing region was essentially free of both dislocations and stacking faults, with densities below the images’ resolutions of 5  106 cm2 and 3  103 cm1 , respectively. The N-face wing region was also free of threading dislocations, though basal plane stacking faults and partial dislocations terminating the faults remained prevalent. N The above results for a-plane GaN LEO films grown with <1100>-oriented stripes have demonstrated that substantial reduction in morphological and structural defects in a-plane GaN may be readily achieved by LEO with HVPE. The reduction in TDD in the overgrown GaN is accompanied by a significant improvement in surface morphology compared to non-LEO planar a-plane GaN films. To achieve uniformly low defect density across a wafer, however, additional steps are required or another approach must be taken. A coalesced (planar) LEO film may be newly patterned with a second mask offset from the original stripe array, such that defects remaining in the window regions are blocked. While effective, this approach requires careful lithographic alignment, and also assumes a sufficiently smooth and planar coalesced surface for second-stage mask deposition and patterning. An alternative is a single-step regrowth technique such as Sidewall LEO (SLEO) [34], first demonstrated using MOCVD. Periodic trenches are etched in a nonpolar GaN film, and regrowth then emerges from trench sidewalls, as shown in Fig. 5.16. In this way, few if any dislocations or stacking faults emerge in the regrown material. There are still potential issues with wing tilt over the masked mesas, as well as pit formation

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Fig. 5.15 (a) Plan-view TEM of coalesced <11N 00> stripes. (b) Corresponding cross-section TEM image with (c) explanation of salient features

at coalescence fronts. However, this technique presents the most effective means of reducing the time and cost required to uniformly lower microstructural defect density in nonpolar GaN thick films.

5.3.2 LEO of m-Plane GaN While the results above demonstrated significant improvement in growth of planar, thick m-plane GaN films, the high microstructural defect densities inherent to these films needed to be reduced to improve device performance. To address this

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Fig. 5.16 Schematic of the SLEO process, in which a GaN film is selectively masked and etched to yield periodic trenches. Regrowth occurs in such a way that nearly all extended defects are blocked in one regrowth step. Courtesy of B. Imer

concern, LEO was developed for m-plane GaN via HVPE. Starting with MBEgrown m-plane GaN or m-plane AlN templates grown on m-plane SiC substrates (both 4H and 6H polytypes), fully coalesced LEO films were produced, achieving comparable improvements in morphology and structural quality to those attained for a-plane GaN LEO. Conventional photolithographic processing and wet etching were performed on SiO2 -coated m-plane GaN templates to create several mask patterns in a similar manner to the approach taken with the a-plane GaN LEO process. Ultimately, N periodic (parallel) stripe masks oriented along the <1120> direction yielded elimination of both dislocations and stacking faults, which will therefore be the focus of the discussion below. Vertical sidewalls were again typical for this stripe orientation, terminated by Cc and c-plane surfaces. In general, both the vertical and lateral growth rates observed were significantly faster than have been observed for c- and a-plane GaN LEO. Vertical growth rates ranged from 40 to 300 m=h at substrate temperatures ranging from 860 to 1;070 ıC. Fully coalesced LEO films were produced by this method, yielding specular surfaces over the full area of the wafers. The direct growth m-plane GaN films described above exhibited a commonly observed striated or “slate” surface morphology. This striated morphology persisted in the window regions of m-plane GaN LEO films, as is shown in Fig. 5.17a. The RMS roughness of this film’s surface was 0.78 nm, which is slightly smoother than was observed for the smoothest standard m-plane GaN films described above.

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Fig. 5.17 Atomic force micrographs of m-plane GaN LEO films taken from (a) the window region of a <112N 0>-oriented stripe, and (b) the Ga-terminated wing of a <112N 0>-oriented stripe

Figure 5.17b shows an AFM image taken from the Ga-terminated wing of an N m-plane LEO GaN stripe grown from <1120>-oriented stripes. The wing surface no longer exhibited the “slate” morphology, but rather exhibited a step-flow surface. The step terraces ranged from 100 to 150 nm across and showed no evidence of being aligned to any particular GaN crystallographic direction. The breadth and orientation of these terraces were likely dictated by the miscut of the underlying substrate. Terrace heights ranged from 0:8 to 1.4 nm, which corresponded to roughly N four to seven atomic layers in the <1100> direction. The roughness of this surface was reduced to 0.53 nm RMS over a 25 m2 area. This superior surface morphology resulted from the lack of extended defects in the overgrown material; indeed TEM imaging indicated that the threading dislocation and stacking fault densities were comparable to the Ga-terminated wings in a-plane GaN LEO films, at 5  106 cm2 and 3  103 cm1 , respectively. The structural quality of the m-plane GaN LEO films was additionally characN terized by XRD. m-plane GaN LEO films grown from <1120>-oriented stripes typically had narrower on-axis reflections compared to non-LEO films, on the order of 1;200 arcsec at D 0 and 3,600 arcsec at D 90ı . The asymmetry in the on-axis rocking curve width was independent of LEO stripe orientation; rather it was a consequence of the in-plane anisotropy of the crystalline mosaic. For comparison, the on-axis FWHM of the MBE GaN templates were typically 1,100 arcsec N reflections was at D 0ı and 7,000 arcsec at D 90ı . No splitting of the 1100 observed when the scattering plane normal was parallel to the LEO stripe direction for either stripe orientation. Thus no wing tilt was observed within the sensitivity of these X-ray measurements, which was limited by the mosaic of the underlying films N reflections, measured at both the and the curvature of the samples. The off-axis 1012 D 0ı and the D 180ı in a skew geometry, ranged from 1,800 to 2,400 arcsec, independent of or LEO stripe orientation.

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5.3.3 LEO of Semipolar GaN Like their a-plane and m-plane counterparts, all of the semipolar GaN films discussed above contained TDs and SFDs too high for demanding optoelectronic device growth and fabrication, i.e., those needed to demonstrate the potential of nonpolar and semipolar devices compared to their high-quality c-plane counterparts. Direct defect reduction via LEO was therefore needed for these semipolar films. As discussed above, TDs and SFs in semipolar GaN films are significantly inclined with respect to the growth direction, since the TD line directions often lay N along <1100> directions, while SFs were on {0001} planes. This complicates the strategy needed for effective LEO defect reduction compared to the methods used N films is given below for a-plane and m-plane films. The approach used for f1122g as illustrative example. N plane, two primary LEO stripe directions were considered, For the f1122g N namely the <0001> and <1100> projections, as shown in Fig. 5.18. The princiN pal advantage of the <1100> direction lies in the fact that defects are largely of one unique inclination, and will be periodically blocked in a fashion similar to their counterparts in a-plane or m-plane LEO films. In addition, the wings will have a “Ga-polar” and “N-polar” character (and therefore differing lateral growth rate), N also analogous to the case of a-plane LEO films with <1100> stripes. Under the same growth conditions used for planar film growth, it was observed N LEO films with <1100> N that the Ga-polar wings in these f1122g stripes indeed had a higher growth rate than the N-polar wings, as shown in Fig. 5.19. The voids N that formed between coalescing stripes were reproducibly terminated by .1N 122/ and N .0001/ planes.

a

(11–22)

[–1100]

[–1–123]

b (11–22)

[–1–123]

[1–100]

Fig. 5.18 Schematics of defect propagation in a GaN f112N 2g film with LEO stripes of two different orientations. After [29]

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Fig. 5.19 SEM cross-sections of coalesced f112N 2g LEO film with <1N 100> stripes of 5 m openings and 15 m mask widths. After [29]

XRD and TEM data demonstrated the effectiveness of defect reduction for this N and off-axis 0002 rocking curves were reduced LEO stripe orientation. On-axis 1122 from 0:39ı and 0:61ı to 0:24ı and 0:18ı , respectively. Plan-view TEM micrographs (not shown) demonstrated fairly uniform TD reduction to 5  106 cm2 over the wing regions, but SF reduction was observed only above areas where triangular voids had formed during coalescence. This was counter to the expectation that additional SFs would not form in “Ga polar” wings, as is the case in a-plane LEO films.

5.4 Conclusions and Future Development This chapter has demonstrated the variety of nonpolar and semipolar GaN thick films that may be grown by HVPE. In all cases, heteroepitaxy on sapphire and other substrates can result in smooth, uniform films, albeit with fairly high densities of threading dislocations and stacking faults. LEO is an effective, direct means of significantly reducing the propagation of such defects via blocking and redirection. In a manner analogous to c-plane LEO, the stripe direction, width, and period all play an important role in the morphology and defect reduction during overgrowth. However, unlike c-plane LEO, the lower in-plane symmetry of the nonpolar and semipolar planes creates both challenges and advantages: mask stripes must often be oriented in a unique direction, but the asymmetry in lateral growth rate that often results presents a means of obtaining higher uniformity in defect reduction. For semipolar films, the highly inclined angle at which threading dislocations propagate relative to the surface normal requires some additional, though not insurmountable, challenges in optimizing LEO stripe direction and width/period. Conventional LEO on either nonpolar or semipolar films is usually effective only at periodically reducing defect density in direct proportion to the mask stripe width and repeat period. “Double” LEO, in which a second mask stripe pattern is typically offset by a half-period from the first layer, can block most remaining dislocations and stacking faults, but this technique is time consuming and relies on first obtaining

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uniformly smooth coalescence of the first layer. Alternate approaches such as SLEO offer the possibility of single-step defect reduction, resulting in more rapid and uniform defect reduction. Future work should focus on such single-step approaches, as they can be quite effective and ultimately more economical for commercial production.

References 1. H.P. Maruska, L.J. Anderson, D.A. Stevenson, J. Electrochem. Soc. 121, 1202 (1974) 2. M. Sano, M. Aoki, Jpn. J. Appl. Phys. 15(10), 1943 (1976) 3. A. Shintani, S. Minagawa, J. Electrochem. Soc. 123(10), 1575 (1976) 4. R. Madar, D. Michel, et al., J. Cryst. Growth 40, 239 (1977) 5. T. Sasaki, S. Zembutsu, J. Appl. Phys. 61(7), 2533 (1987) 6. T. Lei, K.F. Ludwig Jr., T.D. Moustakas, J. Appl. Phys. 74(7), 4430 (1993) 7. W.A. Melton, J.I. Pankove, J. Cryst. Growth 178, 168 (1997) 8. R.J. Molnar, Hydride Vapor Phase Epitaxial Growth of III–V Nitrides, in Semiconductors and Semimetals, Vol. 57. (New York, Elsevier Science and Technology Books, 1999), pp. 1–31 9. P. Waltereit, O. Brandt, et al., Nature 406, 865 (2000) 10. M.D. Craven, S.H. Lim, et al., Appl. Phys. Lett. 81(3), 469 (2002) 11. P. Kung, C. J. Sun, et al., J. Appl. Phys. 75(9), 4515 (1994) 12. B. A. Haskell, F. Wu, et al., Appl. Phys. Lett. 83(8), 1554 (2003) 13. T. Paskova, P.P. Paskov, et al., Phys. Stat. Sol. (A) 183(1), 197 (2001) 14. A. Shintani, S. Minagawa, J. Electrochem. Soc. 123(10), 1575 (1976) 15. R. Molnar, Hydride Vapor Phase Epitaxial Growth of III-V Nitrides, in Semiconductors and Semimetals, vol. 57 (Gallium Nitride II, ed. by T.D. Moustakas and J.I. Pankove, San Diego: Academic Press, 1999). pp. 1–31 16. E.V. Etzkorn, D.R. Clarke, J. Appl. Phys. 82(2), 1025 (2001) 17. T. Böttcher, S. Einfeldt, et al., Appl. Phys. Lett. 78(14), 1976 (2001) 18. H. Marchand, J. Ibbetson et al., MRS Internet J. Nitride Semicond. Res. 3, 3 (1998) 19. M. D. Craven, S. H. Lim, et al., Appl. Phys. Lett. 81(3), 469 (2002) 20. D.N. Zakharov, Z. Lilienthal-Weber, et al., Mat. Res. Soc. Symp. Proc. 798, Y5.28.1–6 (2004) 21. R. Armitage, H. Hirayama, Appl. Phys. Lett. 92, 092121 (2008) 22. T. Wei, R. Duan, et al., Jpn. J. Appl. Phys. 47(5), 3346 (2008) 23. O. Brandt, Y. J. Sun, et al., Phys. Rev. B 69, 165326 (2004) 24. B. Haskell, A. Chakraborty, et al., J. Electr. Mater. 34(4) 357–360 (2005) 25. X. Xu, R.P. Vaudo, et al., J. Cryst. Growth 246, 223–229 (2002) 26. A.E. Romanov, T.J. Baker, et al., J. Appl. Phys. 100, 023522 (2006) 27. X.H. Wu, C.R. Elsass, et al., Appl. Phys. Lett. 72(6), 692 (1998) 28. H. Marchand, J.P. Ibbetson, et al., J. Cryst. Growth 195, 328–332 (1998) 29. T.J. Baker, Ph.D. Dissertation, U.C. Santa Barbara (2006) 30. T.J. Baker, B.A. Haskell, et al., Jpn. J. Appl. Phys. 44(29), L920–L922 (2005) 31. B.A. Haskell, Ph.D. Dissertation, U.C. Santa Barbara (2005) 32. B.A. Haskell, F. Wu, et al., Appl. Phys. Lett. 83(4), 644 (2003) 33. M.D. Craven, S.H. Lim, et al., Appl. Phys. Lett. 81(7), 1201 (2002) 34. B.M. Imer, F. Wu, S.P. DenBaars, J.S. Speck, Appl. Phys. Lett. 88, 061908 (2006)

Chapter 6

High Growth Rate MOVPE K. Matsumoto, H. Tokunaga, A. Ubukata, K. Ikenaga, Y. Fukuda, Y. Yano, T. Tabuchi, Y. Kitamura, S. Koseki, A. Yamaguchi, and K. Uematsu

Abstract In this chapter, growth mechanism of GaN and AlGaN by MOVPE is described in detail. The parasitic reaction that generates particulates in vapor phase is the most probable limiting factor of maximum growth rate of GaN and AlGaN. Both experimental and quantum chemical study of vapor phase reaction between organo-metals and ammonia is described. From the close insight of vapor phase reaction, high flow speed three layered gas injection has been developed. By employing this three-layer gas-injection, GaN was grown with growth rate as high as 28m/h at atmospheric pressure. As long as the present growth conditions are concerned, SIMS and XRD results suggested that the growth rate of 12m/h would be a practical limitation of good quality material in terms of electrical and structural properties.

6.1 Introduction Today, GaN with dislocation density of low 108 cm3 can be grown on sapphire substrates by atmospheric-pressure metal-organic vapor phase epitaxy (MOVPE) by using the conventional two-step growth technique [1]. High-growth-rate technology of GaN by MOVPE is a promising technology for a thick, low-dislocation-density template as an alternative to a native bulk GaN substrate since the device structure can be grown successively in the same reactor. Very recently, Tanaka et al. reported on an InGaN light-emitting diode (LED) on 100-m-thick GaN, which was grown by MOVPE with a growth rate of as high as 56 m=h [2]. Until recently, the growth rate of MOVPE GaN has been limited to a few micrometers per hour in large-scale production equipment at atmospheric pressure. Recent understanding of the obstacles for obtaining high growth rates of GaN by MOVPE can be summarized as follows. To put it simply, high growth rate of GaN is hindered by the generation of particulates in the vapor phase. Creighton observed nanometer-sized particles in an inverted stagnation-point-flow reactor by employing a high-flow-rate condition of organometals [3]. Nanometer-sized particulates were observed on top of a thermal boundary layer by in situ laser scattering. Dauelsberg 119

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et al. reported an experimental and simulation study of GaN MOVPE as a function of thermal gradient in vapor phase and flow speed [4]. They concluded that a large temperature gradient in vapor phase in a shower-head reactor can result in a condensation of nanometer-sized particles and their migration into a relatively cool part of the reactor, followed by removal of precursors from the growing region. They have also suggested that a high flow speed is critical to eliminate particulate generation in a horizontal reactor. Hirako and Ohkawa also reported ŒGaN4 molecules condensing at the entrance region of a horizontal reactor, where there is a large temperature gradient in the vapor phase [5]. High growth rate can be realized by considering above-mentioned problems. In this section, we will describe (1) growth characteristics of AlGaN and GaN by conventional MOVPE, (2) quantum chemical study of vapor-phase reaction, and (3) the result of high-growth-rate GaN by using a high-flow-speed reactor. By employing a high flow speed, we can eliminate the parasitic reaction that hinders a high growth rate.

6.2 Growth Characteristics of AlGaN and GaN by Conventional MOVPE In this section, growth of AlGaN and GaN is described by showing experimental results of a parasitic reaction in the vapor phase. The data are obtained by using horizontal flow reactors at atmospheric pressure. First we will describe AlGaN growth by using a 2-in. single-wafer reactor. The total flow rate was about 40 l/min. Figure 6.1 shows AlN growth rate at atmospheric pressure as a function of growth temperature .Tg / [6]. AlN growth rate sharply dropped at a Tg higher than 500 ı C. This shows that the parasitic reaction involving

Growth Rate of AIN [m m/h]

2.0

1.5

1.0

0.5

0.0

200

400 600 800 1000 Growth Temperature [°C]

1200

Fig. 6.1 AlN growth rate at atmospheric pressure shown as a function of growth temperature .Tg / [6]

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AlGaN Growth Rate [m m/h]

1.2 1.0 0.8 0.6 0.4 0.2 0.0

0

5 10 15 20 Input Flow Rate of TMA [m mole/min]

25

Fig. 6.2 AlGaN growth rate shown as a function of TMA input flow rate at a constant TMG flow rate [6]

trimethylaluminum (TMA) and NH3 requires some excitation energy to initiate the process. Since the upstream gas-phase temperature is far below the susceptor temperature, the onset temperature of the parasitic reaction of the TMA–NH3 system would be less than 500 ıC. Figure 6.2 shows the AlGaN growth rate as a function of TMA input flow rate at a constant trimethylgallium (TMG) flow rate [6]. AlGaN growth rate decreased as TMA input flow rate increased. This implies that intermolecular collision has an important role in the enhancement of the parasitic reaction. The above results can be well explained if we assume that nanoparticles are generated in vapor phase, as proposed by Creighton et al. [3]. The above result can be attributed to the very fast reaction between TMA and NH3 to form oligomers, which easily takes place in the upstream region of the horizontal reactor. In contrast to the TMA and NH3 case, the TMG and NH3 reaction is a reversible reaction: that is, the TMG–NH3 adduct easily breaks to individual molecules by thermal excitation. As a result, adduct chain reaction rarely occurs for the case of TMG and NH3 . However, at high growth rate condition of more than a few 10 m=h, we can see the effect of the parasitic reaction for GaN growth. Actually, Creighton observed nanoparticles by employing the growth condition of ten times dense TMG concentration in vapor phase than that of usual GaN growth rate of a few m/h [3]. They attributed the nanoparticle formation to a radical reaction pathway rather than an adduct reaction. A large temperature gradient above the susceptor also enhances the condensation of GaN molecules at the top surface of the thermal boundary layer [4]. Another important issue is the flow dynamics in the reactor. If there is recirculation flow in the reactor, we will see the growth rate saturation of GaN at the growth rate of a few micrometers per hour. Figure 6.3 shows the growth rate distribution of GaN along the flow direction for various input TMG supply rates [6]. The growth reactor has a 180-mm-diameter platen that can handle 7  200 wafers at a time (TAIYO NIPPON SANSO SR-6000). The wafer susceptor is heated by a resistance heater and can be rotated about one axis. The TMG

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70μmol/min 210μmol/min 420μmol/min 700μmol/min

6 4

2 0

0

20

40

60

100 120 80 Position [mm]

140

160

180

Fig. 6.3 Growth rate distribution of GaN along the flow direction shown for a various input TMG supply rates [6]

input flow rate was varied from 70 to 700 mol=min. The reactor pressure was set at 760 Torr. The typical carrier gas flow rate used was 70 l/min. The NH3 flow rate was kept constant at 15 l/min. A H2 =N2 mixture of 50:50 was used as the carrier gas. The growth temperature was 1;150 ı C. Standard low-temperature GaN buffer technology was used. In this case, organometallics and NH3 were mixed together at the inlet of the reactor. Judging from the wall deposit, there was unexpected recirculation in the reactor. By flow dynamic simulation, it was also confirmed that the reason for the recirculation was the improper design of the upstream flow channel in the flow expansion part. When the TMG input flow rate was below 400 mol=min, the GaN growth rate variation along the flow direction was almost linear. The average growth rate with susceptor rotation was approximately 2 m=h at this TMG flow rate. Note that the growth rate variation of GaN with the TMG supply rate of 700 mol=min is a nonlinear function of the wafer position. The GaN growth rate downstream was almost zero for an input TMG of 700 mol=min, which strongly suggests nanoparticle generation by the CVD mechanism [3]. By improving the flow so that a perfect laminar flow was realized, we have succeeded in achieving a growth rate of more than 7 m=h at the TMG supply rate of 700 mol=min [6]. In Fig. 6.4, GaN growth rate along the flow direction is shown for the reactor with the improved flow design [6]. GaN growth rate linearly decreased downstream. The concept of the reactor is described in the following section on high-flow-speed reactor. Lundin et al. also reported that the growth rate of GaN saturated at about 4 m=h at 800 mbar at a fixed residence time in the growth region [7]. The detailed reaction pathway will be discussed in the next section.

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12 GaN

P=760torr

Growth rate [mm/hr]

10 8 6 4 2 0

0

20

40

60

80

100

120

140

160

180

Position [mm]

Fig. 6.4 GaN growth rate along the flow direction shown for the improved flow design reactor [6]

6.3 Quantum Chemical Study of Vapor-Phase Reaction In this section, we will describe an elementary process which is supposed to be responsible for a parasitic reaction. We will discuss the unimolecular elimination of methane from TMG under excess NH3 condition [8]. First, the coordination interaction between organometallics and NH3 is shown. The M.CH3 /3 molecule (1; each one, respectively, denotes 1a, 1g, and 1i for M D Al, Ga, and In, and so forth) forms a very stable complex with ammonia, .CH3 /3 M W NH3 (2), due to the M–N coordinate bond through the following reaction: M.CH3 /3 .1/ C NH3 ! .CH3 /3 M W NH3 .2/

(6.1)

The stabilization energy of complex formation is the largest in the TMA:NH3 system. It is 23:17 kcal=mol for aluminum, 18:92 kcal=mol for Ga, and 18:22 kcal=mol for indium. They form a coordination bond without an energy barrier. Under excess NH3 condition, TMG forms a stable complex due to its coordination bond with two ammonia molecules, H3 N W .CH3 /3 Ga W NH3 (3g), without a potential energy barrier as follows: Ga.CH3 /3 .1g/ C NH3 ! .CH3 /3 Ga W NH3 .2g/ .CH3 /3 Ga W NH3 .2g/ C NH3 ! H3 N W Ga.CH3 /3 W NH3 .3g/

(6.2) (6.3)

The elimination of a methane molecule can occur in the presence of excess ammonia, by the intramolecular reaction of the complex 3g or the intermolecular collision

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N 2.042

1.290 H

M

1.448 2.540

2.374 C

Fig. 6.5 The geometry of TS1 [8] TS1g

TS4g

14.83

13.04

1g 0.00 + NH3

– CH4

2g –18.92

+ NH3 7g –22.02

1g: Ga(CH3)3 2g: (CH3)3Ga NH3 3g: Ga(CH3)2NH2 7g: H3N Ga(CH3)3 NH3 8g: H3N Ga(CH3)2NH2

3g –22.62

·

· ·

– CH4

+ NH3

·

8g –41.03

Fig. 6.6 Energy diagram of the reaction described in (6.4) [8]

between the complex 2g and an ammonia molecule. H3 N W Ga.CH3 /3 W NH3 .3g/: orŒ.CH3 /3 Ga W NH3 .2g/ C NH3  ! transition state1ŒTS1 ! H3 N W Ga.CH3 /2 W NH2 .4g/ C CH4

(6.4)

In Fig. 6.5, the geometry of TS1 is shown [8]. Here, it is noteworthy that Bergmann et al. observed a fragment of the molecule 3g by Quadrupole Mass Spectrometer spectroscopy [9]. In Fig. 6.6, the energy diagram of the reaction described in (6.4) is shown [8]. It is found that the potential energy barrier is reduced through TS1 in the presence of excess ammonia for each M.CH3 /3 C 2NH3 system. In particular, the potential energy barrier for TMA is reduced to 4.67 kcal/mol (Table 6.1) [8]. Accordingly, it is considered that the gas-phase reaction of TMA proceeds rapidly in the presence

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Table 6.1 Potential energy for different NH3 association numbers and excitation states for M.CH3 /3 and NH3 [8] Ground state (a) Ground state (b) TS of (a) TS of (b) Oligomer .CH3 /3 M W NH3 H3 N W .CH3 /3 M W NH3 with one with two H3 N W M.CH3 /2 NH3 NH3 molecules molecules Al.CH3 /3 23.17 28.86 7:37 4:67 49.86 18.92 22.02 14:83 13:04 41.03 Ga.CH3 /3 In.CH3 /3 18.22 25.83 13:96 7:17 35.26 TS energy decreases with an excess NH3 . All values are in kcal/mol

of excess ammonia. Since the stabilization energy of the coordination by the second ammonia molecule (reaction (6.3)) is relatively small, the mainstream of reaction (6.4) would be the intermolecular reaction between the complex 2 and an ammonia molecule together with the effective energy transfer due to the collision. In Table 6.1, the enthalpy of the energy diagram is summarized. Watwe et al. have studied the gas-phase chemistry of TMG and NH3 by quantum chemical calculation and have shown that the TMGa:.CH3 /2 GaNH2 W NH3 species has the lowest standard free energy activation barrier for methane elimination [10]. If we assume that .CH3 /2 GaNH2 is back-diffusing from downstream, this catalytic reaction will easily occur and proceed to oligomer formation. Once an oligomer starts to grow, nanoparticles would result by an iterative process. In the downstream region, most of the precursors would be consumed in the nanoparticle growth by CVD on the particle [3]. Since the decomposition velocity of Œ.CH3 /3 Ga W NH2 x is determined by ambient gas-phase temperature, the rate of their decomposition is independent of their concentration. Then, the growth rate of Œ.CH3 /3 Ga W NH2 x is proportional to its collision frequency with TMG. Therefore, at some critical TMG input, Œ.CH3 /3 Ga W NH2 x will start to grow. In the particular case shown in Fig. 6.3, the growth rate of GaN showed saturation at approximately 2:5 m=h. We must confine molecules such as Œ.CH3 /3 Ga W NH2 x in the very vicinity of the crystalgrowth surface to suppress nanoparticle generation. When we attempted to grow AlGaN with the same reactor at atmospheric pressure, we failed because of a very strong parasitic reaction [11]. Hirako and Ohkawa calculated elemental processes and spatial distribution of species in each step by computational fluid dynamics. Figure 6.7 illustrates major steps for growing GaN [12]. They suggested that a major reaction pathway is (1) TMG W NH3 adduct formation, (2) methane elimination from TMG W NH3 adduct, followed by another methane elimination, and (3) GaN molecule formation. As a side reaction, higher order molecules are formed. They have also calculated the effect of growth pressure on reaction by-products (Fig. 6.8) [13]. At atmospheric pressure, the major reaction product in the vicinity of the growth surface is the GaN molecule. However, at reduced growth pressure, the major reaction product

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DMGa

TMGa:NH3

H3N:TMGa:NH3

DMGaNH2

H3N:DMGaNH2

MMGaNH2 MMGa:NH3

[DMGaNH2]2 MMGa

MMGaNH

[DMGaNH2]3

[MMGaNH]2

Ga-N

Ga

N2

MMGa(NH2)2 DMGaMMGa(NH2)3

[MMGaNH]3-6 N2H4

[Ga-N]3-6

[Ga-N]2

GaNH2

H3N:MMGa(NH2)2

N

NH

NH2

NH3

GaN Layer

Fig. 6.7 Major steps for growing GaN. The major reaction pathway is (1) TMG:NH3 adduct formation, (2) methane elimination from TMG:NH3 adduct, followed by another methane elimination, and (3) GaN molecule formation [12]. Copyright APEX/JJAP. Reproduced with permission

10–6

Concentration (mol/L)

Ga-N

10–9

DMG:NH2 (DMG:NH2)3

10–12 (DMG:NH2)2

10–15

NH2 10

100

Pressure (kPa) Fig. 6.8 The effect of growth pressure on reaction by-products. At atmospheric pressure, the major reaction product in the vicinity of growth surface is the GaN molecule. However, at reduced growth pressure, the major reaction product is .CH3 /2 Ga W NH2 [13]. Copyright Wiley-VCH Verlag GmbH & Co. KGaA. Reproduced with permission

is .CH3 /2 Ga W NH2 . This might be responsible for higher carbon incorporation for materials grown at low pressures [14]. In Fig. 6.9, the spatial distribution of Œ.CH3 /GaNH4 and ŒGaN4 is shown [12]. We can see condensation of ŒGaN4 at the entrance of the growth region, where there is a large temperature gradient.

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a [MMGaNH]4 700 K

1300 K

Susceptor

0

8×10–9 mol/L

0

5.5×10–8 mol/L

b [Ga-N]4 700 K

1300 K

Susceptor

Fig. 6.9 Spatial distribution of Œ.CH3 /GaNH4 and ŒGaN4 . The condensation of ŒGaN4 at the entrance of the growth region is seen, where there is a large temperature gradient [12]. Copyright APEX/JJAP. Reproduced with permission

6.4 Result of High-Growth-Rate GaN by Using a High-Flow-speed Reactor Figure 6.9 shows three-layer gas-injection reactor, which has an additional gas injection part on top of the conventional two-injection nozzle [15]. As shown in Fig. 6.10, organometallics and NH3 are separately injected and mixed together by molecular diffusion. Under typical flow conditions, the spacing between the gas separators and their position are designed so that organometallics would diffuse laterally to the flow direction and reach the substrate surface just at the inlet of the growth region. Therefore, the mainstream of adduct .CH3 /3 M W NH3 would be isolated by a sheath of gas stream from the heated reactor wall until it reaches the growth region. This sheath effect is useful to avoid undesirable upstream parasitic reaction [6]. This type of gas injection is also useful for reducing the entrance effect in a high-flow-speed reactor. By varying flow rate ratio between the carrier gases, the growth rate near the leading edge of the deposition zone is controlled. When the carrier gas flow of

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Fig. 6.10 Three-layer gas-injection reactor which has an additional gas injection part on top of a conventional two-injection nozzle [15]

the bottom NH3 gas is larger than those of the organometallics, the growth starts at a downstream side of the flow. This is because organometallics must travel through the thicker underlying flowing layer, since the large flow rate results in an expansion of the flow. On the contrary, when the bottom gas flow rate is smaller, the growth starts upstream of the flow. The same effect can be realized by increasing the top NH3 carrier flow rate by compressing the underlying two-channel gas flows [16]. This mechanism is especially useful in case of a centrifugal gas injection planetary reactor with a gas separator configuration. Gas injection flow speed in this type of reactor is extremely high for a large high-flow-speed reactor. As a result of the high flow speed, organometallics cannot reach the wafer surface within a reasonable distance from the gas injection nozzle without the suppression effect of the additional top flow of three-layer gas injection [17]. For high-growth-rate GaN, we have developed a production-scale multiwafer reactor (Taiyo Nippon Sanso Corp., SR23K) [18]. The reactor is capable of growing 10  2 in. or 8  3 in. wafers at a time. The reactor geometry is similar to the planetary reactor in Ref. [17]. We have employed a specifically designed three-flow gas injection with a flow speed of more than 1 m/s at the wafer center. GaN layers were grown at atmospheric pressure. NH3 and TMG were used as precursors. Hydrogen and nitrogen were used as the carrier gas. By using SR23K, GaN was grown at a growth rate of up to 28 m=h. In Fig. 6.11, the growth rate of GaN is shown as a function of normalized TMG supply rate [18, 19]. The maximum growth rate was limited only by the limitation of TMG supply. Surface morphology of all the samples was smooth regardless of the growth rate. Figure 6.12 shows ¨-scan X-ray diffraction full width at half-maximum (XRD FWHM) of (002) and (102) reflection as a function of the growth rate [18, 19]. In this experiment, the NH3 flow rate was fixed at 100 SLM, and the TMG supply rate was changed. The V/III ratio was 300–3,000 according to the growth rate. Each

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Normalized TMG supply rate Fig. 6.11 The growth rate of GaN as a function of normalized TMG supply rate [18] 600

XRD FWHM (arcsec)

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Growth Rate (m m/h) Fig. 6.12 ¨-scan XRD FWHM of (002) and (102) reflection as a function of growth rate [18]

high-temperature-grown GaN was directly deposited on a low-temperature buffer layer. The total thickness of the samples was fixed to about 4 m. The FWHM of symmetric (002) reflection was less than 300 arcsec. The FWHM of asymmetric (102) reflection was 350, 390, and 560 arcsec for the growth rate of 2.4, 12, and 28 m=h, respectively. It is notable that the XRD FWHM is not very much

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Table 6.2 Carbon concentration and sheet resistance of GaN shown as a function of growth rate. NH3 flow rate was kept constant at 100 SLM. Sheet resistance was measured by Lehighton [20] Growth rate .m=h/ 2.4 12 28 V/III ratio 3,000 600 300 2:5  1016 1:5  1017 6:0  1017 Carbon concentration .cm3 / Sheet resistance .=sq/ 1,200 4,500 16,000

affected up to the growth rate of 12 m=h, even though high-temperature GaN is directly grown on a low-temperature buffer layer. It is well established that optimum nucleation density and the resulting lateral growth is important to reduce dislocation density [20]. Trapezoidal facet formation is enhanced during initial growth stage by a relatively low V/III ratio [21]. Density of high-temperature-grown nuclei tends to decrease by low V/III ratio, followed by lateral growth and bending of dislocations [22]. Lateral to vertical growth rate ratio seems smaller in case of the highest growth rate of 28 m=h, which resulted in a broader line width of (102) reflection. Since the V/III ratio affects the electrical property of GaN, carbon concentration and sheet resistance were measured for these samples. Table 6.2 summarizes the carbon concentration and sheet resistance of GaN layers as a function of V/III ratio and the growth rate [19]. At the V/III ratio of 600 (growth rate of 12 m=h), the carbon concentration was 1:5  1017 cm3 . At the V/III ratio of 300 (growth rate of 28 m=h), the carbon concentration was 6  1017 cm3 . Probably due to the carbon contamination, the sheet resistance of GaN was 16; 000 =sq (nominal sheet resistance by Lehighton) for the highest growth rate sample with the V/III ratio of 300. However, it should be noted that the sheet resistance is also affected by the growth process of initial coalescence without regard to the residual carbon concentration. When the growth condition of early coalescence is adopted, the sheet resistance tends to be high. When the retarded coalescence condition is adopted, a current leak channel is often formed near the interface of epitaxial layer and the sapphire substrate. From these results, the usable growth rate range that allows us to dope the film would be lower than 10 m=h.

6.5 Discussion and Summary For comparison with hydride vapor phase epitaxy (HVPE), the reported growth rate of GaN as a function of V/III ratio is plotted in Fig. 6.13 [19]. It is noteworthy that a recently reported GaN growth rate by HVPE [23] lies on the same line as the present work for MOVPE, while old data for HVPE in the literature [24–27] lie below the line. Growth rate data of HVPE of Ref. [23] and the present MOVPE data lie on a common straight line. The growth rate difference of HVPE [23] and the present MOVPE can be explained mostly by the Ga source supply rate, by assuming that the rate-limiting process is diffusion-controlled and the order of NH3 flow rate is comparable. Because the growth rate of the surface-reaction-limited case is

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[23] 100

[27] [26]

MOVPE Tanaka[ 2 ]

[24] [25]

high flow speed MOVPE

HVPE

This work

10

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1.00E+00

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1.00E+02

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1.00E+04

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Fig. 6.13 Reported growth rate of GaN as a function of V/III ratio. The figure is modified from [19] with additional data from [2]

always lower than that of diffusion-limited growth condition, the different growth conditions in Refs [24–27] may partly explain the reason for the scattered data of HVPE in Fig. 6.13. However, because growth temperature in these references is very similar, in the range from 1;000 ıC to 1;050 ıC, growth-equipment-related reasons would be more probable. Incomplete reaction of HCl and metal Ga could be considered. The limitation of the maximum growth rate of GaN by MOVPE partly comes from the reaction between GaN molecules in vapor phase and the stability of resultant particles, since the GaN molecule is supposed to be a major by-product in the vapor phase for atmospheric pressure MOVPE [12]. Adduct reaction can also be responsible for the particulate generation. Adduct process would be important at an upstream region where the decomposition process does not proceed very much in vapor phase [28]. As described in the previous section, when there was flow turbulence at the upstream region, we have observed growth rate saturation. Growth rate saturation is also reported in a case with a long residence time in a reaction zone [7]. There would be some critical residence time for growth rate saturation. In case of a long residence time, small GaN clusters, formed due to any causes, can result in particulate growth under supersaturation of precursors. Since the maximum growth rate in the present report is not the true limit, we can expect a higher growth rate for GaN by increasing TMG supply. Actually, Tanaka et al. reported InGaN LED on 100-m-thick GaN, which was grown by MOVPE with a growth rate of as high as 56 m=h [2]. They have employed a low V/III ratio of 114 at the growth rate of 56 m=h. In Fig. 6.13, we can see that their data lie on the same line of the recent HVPE and MOVPE results. It would be interesting to see whether growth rate of GaN by high-flow-speed MOVPE can reach the criteria of HVPE. Provided that the particulate generation by adduct process expected at the upstream region is well suppressed by proper flow design and gas-phase temperature distribution, we may focus on the stability of GaN clusters in the growth region.

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As a supplementary discussion, it would be interesting to consider the role of HCl in vapor-phase reaction in HVPE. HCl is a reactant to produce GaCl in the source region in HVPE. This is usually almost fully consumed in the source region, but HCl is produced again through the reaction between and NH3 in the growth region. When GaN clusters are formed in vapor phase from some reason, concentration of GaN monomolecules around the cluster would be decreased by consumption through deposition on the cluster surface and the corresponding HCl concentration would be excessive. This would enhance the reverse reaction by HCl to decompose the GaN cluster into GaCl and NH3 . This process provides negative feedback to eliminate GaN cluster formation and growth. It is notable that a high concentration of NH3 can act to stabilize GaN clusters through the reverse reaction of GaN decomposition. Without an efficient etching reaction in vapor phase, the maximum obtainable growth rate of MOVPE GaN would be smaller than that of HVPE GaN. From the above consideration, cluster formation in the vapor phase should be strictly controlled in case of MOVPE. In order to eliminate the possibility of adduct reaction path, low would be recommended, because adduct decomposition is enhanced by excessive NH3 [8]. The actual limitation of the growth rate of MOVPE GaN is open to discussion at present.

References 1. K. Hoshino, N. Yanagita, M. Araki, K. Tadatomo, J. Cryst. Growth 298, 232 (2007) 2. Y. Tanaka, J. Ando, D. Iida, M. Iwaya, S. Kamiyama, H. Amano, I. Akasaki, Phys. Stat. Sol. (C) 5, 3073 (2008) 3. J.R. Creighton, G.T. Wang, W.G. Breiland, M.E. Coltrin, J. Cryst. Growth 261, 204 (2004) 4. M. Dauelsberg, C. Martin, H. Protzmann, A.R. Boyd, E.J. Thrush, J. Kappeler, M. Heuken, R.A. Talalaev, E.V. Yakovlev, A.V. Kondratyev, J. Cryst. Growth 298, 418 (2007) 5. A. Hirako, K. Kusakabe, K. Ohkawa, Jpn. J. Appl. Phys. 44(2), 874 (2005) 6. K. Matsumoto, A. Tachibana, J. Cryst. Growth 272, 360 (2004) 7. W.V. Lundin, E.E. Zavarin, D.S. Sizov, M.A. Sinitsin, A.F. Tsatsul’nikov, A.V. Kondratyev, E.V. Yakovlev, R.A. Talalaev, J. Cryst. Growth 287, 605 (2006) 8. K. Nakamura, O. Makino, A. Tachibana, K. Matsumoto, J. Organometallic Chem. 611, 514 (2000) 9. U. Bergmann, V. Reimer, B. Atakan, Phys. Chem. Chem. Phys. 1, 5593 (1999) 10. R.M. Watwe, J.A. Dumesic, T.F. Kuech, J. Cryst. Growth 221, 751 (2000) 11. H. Tokunaga, H. Tan, Y. Inaishi, T. Arai, A. Yamaguchi, J. Hidaka, J. Cryst. Growth 221, 616 (2000) 12. A. Hirako, K. Kusabake, K. Ohkawa, Jpn. J. Appl. Phys. 44(2), 874 (2005) 13. K. Kusakabe, A. Hirako, S. Tanaka, K. Ohkawa, Phys. Stat. Sol. (C) 1, 2569 (2004) 14. D.D. Koleske, A.E. Wickenden, R.L. Henry, M.E. Twigg, J. Cryst. Growth 242, 55 (2002) 15. K. Uchida, H. Tokunaga, Y. Inaishi, N. Akutsu, K. Matsumoto, T. Itoh, T. Egawa, T. Jimbo, M. Umeno, Mater. Res. Soc. Symp. Proc. 449, 129 (1997) 16. K. Matsumoto, T. Arai, H. Tokunaga, Vacuum 51(4), 699 (1998) 17. C. Martin, M. Dauelsberg, H. Protzmann, A.R. Boyd, E.J. Thrush, M. Heuken, R.A. Talalaev, E.V. Yakovlev, A.V. Kondratyev, J. Cryst. Growth 303, 318 (2007) 18. H. Tokunaga, Y. Fukuda, A. Ubukata, K. Ikenaga, Y. Inaishi, T. Orita, S. Hasaka, Y. Kitamura, A. Yamaguchi, S. Koseki, K. Uematsu, N. Tomita, N. Akutsu, K. Matsumoto, Phys. Stat. Sol. (C) 5(9), 3017 (2008)

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19. K. Matsumoto, H. Tokunaga, A. Ubukata, K. Ikenaga, Y. Fukuda, T. Tabuchi, Y. Kitamura, S. Koseki, A. Yamaguchi, K. Uematsu, J. Cryst. Growth 310, 3950 (2008) 20. P. Fini, X. Wu, E.J. Tarsa, Y. Golan, V. Srikant, S. Keller, S.P. DenBaars, J.S. Speck, Jpn. J. Appl. Phys. 37, 4460 (1998) 21. H. Miyake, A. Motogaito, K. Hiramatsu, Jpn. J. Appl. Phys. 38, L1000–L1002 (1999) 22. D.D. Koleske, A.J. Fisher, A.A. Allerman, C.C. Mitchell, K.C. Cross, S.R. Kurtz, J.J. Figiel, K.W. Fullmer, W.G. Breiland, Appl. Phys. Lett. 81, 1940 (2002) 23. K. Fujito, K. Kiyomi, T. Mochizuki, H. Oota, H. Namita, S. Nagao, I. Fujimura, Phys. Stat. Sol. (A) 205(5), 1056 (2008) 24. Hwa-Mok Kim et al., IPAP Conf. Series 1, 49 25. Sung. S. Park et al., IPAP Conf. Series 1, 60 26. K. Motoki, T. Okahisa, N. Matsumoto, M. Matsushima, H. Kimura, H. Kasai, K. Takemoto, K. Uematsu, T. Hirano, M. Nakayama, S. Nakahata, M. Ueno, D. Hara, Y. Kumagai, A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 40, L140 (2001) 27. Y. Oshima, T. Eri, M. Shibata, H. Sunakawa, A. Usui, Phys. Stat. Sol. (A) 194(2), 554 (2002) 28. A. Thon, T.F. Kuech, Appl. Phys. Lett. 69, 55 (1996)

Part III

Solution Growth Technology

Chapter 7

Ammonothermal Growth of GaN Under Ammono-Basic Conditions ´ ´ ´ R. Doradzinski, R. Dwilinski, J. Garczynski, L.P. Sierzputowski, and Y. Kanbara

Abstract We are presenting some physical and chemical basics of the ammonothermal method of bulk gallium nitride synthesis in ammonobasic regime. Excellent structural parameters and a wide spectrum of electrical properties of truly bulk GaN crystals are revealed. In the considered crystals, a low dislocation density .5  103 cm2 / is attained. At the same time, the crystal lattice is extremely flat, and the rocking curve is very narrow (FWHM D 16 arcsec). Regardless of the crystal size, the radius of lattice curvature is higher than 100 m, whereas in the best crystals it is higher than 1,000 m. Both polar and nonpolar ammonothermal GaN substrates with perfect crystalline properties enable growth of excellent quality, strain-free homoepitaxial layers. The luminescence is dominated by an intensive, perfectlyresolved excitonic structure which is uniform in the whole range of sample surface. In high excitation conditions, a biexciton emission is observed. High PL homogeneity corresponds well with structural and microscopic measurements performed on these layers. The authors are convinced, that due to perfect scalability of the ammonothermal method, large-diameter (above 2 in.) bulk GaN substrates can be employed in mass production. This will make a breakthrough in the manufacturing of high-power GaN-based devices.

7.1 Introduction At present, next generation high-efficiency lighting and high-power electronics are one of the focus areas for energy saving innovative technologies in the world industry. In this context, gallium nitride (GaN) has attracted great attention for its material properties that are useful for applications in short-wavelength optoelectronic and high power electronic devices [1, 2], such as white or colour light emitting diodes, blue laser diodes (LD), UV detectors, and high power–high frequency transistors. However, the currently available devices use GaN deposited by heteroepitaxy. The resulting thin films suffer from large defect concentrations, mainly due to a difference in lattice parameters and thermal expansion coefficients between such 137

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non-native substrates and GaN. Therefore, the development of low-cost optoelectronic and high-temperature electronic devices of sufficiently high efficiency is limited due to a lack of suitable substrates for growing homoepitaxial structures. Under such circumstances, the use of truly bulk GaN substrates would be an ideal solution of this problem. Standard methods of crystallization from melt (Czochralski, Bridgman) are not applicable because of the decomposition of nitrides into metals and N2 . In the high nitrogen pressure method (HNP), this decomposition is inhibited by the use of nitrogen under high pressure [3, 4]. The growth of single crystals is performed in molten gallium and requires both high temperatures (of about 1;500 ı C) and extremely high nitrogen pressures (of the order of 15 kbar). High pressure prevents usage of large growth chambers and nonseeded growth imposes serious limitations on crystal size. Besides the HNP technique, a number of alternative methods of bulk GaN crystallization were developed [5–7]. Satisfactory results were obtained by the growth of thick layer by different vapor-phase-epitaxy techniques - for example hydride vapor phase epitaxy (HVPE) [5]. Substrates produced this way enabled a breakthrough in commercialization of optical storage media based on blue lasers obtained ten years ago [8]. However, the resulting structures still suffer from a large dislocation density originating from the use of non-native seeds. Even after seed separation, the free-standing HVPE GaN is still highly stressed and bowed. Besides HVPE, also other methods of bulk GaN growth were developed. For example, it appeared that a Na-flux method [6] has many advantages (also concerning the crystal quality) when compared to HVPE. However, there are still many problems that need to be resolved, such as poor growth of seeds, unhomogeneity, mosaicity of crystals and poor scalability.

7.2 The Growth Method We propose the ammonothermal (AMMONO) method to overcome the aforementioned obstacles. Main motivation was to use “chemical power” for decreasing high demands of nitride crystallization, similarly to hydrothermal crystallization, where supercritical aqueous solutions are used for recrystallization of oxides (thousands of tons per year are produced with the cost of a few dollars per kilogram [9]). The ammonothermal method is regarded as an analogue of the hydrothermal one, where supercritical ammonia instead of water is used as a solvent and nitrides can be grown instead of oxides. First of all, the ammonothermal method enables growth of highdiameter seeds with perfect crystalline quality. Secondly, it is a very controllable and reproducible process at relatively low temperatures and pressures. Thirdly, it enables excellent scalability with the size of the autoclaves and thus growth of many crystals during one process with minimized material costs (closed system). Fourthly, the use of supercritical ammonia has been proposed to lower the temperature and decrease the pressure during the growth process of nitrides.

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Our group has been developing the ammonothermal method in the so-called ammonobasic environment for sixteen years now. It was initiated in mid-1990s, when R. Dwilinski et al. showed that it was possible to obtain a fine-crystalline GaN by a chemical reaction between gallium and ammonia, in the presence of alkalimetal amides (LiNH2 or KNH2 ) introduced into the reaction zone [10]. The latter compounds played the role of mineralizers, highly increasing reactivity of solution. The processes were conducted at temperatures up to 550 ıC and under pressure of 5 kbar. The resulting GaN crystals were obtained in the form of wurtzite-type microcrystalline powder, with the grain size up to 5 m [10, 11]. On the other hand, very sharp peaks of near band-edge emission accompanied by diminishing of yellow band was observed for samples grown with addition of rare earth elements [12], which suggests high affinity of rare earths to oxygen atoms, leading to reduced concentration of free electrons [13]. At the end of the 1990s, Kolis et al. [14] performed crystal growth of GaN in ammonobasic conditions using ammonia as a solvent and GaN as a nutrient at temperature T D 400 ıC and pressure p D 240 MPa. They obtained transparent GaN platelets or prismatic needles up to 0:5  0:2  0:1 mm3 [15] thanks to mineralizers such as sodium or potassium amide with an additional comineralizer (MX, M D Li, Na, K and X D Cl, I). In the acidic regime, the ammonothermal synthesis from metallic gallium with ammonium halogenide (NH4 X; X D Cl, Br, I) leads to a zincblende-type (cubic) GaN powder [16] or a mixture of hexagonal and cubic phases, depending on the growth temperature. The growth was performed at initial pressure of 70 MPa and thermal gradient 10 K cm1 was applied during heating in the quartz tube. Needle-shaped GaN crystals (with diameter up to a few tens of m and length of several hundreds of m) were also obtained [17] by recrystallization of GaN using NH4 Cl mineralizer at T D 500 ıC and relatively low pressure p D 120 MPa. In these early works the growth and recrystallization were effected spontaneously, thus leading to relatively small dimensions of the crystals. At present, an enormous progress has been made toward enlargement of the size of the ammonothermally grown crystals by taking advantage of chemical transport in temperature gradient. The method was greatly developed by the inventors in AMMONO sp. z o.o. company in collaboration with Nichia Corporation from the beginning of 2000s [18–20].

7.2.1 Physico-Chemical Basics As mentioned earlier, the idea of the ammonothermal technique was inspired by the field of hydrothermal technology, used commercially in quartz mass production. Thus, the scheme of the crystal growth process is as follows: GaN-containing feedstock is dissolved in one zone of the high-pressure autoclave (dissolution zone), then transported by convection in the temperature gradient to the second zone (crystallization zone), where GaN is crystallized on native seeds due to the supersaturation of the solution (Fig. 7.1). In addition, the use of mineralizers is necessary to enhance

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Heater T1 Feedstock Basket

Baffle Seeds

Heater T2

T2 > T1

Fig. 7.1 Schematic idea of the AMMONO-Bulk method

the solubility of GaN in ammonia. They can introduce NH2  ions to the solution in the case of ammonobasic environment of ammonothermal growth, NH4 C ions in case of acidic environment, or they can supply either of them in the case of neutral environment. Lack of the mineralizer implies that the only source of amide and ammonium ions is the dissociation of ammonia (according to reaction  2NH3 ! NHC 4 C NH2 ). Concentration of these native amide and ammonium ions seems negligible, as we could not observe any significant effect of pure ammonia on GaN. The choice of the mineralizer determines the type of environment (basic, acidic or neutral) and therefore the conditions of the ammonothermal growth. In the ammonobasic regime, pure alkali metal or alkali metal amides (LiNH2 ; NaNH2 or LiNH2 ) are used, while in the ammonoacidic regime ammonium chloride NH4 Cl or ammonium iodide NH4 I is generally used. In the case of ammonobasic conditions, implemented by the AMMONO company, precise chemical reactions leading to the synthesis and recrystallization of GaN are not known, because the growth is performed in a closed, inaccessible space of the autoclave. However, due to the fact that formation of gallium amides, i.e., Ga.NH2 /3 and imides, i.e., Ga2 .NH/3 seems to be thermodynamically privileged, some hypothetical series of reactions can be assumed in the AMMONO method. These amides or imides are unstable at growth temperature and are transformed to nitrides easily. It is also possible to obtain complexes of these amides with alkali metals such as KGa.NH2 /4 as temporary compound but it does not change the scheme of reactions leading finally to GaN. In the ammonobasic environment the following scenario is possible: mineralizer (by means of NH2  ions) attacks nitride-containing feedstock in the dissolution zone, leading to soluble metal amide compounds, for example KGa.NH2 /4 or Na2 Ga.NH2 /5 at first temperature and first

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pressure according to the following reactions [21]: KNH2 C GaN C 2NH3 ! KGa .NH2 /4 GaN C 2NaNH2 C 2NH3 ! Na2 Ga .NH2 /5

(7.1) (7.2)

After a convection-induced transport to the crystallization zone, GaN recrystallizes from those soluble intermediates onto GaN seeds present in the crystallization zone at second temperature and second pressure. This process takes place according to the reversed reactions (7.1, 7.2). The typical temperatures and pressures applied are 0.1–0.3 GPa and 400–600 ıC and a proper temperature gradient between crystallization and dissolution zone should be maintained in order to make convection possible. A few groups have presented some important results on GaN synthesis by fluid transport both in ammonobasic [18–25] and ammonoacidic [26–28] conditions. In 2006, Wang et al. have demonstrated single GaN crystals of  10  10  1 mm3 in size using GaN as a nutrient [22]. Spontaneously nucleated precipitates (of maximum size being 10 m) on GaN seeds were observed by Hashimoto et al. [24] if metallic Ga was used as a nutrient in the dissolution zone. D’Evelyn et al. have succeeded in growing mm-sized GaN crystals on HVPE seeds by fluid transport in a high-pressure version of the ammonothermal method [25]. They described their progress in reduction of impurity concentrations, wafering and even fabrication of homoepitaxial LD on ammonothermal GaN substrates. Recently, uniform growth of thick GaN via fluid transport on a 3  4 cm2 HVPE oval-shaped substrate [29] by the ammonobasic route and about 0.5-mm thick single-crystal GaN grown on 1-in. large HVPE seed by ammonoacidic one [30] are presented. In this work, the authors describe growth of truly-bulk GaN monocrystals, up to 2 in. in size in the ammonobasic regime.

7.2.2 Solubility Measurements GaN exhibits good solubility .S / in supercritical ammonia, provided that alkali metal or their compounds are introduced into the reaction zone. The solubility can be defined by the molar percentage: Sm D GaNsolvent W .MeNH2 C NH3 /100% (Me – alkali metal). Figure 7.2 shows the solubility of GaN in supercritical solvent as a function of pressure for temperatures of 400 and 500 ıC in the KNH2  NH3 system with the molar ratio x D KNH2 W NH3 equal to 0.07. The data clearly show that solubility .S / grows monotonously with pressure, but in the whole measured range the solubility values at 400 ıC are higher than those at 500 ı C, while keeping the pressure constant. Formally, this indicates that GaN reveals negative temperature coefficient of solubility .@S=@Tjp < 0/ and positive pressure coefficient .@S=@pjT > 0/ in the investigated range of parameters. Qualitatively similar behavior was observed for a number of other ammonobasic mineralizers, and it seems to be rather a rule than an exception in this regime. Such retrograde solubility was

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T = 400°C

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500°C 3

2

1

0 0

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observed for the first time by the inventors in the early 2000s [18] in ammonobasic conditions, contrary to ammonoacidic conditions, where this temperature coefficient is positive. This unusual feature has important consequences on transport properties of GaN in the ammonobasic solution, and on the autoclave design, as described in the next section. In particular, the negative temperature coefficient implies that in the presence of temperature gradient, the chemical transport of GaN in a soluble form can take place from the dissolution zone having a lower temperature (higher solubility conditions) to the crystallization zone having a higher temperature (lower solubility conditions).

7.2.3 Equipment The autoclave must fulfill some material and construction requirements – it must sustain ammonia under high pressures at elevated temperatures for many days. Proper materials for such a construction appeared to be nickel super alloys, which are characterized by outstanding mechanical properties and high resistivity against chemically aggressive supercritical-ammonia environment. The basic character of the solution is extremely favorable from the point of view of metallic autoclave materials. On the contrary, the acidic solutions have a disastrous effect upon Ni

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superalloys. The partial solution to this problem is an application of tight-sealed noble-metal (e.g., Pt) liners [17], but their cost is high and a non-negligible risk of the autoclave destruction still remains. In the case of ammonobasic mineralizers, supercritical solution is nondestructive for autoclave materials. This feature practically eliminates the risk of rupture of the autoclave which is an extremely important circumstance from the point of view of safety and cost effectiveness. Due to the fact that chemical transport of GaN is directed from the low-temperature zone of the autoclave to the high-temperature zone, and natural convection is a driving force for the mass transport, one has to place the high-temperature zone (with seeds) below the low-temperature zone (with feedstock) in order to obtain an efficient recrystallization process. Such an arrangement is demonstrated schematically in Fig. 7.1, in which we have shown the location of GaN feedstock and GaN seeds. The autoclave is constructed as a closed system in which the dissolution zone and recrystallization zone are separated by a baffle, having at least one opening (Fig. 7.1). The size of the opening(s) should be large enough to allow mass transport between the zones, but should be sufficiently small to maintain a temperature gradient in the reactor (i.e. limit the heat transport). The appropriate size of the opening depends on the size and construction details of the reactor. Two separate heating devices can be applied to two different zones in order to conduct the process of dissolution and crystallization in the temperature gradient.

7.2.4 Seeded Recrystallization The ammonothermal method in its ammonobasic version proved to be highly controllable. It means that it is possible to design a process in which crystallization takes place uniquely on seed crystals, and such a process can be scaled up with the autoclave size. The main issue here is to conduct the process in such a manner, that the supersaturation of the solution in the crystallization zone is not too high. When certain threshold value of supersaturation is exceeded, a spontaneous nucleation of GaN appears which can finally disturb the seeded growth. The factors of high importance are: the temperature-time route, spatial distribution of temperature, and an arrangement of inner-construction elements, such as baffles, crucibles, seed racks etc. Our long experience with the ammonothermal technique indicates usefulness of crystals obtained by ammonothermal method (A-GaN crystals) as seeds for their further multiplication and crystal-size enlargement. Moreover, no deterioration of the quality as a function of crystal thickness was observed, and perfect scalability with the size of the autoclave, as well as the possibility of simultaneous use of many seeds in a high-volume chamber, were also confirmed. It is worth mentioning that the AMMONO team currently does not identify any obstacles to increase the crystal dimensions besides the diameter of the autoclave. The aforementioned features provide a great opportunity to produce polar substrates (exceeding 2-in. diameter) and nonpolar substrates of competitive size. This will require large scale equipment for

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pressures involving supercritical ammonia and manufacturing of substantial amount of high quality seeds. Let us now remind and emphasize serious advantages of the ammonobasic version of the ammonothermal method (the one we use) with respect to its acidic counterpart, in which the temperature coefficient of solubility is positive .@S=@Tjp > 0/, and GaN is transported from the high-temperature zone to the low-temperature zone. First, the basic character of solution is nondestructive on autoclave material. Second, one has to take into account that even Ni  Cr alloys have limited strength and that certain temperature-pressure conditions cannot be exceeded. Therefore, it is favorable to crystallize GaN in the hot zone and take advantage of maximum temperature in the system, as it leads to better growth quality and lower dislocation density. Third, as was reported in [26, 27], the acidic regime suffers from the appearance of cubic-GaN phase at certain conditions, whereas in ammonobasic regime one deals with the purely hexagonal phase, as shown in Sect. 7.3.1. And last but not least, the retrograde GaN transport direction enables to decrease a contamination level of growing GaN with other undesirable phases (e.g., microprecipitates), as they are typically transported to the cold zone of the system.

7.2.5 Doping In the ammonothermal method the bulk monocrystalline GaN can be easily doped by different donors and acceptors by introducing appropriate precursors of dopants into the autoclave and adding them to the feedstock. In general, different donor dopants, acceptor dopants, and magnetic dopants can be selected according to the desired properties of the substrates. Preferably, they are selected, from the group consisting of Si and O, while Zn and Mg are preferred as acceptor dopants. The concentration of dopants can be controlled in order to alter optical and electrical properties of crystals, ranging from highly conductive to semi-insulating [31] (Sect. 7.3). Typically, the concentrations of these dopants range from 1017 to 1021 cm3 and depend on desired final application of the crystal.

7.2.6 Crystal Machining Figure 7.3a shows one of our first 2-in. seed monocrystals, from which 2-in. substrates can be sliced. In Fig. 7.3b we present our typical 1-in. AMMONO-GaN substrate, which is a c-plane oriented and polished wafer, sliced and round-shaped from a larger crystal. Due to a perfect scalability of the ammonothermal method with autoclave diameter we were able to enlarge the size to 1.5-in. (Fig. 7.3b). Very recently, AMMONO Sp. z.o.o. began to increase the population of crystals (Fig. 7.3a) used for the manufacturing of 2-in. wafers. This may facilitate more effective production of optoelectronic devices grown on truly bulk GaN substrates.

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Fig. 7.3 (a) Photograph of 2-in. seed crystal obtained at AMMONO Company, (b) Photograph of 1-in. and 1.5-in. A-GaN wafer manufactured at AMMONO company

For photonic applications non-polar, m-plane GaN substrates are even more promising and interesting (e.g. green lasers) due to lack of internal electric fields. Since now such substrates have been obtained by hydride vapor phase epitaxy (HVPE) method [32, 33] by slicing the c-direction grown bulk boule along the proper direction, yielding about 1-cm square plates. Similar approach was applied to GaN crystals grown by ammonothermal method. In this case the larger size of semi-polar or non-polar GaN substrates can be obtained from much smaller diameter of boule, when compared to 2-in. HVPE GaN. Figure 7.4a shows the photograph of the thick 1-in. GaN crystals, from which a 10:5  19 mm nonpolar substrate can be made (Fig. 7.4b) [34]. Production of semipolar, for example (201) plane oriented wafers, is also achievable. The wafers must be properly polished by mechanical and chemico-mechanical polishing (CMP) in order to obtain an epi-ready surface. Such prepared A-GaN substrates were applied for homoepitaxy by metal-organic vapor phase epitaxy and epilayers of excellent properties were obtained this way. The results of structural and optical investigations of these epilayers are the subject of Sect. 7.4.

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a

b

Fig. 7.4 (a) Photograph of a thick 1-inch A-GaN crystal, (b) typical m-plane A-GaN substrate fabricated from such a crystal.

7.3 Crystal Characterization 7.3.1 Structural Properties Since the AMMONO-Bulk Method enables growth of high-quality seeds, that can be further multiplied and regrown without a noticeable loss of quality, the AMMONO-GaN crystals are characterized by excellent structural properties. All structural data shown below was measured on as-grown AMMONO-GaN crystals of various shapes and dimensions which did not influence their properties. The pure hexagonal phase of ammonothermally grown GaN was confirmed by powder diffractometry of pulverized single crystal (Fig. 7.5) [21]. In addition, we performed X-ray rocking curve measurements on (002) plane of N-face GaN obtained by ammonothermal method. The measurements were conducted using Panalytical X’Pert Pro MRD high resolution with a Cu K’1 line source, U D 40 kV; I D 20 mA, angle resolution 0:0001ı, slit width 0:1  0:1 mm2 for the incident beam and an open-detector mode for the diffracted beam. Excellent crystallinity was manifested by full width at half maximum (FWHM) value as low as 16 arcsec (Fig. 7.6). We also stress that similar value was obtained on asymmetric (014) and (115) planes. Further narrowing of FWHM to about 10 arcsec was achieved by applying a highly monochromatic and perfectly collimated lowdiameter (below 1 m) X-ray microbeam of synchrotron radiation Spring-8 facility located in Hyogo, Japan [21]. This is still very low FWHM value as compared with that obtained for GaN obtained in ammonoacidic conditions directly on HVPE seeds (about 100 arcsec) [35]. The natural cause of such structural properties is the possibility of using native ammonothermal GaN as seeds for further proliferation in order to create more and more new generations of AMMONO-GaN crystals without any loss of quality. Structural properties of AMMONO-GaN were also illustrated by reciprocalspace maps, as shown in Fig. 7.7. The left image corresponds to .Qx ; Qy / crosssection of reciprocal space map for (002) symmetrical reflex of truly bulk A-GaN crystal and the right one, for highly stressed GaN layer grown on sapphire substrate

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25000

20000

15000

10000

5000

0 15

25

35

45 2

55

65

㧛 deg.

Fig. 7.5 X-ray diffraction pattern of pulverized A-GaN crystal 100 μm x 100 μm slit

intensity (a.u.)

150

100

FWHM = 16 arcsec

50

17.34

17.36

17.38

17.40

17.42

omega (degrees)

Fig. 7.6 (0002) X-ray rocking curve of A-GaN crystal

measured for comparison (Fig. 7.7). The general conclusion is that AMMONO-GaN crystals reveal excellent single-crystal structure, without mosaicity and low-angle grain boundaries, as broadening along the Qx direction .FWHM Qx D 0:0008 rlu/ is very small, especially in comparison to broadening of GaN epilayer grown on sapphire. Additionally, one can determine the lattice parameters a, c of AMMONO-GaN monocrystals by solving a set of two equations (7.3)

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Qy·104

Qy·104

AMMONO-GaN

2973

2973

2971

2971

2969

GaN on sapphire

2969 –8.2

Qx·104 8.2

0.0

0.0 Qx·104

–8.2

FWHM Qx = 0.80

8.2

FWHM Qx = 4.67

Fig. 7.7 Reciprocal space map collected for (0002) point collected for A-GaN truly bulk crystal (left panel) and GaN epilayer grown on sapphire (right panel)

1 2 dhkl

D

4 3



h2 C hk C k 2 a2

 C

l2 c2

(7.3)

(hkl – Miller indices) related with two interplane spacings d114 and d015 evaluated from reciprocal space maps collected in asymmetric (114) and (015) reflexes. The procedure was repeated for several measurement points of the same crystal, yielding mean lattice parameters a D .3:18919 ˙ 0:00030/ ; c D .5:18529 ˙ 0:00015/ for intentionally undoped n-type crystals. Semi-insulating crystals reveal lattice parameters a D .3:18896 ˙ 0:00023/ , c D .5:18539 ˙ 0:00014/ and p-type crystals give values a D .3:18871 ˙ 0:00028/ ; c D .5:18551 ˙ 0:00020/ . We stress high resolution of the measurement and very low gradients of lattice parameters across the crystal. The obtained values can be used as reference lattice parameters of unstrained GaN. For comparison, the lattice parameters of high nitrogen pressure grown crystals are a D .3:1881 3:1890 ˙0:0003/ ; c D .5:1846 5:1864 ˙0:0002/ [36]. Measurements of the radius of curvature of the AMMONO-GaN lattice were also performed. For the sake of comparative analysis similar measurements were first performed for a standard free-standing HVPE-GaN. The HVPE-GaN measurement consisted in collecting about a dozen X-ray rocking curves from a (002) plane, measured in colinear points spaced by 5 mm along the wafer. A typical result of such a measurement is presented in Fig. 7.8. Because of the lattice bending, the maxima of diffraction peaks shift systematically on the -axis. This effect reflects the fact that the inclination of (0002) plane changes systematically when moving along the measurement line. For each pair of maxima one can calculate the corresponding radius of curvature. As shown in the table on the right side of Fig. 7.8, the values for HVPE-GaN range between 2 and 12 m. In contrast to that, X-ray rocking curves measured at the AMMONO-GaN sample remained essentially in the same position, as illustrated in Fig. 7.9. When

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counts/s 1400 1200

HVPE GaN R curv = 2 – 12 m

Position Radius of on sample curvature [m]

1000

–25 : –20 –20 : –15 –15 : –10 –10 : –5 –5 : 0 0:5 5 : 10 10 : 15 15 : 20 20 : 25 25 : 30

800 600 400 200 0 17.0

17.2

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17.8

18.0

18.2

18.4

12,26 4,79 3,48 2,64 2,27 2,17 3,36 3,38 2,28 2,10 1,88

18.6 Omega (°)

Fig. 7.8 Family of X-ray rocking curves for HVPE GaN and radius of curvature evaluation counts/s 2500

R curv > 1000 m 2000

1500

1000

AMMONO-GaN

500

0 16.80

16.82

16.84

16.86

16.88

16.90

16.92

16.94 Omega (°)

Fig. 7.9 Family of X-ray rocking curves for AMMONO-GaN and radius of curvature evaluation

calculating the radius of curvature, values higher than 1,000 m were obtained. Although such an excellent result cannot be treated as representative for the population of AMMONO-GaN crystals, it was confirmed in many cases, and the averaged value for the whole population is of the order of hundred meters. These data prove

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HVPE-GaN 64 x 64 mm

CL (420 nm) (420nm)

CL (420 nm) DSD = 3 x106 /cm2

DSD unmeasurable

Fig. 7.10 Cathodoluminescence result for HVPE-GaN (left panel) and AMMONO-GaN (right panel) 1 x 1 mm

EPD = 5 x 103 /cm2

Fig. 7.11 Etch pits in AMMONO-GaN

that crystal lattice of AMMONO-GaN crystals is extremely flat, leading to a very low level of the built-in stress. In order to estimate the surface dislocation density, cathodoluminescence (CL) imaging was performed on both free-standing HVPE-GaN and AMMONO-GaN samples. The representative pictures are shown in Fig. 7.10, where the detectedlight wavelength is 420 nm and the picture area is 64  64 m2. The dark spot density (DSD) of the HVPE material (left panel) can be estimated at the level of 3  106 cm2 . On the contrary, in the AMMONO-GaN sample (right panel) one can hardly notice any dark spot and therefore estimation of DSD was not possible. Finally, we performed etching experiments in molten KOH (400 ıC, 5 min) in order to determine the dislocation density in AMMONO-GaN. The results of 1 mm2 of surface area are presented on differential interference contrast (DIC) microscopy photograph in Fig. 7.11. In general, selective etching reveals hexagonal pits resulting from increased etching rate on dislocations. All defects were revealed in 400 ı C,

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since no new defects appeared by etching in higher temperature 520 ıC. The value of etch pit density (EPD) is estimated at the level of 5  103 cm2 , which is an outstanding result as compared with typical dislocation density of HVPE material (of the order of 106 cm2 ).

7.3.2 Optical Properties Photoluminescence (PL) is one of the best tools to identify the existence of specific defects and their participation in different radiative recombination processes. In this respect, microphotoluminescence (-PL) experiments were performed at temperature 4.2 K using a 3.813 eV line of He  Cd laser as an excitation source. The laser spot was focused on surface using a microscopic reflection with the magnification of 25. Resulting excitation spot size was estimated to be 10 m in diameter. The -PL spectra were collected at an as-grown N-side of AMMONO-GaN crystal at temperature T D 4:2 K in wide spectral range of 2.0–3.5 eV. The results of -PL are shown in Fig. 7.12. The luminescence is dominated by strong band edge emission, consisting mainly of neutral donor bound exciton .D0 X/ recombination (3.4716 eV) [37] together with its single and double

60000

D0 X

T=4.2K

Intensity (Counts)

50000

40000

D0X-LO 30000

20000

D0X-2LO

10000

D-A

0 2.2

2.4

2.6

2.8

3.0

3.2

3.4

Energy (eV) Fig. 7.12 -photoluminsescence spectra for as grown AMMONO-GaN surface measured at T D 4:2 K. The peaks are marked in the following way: D0 X – exciton bound to neutral donor, D0 X  LO: optical phonon replica of exciton bound to neutral donor line, D0 X–2LO: two optical phonons replica of exciton bound to neutral donor line, D–A: donor–acceptor transition

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optical phonon replica .D0 X–LO; D0 X–2LO/ and donor-acceptor pair transition (D–A, 3.283 eV). Intense excitonic emission is accompanied by diminishing of typical of GaN, parasitic yellow luminescence .2:2 eV/, attributed mainly with Ga vacancy acceptors [38]. This observation proves high optical quality and unique purity of the studied material, enabling effective excitonic recombination.

7.3.3 Electrical Properties The carrier concentration of the wafers can be controlled by appropriate doping [39], as mentioned in Sect. 7.2.5. Both n-type .n  1018  1019 cm3 , D 103  102  cm/, p-type .p  1018 cm3 , D 101 –102  cm/ and semi-insulating .  106  cm–1012  cm/ substrates can be grown via the ammonothermal method [39], as measured by both Hall effect experiments and contactless methods. Tuning of the electrical properties suggest, that truly bulk AMMONO-GaN crystals may find application in both optoelectronics (highly conductive substrates) and high power electronics (semi-insulating substrates for high electron mobility transistors production).

7.4 Homoepitaxy on Ammonothermal GaN The epilayers were deposited by metalorganic chemical vapor phase deposition (MOCVD) using a RF heated AIXTRON AIX-200 low pressure horizontal reactor on both c-plane (Ga face) and m-plane oriented substrates. Trimethylgallium and ammonia were used as Ga and N precursors, respectively. The growth procedure was started by annealing the substrate for 5 min at high temperature .1; 100 ıC/ in a H2 C NH3 atmosphere at a total reactor pressure of 350 mbar. For all samples grown on A-GaN substrates the ammonia .NH3 / was kept open from the moment the temperature in the reactor reached 500 ıC till the end of the process when the grown sample was cooled down to 500 ı C. This ammonia gas flow was kept constant using 2 standard liters per minute (slm) till the end of the growth process. This step was used to prevent first GaN substrate and then epitaxial layers from decomposition [40]. After the annealing, the growth time was set such as to obtain the 1–2 m thick GaN layer. The growth temperature was 1;170 ıC for the whole process and was carried out at a total reactor pressure of 50 mbar. The epilayers were also characterized by various techniques: X-ray diffraction, defect selective etching (DSE), -PL, PL, and reflectance. In order to check the overall structural quality of GaN epilayers grown on A-GaN polar substrates, the X-ray rocking curves mode measurements were performed. Representative rocking curves measured on (002) plane of epilayer grown on n-type and SI substrate are shown on Fig. 7.13. The resulting FWHM values (22–24 arcsec) are, to our knowledge, outstanding results, that strongly indicate the

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(a) 2 μm MOCVD GaN/ n-type A-GaN substrate Intensity (counts)

1000 100 μm x 100 μm slit

FWHM = 22 arcsec 500

0

(b) 2 μm MOCVD GaN/ SI A-GaN substrate Intensity (counts)

1000

100 μm x 100 μm slit

FWHM = 24 arcsec 500

0 –0.04

-0.02

0.00

0.02

0.04

angle (deg) Fig. 7.13 X-ray rocking curves of the GaN epilayers grown on n-type (a) and SI (b) A-GaN substrates for the symmetric (002) reflections

excellent crystalline properties of the studied material. This result is a direct consequence of the high crystalline quality of A-GaN substrates adapted to MOCVD growth, described in Sect. 7.3.1. All optical date are collected in the band-edge emission region. Presented in Fig. 7.14 is the -PL spectra for an epilayer grown on SI polar substrate which is dominated by very intensive, perfectly-resolved excitonic emission. This structure consists of both lines related to emission of free excitons .FXA ; FXB /, and very sharp and narrow lines (of the FWHM value of the order of 0.3 meV) related to excitons bound to neutral acceptor .A0 X/ and both excitons  A and B bound to different neutral donors D1 0 XA ; D2 0 XA ; D3 0 XA ; D0 XB . Exact identification of individual donors involved with the latter four transitions requires further research (for example magnetooptic measurements in high magnetic fields). Two kinds of free excitons .FXA ; FXB / are formed with participation of holes originating from crystal field and spin-orbit split valence band. The energies of all the observed transitions are the following: 3.4779 eV, 3.4831 eV, 3.4717 eV, 3.4730 eV,

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0

AX

T=4.2 K

0

FWHM=0.3 meV

D1XA 0

Intensity (Counts)

10000

D3XA

SI substrate

1000

0

D2XA FXA 0 D XB

TES

FXB 100

10

1 3.44

3.46 Energy (eV)

3.48

Fig. 7.14 -PL spectra of GaN epilayer grown on SI A-GaN substrate measured at T D 4:2 K. The peaks are marked in the following way: FXA – free exciton A, FXB – free exciton B, D1 0 XA ; D2 0 XA ; D3 0 XA – exciton A bound to different neutral donors, D0 XB – exciton B bound to neutral donor, A0 X – exciton bound to neutral acceptor, TES – two electron satellite

3.4707 eV, 3.4754 eV, 3.4663 eV for FXA ; FXB ; D1 0 XA ; D2 0 XA ; D3 0 XA ; D0 XB ; A0 X peaks, respectively. They are very close to those reported for unstrained homoepitaxial GaN (3.4776 eV, 3.4827 eV for free excitons A and B recombination [41] and 3.4719 eV and 3.4668 eV for exciton bound to donor and acceptor emission [37]. In addition to excitonic structure we observed two electron transition (TES) – a recombination of exciton bound to donor that leaves the donor in the excited state [42]. We also stress that the optical spectra is reproducible in the entire sample surface range. In other words, the optical properties are highly homogeneous independently on excitation place (Fig. 7.14). High PL homogeneity corresponds well with structural and microscopic measurements performed on these layers. Moreover, the width of bound excitons peaks .FWHM  0:3 meV/ is one of the smallest ever reported for homoepitaxial GaN. It proves that ammonothermal GaN substrates with perfect crystalline properties enable the growth of excellent quality, strain-free homoepitaxial layers. Additional PL measurements were also performed at various impulse excitation conditions in a standard configuration using a Si multichannel detector, 0.55 m length monochromator and YAG laser emitting at 266 nm with an average output power of 2 mW (0:1 J in one 10 ns pulse). The excitation beam was defocused to a diameter of 2 mm and its intensity was changing from 0.01 to 2 mW. The same experimental set-up was used to measure reflectance spectra. The sample was illuminated by white light from a halogen lamp (100 W) at near normal incidence.

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For low temperature measurements the samples were mounted on a cold finger in a continues-flow liquid helium cryostat. Those additional PL measurements performed under pulsed laser excitation (Fig. 7.15) seem to corroborate the above results. In this case the bound exciton recombination is saturated and suppressed in the scale of the spectra due to the limited number of donors and acceptors. Instead, the biexciton emissions (marked by XX1 ; XX2 ) appear and dominate the PL spectra together with free exciton .FXA ; FXB / recombination under high excitation conditions (Fig. 7.15). Exact study of biexciton emission is beyond the scope of this chapter and will be published elsewhere. In the reflectance spectra (Fig. 7.16a) two dispersion lines .FXA ; FXB /, related to free excitons A and B described earlier, are clearly visible. In addition, there is a third one, a bit weaker, assigned as FXC , that consists of the electron and hole originating from the third remaining valence subband resulting from crystal field and spin-orbit splitting. The energy position of the dispersion lines FXA and FXB are in reasonable agreement with PL and -PL results and literature data [41]. More thorough analysis of reflectance and PL spectra will be a subject of an upcoming paper. In case of GaN epilayers deposited on nonpolar m-plane A-GaN substrates it should be expected that the FXA exciton will be completely polarized in the epilayer plane [34]. Indeed, the studied films exhibit intrinsic narrow exciton lines, which are very sensitive to the optical selection rules typical for hexagonal

GaN epilayer on n-type GaN substrate 1000

XX1 P max

PL intensity (arb. u.)

10

FXA FX B

XX2

100

T=10K

P min

1

0.1

0.01

1E-3 3.45

3.46

3.47

3.48

3.49

Energy (eV) Fig. 7.15 PL spectra measured for GaN epilayers grown on n-type A-GaN substrate collected at T D 10 K at various excitation conditions. The peaks are marked in the following way: FXA – free exciton A, FXB – free exciton B, XX1 ; XX2 – biexciton lines

T=10K

B

FXC

XX

c-plane

(b)unpolarized

FXA

FXB

FXC

E^ c E II c

I m-plane

3.42

3.44

3.46

3.48

Reflectance (arb. u.)

PL intensity (arb. u.)

FXA FX

(a)

Reflectance (arb. u.)

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3.50

Energy (eV) Fig. 7.16 (a) Reflectance and PL spectra of GaN epilayers grown on n-type polar A-GaN substrate collected at T D 10 K, (b) Polarized reflectance spectra measured at 10 K for GaN homoepitaxial layer grown on m-plane A-GaN substrate. Polarization configurations are ¢ .k?c; Ejjc/ and   .k?c; E?c/. The lines are marked in the following way: FXA – free exciton A, FXB – free exciton B, FXC – free exciton C, XX – biexciton, I – bound exciton line

symmetry – the free exciton A line completely disappears in Ejjc configuration, proving truly nonpolar character of m-plane oriented A-GaN crystals (Fig. 7.16b). The crystals have an ideal hexagonal symmetry without any local structural imperfections which could destroy this symmetry and change the optical selection rules. Again, observation of both free excitons and biexciton transitions is a fingerprint of high quality studied material. High quality of homoepitaxial layers deposited on polar c-plane A-GaN substrates was also confirmed by modulation spectroscopy: photoreflectance (PR) and contactless electroreflectance (CER) measurements. We note that the PR technique has never been applied to study homoepitaxial films because of the limiting access to bulk GaN substrates. The PR spectra (Fig. 7.17) exhibits the following

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0.03

FXA

0.02

FXB

T = 10 K FXC

0.01

DR/R

157

0.00 –0.01 –0.02

GaN(2μm)/A-GaN

–0.03

Exp. Fit

–0.04 3.47

3.48

3.49

3.50

Energy (eV) Fig. 7.17 Photoreflectance spectra from a GaN epilayer grown on polar c-plane A-GaN substrate. Solid line represents Lorentzian line fit. The lines are marked in the following way: FXA – free exciton A, FXB – free exciton B, FXC – free exciton C

features: (1) Typical three resonances corresponding to intrinsic free exciton transitions (FXA, FXB and FXC) of the energies evidently lower (3.4760 eV, 3.4817 eV, 3.4991 eV) than those typical for GaN layers grown on sapphire substrates (3.483 eV, 3.491 eV, 3.512 eV, respectively) [43] were observed, (2) Exciton linewidth for the homoepitaxial layer are significantly smaller than those for the heteroepitaxial layer, the observed transitions are also about 20–50% narrower than the exciton transitions reported in [41] measured by reflectivity for homoepitaxial layers grown on HNP crystals. Such narrow resonances and homogeneous broadening of excitonic transitions mean that no residual strain exists in GaN epilayer. Very narrow CER resonances .15 meV/, typical of high quality material, have been observed for considered epilayers, while broad CER resonances typical of band-to-band absorption in bulk material were recorded for A-GaN substrates [44]. It confirms the excellent usefulness of ammonothermal GaN substrates for homoepitaxy.

7.5 Conclusions The results of measurements performed for bulk GaN single crystals obtained by the AMMONO-Bulk Method, such as: (a) FWHM value of (0002) X-ray rocking curve of about 16 arcsec, (b) the radius of curvature of the crystal lattice of the order of 102 –103 m, and (c) EPD of the order of 5  103 cm2 (d) wide spectrum of electrical properties, all unambiguously prove excellent crystalline quality of the grown material. The excellent structural properties of crystals, combined with advantages of the method itself (seed-multiplication, repeatability, scalability), allow us to claim that AMMONO-Bulk Method is an extremely promising technology for fabricating

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high-quality, large diameter GaN substrates. Confirmed high scalability of the method will play a crucial role in further improving its cost-effectiveness when implemented on an industrial scale. Moreover, as no deterioration of the quality within crystal thickness was observed, the growth of sufficient sized non-polar, mplane substrates was succeeded. Such substrates eliminate the problem of internal piezoelectric field, which reduces efficiency of devices and are for example essential for achieving green lasers. Structural and optical properties of GaN MOCVD layers grown on ammonothermal GaN substrates have been studied by X-ray diffraction, defect selective etching, PL, -PL microphotoluminescence, reflectance, and modulation spectroscopy. We have succeeded in performing homoepitaxy by MOVPE resulting in high quality, strain free GaN epilayer, which was documented by the small width of bound exciton peaks, observation of free excitons and biexciton emission (and absorption), highly homogeneous luminescence properties, and small FWHM of X-ray rocking curve. Presented results indicate that both polar and nonpolar (m-plane) A-GaN substrates may be excellent for homoepitaxy. Moreover, controllable conductivity of A-GaN substrates (from n-type to SI) can be very useful in fabrication of various devices, like LD operating in blue light range, UV detectors, or high electron mobility transistors (HEMTs). Acknowledgements This project was partially supported by the European Union fund “Sektorowy Program Operacyjny Wzrost Konkurencyjno´sci Przedsie¸biorstw – Program 1.4.1” (Europejski Fundusz Rozwoju Regionalnego). The authors acknowledge the following colleagues for their contribution to this work: Dr. Robert Kucharski and his team for machining A-GaN crystals in order to prepare samples for epitaxy and measurements, Tomasz Tracz and his team for growing bulk A-GaN crystals, dr Mariusz Rudzi´nski for performing homoepitaxy, Arkadiusz Puchalski and Dr. Jarosław Serafi´nczuk for X-ray measurements, Dr. Andrzej Wysmołek for -PL measurements, and Dr. Robert Kudrawiec for PL, reflectivity, PR measurements.

References 1. See, for example, S. Nakamura, G. Fosol, The Blue Laser Diode (Springer, Berlin, 1998) 2. See, for example, H. Morkoc, Nitride Semiconductors and Devices (Springer, Heidelberg, 1999) 3. S. Porowski, Mater. Sci. Eng. B 44, 407 (1997) 4. I. Grzegory, M. Bo´ckowski, B. Łucznik, S. Krukowski, Z. Romanowski, M. Wróblewski, S. Porowski, J. Cryst. Growth 246, 177 (2002) 5. B. Łucznik, B. Pastuszka, I. Grzegory, M. Bo´ckowski, G. Kamler, E. Litwin-Staszewska, S. Porowski, J. Cryst. Growth 281, 38 (2005) 6. H. Yamane, M. Shimada, S.J. Clarke, F.J. DiSalvo, Chem. Mater. 9, 413 (1997) 7. A. Denis, G. Goglio, G. Demazeau, Mater. Sci. Eng. R 50, 167 (2006) 8. S. Nakamura, M. Senoh, S. Nagahama, N. Iwasa, T. Yamada, T. Matsushita, H. Kiyoku, Y. Sugimoto, T. Kozaki, H. Umemoto, M. Sano, K. Chocho, Appl. Phys. Lett., 72, 211(1998) 9. K. Byrappa, M. Yoshimura, Handbook of Hydrothermal Technology (Noyes, NJ, USA, 2001) 10. R. Dwilinski, A. Wysmolek, J. Baranowski, M. Kaminska, R. Doradzinski, J. Garczynski, L. Sierzputowski, Acta Physica Polonica 88(5), 833 (1995)

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11. R. Dwilinski, J. Baranowski, M. Kaminska, R. Doradzinski, J. Garczynski, L. Sierzputowski, Acta Physica Polonica 90(4), 763 (1996) 12. R. Dwilinski, R. Doradzinski, J. Garczynski, L. Sierzputowski, J. Baranowski, M. Kaminska, Mater. Sci. Eng. B 50, 46 (1997) 13. M. Palczewska, B. Suchanek, R. Dwilinski, K. Pakula, A. Wagner, M. Kaminska, MRS Internet J. Nitride Semicond. Res. 3, Article 45 (1998) 14. J.W. Kolis, S. Wilcenski, R.A. Laudise, Mater. Res. Soc. Symp. Proc. 495, 367 (1998) 15. D.R. Ketchum, J.W. Kolis, J. Cryst. Growth, 222, 431 (2001) 16. A.P. Purdy, Chem. Mater. 11, 1648 (1999) 17. A. Yoshikawa, E. Ohshima, T. Fukuda, H. Tsuji, K. Oshima, J. Cryst. Growth 260, 67 (2004) 18. R.T. Dwilinski, R.M. Doradzinski, J.S. Garczynski, L.P. Sierzputowski, Y. Kanbara, Polish Patent Application no. P-347918 (06.06.2001) 19. R.T. Dwilinski, R.M. Doradzinski, J.S. Garczynski, L.P. Sierzputowski, Y. Kanbara, International Patent Application no. PCTnIB02n04185 (17.05.2002) 20. R.T. Dwilinski, R.M. Doradzinski, J.S. Garczynski, L.P. Sierzputowski, Y. Kanbara, United States Patent no. 6,656,615 B2 (02.12.2003) 21. R. Dwilinski, R. Doradzinski, J. Garczynski, L.P. Sierzputowski, A. Puchalski, Y. Kanbara, K. Yagi, H. Minakuchi, H. Hayashi, J. Cryst. Growth 310, 3911 (2008) 22. B. Wang, M.J. Callahan, J. Cryst. Growth 291, 455 (2006) 23. T. Hashimoto, K. Fujito, B.A. Haskell, P.T. Fini, J.S. Speck, S. Nakamura, J. Cryst. Growth 275, e525 (2005) 24. T. Hashimoto, K. Fujito, R. Sharma, E.R. Letts, P.T. Fini, J.S. Speck, S. Nakamura, J. Cryst. Growth 291, 100 (2006) 25. M.P. D’Evelyn, H.C. Hong, D.-S. Park, H. Lu, E. Kaminsky, R.R. Melkote, P. Perlin, M. Leszczynski, S. Porowski, R.J. Molnar, J. Cryst. Growth 300, 11 (2007) 26. A.P. Purdy, Cryst. Growth Des. 2(2), 141 (2002) 27. A.P. Purdy, A.C. Cas, N. Muratore, J. Cryst. Growth 252, 136 (2003) 28. Y. Kagamitani, D. Ehrentraut, A. Yoshikawa, N. Hoshino, T. Fukuda, H. Tsuji, K. Oshima, Jpn. J. Appl. Phys. 45, 4018 (2006) 29. T. Hashimoto, K. Fujito, M. Saito, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 44, L1570 (2005) 30. T. Fukuda, D. Ehrentraut, J. Cryst. Growth 305, 304 (2007) 31. R.T. Dwilinski, R.M. Doradzinski, J. Garczynski, L.P. Sierzputowski, Y. Kanbara, International Patent Application no. PCTnPL2005n000036 (10.06.2005) 32. T. Paskova, R. Kroeger, D. Hommel, P. P. Paskov, B. Monemar, E. Preble, A. Hanser, N.M. Wiliams, M. Tutor, Phys. Stat. Sol. (C) 4, 2536 (2007) 33. K. Fujito, K. Kiyomi, T. Mochizuki, H. Oota, H. Namita, S. Nagao, I. Fujimura, Phys. Stat. Sol. (A) 205, 1056 (2008) 34. R. Kucharski, M. Rudzi´nski, M. Zaja¸c, R. Doradzi´nski, J. Garczy´nski, L. Sierzputowski, R. Kudrawiec, J. Serafi´nczuk, W. Strupi´nski, R. Dwili´nski, Appl. Phys. Lett. 95, 131119 (2009) 35. D. Ehrentraut, Y. Kagamitani, C. Yokoyama, T. Fukuda, J. Cryst. Growth 310, 891 (2008) 36. M. Leszczy´nski, H. Teisseyre, T. Suski, I. Grzegory, M. Bo´ckowski, J. Jun, S. Porowski, K. Pakuła, J.M. Baranowski, C.T. Foxon, T.S. Cheng, Appl. Phys, Lett. 69, 73 (1996) 37. K. Pakuła, A. Wysmołek, K.P. Korona, J.M. Baranowski, R. Ste¸pniewski, I. Grzegory, M. Bo´ckowski, J. Jun, S. Krukowski, M. Wróblewski, S. Porowski, Sol. State Commun. 97, 919 (1996) 38. J. Neugebauer, C. van de Walle, Appl. Phys. Lett. 69, 503 (1996) 39. R.T. Dwilinski, R.M. Doradzinski, J. Garczynski, L.P. Sierzputowski, Y. Kanbara, International Patent Application no. PCTnPL2005n000036 (10.06.2005) 40. D.D. Koleske, A.E. Wickenden, R.L. Henry, M.E. Twigg, J.C. Culbertson, R.J. Gorman, Appl. Phys. Lett. 73, 2018 (1998); Erratum Appl. Phys. Lett. 75, 1646 (1999).

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Chapter 8

A Pathway Toward Bulk Growth of GaN by the Ammonothermal Method Tadao Hashimoto and Shuji Nakamura

Abstract In this chapter, research progress of bulk GaN growth by the ammonothermal method at the University of California, Santa Barbara is reviewed. Starting from a fundamental study on chemistry by powder synthesis, ammonothermal GaN growth process has been developed through solubility measurement, growth from metallic Ga nutrient and growth from polycrystalline GaN nutrient. Based on accumulated knowledge on the growth process, polyhedron-shaped GaN bulk crystals were achieved. Future challenges are presented through characterization of GaN wafers sliced from the bulk crystals.

8.1 Introduction Until the early 1990s, many researchers in the field of compound semiconductors aspired to become the “first” to realize blue light emitting diodes (LEDs) and blue laser diodes (LDs). Despite the earlier achievements of blue LEDs by silicon carbide [1] and blue LDs by II–VI materials [2], commercially reliable blue/green/white LEDs and blue LDs have been achieved with GaN-based solid solutions. After the successful demonstration of InGaN blue LEDs [3] and blue LDs [4], many crystal growers’ interest has shifted toward growth of bulk GaN crystals. Although not many material scientists are privileged to pursue this “second dream” due to limited access to special growth apparatus, the entire body of nitride research has longed for realization of GaN wafers and bulk GaN crystals. This is because GaN or AlN substrates provide a direct and ultimate solution for fundamental growth problems arising from heteroepitaxy (e.g., on sapphire). Growth of bulk GaN, however, is extremely challenging due to the high melting temperature and high equilibrium N2 pressure at high temperature. Several growth techniques based on a Ga melt have been investigated to grow bulk GaN crystals [5–8]. Nevertheless, these methods encountered difficulties to produce large enough bulk GaN crystals which can be sliced into 2 in. wafers. In the meanwhile, hydride vapor phase epitaxy (HVPE) was developed to take a quasi-bulk approach to fabricating GaN wafers. GaN wafers by HVPE made a significant contribution to mass 161

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production of blue LDs for blu-ray DVDs; however, cost reduction is one of the most challenging issues in the quasi-bulk approach. Considering a rapid decrease in the market prices of blue and white LEDs grown on sapphire or SiC, a bulk growth method must be expandable to a volume production. Ammonothermal growth, a solvothermal method using NH3 as a fluid, is a promising candidate owing to its potential scalability. Ammonothermal growth is regarded as a highly expandable method since it is an analogue of hydrothermal growth widely used for the quartz mass production. The recent demonstration of bulk GaN grown by the ammonothermal method [9] opened up a new substrate choice for nitride-based LEDs. For instance, LEDs fabricated on nonpolar or semipolar bulk GaN substrates have already realized performance comparable to state-of-the-art c plane LEDs [10, 11]. Nitride-based solid-state lighting which requires even brighter LEDs is expected to be realized with low-cost and high-quality nonpolar/semipolar GaN substrates sliced from ammonothermally grown bulk GaN. Since Peters reported the AlN synthesis in supercritical NH3 in 1990 [12] several groups have reported key achievements in the ammonothermal growth of GaN [9, 13–27]. Ammonothermal activity at the University of California, Santa Barbara (UCSB) started in July 2000. At that time, there were only a few publications on the ammonothermal growth of nitride: AlN synthesis by Peters [12], microcrystalline GaN growth in ammonobasic conditions [13–15], and microcrystalline GaN growth in ammonoacidic conditions [16]. As our group had no background in highpressure technology, we explored our own pathway toward bulk GaN growth with limited landmarks. In this chapter, we will trace our research findings until bulk GaN crystals were obtained by the ammonothermal method.

8.2 Impact of Mineralizer on Ammonothermal Synthesis of GaN A mineralizer is an ionic additive to increase solubility of solute in solvothermal growth. The choice of the mineralizer is very important since it determines the chemistry of crystal growth. The type of mineralizers can be categorized into three groups in terms of the resulting acidity of the fluid. Since NH3 dissociates into NHC 4 C and NH 2 , mineralizers that generate NH4 are called acidic mineralizers whereas those which generate NH 2 are called basic mineralizers. Mineralizers which do not  dominantly generate either NHC 4 or NH2 are called neutral mineralizers. One of the major impacts of the mineralizer in the ammonothermal synthesis of GaN is its influence on phase selection. When a basic mineralizer is used, wurtzite GaN is preferentially formed at temperature as low as 600 ıC, whereas mixed phase (i.e., wurtzite and zinc-blende) GaN tends to be synthesized with acidic and neutral mineralizers. Another important factor is its chemical compatibility with autoclave materials. Acidic NH3 severely corrodes steel, stainless steel, and Ni  Cr based superalloys whereas basic and neutral NH3 does not attack most of Ni  Cr based superalloys. To avoid corrosion of the autoclave and related components (e.g., valves

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and pressure transducers) in acidic ammonothermal growth, all wet surface must be protected with Pt-based liner. Synthesis of microcrystalline GaN in autoclaves made of NiCr based alloy with various mineralizers is summarized in Table 8.1. Detailed experimental procedure and method for calculating the ratio of phases are explained elsewhere [22]. When a mineralizer was not added, Ga metal did not react with NH3 . The synthesized powder did not show any definite crystal facets under scanning electron microscope (SEM) observation. Among neutral mineralizers, only KI, NaI, and LiCl yielded an appreciable amount of GaN powder. The powder yield appeared to decrease at higher concentration of KI, although the mechanism of this is unclear. The powder yield with KI was about 1–2 wt% and the powder yield with NaI was slightly higher (3.3 wt%). Among neutral mineralizers in this study, LiCl resulted in the highest yield (16.7 wt%). Considering that KCl and NaCl did not yield GaN, Li seems to enhance GaN formation compared to K and Na. In acidic conditions, powder yield for NH4 Cl; NH4 I, and GaI3 was 6.3 wt%, 3.3 wt%, and 22 wt%, respectively. GaI3 mineralizer resulted in very high powder yield, because GaN was formed not only from the Ga nutrient but also from GaI3 . In the case of NaNH2 (a basic mineralizer), the powder yield was even higher .>20 wt%/. However, in some occasions GaN was not synthesized (Table 8.1). When GaN powder was not formed in basic conditions, the Ga charge became a yellowish powder which dissolved in water with an accompanying emission of NH3 . Judging from this emission, the yellowish powder is thought to be some form of amide, possibly NaGa.NH2 /4 [28, 29]. Therefore, under basic conditions, it is surmised that GaN was synthesized from further ammonolysis of gallium-containing amide. This implies that the reaction path in basic conditions is different from that in neutral and acidic conditions. Figure 8.1 shows X-ray 2  ! scans of the synthesized GaN powders. GaN powders synthesized with NH4 Cl (acidic), NaI (neutral), and KI (neutral) showed diffraction from zinc-blende GaN (zb-GaN) as well as wurtzite GaN (wz-GaN). In contrast, GaN powder synthesized with NaNH2 CNaI (basic) at 525 ı C did not show a clear peak from zb-002 diffraction. For GaN powder synthesized at higher temperature and pressure .600 ıC; 68:3 MPa/ with NaNH2 C NaI, the diffraction pattern showed phase-pure wz-GaN. The quantified zb/wz ratios are shown in Table 8.1. At 525 ı C the zb/wz ratios are about one order of magnitude lower for the basic condition than acidic and neutral conditions. As calculated in [22], the Gibbs free energy of formation for wz-GaN is smaller than that of zb-GaN by small energy difference: only 720 J/mol (8 meV) at 600 K and 1,310 J/mol (14 meV) at 800 K. In acidic and neutral conditions, higher zb/wz ratio at higher formation rate (Table 8.1) indicates that there is a kinetic mechanism which enhances zb-GaN formation at faster reaction. In the case of basic conditions, the ratio of zb-GaN to wz-GaN was about one order of magnitude lower at 525 ıC, and phase-pure wz-GaN was obtained at 600 ıC, at which temperature even vapor-phase epitaxy on c-plane sapphire experience inclusion of zinc-blende phase [30, 31]. These results indicate that the phase selection in the ammonothermal synthesis of GaN is primarily influenced by kinetics related to mineralizer type.

0:5

0:5

0:5

0:5

2:96

3:29

1:1

2:96

NaNH2

NaNH2

NaNH2

NaNH2

600

600

525

525

525

0:5

Basic mineralizers 1:77 NaNH2

525 525 525 525 525 525 525 525 525

525 525 525

3 3 3 3 3 3 3 3 3

525

Temp. (C)

Acidic mineralizers 2:83 NH4 C1 NH4 I 2:83 GaI3 2:83

3

Ga (g)

Neutral mineralizers KCl 2:87 KBr 2:87 KI 2:87 KI 2:87 KI 2:87 NaCl 2:87 NaBr 2:87 NaT 2:87 LiCl 2:87

NH3 (mol)

Without mineralizer 2:87

Mineralizer

68:3

29:3

56:9

53:1

40

34:5 32:8 32:4

32:8 31:7 32:4 32:4 32:8 32:4 32:6 32:4 33:1

33:8

Pressure (MPa)

18

48

24

24

37

7 7 7

7 7 7 7 7 7 7 7 7

19

Time (h)

Process Conditions

1

1

1

1

1

0:7 0:7 0:7

0:7 0:7 0:14 0:7 1:4 0:7 0:7 0:7 0:7

0

%mol of Mineralizer to NH3

Table 8.1 Summary of GaN synthesis with various mineralizers

140

no yield

no yield

no yield

100

190 100 660

No yield No yield 70 50 20 No yield No yield 110 480

No yield

GaN Powder Weight (mg)

28.00%

no yield

no yield

no yield

20.00%

6.33% 3.33% 22.00%

No yield No yield 2.33% 1.67% 0.67% No yield No yield 3.67% 16.00%

No yield

GaN Powder Yield (wt%)

Water soluble yellowish powder Water soluble yellowish powder Water soluble yellowish powder Pure wz-GaN

Mostly wz-GaN

wz, zb-GaN wz, zb-GaN wz. zb-GaN

N/A N/A wz. zb-GaN wz. zb-GaN wz, zb-GaN N/A N/A wz. zb-GaN wz, zb-GaN

N/A

Product

Result

0

0

N/A

N/A

0.66

0.41 0.21 1.14

N/A N/A 0.26 0.21 0.22 N/A N/A 0.33 0.76

N/A

zb/wz Ratio

With comineralizer of 0.1% N With comineralizer of 0.1% N With comineralizer of 0.1% N With comineralizer of 0.1% N With comineralizer of 0.1% N

Ga metal left unreacted

Comments

164 T. Hashimoto and S. Nakamura

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A Pathway Toward Bulk Growth of GaN by the Ammonothermal Method

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Fig. 8.1 X-ray 2  ! scans of GaN powders synthesized with various mineralizers

In ammonothermal growth of GaN, growth temperature is limited by the creep temperature of autoclave materials, which is approximately at 600 ı C. Therefore, the phase selection nature of basic mineralizer is beneficial to avoiding potential phase mixing in grown crystals. In addition, chemical compatibility of basic NH3 with Ni  Cr based superalloys is advantageous over acidic NH3 because the autoclave price becomes enormous when the entire wet surface is covered with a Pt-based liner.

8.3 Solubility of GaN in Ammonobasic Solutions Dissolution of a nutrient into a supercritical fluid is the first step to attain successful crystal growth in solvothermal growth. It is reported that reasonable solubility for hydrothermal crystallization is 2–5 wt% [32]. To achieve high solubility and optimum supersaturation for crystal growth, detailed solubility data are indispensable. Although comprehensive data on solubility of GaN in supercritical NH3 are not available for either basic or acidic conditions, a few measurement data have been presented thus far [17, 20, 23, 33, 34]. As Peters implied in his paper on synthesis of AlN in supercritical ammonobasic solutions [12], GaN shows retrograde solubility in supercritical ammonobasic solutions. At UCSB, we also observed an indication of retrograde solubility in the early stage of growth experiment [20], and confirmed it with a well-controlled weight loss method [35] as shown below. In the weight loss measurement presented in this section, the following points were taken into consideration to minimize measurement errors: (1) sufficient amount

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.5 g/ of polycrystalline GaN was loaded in a small-volume autoclave .22 ml/; (2) the small-volume autoclave was heated in a one-zone furnace to minimize temperature gradient in the autoclave; (3) polycrystalline GaN granules with large grain size (2–3 mm) were used; (4) NH3 was released at elevated temperature followed by rapid air-cooling to avoid precipitation of dissolved GaN on the granules; (5) GaOx and metallic Ga were removed from the surface of polycrystalline GaN. HVPE-grown polycrystalline GaN granules were first baked in an H2 =N2 ambient at 1;000 ı C for 4 h to remove oxides on the crystal surface. After H2 =N2 baking, the GaN granules were dipped in a 50% HCl aq. solution for 3 h to remove residual metals on the crystal surface. Then, the GaN granules were rinsed with DI water and dried in an oven at 120 ı C for more than 8 h. The pretreated GaN granules placed in a Ni-basket and NaNH2 (1.5% mol to NH3 ) were loaded in the autoclave and NH3 was filled. Then, the autoclave was heated for 120 h followed by rapid release of high-pressure NH3 and immediate air-cooling. The GaN granules were rinsed with isopropanol and DI water to remove Na and Na compounds. Before weighing, the GaN granules were dipped in a 50% HCl aq. solution for 3 h, rinsed with DI water, and dried in an oven at 120 ı C for more than 8 h to remove residual Ga and other metallic compounds. The surface of the GaN granules was characterized with SEM and electron-dispersive X-ray analysis (EDX). The total weight of the GaN granules was carefully weighed before and after the process with an electric balance with a resolution of 0.1 mg. The weight loss of the GaN granules was determined by subtracting the final weight from the initial weight. The possible measurement errors are weighing errors of GaN granules due to the resolution of the balance .˙0:2 mg/, errors of the GaN weight due to surface contamination of the granules .C0:5 mg/, and weighing errors of NH3 due to resolution of the balance .˙0:1 mg/. The upper and lower limit of the measurement errors were calculated for all dissolution values or solubility values, and presented as error bars in the figures. Figure 8.2a shows a surface of a GaN granule after H2 =N2 baking observed by SEM and Fig. 8.2b is a corresponding EDX spectrum collected from the entire area of the SEM image. No O was detected by EDX analysis indicating that the surface was effectively reduced by H2 =N2 baking. White particles on the surface were also confirmed to be O-free by EDX. Weak N peak indicates that the surface of the GaN granule was Ga rich, although no evident Ga droplet was observed by SEM. Aluminum was also detected indicating possible physical transport of Al from Al2 O3 reactor tube during H2 =N2 baking. Figure 8.3a shows a surface of a GaN granule after HCl dip observed by SEM and Fig. 8.3b is a corresponding EDX spectrum collected from the entire area of the SEM image. The enhanced N peak and the absence of Al peak in the EDX spectrum indicate effective removal of metallic Ga and Al from the surface. The crater-like holes observed by SEM are probably due to etching of Ga islands on the surface. Baking in H2 =N2 followed by HCl dip is an important pretreatment to minimize measurement error in the weight loss method. A surface treatment of the GaN granules after run was also important to minimize measurement error. Figure 8.4a shows an SEM image of a GaN granule surface before weighing. The granules were rinsed with isopropanol and DI water, dipped in HCl, followed by a rinse with DI water. Although a wide area EDX spectrum

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Fig. 8.2 (a) A surface of a GaN granule after H2 =N2 baking observed with SEM. (b) An EDX spectrum collected from the entire area of the SEM image

(Fig. 8.4b) did not show any other peak than Ga and N, there were many white dots observed on the surface. The focused EDX measurement on the white dot revealed that the dot contained O and Cl (Fig. 8.4c). These white dots were probably Ga  O  Cl complex formed from gallium amide or gallium imide compounds during the surface treatment after run. No evidence of etching on the GaN surface was observed after HCl dip. To estimate solubility of GaN in supercritical ammonobasic solutions, weight loss was measured as a function of time at 450 ıC and 650 ı C. The NH3 fill factors at each temperature were determined so that the pressure at high temperature became 76 ˙ 12 MPa .11;000 ˙ 1;800 psi/. The wide pressure allowance was due to pressure increase or a minor NH3 leak during the run. As shown in Fig. 8.5, weight loss saturated after about 120 h at 450 ıC. This means that the supercritical ammonobasic solution was saturated with GaN after 120 h. From the amount of loss, the solubility of GaN in this solution was estimated to be 0.11 wt%. On the other hand, the saturation occurred in about 26 h at 650 ı C. This means that a kinetic process in dissolving was enhanced at higher temperature. Based on this experiment, the duration of 1 run was determined to be 120 h (5 days) for the rest of the experiments. Figure 8.6 shows solubility of GaN in supercritical ammonobasic solution as a function of temperature. Retrograde solubility was observed at temperature higher than 600 ıC in this experiment.

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Fig. 8.3 (a) A surface of a GaN granule after HCl dip observed with SEM. (b) An EDX spectrum collected from the entire area of the SEM image

There are three possible reasons for retrograde solubility of GaN in supercritical ammonobasic solutions. The first possible reason is dissociation of NH3 at high temperature. The solubility decreased at higher temperature range .600–650 ı C/ at which NH3 might start to dissociate into N2 and H2 [36]. Although the equilibrium temperature and pressure for the dissociation of supercritical NH3 cannot be calculated due to lack of thermodynamic data at high temperature and high pressure, pressure increase observed during the typical ammonothermal growth indicates possible dissociation of NH3 . If NH3 dissociates into N2 and H2 , the total amount of NH3 to dissolve GaN decreases, resulting in underestimation of solubility. The second possible reason is formation of gallium amide or gallium imide compounds. As shown in the previous section, formation of gallium amide or gallium imide compounds is more favored than formation of GaN at temperature below 600 ıC. This means that more of GaN might be transformed into gallium amide or gallium imide at lower temperature, resulting in higher solubility of GaN at lower temperature. The third possible reason is the effect of entropy in dissolution process. The standard free energy of solution, G is expressed as follows: G D H  T S; where H and S are changes in enthalpy and entropy associated with dissolution, and T is temperature. In dissolution of ionic salts, an entropy contribution to the free

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A Pathway Toward Bulk Growth of GaN by the Ammonothermal Method

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Fig. 8.4 (a) An SEM image of a GaN surface after run. (b) An EDX spectrum collected from the entire area of the SEM image. (c) An EDX spectrum collected from the white dot observed in the SEM image

energy cannot be neglected since the enthalpy changes associated with dissolution process are usually comparatively small [37]. It is reported that the entropy of solution is exceedingly negative when the anion and cation that would be formed on dissociation are both triply charged [37]. This might be the case in GaN dissolution, i.e., the entropy of GaN dissolution into NH3 is negative, resulting in higher free energy at higher temperature. The situation where entropy becomes negative on dissolution can be explained as follows: If GaN generates Ga3C on dissolution, three 3C NH 3NH 2 molecules are attracted, forming an ordered structure of Ga 2 complexes. Therefore, the randomness of the NH3  GaN system without Ga3C 3NH 2 complexes becomes higher than that with Ga3C 3NH 2 complexes.

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DISSOLUTION (wt%)

0.300%

NaNH2 1.5%, 450°C NaNH2 1.5%, 650°C

0.200%

0.100%

0.000%

0

100 200 TIME (hour)

300

Fig. 8.5 Weight loss of GaN granules as a function of time. The pressure was 76 ˙ 12 MPa .11;000 ˙ 1;800 psi/ and the mineralizer was 1.5%mol NaNH2

Fig. 8.6 Solubility of GaN in supercritical ammonobasic solutions as a function of temperature. The pressure was 76 ˙ 12 MPa .11;000 ˙ 1;800 psi/ and the mineralizer was 1.5%mol NaNH2

Currently, there is no evidence determining which of the three factors has a major influence on the retrograde solubility of GaN. Further detailed study must be carried out to answer this question.

8.4 Seeded Growth of GaN with Metallic Ga Nutrient To apply the ammonothermal method to bulk GaN growth, it is necessary to grow GaN on seed crystals via fluid transport. As shown in the previous section, the solubility of GaN in supercritical NH3 with 1.5 mol% NaNH2 was at most 1.2 wt%. In addition, it took a few days to saturate GaN in supercritical NH3 . Given this low solubility and slow kinetic process, the growth rate of GaN using GaN nutrient was expected to be extremely low for low mineralizer concentration. In the early stage of experiment at UCSB, there was a limitation in the mineralizer concentration due to clogging of high-pressure tubing. Therefore, growth of GaN using metallic

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Ga nutrient was attempted first. In this section, successful growth of GaN on HVPE-grown GaN seed via fluid transport of metallic Ga nutrient is reviewed. Seeded growth was carried out in an autoclave made of Ni  Cr-based superalloy. The inner room of the autoclave was divided into two regions with a baffle. Considering the retrograde solubility, the Ga nutrient was placed at a lower-temperature zone (an upper zone) and single-crystalline GaN seed crystals were placed at a higher-temperature zone (a lower zone). 1.5 mol% NaNH2 was added to NH3 as a mineralizer. The NH3 pressure was about 150–200 MPa. The surface morphology and the thickness of the GaN films grown on the seeds were characterized by SEM. Selected single-crystal samples were characterized by Philips high-resolution fouraxis X-ray diffractometer using the Cu K’ line monochromated with a four-crystal Ge (002). The X-ray tube was operated at 45 keV and 40 mA. The sample was evaluated by measuring both on-axis and off-axis rocking curves. The line-width of the 0002 reflection was used for evaluating dislocations with a screw component N reflection was used to evaluate dislocations with an and the line-width of the 2021 edge component. The microstructure of the films was characterized by transmission electron microscopy (TEM: FEI Tecnai 20) operated at 200 kV. TEM samples were prepared for observation nearly along the <10 1N 0> axis by mechanical polishing followed by ArC ion milling (Gatan PIPS). For determination of polarity, convergent beam electron diffraction (CBED) was observed when the samples were characterized by TEM. The experimental CBED patterns were compared with patterns simulated by the electron microscopy simulation (EMS) software package with 33 zero-order reflections [38]. All TEM and CBED characterizations were carried out by Dr. Feng Wu. Figure 8.7a, b are cross-section of GaN films grown with 100 g of Ga nutrient on Ga-face and N-face, respectively. The temperatures measured inside the autoclave in the seed zone and the nutrient zone were 550–565 ı C and 520–530 ı C, respectively. After 19-day growth, 17 m-thick GaN and 45 m-thick GaN were grown on the Ga-face and the N-face, respectively. The resulting growth rates were 0:9 m/day for the Ga-face and 2:4 m/day for the N-face. As explained below, the actual growth happened within approximately 1 day. The film on the Ga-face consisted of hexagonal columns as shown in Fig. 8.7c, whereas the N polar surface was featureless as shown in Fig. 8.7d. A cross-sectional TEM observation of the film grown on the N-face revealed numerous voids and defects at the interface between the seed and the film. However, at the surface of the film no void was observed and defects were greatly reduced. The TEM images taken with g D 000 2N and g D 1 1N 00 showed that the majority of the dislocations were mixed-character dislocations. It was also confirmed that there was no evidence of basal-plane stacking faults in the surface region. The dislocation density estimated from the cross-sectional TEM images was in the low-109 cm2 level. The TEM images of this sample can be found in [39]. Due to the voids at the interface, the grown film was easily separated from the seed crystal. The free-standing GaN platelet separated from the N-polar surface of the seed crystal was characterized by X-ray rocking curve (XRC). The FWHM of XRC from 0002 reflection and 20 2N 1 reflection were 3,400 arcsec and

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Fig. 8.7 (a, b) Cross-sectional SEM images, and (c, d) plan-view SEM images of the GaN films grown on a GaN seed crystal with 100 g of Ga. (a, c) the Ga-side, (b, d) the N-side

6,300 arcsec, respectively. The high value of FWHM is considered to originate from poor structural quality near the interface between the film and the seed crystal. Hall measurement was carried out on the free-standing GaN platelet. The electron con1 centration and mobility were 5:4  1019 cm3 and 73 cm2 Vs , respectively. The high carrier concentration is probably due to high impurity concentrations and high defect density. One of the greatest advantages of solvothermal growth is its excellent uniformity. Since the convective flow of the solvent during the growth is typically high enough to attain low distribution of temperature within an independent region (i.e., a nutrient region or a seed region), growth uniformity is excellent compared to other growth methods. Growth shown in Figs. 8.8 and 8.9 demonstrate the excellent uniformity of GaN film grown on 3 4 cm2 oval shaped seed. The growth temperature was 675 ıC and the nutrient temperature was 625 ı C. Note that these temperatures are values measured on the external wall of the autoclave. The temperature inside the autoclave was estimated to be about 50–100 ıC lower than these values. Growth was carried out for 1 day with 40 g of Ga nutrient. The surfaces of the grown films on the Ga-face and N-face are shown at the top left corner of the Figs. 8.8 and 8.9, respectively. The grown films showed yellowish color and the Ga-face was relatively rough whereas the N-face showed a specular surface. The surface morphology at various points

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Fig. 8.8 A photograph of GaN film grown on a large-sized GaN seed crystal (Ga-face) and surface morphology at various points observed with SEM

Fig. 8.9 A photograph of GaN film grown on a large-sized GaN seed crystal (N-face) and surface morphology at various points observed with SEM

on the Ga-polar surface was observed with SEM as shown in Fig. 8.8. The Ga-polar surface was filled with numerous pits whose sizes ranged from submicron to 10 m. However, the surface morphology was uniform over the entire growth area, indicating the excellent scalability of the ammonothermal method. Similarly, the surface morphology at various points on the N-polar surface was observed with SEM as shown in Fig. 8.9. The morphology of N-polar surface was featureless except for scattered dots of GaN. The N-face was also very uniform over the entire growth area. This growth demonstrated excellent uniformity of ammonothermal growth. Despite these achievements, there was a critical problem in the growth with a metallic Ga nutrient. Figure 8.10 shows the relationship between the thickness of the film and the growth time when metallic Ga was used as a nutrient. As clearly shown in the figure, the thickness did not increase after 12 h. We have achieved maximum growth rate of 88 m per day for optimized 12-h growth; however, the growth did not continue for extended time. This was because the metallic Ga nutrient was

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Fig. 8.10 Thickness of grown films as a function of growth time in the case of a metallic Ga nutrient

transformed into GaN in the nutrient crucible during growth. Once the metallic Ga is converted to GaN, the dissolution of Ga-containing species into NH3 decreased significantly and GaN ceased to grow. Another problem associated with metallic Ga nutrient is that GaN film started to grow at very low temperature .200–300 ıC/ due to high reactivity of metallic Ga in the supercritical ammonobasic solution resulting in poor growth interface. It is challenging to control the growth initiation in the closed system of ammonothermal growth. Due to these critical issues, we decided to switch to polycrystalline GaN nutrient.

8.5 Seeded Growth of GaN with Polycrystalline GaN Nutrient In this section, seeded growth of GaN by the ammonothermal method with polycrystalline GaN nutrient is reviewed. As explained in the previous section, GaN growth from Ga nutrient continued only for about 12 h. By utilizing polycrystalline GaN nutrient with increased mineralizer concentration (3.9–4.2 mol% of NaNH2 ), continuous growth up to 90 days was achieved as shown in Fig. 8.11. Improved microstructure in the crystal grown for 50 days was confirmed with TEM as shown in Fig. 8.12a, c. The cross-section samples were fabricated with focused ion beam (FIB). It was confirmed with an optical microscope that the location of the FIB hole (i.e., the place where the cross-sectional sample was obtained) was on the growth interface as shown in Fig. 8.12b. Cross-sectional TEM images were observed under bright field condition with g D 11 2N 0 and g D 0002. Figure 8.12a shows cross-sectional TEM micrographs at the interface on the Ga-polar

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Fig. 8.11 Thickness of grown crystal along Cc, c and a direction as a function of growth duration

side. The interface could not be identified and no dislocation was observed. Figure 8.12c shows cross-sectional TEM micrographs on the N-polar side. Unlike the Ga-polar side, a clear interface was observed on the N-polar side. One void (indicated as “V” in the image) and one mixed-character dislocation propagating horizontally (indicated as “D” in the image) were identified. To estimate the threading dislocation density (TDD), plan-view TEM images were observed under bright field condition with g D 11 2N 0. There were no dislocations observed on the Ga-face and only several dislocations were observed on the N-face. The TDD on the Ga-face was estimated to be less than 1  106 cm2 and TDD on the N-face was estimated to be 1  107 cm2 . Despite the excellent microstructure, the grown crystals contained multiple grains, which deteriorated FWHM of XRC as summarized in Table 8.2. The FWHMs of XRC from the Ga-polar surface of the 40-m crystal (50-day growth) were approximately 290 arcsec and 110 arcsec for 0002 and 20 2N 1 reflections, respectively. Although these numbers are comparable to the FWHMs for the seed crystals, the additional 90-day growth (total thickness of 160 m) resulted in higher FWHM values of 1,400 arcsec and 650 arcsec for 0002 and 20 2N 1 reflections, respectively. Similarly, the FWHMs of XRC from the N-polar surface of the 180-m crystal (50-day growth) were approximately 840 arcsec and 490 arcsec for 0002 and 20 2N 1 reflections, respectively and the additional 90-day growth (total thickness of 430 m) resulted in FWHMs of 1,400 arcsec and 570 arcsec for 0002 and 20 2N 1 reflections, respectively. It is surmised that GaN crystal grew under strain on HVPEgrown GaN seed, of which the crystal lattice is bowed and after reaching certain thickness the stress is relaxed resulting in increase in FWHM. Figure 8.13 is a direct observation of the grain boundary on Ga-polar surface with Nomarski micrograph after selective etching in a hot .160 ı C/ H3 PO4 solution. The detail of the etching is explained elsewhere [26, 27]. As indicated with an arrow, the grain boundary was clearly observed on the surface. The dimension of each grain was in the range of 500 m.

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Fig. 8.12 (a) Cross-sectional TEM images of the interface on the Ga-face with g D 0002 and g D 112N 0. No dislocation was observed. The interface was not identified with TEM. Note that the horizontal line is not an interface but damage during sample preparation. (b) Comparison of cross-sectional surface before and after the FIB process observed with a Nomarski microscope. It was confirmed that the “FIB holes” lie on the interface for both N-polar side and Ga-polar side. (c) Cross-sectional TEM images of the interface on the N-face with g D 0002N and g D 112N 0. One void (indicated as “V”) and one dislocation propagating horizontally (indicated as “D”) were observed. The interface was clearly identified as indicated with lines in the picture

Preliminary SIMS measurement was carried out for the crystal grown for 50 days. The crystal on the Ga-polar side contained 2  1019 cm3 of oxygen and 1  1018 cm3 of carbon. The crystal on the N-polar side contained 6  1019 cm3 of oxygen and 4  1018 cm3 of carbon. In addition, other impurities such as Li, Na,

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Table 8.2 FWHMs of XRC obtained from GaN crystals grown by the ammonothermal method with a polycrystalline GaN nutrient Sample 0002 reflection (arcsec) 202N 1 reflection (arcsec) Ga-face HVPE-grown seed crystal 280 90 40 m (50 days) 290 110 160 m (90-day regrowth 1;400 650 on 50-day crystal) N-face HVPE grown seed crystal 360 100 180 m (50 days) 840 490 430 m (90-day regrowth 1;400 570 on 50-day crystal)

Fig. 8.13 A Nomarski micrograph of the Ga-face after selective etching. The arrow indicates a grain boundary

Mg, Al, Ni, Cr, and Mn were detected. Reduction of impurities will be one of the future challenges in the ammonothermal method.

8.6 Growth of Bulk GaN Crystals and Sliced Wafers GaN in bulk shape has been obtained in 82-day growth with increased growth rate. Figure 8.14 shows bulk GaN crystals. The crystals showed dark brownish color. We surmise that the origin of the color center is a combination of impurities and point defects. The crystals were in polyhedron shape with a flat hexagonal bottom and sidewalls. The flat hexagonal surface was c plane and the sidewalls were m planes. The Ga-polar surface was decorated with semipolar facets. The height (i.e., thickness) of the crystal was 4.5 mm and 5.1 mm for the crystals in Fig. 8.14a (Sample A), b (Sample B), respectively. The grown thickness along c directions, which represents the addition of grown thicknesses along Cc and c directions was

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Fig. 8.14 Bulk GaN crystals grown by the ammonothermal method. (a) Sample A, (b) Sample B

approximately 4.1 mm and 4.6 mm for Sample A and Sample B, respectively. The lateral dimension of Sample A and Sample B was approximately 5.7 mm (pointto-point) and 7 mm (the longest side), respectively. Since the polycrystalline GaN nutrient had been completely consumed when the autoclave was open, we suspect the actual growth duration was less than 82 days. Therefore, the average growth rate along the c direction is estimated to be more than 50 m per day for Sample A and 55 m per day for Sample B. Since Sample B was placed at a lower position than Sample A in the autoclave, where the temperature was higher, the growth rate was higher due to larger temperature difference. One crystal was sliced to obtain 400-m-thick c plane wafers as shown in Fig. 8.15. The picture on the left in Fig. 8.15 shows the side view of the crystal before slicing. The parallel lines and the rectangle represent the approximate wire positions and the seed position, respectively. As clearly seen in Fig. 8.15, the growth nature was strongly dependent on growth directions. The wafer #1 has a relatively transparent rectangular region of the seed and laterally grown regions. The growth rate along a directions was much faster than that along m directions. The dimensions of the grown crystal along a directions and m directions were 1.5 and 0.3–0.5 mm, respectively. As seen in the wafer #2, the crystal grown along c direction was

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Fig. 8.15 Photographs of c plane wafers. The left picture is a crystal before slicing and right pictures are sliced wafers. The parallel lines represent the approximate wire positions and the rectangle represents the approximate seed position. The sliced wafers were labeled from top to bottom as #1 to #4

Fig. 8.16 Photographs of m plane wafers. The left picture is a crystal before slicing and right pictures are sliced wafers. The parallel lines represent the approximate wire positions and the trapezoid represents the approximate seed position. The arrows on the wafers represent the lateral interface between the seed and the grown crystal. The sliced wafers were labeled from top to bottom as #1 to #7

significantly colored. Also, one can see that the dark region grew laterally as crystal grew along c direction, indicating the growth along c direction was dominant over domains growing along a or m directions. Another crystal was sliced to obtain m plane wafers as shown in Fig. 8.16. The picture on the left in Fig. 8.16 shows the top view of the crystal before slicing. The

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Fig. 8.17 XRD rocking curves from 10 1N 0 reflection of the m plane wafers shown in Fig. 8.16. The wafer numbers correspond with those in Fig. 8.16

parallel lines and the trapezoid represent the approximate wire positions and the seed position, respectively. The arrows in the pictures on the right represent the lateral boundary between the seed platelet and the grown crystal. Again, one can clearly see growth anisotropy: the growth along Cc direction shows more grain structure with f10 1N 1g facets, the growth along c direction shows the darkest color and the growth along a direction shows the lightest color merging with domains growing along Cc and c directions. Since the wafers #1 and #7 consist of laterally grown domains, they showed better structural quality as shown later. The grown thickness along Cc and c directions was 2.5 and 2.1 mm, respectively. Figure 8.17 shows XRC from 10 1N 0 reflection of the m plane wafers shown in Fig. 8.16. The XRC indicates the multiple grain structure. As discussed above, the wafer #1 which predominantly consists of a laterally grown domain showed the narrowest FWHM of 220 arcsec. The wafer #7 also showed relatively narrow FWHM of 420 arcsec, although the rocking curve is not shown in Fig. 8.17. This implies that the degradation of the crystal quality arises from insufficient structural quality of the HVPE-grown seed crystals.

8.7 Summary Starting from powder synthesis, we have developed our own technology of ammonothermal growth at UCSB. Although there still remain a lot of uninvestigated matters in this method, the bulk crystals with three-dimensional polyhedron shape in appreciable size demonstrate the high feasibility of ammonothermal growth. The obtained growth rate was in the practical range for commercialization of ammonothermal growth. At this moment, multiple grains and coloration are the major challenges in our research; however, these problems will be solved in the near future by

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engineering. We believe that the ammonothermal growth is the enabling method to realize our “dream” bulk crystals of GaN. Acknowledgements This research was conducted as a part of Nakamura Inhomogeneous Crystal Project of Exploratory Research for Advanced Technology (ERATO) program funded by Japan Science and Technology Agency (JST). The authors appreciate the long-term and flexible funding of JST. The authors thank Professor James S. Speck and Dr. Feng Wu at UCSB for TEM observation and fruitful discussions. Also, the authors thank Mr. Kenji Fujito and Mr. Makoto Saito at Mitsubishi Chemical Corporation for participating in the project as visiting researchers. In addition, the authors acknowledge Mitsubishi Chemical Corporation for supplying HVPE-grown GaN platelets. The authors are also thankful to Takatori Corporation and Yasunaga Wire Saw Systems Corporation for slicing the crystals with wire saws.

References 1. R.W. Brander, R.P. Sutton, J. Phys. D (Appl. Phys.) 2, 309 (1969) 2. M.A. Haase, J. Qiu, J.M. DePuydt, H.C. Cheng, Appl. Phys. Lett. 59, 1272 (1991) 3. S. Nakamura, T. Mukai, M. Senoh, Jpn. J. Appl. Phys. Part I 32, L16 (1993) 4. S. Nakamura, M. Senoh, S. Nagahama, N. Iwasa, T. Yamada, T. Matsushita, H. Kiyoku, Y. Sugimoto, Jpn. J. Appl. Phys. Part II 35, L74 (1996) 5. S. Porowski, MRS Internet J. Nitride Semicond. Res. 4S1, G1.3 (1999) 6. T. Inoue, Y. Seki, O. Oda, S. Kurai, Y. Yamada, T. Taguchi, Phys. Stat. Sol. (B) 223, 15 (2001) 7. H. Yamane, M. Shimada, T. Sekiguchi, F.J. DiSalvo, J. Cryst. Growth 186, 8 (1998). 8. F. Kawamura, M. Morishita, K. Omae, M. Yoshimura, Y. Mori, T. Sasaki, Jpn. J. Appl. Phys. 42, L879 (2003) 9. T. Hashimoto, F. Wu, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 46, L889 (2007) 10. M.C. Schmidt, K.C. Kim, H. Sato, N. Fellows, H. Masui, S. Nakamura, S. DenBaars, J.S. Speck, Jpn. J. Appl. Phys. 46, L126 (2007) 11. A. Tyagi, H. Zhong, N.N. Fellows, M. Iza, J.S. Speck, S.P. DenBaars, S. Nakamura, Jpn. J. Appl. Phys. 46, L129 (2007) 12. D. Peters, J. Cryst. Growth 104, 411 (1990) 13. R. Dwilinski, A. Wysmolek, J. Baranowski, M. Kaminska, R. Doradzinski, J. Garczynski, L. Sierzputowski, H. Jacobs, Acta Phys. Polonica A 88, 833 (1995) 14. R. Dwilinski, R. Doradzinski, J. Garczynski, L. Sierzputowski, M. Palczewska, A. Wysmolek, M. Kaminska, MRS Internet J. Nitride Semicond. Res. 3, 25 (1998) 15. D.R. Ketchum, J.W. Kolis, J. Cryst. Growth 222, 431 (2001) 16. A.P. Purdy, R.J. Jouet, C.F. George, Cryst. Growth Des. 2, 141 (2002) 17. R.T. Dwilinski, R.M. Doradzinski, J. Garczynski, L.P. Sierzputowski, Y. Kanbara, U. S. Patent 6656615 B2 (2003) 18. M.J. Callahan, B. Wang, L.O. Bouthillette, S.Q. Wang, J.W. Kolis, D.F. Bliss, Mater. Res. Soc. Symp. Proc. 798, Y2.10 (2004) 19. A. Yoshikawa, E. Ohshima, T. Fukuda, H. Tsuji, K. Oshima, J. Cryst. Growth 260, 67 (2004) 20. T. Hashimoto, K. Fujito, B.A. Haskell, P.T. Fini, J.S. Speck, S. Nakamura, J. Cryst. Growth 275, e525 (2005) 21. T. Hashimoto, K. Fujito, M. Saito, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 44, L1570 (2005) 22. T. Hashimoto, K. Fujito, R. Sharma, E.R. Letts, P.T. Fini, J.S. Speck, S. Nakamura, J. Cryst. Growth 291, 100 (2006) 23. M. Callahan, B.G. Wang, K. Rakes, D. Bliss, L. Bouthillette, M. Suscavage, S.Q. Wang, J. Mater. Sci. 41, 1399 (2006)

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24. B. Wang, M.J. Callahan, K.D. Takes, L.O. Bouthillette, S.-Q. Wang, D.F. Bliss, J.W. Kolis, J. Cryst. Growth 287, 376 (2006) 25. Y. Kagamitani, D. Ehrentraut, A. Yoshikawa, N. Hoshino, T. Fukuda, S. Kawabata, K. Inaba, Jpn. J. Appl. Phys. 45, 4018 (2006) 26. T. Hashimoto, F. Wu, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 46, L525 (2007) 27. T. Hashimoto, F. Wu, J.S. Speck, S. Nakamura, Nat. Mater. 6, 568 (2007) 28. P.P. Molinie, R. Brec, J. Rouxel, P. Herpin, Acta Crystallogr. B29, 925 (1973) 29. H. Jacobs, B. Nocker, Zeitschrift für Anorganische und Allgemeine Chemie 619, 381 (1993) 30. X.H. Wu, D. Kapolnek, E.J. Tarsa, B. Heying, S. Keller, B.P. Keller, U.K. Mishra, S.P. DenBaars, J.S. Speck, Appl. Phys. Lett. 68, 1371 (1996) 31. A. Munkholm, C. Thompson, C.M. Foster, J.A. Eastmen, O. Auciello, G.B. Stephenson, P. Fini, S.P. DenBaars, J.S. Speck, Appl. Phys. Lett. 72, 2972 (1998) 32. A.A. Ballman, R.A. Laudise, Hydrothermal Growth, The Art and Science of Growing Crystals, ed. by J.J. Gilman (Wiley, New York 1963), p. 231 33. T. Hashimoto, M. Saito, K. Fujito, F. Wu, J.S. Speck, S. Nakamura, J. Cryst. Growth 310, 311 (2008), note that the non-polar plane is mislabeled as m-plane which should be corrected as a-plane 34. D. Ehrentraut, Y. Kagamitani, C. Yokoyama, T. Fukuda, J. Cryst. Growth 310, 891 (2008) 35. K. Byrappa, M. Yoshimura, Handbook of Hydrothermal Technology, Chap. 4 (Noyes, London, 2001) 36. H. Jacob, S. Schmidt, High-pressure ammonolysis in solid-state chemistry, Current Topics in Materials Science, Vol. 8, Chap.5, ed. by E. Kaldis (North-Holland, Amsterdam, 1982) 37. D.A. Johnson, Some thermodynamic aspects of inorganic chemistry, 2nd edn., Chap. 5 (Cambridge University Press, Cambridge, 1968, 1982) 38. P.A. Stadelmann, Ultramicroscopy 21, 131 (1987) 39. T. Hashimoto, K. Fujito, F. Wu, B. A. Haskell, P. T. Fini, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 44, L797 (2005)

Chapter 9

Acidic Ammonothermal Growth Technology for GaN Dirk Ehrentraut and Yuji Kagamitani

Abstract The ammonothermal technique for the crystal growth of bulk GaN is presented. General aspects are discussed and the use of acidic mineralizer is focused on. This method is a recent development to challenge the problem of latticematched substrate crystal for future group-III nitride device technology. Chemical and physical properties are discussed and an outlook to the future prospective and technological requirements is given.

9.1 Introduction The accessibility of lattice and thermally matched substrates is typically a great benefit in the design and fabrication of optimized device structures providing high efficiency and lifetime [1]. The hexagonal phase of the wide bandgap material gallium nitride (GaN; space group P63 mc, with lattice parameters a D 3:189 Å and c D 5:186 Å), to which we refer as GaN unless otherwise mentioned, is of significant importance for numerous applications in the big market of electronics and optoelectronics because of its excellent properties (Eg D 3:4 eV, breakdown 1 voltage 3 MV cm1 , electron mobility 1; 000 cm2 V s1 , and high saturation velocity); see, for example, Chap. 1 in this book. It is widely recognized that providing high-quality GaN substrates at a reasonable price is the key toward further development of next-generation electronic and optoelectronic devices. Different approaches have been chosen in the extensively employed vapor phase technologies like hydride vapor phase epitaxy (HVPE) or molecular beam epitaxy (MBE) to lessen the problem of lacking lattice and thermally matched substrates by some means. Templated single crystalline ’-Al2 O3 (sapphire) substrate of different crystallographic orientations (polar, nonpolar) is the most prominent one realized to grow GaN crystals of about 10 mm thickness along the h0001i axis [2]. Other mateN 6H-SiC [4, 5], (111) and rials of choice involve (111)GaAs [3], (0001) and .0001/ N (001)Si [6], .100/”-LiAlO2 [7], (0001) and .0001/ ZnO [8, 9] to mention but a few. The GaN crystals grown on these substrates yet contain a large amount of crystal defects like grain boundaries, threading dislocations, stacking fault, etc. Whatever 183

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be the foreign substrate material, it is the thermal and lattice mismatch between substrate and GaN that causes the formation of crystal defects in high densities, strain and subsequently bowing of the final product epi-ready GaN wafer. What is more, the substrate must often be removed if realization of a transparent and electrically conductive back contact is considered. Free-standing GaN wafers of two inches in size (minimum requirement) and low defect concentration around 104 cm2 are in vital demand in the industry. The answer to the problem could lie in the growth of GaN by a process from liquid solution. Growth from the solution is generally attractive in that grown crystals are of high structural perfection and high throughput can be achieved. Quartz .’-SiO2 / and more recently ZnO have paved the way for the solvothermal growth technique [10]. More specifically, the solvothermal growth technique basically employs a polar solvent to dissolve and successively form metastable products with the solute. Later, after transport to the growth zone, under slightly different temperature conditions, crystallizes the desired phase of GaN on a seed of similar crystal structure. Mineralizers are essential to amplify the solubility of the solute in the solvent. Since a closed system is utilized, i.e., exchange of matter with the ambient is impossible, the solvent takes over a supercritical (SC) state, which also improves the solubility of the nutrient. Basically, acidic and basic mineralizers are employed in the ammonothermal growth of GaN. Whereas the latter are known for its retrograde solubility (also compare Chaps. 7 and 8 of this book), acidic mineralizers can be utilized under classic conditions, i.e., crystal growth occurs in the cooler zone of an autoclave or upon lowering system temperature. The solvothermal crystal growth technology, if compared to the vapor phase growth technologies, is characterized by the following merits: (1) the operation near the thermodynamic equilibrium for a given temperature–pressure range. This leads to generate an extremely low supersaturation in the solution and consequently a high crystallinity can be expected. (2) The temperature gradient at the growing interface is practically zero which results in stable growth conditions. (3) Scalability: large crystal quantities can be controlled over long process time, thus enabling a high throughput. Growth processes employing multiple seeds (around 2,000 quantity) are applied for the low temperature phase of quartz, ’-SiO2 . (4) Environmentally benign conditions for production and capability for complete recycling of the used solution. (5) This technology is used for a long time, i.e., >60 years to mass-produce crystals like ’-SiO2 . Like with most of the regularly used technologies, there are certain limitations in that the incorporation of constituents or reaction products from the solution on a mesoscopic scale may occur if the temperature inside the growth vessel is not controlled precisely. The homogeneous incorporation of dopants over long growth cycles may be hard to manage since the reaction rate of the host crystal to be grown and the dopant under similar conditions usually differ. This effect will result in graded crystals, i.e., changes in the concentration of the dopant along the growth direction of a crystal facet. It should be known here that the process performed at the vicinity to the thermal equilibrium does not allow the growth of thermodynamically unstable solid solutions. This in turn limits the fabrication of lattice-matched

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alloyed substrates, but is a general feature for growth techniques from the liquid solution. However, the main purpose for solvothermal growth technologies lies in the fabrication of crystals of superior crystallinity to slice substrates for use in the following step of device fabrication, i.e., the main purpose is to provide lattice-matched, low-strain substrates of large size at reasonable production cost. Almost naturally, a high-throughput technique like the hydrothermal (ammonothermal) growth employing a solution indeed holds great promise to satisfy the market demand. As side note, an isostructural compound to GaN is the hexagonal zinc oxide (ZnO; space group P63 mc, with lattice parameters a D 3:253 Å and c D 5:213 Å), which is the first semiconductor crystal being fabricated by the hydrothermal technique: 3 in. size wafers have already been demonstrated [10]. However, the situation in the ZnO device technology is opposed to that in the GaN device technology such that ZnO wafers of high crystallinity are available but the device technology seemingly needs more progress before ZnO-based light emitters (LEDs) make their way to the market. Contrary, in case of GaN the device technology is already a multibillion US dollar business but high-quality and cheap GaN substrates are not yet available. This chapter comprehensively reports on the growth of GaN single crystals by the acidic ammonothermal path, including detailed analysis of important growthrelated parameters like solubility and chemistry of the solution, and will start with the historical developments from the early beginning.

9.2 Brief History of the Ammonothermal Growth Technique of GaN The ammonothermal technique is the most recent among the methods to produce bulk GaN crystals for future wafer fabrication as laid out in the book. The HVPE technique is currently the most widely used one (Chaps. 2–5). Other techniques with potential comprise the high-pressure solution (HPS, Chap. 10) growth and the Na-flux method (Chap. 11). Early experiments using SC ammonia have been focused on the materials synthesis, specifically of complex nitrides [11]. In the 1990s, nano- and microcrystals of AlN and GaN have been synthesized by the ammonothermal approach [12–15] and the potential of the method as a prospective candidate for bulk GaN crystal growth had been realized. It was, however, not until around the year 2000 that free-standing millimetersized GaN crystals have been grown by the ammonothermal technique: Ketchum and Kolis have grown 0:5  0:2  0:1 mm3 crystals in a basic ammonothermal environment over 10-day cycles at 400 ı C and 240 MPa [16]; Callahan et al. have used polycrystalline and HVPE-GaN as seed crystals to grow thick films using basic mineralizer and system pressure around 100–300 MPa [17]; Wang et al. have reported in 2006 the growth of GaN crystals up to 10  10  1 mm3 at growth rates 50 m

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Fig. 9.1 First 2-in. size GaN produced by the acidic ammonothermal route grown from the autoclave shown in Fig. 9.4. HVPE GaN crystals served as seeds

per day [18]. Dwilinski et al. and Hashimoto et al. have reported their huge efforts over the past years in the basic ammonothermal technology; for details see Chaps. 7 and 8. Acidic mineralizers have been employed recently by Purdy et al., Yoshikawa et al., and Kagamitani and Ehrentraut et al. to synthesize free-standing hexagonal GaN crystals and homoepitaxial films [19–21]. More recently, the first successful trial to use a 2-in. size HVPE-GaN substrate to produce ammonothermally grown GaN was published [22]. Figure 9.1 shows a roughly 2-in. size GaN crystal, the first ever grown under acidic ammonothermal conditions. The acidic ammonothermal technique is by far the most juvenile, though swiftly developing, technology with substantial prospect to fabricate bulk size GaN crystals.

9.3 Growth Technology A closed system, so-called autoclave, is employed rugged enough to bear high pressure under elevated temperature. Important is the choice of a proper alloy for the autoclave to withstand the chemically harsh conditions over long term. In the case of the hydrothermal method, SC water .H2 O/ is used as solvent and Fe-based alloys are mostly chosen as for the autoclave material. Fe-based alloys have excellent mechanical properties; however, owing to corrosion they cannot be used under the conditions of SC NH3 C acidic mineralizers. Therefore,

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corrosion-resistant Ni-based alloys (alloy 625, Rene 41, Nimonic 90, etc.) may be used as material of choice for the main body of the autoclave. Even for the aforementioned Ni-based alloys, though, little corrosion may occur depending on the acidity of the solution. Autoclaves for acidic ammonothermal growth processes are therefore equipped with an inner liner made of chemically inert platinum (Pt) to fully prevent the corrosion of the autoclave and successively incorporation of impurities into the growing crystal. The baffle plate, wire to fix seed crystals, and the basket to contain the precursor are all made of Pt. Figure 9.2 reveals a typical autoclave setup. A baffle serves to separate the volume into dissolution zone with higher temperature and growth zone with slightly lower temperature. Ga metal and/or polycrystalline GaN or other Ga-contained materials can be used as precursor. Gallium metal (Ga; 99.9999% purity) had been used initially as the precursor and was gradually replaced by polycrystalline GaN. The latter is a by-product obtained from the synthesis of hexagonal GaN by HVPE. The ratio of the molar concentration of mineralizer and precursor, calculated to .1  x/ GaN C x Ga, was up to 0.15. NH4 X (X D Cl, Br, I) is added as mineralizer. The mineralizers are of 99.999% phase purity and have been used in the molar ratio of mineralizer and ammonia in the range  102 . GaN single crystals are prepared as seed. Precursors and mineralizes are placed in the high-temperature zone and the seed crystals are mounted in the low-temperature zone. The autoclave is then sealed and thereafter vacuum pumped and successively flooded with nitrogen to remove oxygen. Generally, several methods maybe used to fill the autoclave with ammonia. In the high-pressure injection system, ammonia is injected into the vessel by a

Fig. 9.2 Photograph and sketch of an autoclave for the acidic ammonothermal growth of GaN at the research scale

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high-pressure injection syringe. However, dirt-like oil in the cylinder can mix in the container, thus polluting the system. Another, yet better method is to process under low temperatures, so-called cooling introduction. This is suitable for small autoclaves. The autoclave is vacuum-pumped and cooled below the boiling point Tb of NH3 (under atmospheric pressure Tb D 33:35 ı C, melting point Tm D 77:7 ı C). Then, ammonia of >99:999% purity is flushed into the autoclave in a controlled way using a mass flow controller. The ammonia in the autoclave takes over the liquid state. Typically, we worked with the fill ranging 40–70% of NH3 . Already well working for small-size autoclaves, this filling technique is currently under development for large-scale autoclaves to hold 2-in. crystals. The final step prior growth is heating up the autoclave in a controlled way. We have performed the GaN growth in the temperature range T D 360–550 ı C under pressures p D 130–170 MPa. The temperature is kept constant during the growth process.

9.4 Chemistry of the Solution and Growth Mechanism 9.4.1 Solubility When compared to H2 O; NH3 is a weaker solvent for ionic substances since its dielectric constant ", a measure of the energy of solvation compared to the lattice energy, is smaller, i.e., " D 16:5 for NH3 compared to 80 for H2 O. This will result in the formation of low-crystallinity phases; however, working at much higher pressures may lead to partly overcome this problem as " increases at higher density of the solvent [11]. Anyway, it has become obvious that mineralizers are needed to enhance the solubility of GaN while processing under reasonable conditions. It has been shown in the case of GaN that the solubility strongly depends on the chemical environment in the ammonothermal solution. Retrograde solubility was reported for basic mineralizers like KNH2 [23], consequently, the seed crystals have to be assembled in the hot zone to achieve crystal growth while the precursor is placed in the colder zone of the autoclave where the solubility is highest. Contrary, the acidic mineralizers NH4 X in SC NH3 provides a positive solubility coefficient [24]. The precursor is placed in the hotter zone and the GaN seed crystals in the cooler zone inside the growth vessel. An advantage of the acidic mineralizer NH4 Cl over basic mineralizers such as KNH2 or NaNH2 is the high solubility in NH3 around room temperature where 124 g NH4 Cl .24:8 ı C/, 3.6 g KNH2 .25 ı C/, and 0.163 g NaNH2 .20 ı C/ per 100 g NH3 were reported [25]. Thus, a larger amount of GaN would be dissolvable in ammonia containing an acidic mineralizer in comparison to a basic mineralizer. We have worked on the measurement of the solubility of GaN under acidic ammonothermal conditions, in particular with NH4 Cl, and different molar concentrations c of the mineralizer. A high-pressure cell of 10 mL volume has been

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a Solubility Gan [mol]

0.25

189

GaN: Molar ratio NH4Cl/NH3 = 0.0127; 0.032; 0.127 AIN: Molar ratio NH4Cl/NH3 = 0.0318

0.20

0.127

0.15 0.10 contains 8 mol% NH4I

0.05

0.032 AIN 0.0127

0.00 250

b

350 400 450 500 Temperature [°C]

550

Molar ratio mineralizer/NH3 = 0.31– 0.34

0.008 Solubility GaN [mol]

300

T = 431– 437 °C, p = 95 –100 MPa

0.006

NH4 Cl1-x Xx with: X=l X = Br

0.004 0.002 0.000 0.0

0.6 0.8 1.0 0.2 0.4 Molar content x in mineralizer NH4Cl1-x Xx

Fig. 9.3 Solubility of GaN (a) and AlN for different concentrations of the mineralizer NH4 Cl with respect to 1 mol NH3 and (b) for mixed mineralizer NH4 Cl1x Xx with X D Br, I

employed at pressures around 100 MPa [24, 26]. Experiments were conducted for 120 h under isothermal conditions so as to avoid partial recrystallization of the dissolved GaN. Very similar to the ammonothermal growth process for GaN, polycrystalline HVPE GaN has been used as precursor. In Fig. 9.3a is shown the solubility of GaN in NH4 Cl-containing NH3 . It gets clear that an increase of CNH4Cl as well as temperature results in an increase of CGaN in the solution. Also shown in Fig. 9.3a is a data set of the solubility of the isostructural AlN under comparable conditions, i.e., similar T and molar ratio NH4 Cl=NH3 about 0.032. AlN shows a somewhat higher solubility than the GaN. Increasing the acidity of the solution by adding NH4 I lowers the solubility of AlN. Derived from the single exponential growth of CGaN over T D 250–600 ı C, we have calculated the energy of formation HF D 15:9 kcal mol1 [24]. It was found that this value is independent of the concentration of mineralizer in NH3 and will provide a tool to calculate the overall energy needed to control the solubility of GaN,

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such way controlling the growth of GaN under acidic ammonothermal conditions for a given system. In Fig. 9.3b is shown the solubility of GaN when the mixed mineralizer NH4 Cl.1x/ Xx with X D Br, I is employed. Increasing x results in a lowered solubility of GaN. Utilizing a fraction of NH4 Br gives a bit higher solubility than making use of NH4 I, but does not differ remarkably when the pure mineralizer NH4 Br or NH4 I, i.e., x D 1 is used.

9.4.2 Growth Rate and Chemistry of the Solution In general, the growth rate is firstly depending on the polarity, i.e., electrical charge and density of the ion attached to the crystal surface. The second contribution comes from the abundance of growth species drawn from the solution in front of the growing face of a crystal. Unlike the case of say NaCl dissolved in H2 O, no clear features exist from the chemical constitution of the acidic ammonothermal system at high pressure and temperature. Wang and Callahan [23] have recently summarized from the literature the results on the synthesis of GaN. They reported that the most favorable conditions for the growth of GaN under basic ammonothermal conditions with either purely KNH2 or mixed KNH2 and KN3 as mineralizers are at temperatures around 550 ıC in the hot zone and 172–310 MPa pressure. The ammonia fill must be >60%. In the case of acidic mineralizers such as NH4 Cl, the formation of ammonium chlorogallates being similar to .NH4 /3 GaCln is likely to happen, quite in analogy to ammonium hexafluorogallate, .NH4 /3 GaF6 [27]. The ammonolysis of the latter compound tends to produce GaN at about 400 ı C. As for acidic ammonothermal conditions, pentaaminechlorogallium (III) dichloride ŒGa.NH3 /5 ClCl2 has recently been suspected to be an effective precursor for the growth of GaN. This specie has been synthesized at 840 K [28], i.e., it is stable under the temperature conditions of acidic ammonothermal growth of GaN. The structure is described as cationic [Ga(NH3 )5 Cl]2C octahedra which are surrounded by distorted cubes of Cl anions. The later are weakly bonded and therefore relatively easily disconnect from the cationic octahedra in a given environment. Besides, the free Cl anions might easily form NH4 Cl, a mineralizer molecule again with the NH4C derived from the autoprotolysis of NH3 , accordingly:  2NH3 $ NHC 4 C NH2

(9.1)

If we now assume that the double positively charged ŒGa.NH3 /5 Cl2C does exist in the solution of SC NH3 and NH4 Cl, it might be the reason for the higher growth N face in comparison to the (0001) face. Quite in analogy to the rate of the .0001/ hydrothermal ZnO system with basic mineralizers like LiOH and KOH [10], positively charged growth species like ŒGa.NH3 /5 Cl2C are likely to be attracted by the N face rather than the positively charged (0001) face. We negatively charged .0001/

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have earlier noted that only one process is controlling the chemical equilibrium since the Arrhenius plot from the solubility chart clearly shows a linear slope. More investigations are required to really understand whether ŒGa.NH3 /5 Cl2C or another, perhaps very similar, complex molecule is the key to the chemical equilibrium in the reactions of dissolving and depositing GaN under acidic ammonothermal conditions. The estimation of the maximum stable growth rate Vmax , at which high crystal quality may still obtain, gives an idea of whether a solution growth method is of principal interest in terms of future production process. Above Vmax entrapment of inclusions, columnar growth mechanism and formation of related crystal defects are most likely to happen. Thus far, nothing is known for GaN grown under ammonothermal condition, even ZnO is not well treated this way. The concept of the maximum stable growth rate basically derived from Nernst’s principle [29] and later modified by Carlson [30]:  Vmax D

0:214Du 2 Ce2 Sc1=3 2 L

1=2 :

(9.2)

Here, the parameters are: D – diffusion coefficient, u – solution flow rate,  – thickness of the diffusion boundary layer in front of the crystal face, Ce – equilibrium solute concentration obtained from the solubility, – density of the fluid, L – length of the crystal surface, Sc – Schmidt number. The Sc gives the ratio of friction and mass transport by diffusion, and can be calculated: Sc D = D

(9.3)

Playing the parameters by means of  D 4:6  105 Pa s; .673 K/ D 0:39 g cm3 ; 1 D D 2104 cm2 s , calculated after [31], we would obtain Sc D 5:77. Generally, Sc D 0:2–3 for gas mixtures and 100–1,000 for mixtures of liquids. Our calculated value of Sc D 5:77 suggests that we are dealing with a gas phase rather than purely liquid phase. The calculation of Vmax is based on the bulk transport of the solute and therefore only realistic for macroscopic crystal length. As can be seen from (9.2), higher supersaturation would yield a higher Vmax . This fact has been proven experimentally [32]. As shown above, increasing CNH4Cl in NH3 from almost 0 to >1 mol% gives rise to an improved growth rate up to >50 m per day. Back to reality, very high flow rates are not realistic in closed systems without forced convection like stirring. Consequently, the determining parameter in the calculation is the maximum solution flow rate which should be established in a sustainable way. The effect of the other parameters of the equation is basically through their temperature dependences, and is therefore relatively small. Typical cases of flux growth does show a Vmax D 72–180 m h1 (1;700–4; 320 m per day) [33]. In the production of large size hydrothermal ’-SiO2 and ZnO crystals, a growth rate of 500–1,000 and 200–300 m per day, respectively, is

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typically observed for the (0001) face to produce the large-size, high-quality crystals [10]. Now, further playing the parameters for the case of ammonothermal growth of 1 GaN, thus making following assumptions that D D 2104 cm2 s ; u D 10 cm s1 as obtained from simulation [34], D0:1 cm; Ce .773 K/D1 (i.e., 1 mol GaN/1 mol NH4 Cl), Sc D 5:77; .673 K/ D 0:39 g cm3 , and L D 1 cm; Vmax D 15:48 m h1 (or 371:5 m per day) for a crystal of 1 cm in length. Accordingly, a GaN seed of two inch diameter could be grown at Vmax D 164:6 m per day. 1 Uncertainties are due to D (ranging 104  6  104 cm2 s );  was approximated, and Sc was calculated on considered D. D and  however roughly compensate so that Sc and also Ce play the central role in the calculation of Vmax . Main increase of Vmax will come from an improved flow rate u, which in turn reduces , which then allows to increase Ce . This can be achieved by a suitable temperature gradient T between the dissolution and growth zone as shown for the case of SC H2 O [35]. Increasing the T from 15 to 25 K almost doubled the main flow rate. We have found experimentally that a higher flow rate can easily be established in larger autoclave like the one shown in Fig. 9.4. A higher T can be established which causes the flow rate going up. The growth rate does have a high impact on the overall economy of a crystal growth process and strongly determines the price of wafers cut from a crystal. It is known that the growth of a single crystal from a solution is a much slower process

Fig. 9.4 A large size autoclave, mantled by the heater assembly, capable of growing 2-in. size GaN crystals

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than say growth from the melt and the reasons for that have been discussed extensively in the various literatures, e.g., [36]. If, however, multiple crystals are grown during a growth cycle then the growth from a solution gets much more efficient in terms of total crystal yield. The growth rate as determined from experiments is a function of the concentration of the mineralizer and the absolute temperature in the autoclave. A growth rate exceeding 50 m per day for long-term growth cycles requires 1 mol% NH4 Cl per 1 M NH3 and higher growth rates close to 100 m per day were already achieved in very recent experiments. A stable growth rate around 80–100 m per day would already be satisfying for a future mass production of GaN by the ammonothermal route by making use of large size autoclaves for hundreds of GaN seed crystals per growth cycle [23]. It is interesting to note that the growth rates of GaN obtained for the acidic and basic ammonothermal system are very similar, i.e., the differences in the chemistry behind does not prefer one method over the other in terms of the growth rate.

9.4.3 Effect of Acidity on Formation of GaN The mineralizer NH4 X (X D Cl, Br, I) as employed show an increase in the acidity from X D Cl toward X D I, i.e., the ligands X are determining the concentration of free NH 2 in the solution. A very welcomed effect of increasing the acidity by switching to NH4 Br or even better NH4 I is the increase of the yield of formed GaN, indicated by the quantity of used up GaN precursor, as was recently demonstrated by Ehrentraut et al. [37], see Fig. 9.5. Whereas employing NH4 Cl does not substantially change the yield over increasing the reaction temperature from 360 to 550 ıC, the mineralizers NH4 Br and NH4 I do, leading to raise the yield at T  450 ıC. This might be favorable for an increased growth rate; however, much work has still to be performed to apply this phenomenon. What else can be seen in Fig. 9.5 is that phase stability to the favor of the hexagonal phase is better achievable at T  500 ı C when NH4 Br and NH4 I are used, quite in agreement to the work reported by Purdy et al. [19]. The increasing acidity on the other hand goes along with an increased quantity of self-nucleated cubic GaN, space group F 43N m, crystals as evidenced by powder X-ray scanning [37]. This is strictly true only for the growth of GaN without providing a seed crystal of hexagonal GaN. The admixture of up to 20 mol% NH4 Br or NH4 I to NH4 Cl is still gaining yields around 95% while the hexagonal phase of GaN is still preserved. Figure 9.6 shows that the (200) reflection around 2 D 40:416ı due to cubic GaN, according to the JCPDS International Centre for Diffraction Data card no. 88–2,364, is hardly visible from the powder X-ray scan on self-nucleated GaN. First experiments have evidenced that a higher growth rate of hexagonal GaN on HVPE GaN seed crystals was obtained [37].

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Fig. 9.5 Yield of self-nucleated GaN in terms of quantity of dissolved GaN precursor and phases as evidenced by XRD. (After [37]). The h and c indicates a purely hexagonal and hexagonal with cubic phase, respectively, as revealed by XRD measurements

Fig. 9.6 Phase purity of self-nucleated GaN crystals grown from mixed mineralizers .1x/ mol% NH4 Cl=x mol%NH4 Br and .1  x/ mol% NH4 Cl=x mol%NH4 I. x D 20, 30, and 50 is denoted by blue, green, and red graphs, respectively

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The company of oxygen in the system at high levels at >1018 cm3 causes the formation of crystalline gallium oxide (Ga2 O3 , space group R 3N c) as precipitates if unseeded growth is performed.

9.5 Properties of Ammonothermal GaN Characterization of produced GaN crystals was performed using a Nomarski differential interference microscopy (NDIM), scanning electron microscope (SEM; JEOL JSM-7000F), energy-dispersive spectroscopy (EDS; JEOL JSM-7000F, 15 kV, 1 nA, detection range 0–20 keV, accumulation time 60 s) and secondary electron mass spectroscopy (SIMS; primary beam species CsC and O2C at 5 kV and around 200– 400 nA) to identify elements and their concentrations in comparison to our standard reference material, HVPE-GaN from Mitsubishi Chemical Corporation. Powder X-ray diffractometry (XRD; Rigaku RINT-2000, 40 kV and 40 mA) was employed using CuK’1 radiation in combination with a Ge (220) channel monochromator (12 arcsec beam divergence, scan speed 0:01ı min1 , step width 104ı ) to evaluate the crystallinity. The internal standard reference material for XRD was silicon (JCPDS-International Centre for Diffraction Data card no. 27–1,402). Steadystate photoluminescence (PL) was performed using a He–Cd laser (Omnichrome 3056-M-A01:  D 325 nm; Pout D 10 mW) for excitation and a cryostat (Daikin UV202CL) for temperature control. Spectra were obtained by a photomultiplier using a monochromator (Jobin Yvon Spex HR320: 1,200 lines/mm gratings, 320 mm focal length). The first crack-free ammonothermal GaN crystal has been grown on a 1 cm2 large (0001) HVPE-grown GaN seed crystal with a mechanically untouched surface around the year 2005 as shown in Fig. 9.7. The grown GaN crystal in Fig. 9.7 after the experiment was completely covered with a film mainly containing the mineralizer. This film could have been removed easily upon treatment with distilled water and 2-propanol. The quality of the surface of a seed crystal has a tremendous impact on the nucleation and successive growth of the GaN [21,37]. A mechanically untouched surface would certainly provide the best surface condition with no damage. However, proper CMP or purely chemical etching if one gets very smooth surfaces without macrosteps has shown to deliver a satisfying quality [21]. A similar observation has been made in the growth on m-plane seed crystals as we will discuss later. Figure 9.8 provides some images of a c-plane ammonothermal crystal grown on a HVPE GaN seed crystal. Here, the result was obtained from remodeled heater system which allowed a better control of the overall thermal conditions of the autoclave. Although the surface of the crystal has not been polished, just as-grown, a highly transparent ammonothermal crystal was obtained. The thickness was around 15 m for both faces, Ga and N-polar. The XRC FWHM from the (0002) reflection in Fig. 9.9 is very homogeneous with values around 108 and 339 arcsec for the Ga-polar and N-polar face, respectively. Such substrates can certainly be

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Fig. 9.7 Image of the first thick GaN crystal grown under acidic ammonothermal conditions on a mechanically untreated HVPE GaN crystal

Fig. 9.8 Optical micrographs of high-quality ammonothermal GaN grown on HVPE GaN seed shown in .a/ reflection, .b/ transmission, .c/ NDIM from Ga-polar face, and .d / NDIM from N-polar face. Reprinted with permission. Copyright (2009), Springer-Verlag

used for successive deposition of simple device structures. This work is currently progressing. It has recently been taken into consideration [21] that cobble patterning of an as-grown (0001) surface might be related to formed dislocations, quite in analogy to quartz. It can be seen from Fig. 9.8c, d that the density of this hillocks is N than for the (0001) face. Hashimoto et al. about three times larger for the .0001/

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Ga polar face

Intensity [a.u.]

1000

100

10

1 17.15

17.20 17.25 17.30 17.35

17.40 17.45

N polar face

Intensity [a.u.]

10000 1000 100 10 1 16.0

16.5

17.0

17.5

18.0

18.5

Fig. 9.9 XRC from the (a) Ga and (b) N-polar face of an ammonothermally grown GaN crystal

[38] have confirmed by transmission electron microscopy (TEM) that the threading dislocation density is higher for the N-polar face, i.e., about 106 and 107 cm2 for the Ga and N-polar face, respectively. The cross-sectional SEM in Fig. 9.10a shows about 100 m GaN had been grown on both polar faces, and pyramidal facets already had commenced to form. Zooming into the area of such formed interface between the HVPE seed and ammonothermal GaN crystal (Fig. 9.10b) from a typical short-term growth reveals a high degree of flatness and the cracking pattern just continue from the HVPE into the ammonothermal material. Voids have not been found. Figure 9.10c features the interface region N HVPE seed. The of a grown ammonothermal GaN crystal on the nonpolar .1010/ growth of GaN structures on nonpolar substrates draw a lot of attention recently since it has been demonstrated that improved internal quantum efficiency can be achieved because of the lack of an electrostatic field [39].

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Fig. 9.10 Cross-sectional SEM (a) showing the Ga-polar and N-polar face, both about 100 m thick, including commencing formation of pyramidal facets and featuring the flat interface region between (b) the (0001) HVPE seed and the ammonothermal crystal and (c) the .101N 0/ HVPE seed and the ammonothermal crystal

Fig. 9.11 Optical micrographs of high-quality ammonothermal GaN grown on m-plane HVPE GaN seed shown in (a) reflection, (b) transmission, (c) NDIM at higher magnification displaying growth features because of mechanical damaging and (d) NDIM revealing a very flat surface. Reprinted with permission. Copyright (2009), Springer-Verlag

N We have used .1010/, or m-plane, HVPE seed crystals which are available in 2 sizes of up to 1 cm and excellent quality with the (0002) XRC FWHM around 100 arcsec and a dislocation density of 2  105 cm2 [40]. Figure 9.11 demonstrates the result from the ammonothermal growth. A very smooth surface morphology is obtained. It must however be noted here that damage from slicing is still sometimes visible under NDIM observation (Fig. 9.11c). We have recently demonstrated the acidic ammonothermal growth on a 1in. seed crystal [26] and have shown in Fig. 9.1 the first result of the acidic ammonothermal growth of GaN on a 2-in. large HVPE (0001) GaN seed under similar conditions as above, i.e., temperature and pressure around 500 ıC and 100 MPa, respectively. The curvature radius of the seed crystal was around 9 m and the miscut orientation ˙0:15ı . The autoclave used for the process measures about 100  1; 000 mm2 in inner diameter  inner length, presently the largest one in use for theammonothermal growth of GaN from acidic mineralizer, see Fig. 9.4a. Although the crystal quality is still not satisfactory, in spite of the few attempts

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however this result is very promising in so far that we have shown the practical feasibility of growing large size GaN crystals under moderate process temperature and pressure conditions. This is an important issue for future industrial production. It should be noted here that strain-free (no bending) and large size, i.e., 2 in. GaN seed crystals are yet not available. Only Dwilinski et al. have recently shown the growth of large size, extremely flat and low-strain GaN crystals by using their own GaN crystals as seeding material, see Chap. 7. Impurities in a crystal are a major concern and the sources for them are quite manifold when autoclave technology is employed: the autoclaves, gas pipes, baffle and seed crystal holder in addition to precursor, mineralizer and the ammonia itself are contributing sources. In summary of the many experiments we have carried out thus far, we note levels of impurities such as Cr, Fe, and Ni of the order around 1015 –1018 cm3 ; 1017 –1020 cm3 , and 1017 –1020 cm3 , respectively. Platinum from the inner liner contributed with maximum 1016 cm3 . Silicon (Si) and oxygen (O) ar detected at levels around 1019 cm3 and 1018 –1020 cm3 , respectively. The values for Si and O are somewhat comparable to ammonothermal GaN fabricated from basic mineralizers where around 1019 cm3 for both elements [41] and around 1018 cm3 Si and 1020 cm3 O were reported [23, 41]. In all cases, it was found that the N-polar face contains higher amounts of impurities, quite in agreement with ZnO where the O-polar face easily traps impurities [10]. Also, the lower levels are the result of advanced growth technology. PL measurements were made to evaluate the optical quality of some of the ammonothermal GaN samples [24, 42, 43]. A general observation made by PL is that the GaN nucleated on the N-polar face seems to be of higher quality than the GaN nucleated on the Ga-polar face. Moreover, we have recently reported on the strong improvement of the optical quality of acidic ammonothermal GaN and shown that an ammonothermal GaN samples is similar to standard HVPE-grown GaN wafer [24]. In Fig. 9.12, the PL spectrum obtained at 10 K from ammonothermal GaN is compared with that from HVPE-grown GaN. Emission from the exciton bound to neutral donor .D0 X/ line of the hexagonal structure at 3.472 eV was observed from the HVPE GaN and at lower intensity also from the ammonothermal GaN. However, an additional peak at 3.357 eV was sometimes found which was assigned to the Y4 line [44]. Temperature dependence of this line however shows characteristics similar to the D0 X emission from 2H -GaN and to the reported exciton bound to point defects trapped at threading dislocations. It was therefore suggested that the Y4 line might be related to another crystal structure when comparing the intensity distribution of the (0006) !  2 XRD satellite peak and the Y4 emission. From bandgap energy consideration, the Y4 line could be linked to the D0 X emission from the 6H structure of GaN which is a metastable phase [43]. A deep level luminescence peaking at 1.93 eV, red luminescence, not shown in the figure, was observed from the ammonothermal GaN crystal at 10 K [42]. The characteristics of the temperature dependence of the PL intensity and TR-PL suggest that the high oxygen concentration of the ammonothermal GaN may be the origin of

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Fig. 9.12 Photoluminescence at 10 K from the Ga and N-polar face from a high-quality HVPE and an ammonothermal crystal

this observed red luminescence, quite in good agreement to the high oxygen levels in other GaN samples.

9.6 Prospects and Future developments for Ammonothermal GaN For the first time, a high-quality, free-standing 1 in. GaN wafer had been shown in 2007 by the Ammono Company [41]. Since then, the ammonothermal growth of GaN is being considered a technique competitive to the other, a more established vapor phase technique to grow GaN crystals. Fukuda & Ehrentraut have recently compared the developments in the ammonothermal technology to precedent developments in ZnO and SiO2 [26]. Supposing a comparable technological time route for scaling-up of the autoclave, GaN seed crystal size, and optimization of the process regime including a better understanding of the rate-determining process, one could expect that commercially available wafers of ammonothermal GaN of 2 in. in size may soon be available. The technological developments in the hydrothermal growth of ZnO were purely based on the achievements from hydrothermal growth of SiO2 . In contrast, the ammonothermal growth of GaN is very much profiting from both SiO2 and ZnO

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and problems such as the design of a large size autoclave including choice of the alloy, establishing of the thermal gradient or the process control are much better understood. Furthermore, numerical simulation of mass and heat transport under ammonothermal ambient to optimize growth conditions has seen quite some progress recently [45]. We therefore may expect the growth of large quantities of large size .2 in:/ GaN by the ammonothermal technique within the near future [26]. By comparison of the growth rates of some semiconductor crystals with those of SiO2 and ZnO, it gets clear that the melt growth is clearly dominant in terms of growth rates, e.g., up to 60 mm h1 for Si. Hydrothermal ZnO, however, grows with about 10 m h1 only. The target number for GaN would be about half that of ZnO; however, only 2 m h1 is currently achieved. The presently low growth rate, roughly 10% of that of quartz, seems to impede from the application of ammonothermal growth technique to produce GaN in an economic fashion. However, several factors need to be taken into account: (1) Firstly, scalability is a main advantage strongly speaking for the ammonothermal rather than for any of the other above mentioned technologies to produce GaN crystal. Analogous to SiO2 , i.e., using 1,000 seeds to grow 1,000 crystals from one autoclave at the same time would drastically bring down the price per one GaN wafer. Also, the size of the crystal to be grown is basically limited mostly by the size of the autoclave inner diameter. Taking all the mentioned points into account, we have recently [26] estimated that the price for one ammonothermal GaN wafer is comparable to that of InP, despite the low growth rate and high-pressure equipment. (2) Secondly, the growth rate of SiO2 in the early work of Spezia [46] was estimated to be 1 m h1 [47]. However, the growth rate of SiO2 has been improved by a factor of about 20 through better understanding of the rate determining parameters along with better knowledge of the chemistry in the hydrothermal solution. Right now we are facing a similar situation in the ammonothermal growth of GaN. (3) Thirdly, the period that ammonothermal growth technique is seriously under development is pretty short, about 15 years [13]. Moreover, only a very few groups have been working on the subject. This is in contrast to the melt growth of now dominating semiconductors like Si, Ge and to a lesser amount GaAs, which is the object for scientific investigation for up to half a century already by an incomparably larger community of researchers as well as monetary funding. The major tasks for the successful ammonothermal growth of GaN have recently been taken into account by Fukuda & Ehrentraut [26]: Profound knowledge of the solubility of the GaN feedstock and the mineralizers; availability of large-size, highquality GaN seeds; stable feeding with the GaN feedstock to achieve long-term growth; scaling-up of the size of present autoclaves for 2 in: crystals. N N N and Next, semi- and nonpolar oriented GaN wafers, i.e., f1011g; f1122g; f1010g, N f1120g will be fabricated from large-size GaN crystals, which will allow the fabrication of device structures free of electrostatic fields, therefore yielding an improved quantum efficiency [39, 48]. The dislocation density of GaN wafers must be decreased to 103 cm2 which is attainable as demonstrated for the case of LPE-grown GaN [49] and more recently also for the basic ammonothermal growth technique of GaN [41]. Large GaN seed

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crystals of highly crystalline perfection, i.e., unstrained with no bending would then become available. If we look about 5–10 years ahead from now, we believe that a countable fraction of the GaN wafers for devices is likely to be produced from GaN crystals grown by the ammonothermal growth technique [26].

References 1. T. Koyama, S.F. Chichibu, J. Appl. Phys. 95, 7856 (2004) 2. Mitsubishi Chemical Corporation, Tokyo, Japan http://www.m-kuyaku.co.jp 3. K. Motoki, T. Okahisa, N. Matsumoto, M. Matsushima, H. Kimura, H. Kasai, K. Takemoto, K. Uematsu, T. Hirano, M. Nakayama, S. Nakahata, M. Ueno, D. Hara, Y. Kumagai, A. Koukitu, H. Seki, Jpn. J. Appl. Phys. 40, L140 (2001) 4. T. Sasaki, T. Matsuoka, J. Appl. Phys. 64, 4531 (1988) 5. V. Härle, B. Hahn, H.-J. Lugauer, S. Bader, G. Brüderl, J. Baur, D. Eisert, U. Strauss, U. Zehnder, A. Lell, N. Hiller, Phys. Stat. Sol. A 180, 5 (2000) 6. Y. Honda, Y. Kawaguchi, Y. Ohtake, S. Tanaka, M. Yamaguchi, N. Sawaki, J. Cryst. Growth 230, 346 (2001) 7. P. Waltereit, O. Brandt, M. Ramsteiner, R. Uecker, P. Reiche, K.H. Ploog, J. Cryst. Growth 218, 143 (2000) 8. Z. Sitar, M.J. Paisley, B. Yan, R.F. Davis, Mater. Res. Soc. Symp. Proc. 162, 537 (1990) 9. E.S. Hellman, D.N.E. Buchanan, D. Wiesmann, I. Brener, MRS Internet J. Nitride Semicond. Res. 1, 16 (1996) 10. D. Ehrentraut, H. Sato, Y. Kagamitani, H. Sato, A. Yoshikawa, T. Fukuda, Prog. Cryst. Growth Charact. Mater. 52, 280 (2006) 11. H. Jacobs, D. Schmidt, High-pressure ammonolysis in solid-state chemistry, In Current Topics in Materials Science, vol. 8, ed. by E. Kaldis (North-Holland, Amsterdam, 1982), p. 379 12. D. Peters, J. Cryst. Growth 104, 411 (1990) 13. R. Dwili´nski, A. Wysmolek, J. Baranowski, M. Kami´nska, R. Doradzi´nski, J. Garczy´nski, L.P. Sierzputowski, Acta Phys. Pol. 88, 833 (1995) 14. J.W. Kolis, S. Wilcenski, R.A. Laudise, Mater. Res. Soc. Symp. Proc. 495, 367 (1998) 15. A.P. Purdy, Chem. Mater. 11, 1648 (1999) 16. D.R. Ketchum, J.W. Kolis, J. Cryst. Growth 222, 431 (2001) 17. M.J. Callahan, B. Wang, L.O. Bouthillette, S.-Q. Wang, J.W. Kolis, D.F. Bliss, Mater. Res. Soc. Symp. Proc. 798, Y2.10.1 (2004) 18. B. Wang, M.J. Callahan, K.D. Rakes, L.O. Bouthillette, S.-Q. Wang, D.F. Bliss, J.W. Kolis, J. Cryst. Growth 287, 376 (2006) 19. A.P. Purdy, R.J. Jouet, C.F. George, Cryst. Growth Des. 2, 141 (2002) 20. A. Yoshikawa, E. Ohshima, T. Fukuda, H. Tsuji, K. Oshima, J. Cryst. Growth 260, 67 (2004) 21. Y. Kagamitani, D. Ehrentraut, A. Yoshikawa, N. Hoshino, T. Fukuda, S. Kawabata, K. Inaba, Jpn. J. Appl. Phys. 45, 4018 (2006) 22. D. Ehrentraut, K. Kagamitani, T. Fukuda, F. Orito, S. Kawabata, K. Katano, S. Terada, J. Cryst. Growth 310, 3902 (2008) 23. B. Wang, M.J. Callahan, Cryst. Growth Des. 6, 1227 (2006) 24. D. Ehrentraut, K. Kagamitani, C. Yokoyama, T. Fukuda, J. Cryst. Growth 310, 891 (2008) 25. D. Nicholls, Inorganic Chemistry in Liquid Ammonia, in Topics in Inorganic and General Chemistry, Monograph 17, ed. by R.J.H. Clark (Elsevier, Amsterdam, 1979) 26. T. Fukuda, D. Ehrentraut, J. Cryst. Growth 305, 304 (2007) 27. M. Roos, J. Wittrock, G. Meyer, S. Fritz, J. Strähle, Z. Anorg. Allg. Chem. 626, 1179 (2000) 28. H. Yamane, Y. Mikawa, C. Yokoyama, Acta Crystallogr. E 63, i59 (2007) 29. W. Nernst, Z. Physik. Chem. 47, 52 (1904)

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30. A. Carlson, in Growth and Perfection of Crystals, ed. by R.H. Doremus, B.W. Roberst, E. Turnbull, (Wiley, New York, 1958), p. 421 31. C.-H. He, Y.-S. Yu, Ind. Eng. Chem. Res. 37, 3793 (1998) 32. D. Ehrentraut, Y. Kagamitani, Ammonothermal growth of GaN, IWBNS-V, Itaparica, Salvador, Brazil, September 24–28, 2007, Workshop Program and Abstracts 33. H.J. Scheel, D. Elwell, J. Cryst. Growth 12, 153 (1972) 34. Y. Masuda, Private communication 35. K. Nagai, J. Asahara, J. Jpn. Assoc. Cryst. Growth 27, 68 (2000) 36. A.A. Chernov, in Modern Crystallography III: Crystal Growth, Springer Series in Solid-State Sciences, vol. 36 (Springer, Berlin, Heidelberg, New York, Tokyo, 1984) 37. D. Ehrentraut, N. Hoshino, Y. Kagamitani, A. Yoshikawa, T. Fukuda, H. Itoh, S. Kawabata, J. Mater. Chem. 17, 886 (2007) 38. T. Hashimoto, F. Wu, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 46, L525 (2007) 39. P. Waltereit, O. Brandt, A. Trampert, H.T. Grahn, J. Menniger, M. Ramsteiner, M. Reiche, K.H. Ploog, Nature 406, 865 (2000) 40. F. Orito, K. Katano, S. Kawabata, Y. Kagamitani, D. Ehrentraut, C. Yokoyama, H. Yamane, T. Fukuda, Growth strategy for gallium nitride bulk crystals, 5th International Workshop on Bulk Nitride Semiconductors (IWBNS-V), Sept. 24–28, (2007), Itaparica, Bahia, Brazil, Workshop Program and Abstracts 41. R. Dwili´nski, R. Doradzi´nski, J. Garczy´nski, L.P. Sierzputowski, A. Puchalski, Y. Kanbara, K. Yagi, H. Minakuchi, H. Hayashi, J. Cryst. Growth 310, 3911 (2008) 42. K. Fujii, G. Fujimoto, T. Goto, T. Yao, Y. Kagamitani, N. Hoshino, D. Ehrentraut, T. Fukuda, Phys. Stat. Sol. (A) 204, 3509 (2007) 43. K. Fujii, G. Fujimoto, T. Goto, T. Yao, Y. Kagamitani, N. Hoshino, D. Ehrentraut, and T. Fukuda, Phys. Stat. Sol. (A) 204, 4266 (2007) 44. M.A. Reshchikov, H. Morkoc, J. Appl. Phys. 97, 061301 (2005) 45. Q.-S. Chen, V. Prasad, W.R. Hy, J. Cryst. Growth 258, 181 (2003) 46. G. Spezia, Atti. Accad. Sci. Torino 35, 95 (1900) 47. F. Iwasaki, H. Iwasaki, J. Cryst. Growth 237–239, 820 (2002) 48. A. Chakraborty, T.J. Baker, B.A. Haskell, F. Wu, J.S. Speck, S.P. DenBaars, S. Nakamura, U.K. Mishra, Jpn. J. Appl. Phys. 30, L945 (2005) 49. Y. Mori, T. Sasaki, Oyo Buturi 75, 529 (2006), in Japanese

Part IV

Flux Growth Technology

Chapter 10

High Pressure Solution Growth of Gallium Nitride Michal Bo´ckowski, Pawel Strak, ˛ Izabella Grzegory, and Sylwester Porowski

Abstract State-of-the-art high pressure solution (HPS) growth of gallium nitride is presented in this chapter. The spontaneous crystallization and seeded growth are described in detail from the technological point of view. Two classes of pressuregrown crystals and their use as substrates for blue laser diodes (LDs) are demonstrated.

10.1 Introduction The melting temperature of gallium nitride (GaN) is extremely high 2; 493 K- and was determined with very good precision by Utsumi et al. [1] in 2003. The nitrogen pressure needed for congruent melting of GaN is 6 GPa. Before 2003, people had known that the melting parameters were high from theoretical approximation [2]. Since it was impossible to crystallize GaN from the melt, other crystal growth methods were developed. One of them was crystallization from the high temperature liquid solution of gallium under high nitrogen pressure. According to the definition given by Elwell and Scheel [3], in the high temperature solution growth the constituents of the material to be crystallized are dissolved in a suitable solvent and crystallization occurs as the solution becomes critically supersaturated. The supersaturation can mainly be promoted by cooling the solution or by a transport process in which the solute flows from a hotter to a cooler region [3]. Obviously, the main advantage of the solution growth is that the growth temperature is lower than that required for growth from the melt. This advantage often results in a better crystal quality with respect to point defects, dislocation density, and low angle grain boundaries, compared to crystals grown directly from their melts [3]. The main disadvantages of the high temperature solution growth method are associated with microscopic and macroscopic inclusions of solvent or impurities in the crystals, nonuniform doping, and a slow growth rate [3]. The high temperature solution growth of gallium nitride proceeds at temperature of the order of 1,800 K and at nitrogen pressure up to 2 GPa. The molecular nitrogen being under pressure is dissociated on the hot liquid gallium surface and the atomic nitrogen is dissolved in 207

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the liquid metal. Then, the convective flow of gallium, due to applied temperature gradients, is exploited in order to provide atomic nitrogen to the supersaturated zone of the solution, where the GaN crystallization takes place. Since the high temperature solution growth of GaN requires high nitrogen pressure, gallium nitride has always been in the field of interest of the Institute of High Pressure Physics Polish Academy of Sciences (formerly High Pressure Research Center PAS). In 1984 Karpinski et al. [4] determined the equilibrium nitrogen pressure–temperature curve for gallium nitride in the temperature range up to 2,000 K. At the beginning of the nineties, on the wave of nitrides popularity, the thermodynamic background of the system Ga–GaN–N2 and thermodynamic properties of GaN have been determined and formulated [5]. By the end of 1995 small GaN platelets (up to 2 mm) had been obtained. In 1996, 1 cm2 and 100 m thick GaN crystal (platelet) of perfect structural quality and low dislocation density .102 cm2 / was obtained for the first time. Based on this success, relatively fast development of the Crystal Growth Laboratory was observed. The quantity of high pressure (HP) reactors (Institute’s home-made apparatus) increased several times during a few years. The interior diameter, and thus the working volume of the HP chamber increased too. A small pilot production of GaN crystalline platelets was started thanks to the Polish Government Project [6]. In parallel, the epitaxial techniques (MOCVD and MBE) and laser processing technology were created and implanted to the Institute. In 2002, the first blue laser diode (LD) was built on the pressure grown GaN crystals by the MOCVD method [7]. Three years later, GaN-based blue LDs were obtained by plasma-assisted MBE [8]. The strategic Polish Government Program “Blue Optoelectronics” [9] helped develop the nitrides technology in the Institute of High Pressure Physics and created a spin-off laser company called TopGaN Ltd. The technology of the blue lasers diodes made by TopGaN has just been based on homoepitaxial growth on pressure-grown gallium nitride crystals. The idea of the homoepitaxy has always been the base for the nitrides technology in the Institute of High Pressure Physics. Unfortunately, the spontaneous crystallization by the HP technique has not allowed growth of bigger than 3 cm2 GaN crystalline platelets. Therefore, hydride vapor phase epitaxy (HVPE) was introduced in the Institute. The new idea has been to combine the HVPE with the HP technique in order to create 2 in., low dislocation density crystals with good structural quality and appropriate physical properties to be the substrates. In this chapter, this idea is partially described. Partially, since the use of HVPE-GaN crystals as seeds for the HP crystallization is just shown. The use of HP-GaN crystals as seeds for the HVPE technique is presented in detail in Chap. 3 of this book. The chapter starts with the brief description of the high pressure solution (HPS) growth method and its thermodynamic and kinetic aspects. Then, the HP experimental setup and the spontaneous crystallization are described. Physical properties of the pressure grown crystals (platelets and needles) are presented. Next part of the chapter is focused on a seeded crystallization under pressure. The results of the liquid phase epitaxy (LPE) technique are shown. The relatively new approach of seeded growth, with convective flow of gallium under control, is described in details. At the

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end of this chapter the application of the pressure grown GaN crystals as substrates for the LDs is briefly reported.

10.2 Growth Method The HPS method is temperature gradient growth method based on direct reaction between gallium and nitrogen at high temperature and high nitrogen pressure. Figure 10.1 presents schematic illustration of the crystallization by the HPS method. Nitrogen molecules dissociate on gallium surface and dissolve in the metal. Therefore, the crystals are grown from the solution of atomic nitrogen in liquid gallium. Supersaturation, the driving force for the crystallization process, is created by the application of temperature gradient along the liquid gallium. Thus the HPS method consists of three stages: (1) dissociative adsorption of nitrogen on liquid metal surface, (2) dissolution and transport of nitrogen atoms from the hot end of the solution to the cooler one, and (3) crystallization process. Due to the temperature gradient in the system, the atomic nitrogen is transported from the hot end of the solution to the cooler part. It has been shown that the convection mechanism plays the dominant role in the transport phenomenon [10]. The nitrogen in the solution is uniformly distributed due to the strong convection flow. Thus, the solution at the cold zone must be supersaturated. It means that the excess nitrogen concentration related to the equilibrium value for a given temperature exists at the cold zone of the liquid metal. Hence, GaN crystal growth takes place at this zone.

crucible N2 molecules under pressure Dissociative adsorption of N2 on liquid metal surface Dissolution and transport to the crystallization zone

ΔT

Ga GaN

Spontaneous crystallization

T

Fig. 10.1 Schematic illustration of the crystallization process in a temperature gradient by the high pressure solution growth method; three basic stages of the method are marked

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10.2.1 Thermodynamic and Kinetic Aspects of HPS Growth As mentioned in Sect. 10.1, the equilibrium pressures of N2 for the Ga–GaN system have been determined by Karpinski et al. [4]. The direct synthesis and decomposition experiments performed by gas pressure technique for nitrogen pressures up to 2 GPa were practiced. The curve following from these data is shown in Fig. 10.2a. Figure 10.2b shows, in turn, the N solubility data resulting from the annealing of Ga in N2 atmosphere at the three phase Ga–GaN–N2 equilibrium conditions for the temperature range from 1,273 to 1,973 K. The solid line has been calculated in the ideal-solution approximation. The nitrogen concentrations are not too high; they are below 1 at% [5]. This curve allows us to determine the maximum of the relative supersaturation reached in the solution. If the convective transport in the solution is faster than the crystallization, the relative supersaturation can be expressed as: x.T2 /  x.T1 / ; x.T1 /

D

(10.1)

where  represents the relative supersaturation and x.T2 /  x.T1 / is the difference between equilibrium nitrogen concentration at hot and cold zone of the solution. The determination of  is also schematically presented in Fig. 10.2b. The analysis of thermodynamic properties of GaN and the system of its constituents explains the role of the use of HP [5]. It is a factor increasing thermodynamical potential of GaN constituents. It makes the GaN crystal stable at higher temperatures necessary for crystallization. The pressure, however, is also important for kinetics of the GaN synthesis. As mentioned, one of the stages of GaN synthesis from its constituents is dissolution of nitrogen in hot liquid gallium. This process was analyzed by quantum mechanical calculations [11]. The nitrogen molecule approaching the metal surface is repelled by the metal which results in a potential

a

temperature, K 2500 2250 2000 1750

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b 1900 1.45 GPa

10000 pressure limit of 2 GPa

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GaN Ga+1/2N2

100 0.0004

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1800 T2 1700 T1 1600 1500 1400

0.0005

0.0006 1/T, K–1

0.0007

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1.21 GPa

1.01 GPa 0.81 GPa 0.73 GPa 0.6 GPa

0.3 GPa

x(T2)-x(T1) x(T1)

x(T2)

0.001 0.002 0.003 0.004 0.005 XN, at.fr.

Fig. 10.2 (a) p  T equilibrium curve for GaN determined by Karpinski et al. [4]; (b) liquidus curve for Ga–GaN system [5]; determination of the relative supersaturation in the solution

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barrier. If the molecule has enough energy to overcome the potential barrier, it comes closer to the gallium surface and dissociates into two atoms, forming new bonds with the metal. The potential barrier is lower than the bonding energy in the nitrogen molecule. However, its value of 3.5 eV seems to be quite high. Therefore, the density of the interacting gas (thus its pressure) is of crucial importance. Here, it should be pointed out that for oxygen interacting with gallium, there is no potential barrier [12]. One can always expect the oxygen atoms in liquid gallium. This fact is very important for the properties of GaN crystals grown by HPS method, what will be shown later. The high potential barrier for nitrogen dissociation on the liquid gallium surface also suggests that this dissociation process is kinetically controlled even for relatively high temperatures. Determining the rate of dissociation at 2 GPa, Krukowski et al. [11] showed that for gallium nitride the rate of dissociation is not fast. For effective synthesis (10 mg cm2 during 100 h), the temperature should be higher than 1,500 K. Comparing this last result with the equilibrium p  T curve for GaN (see Fig. 10.2a), it may be concluded that only application of the high nitrogen pressure allows the synthesis of gallium nitride from its constituents.

10.2.2 Experimental One of the typical experimental systems for crystal growth of GaN at nitrogen pressure is shown in Fig. 10.3. This system, consisting of two HP chambers shown in the photograph, is connected to the central pressure line where 1 GPa of nitrogen can be obtained by a set of membrane compressors and intensifiers. Vertically positioned, technological gas pressure chambers (reactors) of internal diameters of 4, 6, or 10 cm, allowing crucibles with the working volume of 10, 25, and 150 cm3 respectively, are connected by capillaries to the special gas chambers with compressors, which serve as reservoirs of nitrogen and allow stabilizing pressure in the chambers during experiments. The multizone cylindrical graphite furnaces, capable of reaching temperatures up to 2,000 K, are placed inside the gas pressure reactors. The HP chambers are equipped with additional systems necessary for annealing in vacuum, cooling of the reactor, electronic stabilization and programming of pressure and temperature. To register the temperature during the crystal growth experiments, PtRh6%–PtRh30% thermocouples are used. They are arranged along the furnace and coupled with the input power control electronic systems. The pressure is measured by manganine gauges positioned in the low temperature zone of the HP systems. The pressure and temperature are stabilized with an accuracy of ˙1 MPa and ˙0:1 K, respectively. The typical crystal growth experiment is that the metal (Ga) placed in the graphite crucible is heated in the furnace inside the HP chamber with a constant rate to given axial and/or radial temperature profiles. Then, the system is annealed at these conditions under high nitrogen pressure for 100–500 h. After that, the furnace is cooled down at a constant rate, the system is decompressed, and the metal with

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the set of membrane compressors and intensifiers

to the next high pressure system

2m

Fig. 10.3 One of the experimental system for crystal growth of GaN at N2 pressure; two visible high-pressure chambers are connected to the central pressure line. The photograph at the bottom left corner of this figure shows a furnace being introduced into the reactor

GaN crystals inside is removed. The charge from the crucible is etched with boiling HNO3 =HCl acid solution to extract the crystals. During the crystallization process, the pressure and the temperature in the whole gallium sample correspond to the nitride stability range. Thus, the highest measured temperature is lower than the equilibrium temperature for a given nitrogen pressure at three phase Ga–N2 –GaN equilibrium conditions. The gallium nitride crystals grow spontaneously from the wall of the crucible. They are perpendicular to the wall surfaces and distributed randomly in the supersaturated cold zone of the solution. The size of this zone depends on the applied axial temperature profile and temperature distribution in the liquid gallium. Two types of the axial temperature gradient configurations are used: positive (direct) and negative (reversed). They are schematically shown in Fig. 10.4a, b. The temperature distribution is evaluated by three, four, or five thermocouples positioned in the crucible wall. Three types of the axial temperature profiles – linear, convex, and concave are observed. Figure 10.4c represents the radial temperature distribution on the gallium top and crucible bottom. The radial gradient seems to be very important since it drives the convective flow in the liquid gallium. A few experiments with thermocouples positioned close to the center of the gallium top surface were carried out. These experiments showed that the radial temperature distribution on the gallium top was flat, as is shown in Fig. 10.4c. At the bottom of the crucible the temperature has always been measured at the sides and in the center. Three radial temperature profiles represented in Fig. 10.4c by curves 1, 2, and 3 have been detected.

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a

linear

positive (direct)

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z

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z

z

Ga

T

b

T

linear

T

convex

negative (reversed) z

concave

z

z

Ga

T

T

T

c radial temperature distributions on the crucible′s top and bottom

T

Ga

1 2 3 r Fig. 10.4 (a) Positive (direct) temperature gradient configuration with three possible temperature profiles; (b) negative (reversed) temperature gradient configuration with three possible temperature profiles; (c) radial temperature distributions on the crucible’s top and bottom; at the bottom of the crucible three temperature profiles can be obtained; filled circle denotes the positions of thermocouples

For seeded crystallization, a seed crystal is immersed in the gallium and back melted by the metal [13]. Then, the procedures are similar to that described above. The details of seeded and spontaneous growth configurations and procedures are presented separately in the corresponding subsections of this chapter.

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10.3 Spontaneous Crystallization by HPS Growth Method The crystallization without seeds in the solution is called here the spontaneous crystallization. The dominating morphological form of GaN crystals grown by the spontaneous method is a hexagonal platelet. In order to crystallize big (1 cm2 or more) platelets the positive temperature gradient configuration with the convex temperature profile is usually applied. The maximum temperature of the solution is about 1,750 K at nitrogen pressure of 1 GPa. Typical duration of a crystal growth run is 100 h. A difference between the temperatures measured at the hottest and the coldest places in the crucible (a supercooling) is always lower than 50 K. The average temperature gradient along the z-axis of the crucible attains 10 K/cm. The typical relative supersaturation determined from (10.1) varies from 5% to 20%. The radial temperature distribution is represented by the curve 3 from Fig. 10.4c, i.e., the temperature at the center of the crucible is lower (several degrees) than those measured at the sides. In this configuration the crystals are grown at the bottom of the crucible and on the crucible wall close to the bottom. Up to five 1 cm2 platelets can be obtained during one crystallization run in the reactor of internal diameter of 6 cm. Other crystals are smaller. An increase in the crucible’s volume - thus crystallization in bigger reactors - does not increase the size of the platelets. Only number of the crystals grows up. The linear or concave temperature distributions allow increasing the supersaturated zone in the solution and then the crystals can grow on the crucible wall, relatively far from the bottom. It does not mean though that the crystals are bigger. They are just distributed in the crucible in a different way. A significant increase in the temperature gradient (to 50 K/cm) does not lead to a significant acceleration of the growth rate but to a change of the growth mode. Then the needles become the dominating morphological form of GaN crystals grown by the spontaneous method.

10.3.1 Habit and Morphology of the Crystals Figure 10.5a presents a scheme of an ideal GaN platelet. Its large hexagonal surfaces correspond to {0001} polar crystallographic planes of the wurzite structure. ˚  ˚  The N planes and semipolar 1011 N and side faces of the crystals are nonpolar 1010 ˚  N planes. For the real GaN platelet (see Fig. 10.5b) these last planes are not 1012 well developed. The crystals in the form of hexagonal platelets grow slowly at a rate N lower than 0.1 mm/h into <1010> directions (perpendicular to the c-axis). They are single crystals, slightly gray or transparent, very often with flat mirror like faces. The maximum lateral size of the platelets is 3 cm2 whereas the thickness is about 150 m. The tendency for unstable growth is stronger for one of the polar {0001} faces of the platelets. The morphological features like macrosteps, periodic inclusions of solvent, or cellular growth structures are observed on this side. The opposite surface is mirror-like and often atomically flat. For crystals grown without an intentional doping, the unstable surface always corresponds to the Ga-polar

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Fig. 10.5 Two habits of GaN crystals grown by the HPS method (a) ideal GaN wurzite platelet; (b) real GaN pressure grown platelet, N-side as grown; grid 1 mm; (c) ideal GaN wurzite needle; (d) real GaN pressure grown needle

(0001) face of GaN. The polarity of the crystal surfaces has been identified by etching in hot alkali aqua solutions since the Ga-polar surface is inert to etching whereas the N-polar one etches very well. The method was calibrated by convergent beam electron diffraction (CBED) and X-ray photoelectron spectroscopy (XPS) measurements [14–16]. Figure 10.5c represents the second possible form of the pressure grown crystals – a needle. The real crystal is presented in Fig. 10.5d. The needles are elongated into c-direction, with diameter of about 1 mm and up to 10 mm in length. It is clearly seen in Fig. 10.5d that only some facets of the needle are well formed and grown. The others are not well developed or they do not simply exist. Some of them collapse and then disappear. Thus, the GaN needles are strongly unstable crystals. They are dark due to the presence of microdefects like dislocation loops and gallium inclusions being the result of the fast and unstable growth in the c-direction [17]. These crystals are often hollow from inside. However, they have very well developed faces of orientations different from the <0001> ones, which give an opportunity to check the growth behavior in these directions using the needles as seeds (see Chap. 3).

10.3.2 Physical Properties of the Crystals As was already mentioned, for oxygen interacting with Ga, there is no potential barrier for dissociation. Therefore, even traces of this impurity in the growth system are sources of the unintentional oxygen doping of GaN. Consequently, the HP-GaN crystals (both platelets and needles) are strongly n-type with uniform free electron

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concentration of about 5  1019 cm3 (metallic conductivity) and mobility of about 60 cm2 =Vs [18]. These free carriers can be fully eliminated by Mg acceptor added into the growth solution [19]. The presence of the native point defects in the crystals (platelets) was checked by positron annihilation measurements [20]. High concentration of Ga vacancies .VGa / was found in the conductive crystals in contrast to the Mg-doped samples where VGa were not observed. This agrees with the theoretical prediction that the formation energy of VGa decreases with an increase of the Fermi level energy [21]. It suggests that the creation of these defects is thermodynamically controlled. The difference in the PL spectra of the conductive (strong yellow emission) and Mgdoped crystals (no yellow emission, but blue Mg-related signal) revealed that VGa are involved in strong yellow luminescence in GaN. The structure of the pressure grown GaN crystals has been studied by X-ray diffraction [22]. The crystallinity of the conductive platelets depends on the size of the crystal. The full widths at half maximum (FWHM) for the reflection 002 are of the order of 0.02 degree for 1–3 mm crystals. For larger platelets the rocking curves often split into a few peaks showing the presence of low angle boundaries separating grains of a few millimeters in size. Misorientation between grains increases monotonically from end to end of the crystal [21]. The crystallinity of the needles is similar to that of the platelets [17].   It has been shown by TEM examination [14, 23] that the N-polar 0001N surface of the n-type pressure grown GaN platelets is often atomically flat and that the crystals with this surface are practically free of extended defects. On the opposite, rough surface, a number of extended defects like stacking faults, dislocation loops, and Ga microprecipitates are observed. The relative thickness of this part usually consists of 10% of the entire thickness of the platelet. In order to determine dislocation densities in GaN crystals, the defect selective etching (DSE) method has been developed [24]. It has been shown that etching in molten KOH–NaOH eutectic (at 723 K) reveals dislocations in GaN pressure-grown single crystals. The etch pit density in HP-GaN (estimated by counting the etch pits in a given area) is very low, of the order of 102 cm2 . More detailed description of the properties of the pressure grown GaN crystals can be found in [25].

10.4 Seeded Growth by HPS Method The main disadvantage of spontaneous crystallization by the HPS method is poor reproducibility of the size and distribution of GaN crystals in the crucible. Better control of these two parameters may be achieved by growth with seed crystals introduced intentionally into the solution. Two techniques have been developed in the Crystal Growth Laboratory of the Institute of High Pressure Physics: LPE and seeded growth with convective flow of gallium under control.

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10.4.1 Liquid Phase Epitaxy in the c-Direction on Various Substrates The LPE technique, i.e., deposition of GaN on the substrate to force the growth in particular direction has been examined at two types of gradient configurations: positive and negative and with five kinds of substrates: HP-GaN, HVPE-GaN, SiC, sapphire, and sapphire/GaN MOCVD templates. The experiments have been carried out in three brands of the HP reactors, what allowed us to use three sizes (in diameter) of substrates: 1 cm, 1 in., and 2 in., respectively. The substrate has been placed at the bottom or at the top of the crucible. A thin tube has always been located in the crucible in such a way that the edges of the seed have not been exposed to the liquid gallium. In this way, the edge nucleation has been eliminated. The crucible with a baffle plate positioned close to the seed has been used. This baffle plate has allowed obtaining a flat crystallization front on the substrate and consequently maintaining a flat GaN surface during a long crystallization run [26]. The crystallization without the baffle has always led to nonuniform crystallization front,and thereby formation of the growth hillock at the center of the substrate [13]. This last phenomenon has perturbed the crystallization process. The crystallization process has also been perturbed if the nitrogen polar surface of the substrate has been exposed to the liquid gallium (case of HP-GaN and HVPE-GaN). Then, the presence of several growth centers has been observed. The simultaneous nucleation and growth of randomly oriented crystals have been noticed, too [13]. In turn, on the gallium polar surface (case of HP-GaN, HVPE-GaN and sapphire/GaN MOCVD template) the growth mode has been stable and proceeded by the steps propagation from the center of the sample to its sides [27]. The LPE crystallization on pure sapphire has been dominated by the polycrystalline GaN growth [28]. On the 4H-SiC substrates the main mechanism of crystallization has been the formation of many growth hillocks, coalesced by small steps propagating from each of the hillock [28]. The permanent feature of GaN grown on SiC has been that the GaN has totally been cracked. These cracks have been formed during cooling procedure, just after the main growth process, due to the difference in thermal expansion coefficients between GaN and SiC at low temperature [28]. It is worth noting that cracking has never been observed for the growth on HP-GaN, HVPE-GaN, or sapphire/GaN MOCVD template. After HP LPE on the sapphire/GaN MOCVD template the bending of the sample has just been observed. It has been found that for the same temperature gradient and supercooling the growth rate in negative (reversed temperature gradient) configuration is twice as fast [27]. It could indicate that the growth rate is governed by a nitrogen transport mechanism. In the configuration with reverse temperature gradient, the convection flow is always stronger. Therefore, for such configuration the crystal can be grown in relatively higher supersaturation compared to the direct gradient system. It has been shown however that there are two factors responsible for GaN growth in the c-direction and these factors are not depended on the configurations or the kind of the substrates used. For a short time .30 h/, the growth rate is governed mainly

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by nitrogen transport to the crystallization zone, and the observed average growth rate can attain 10 m=h. After a longer time .100 h/, even at high temperature gradients, the surface kinetic factor becomes more important and the average rate decreases to 1 m=h. In our opinion, the macrosteps and terraces have been created on the growing crystal surface which finally hindered or even stopped the growth [26]. However, Hussy et al. [29] have proposed a different explanation. During LPE at low pressure solution growth process, it has been observed that at relatively high supersaturation the parasitically growing GaN consumed nearly all nitrogen which has been supplied in the solution. The amount of nitrogen that has been transported through the whole solution volume to the seed has been decreasing with increasing process time. Therefore, the growth rate of the seed has decreased. The analysis of the GaN mass crystallized at configuration with and without seed has confirmed that the growth rate on the seed is governed by surface kinetics. The GaN mass crystallized at the bottom of the crucible without seed was bigger than the mass crystallized on the seed (with no parasitic growth observed) under the same experimental conditions [28]. Thus, not all nitrogen atoms from the solution reaching the seed surface were adsorbed there. Figure 10.6a, b present a general view of the samples deposited on HP-GaN and sapphire/GaN MOCVD templates, respectively. The average dislocation density in

Fig. 10.6 (a) General view of 100 m thick GaN layer deposited on HP-GaN; (b) general view of 100 m thick GaN layer deposited on 2 in. MOCVD GaN/sapphire template; (c) general view of 100 m thick GaN layer deposited on HVPE-GaN – grid 1 mm; (d) Nomarski microphotograph of the surface after DSE. Homogenous distribution of etch pits is visible

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GaN crystallized on the HP-GaN is always as low as 102 cm2 . The dislocation density in the pressure-grown material on sapphire/GaN templates is of the order of 3–5  107 cm2 , for layers thicker than 40 m. This value is more than one order of magnitude lower than in the outgoing templates where the dislocation density varies from 8  108 to 109 cm2 . Lower defect density can be obtained if the HVPE-GaN is used as a substrate. Then, the dislocation density in the starting material is of the order of 5  106 cm2 . For 40 m thick GaN layer the defect density is 2  106 cm2 . For the layer 250 m thick the etch pit density attains even 8  105 cm2 . Figure 10.6c,d present newly deposited material on HVPE free-standing crystal and its surface with etch pits, respectively. Two kinds of etch pits are visible: smaller are associated with edge dislocations, larger associated with screw or mixed-type dislocations. The number of edge dislocations is bigger by a factor of two. The new deposited material on HP-GaN, HVPE-GaN, SiC, or sapphire/GaN template is always n-type with free electron concentration of about 5  1019 cm3 and mobility of 60 cm2 =V s. It also shows strong yellow luminescence. The FWHM of GaN grown on the sapphire/GaN MOCVD template consists of a few narrow peaks, suggesting that the new deposited GaN is composed of small grains of about 0.1– 0.2 mm, misoriented by a few angular minutes. The curves have relatively large total width due to the significant bending of the sapphire/GaN system [26]. The FWHM of GaN grown on the HVPE-GaN is of the order of 0:1–0:15ı for the reflection 002 and the beam dimension of 3  10 mm2 . The 60 m thick pressure grown material, cut from the HVPE-GaN seed for the reflection 002 and beam 0:310 mm2 , showed the FWHM of 0:04ı and the bending radius of 30 cm. In turn, during the LPE runs on pressure grown GaN platelets the bending phenomenon is never observed and the rocking curve is always very narrow .0:01–0:02ı /, which suggests the best structural quality of the deposited material.

10.4.2 Modeling of the Convective Transport in Gallium for LPE Growth (Time Independent Solutions) The finite element calculation has been used for modeling the convective transport in the solution, during LPE process in positive temperature gradient configuration, especially for analyzing the effects of the crucible wall, seeds, and baffle. The LPE configuration with seed of diameter of 1 in. and thickness of 400 m has been selected for the modeling. The relatively small volume of Ga used in this configuration .5 cm3 / has resulted in its small Reynolds number which has assured convergences of the solutions of the calculations. The shortness of the crucible has permitted the obtaining of symmetrical temperature distribution in it. Thus, the radial temperature distribution at the bottom of the crucible could be approximated as parabolic and the vertical temperature distribution as linear (as is shown in Fig. 10.7a). The radial temperature distribution on the crucible’s top has been assumed as constant. The analysis of the convective transport in the gallium at twodimensional approximation has been carried out with FIDAP 8.5 supplied by Fluent Inc. The convectional flow velocity in the gallium and the temperature distribution

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a

N2

z

T(r) 1723 1673

1623

1623 1593

T substrate T(r)

b

Fig. 10.7 (a) Scheme of the experimental LPE configuration with the baffle plate; the temperature values and assumed radial and axial temperature distributions are marked; (b) stream lines in the gallium for the presented configuration

in the liquid, in the seed, and in the crucible wall have been determined. One has considered Ga as the incompressible, viscous and Newtonian fluid, where density changes have been produced only by temperature changes. The boundary conditions have been fixed by temperatures measured during the LPE process. They are presented in Fig. 10.7a. The mesh composed of 41,000 elements has been introduced. Assuming the laminar flow in the gallium the following set of equations has been solved in order to find stationary and time independent solutions: 1. continuity equation for Ga r  .Eu/ D 0;

(10.2)

where uE is a velocity vector for the given mesh’s element. 2. momentum balance equation for Ga

0 uE.r  uE/ D rp C r 2 uE  0 ˇT0 .T  T0 /g; E

(10.3)

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where T0 is the arbitrary chosen reference temperature .T0 D 1; 673 K/; 0 is the gallium density at T0 ;  the viscosity, ˇT0 the gallium thermal expansion coefficient at T0 ; T the temperature, gE the gravity vector and p pressure in the liquid metal given by the following formula: 1 p D  r  uE "

(10.4)

where " is a penalty parameter equal to 107 [30]. 3. energy balance equation for Ga,

0  cp .Eu  rT / D r  .krT /

(10.5)

where cp is the specific heat capacity of gallium, and k is its thermal conductivity coefficient. Since uE D 0 in the crucible and seed, for graphite and seed the following equation has been taken: r  .ki rT / D 0; (10.6) where ki represents the thermal conductivity coefficient for graphite or the seed respectively. The solution of the equations presented above has allowed us to determine: uE and T and present them graphically as the stream lines, velocity vectors, and isotherm lines for the examined system (graphite crucible, seed, and first of all the liquid gallium). Some of the material properties of gallium, graphite, and seeds used in the simulations are presented in Table 10.1. Since the temperature dependence of the specific heat of liquid gallium is not strong, one has considered cp as constant for the temperature range used [31]. However the thermal conductivity of gallium has been determined for the given temperature from the following equation [31]: k0 D 22:3 C 1:8  .T  273:15/0:5 :

(10.7)

The gallium viscosity has been approximated by the relation [31]:  .T / D 1:729 

T  273 350

n ;

(10.8)

Table 10.1 Properties of gallium, graphite GaN, and sapphire used in the simulations Density Specific Thermal Thermal expansion heat conductivity coefficient Gallium Graphite GaN

5;360 kg=m3 [31] 2;300 kg=m3 [32] 6;100 kg=m3 [33]

374 J/kgK [31] 202 J/kgK [33] 500 J/kgK [34]

See the text 30 W/mK [32] 24 W/mK [35]

0:11  103 1=K [31] – –

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where n.T / D 1:135  0:41  103 .T  623/ C 0:80  107 .T  623/2:

(10.9)

In general, the modeling of the convective transport in the solution has allowed us to demonstrate that the measurements of the temperature distribution in the wall and bottom of the crucible have not reflected the real temperature distribution in the liquid gallium [27]. For the given temperature distribution the convective rolls’ behavior has been totally different for the gallium filling the crucible and the gallium being free (crucible neglected). It has also been shown that the substrate in the crucible can change the order of the convectional flow in the solution [36]. It depends on its thermal conductivity coefficient and a relation of this coefficient to the thermal conductivity of the crucible material. Since the temperature gradients realize in the material of the lowest coefficient of thermal conductivity, the seed of the coefficient lower than for graphite (case of sapphire) can cut off the influence of the radial gradient (controlled in the bottom) for the mass of liquid gallium. For the same reason an introduction of the graphite material to the gallium (case of the baffle) also perturbs and changes the roll’s and isotherm lines’ distributions in the liquid. Figure 10.7b shows the streamlines for the configuration with the baffle plate and the GaN crystal as seed. One can see two pairs of the convectional rolls above the baffle. They are separated by the graphite stick. Below the baffle two rolls are observed. There, the gallium flows along the baffle to the center, then comes downwards and backs along the substrate to the sides of the crucible. The maximum velocity of the convectional flow above the baffle is 1 mm/s. Under the baffle the convection is very weak; the maximum flow velocities in two rolls achieve 0.2 mm/s. The nitrogen is transported to the baffle by relatively strong convection, then the N atoms are dispersed on the baffle but by the sides (opening areas) they are supplied below the baffle to the seed region, where their velocity is low and thus by very weak convection process they are transported to the seed. It seems that due to very low convectional flow velocity the growth is macroscopically stable but the rate as slow as 1 m=h.

10.4.3 Seeded Growth with Convective Flow of Gallium Under Control Bulk GaN crystallization may take place when a properly prepared GaN seed is wholly immersed in the liquid Ga:N solution. Then, the growth can proceed in any of the available crystallographic directions. Additionally, the control of the convective flow of gallium can give a proper nitrogen transport to a supersaturated zone of the solution and therefore stable and relatively fast GaN growth. Figure 10.8a presents one of typical experimental configurations used for the bulk growth by HPS method: a graphite crucible with special graphite baffles system, tube, seeds, thermocouples positions, and temperature values measured during crystallization runs.

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a 1623 GaN seed 1623 1630

1643 1673

Fig. 10.8 (a) Scheme of typical experimental configurations used for the bulk growth by HPS method: a graphite crucible with special graphite baffles system, tube, seeds, thermocouples positions, and temperature values expected during typical crystallization runs; (b) typical HVPE seed; in its upper part the special slot allowing the hanging of this crystal in the crucible is visible; (c) photograph of the real experimental configuration; four hung seeds, the baffle, and the system of tubes are visible

This is a two-zone reversed temperature gradient configuration; the hot part of the solution is at the lower part of the crucible and the cooler part is above. The convective flow in the liquid gallium is provided and controlled by the radial temperature gradient. The seeds have been placed in the Ostwald–Miers zone of the solution, supersaturated metastable zone in which crystallization is possible, but spontaneous

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nucleation does not take place. The search for this zone for presented configuration is described in details in [37]. The HVPE crystals have been used as the seeds. Their surfaces have been prepared by mechanical polishing and reactive ion etching. The crystals have had a special slot (see Fig. 10.8b), which allowed their hanging and holding by the graphite sticks, at fixed places in the liquid solutions. The use of the HP reactor of the internal diameter of 10 cm has allowed placing in the crucible several, bigger than 1 in., seeds in one crystallization run. The photograph of the configuration described above is shown in Fig. 10.8c. The seeds have been positioned perpendicular to the crucible radius, although they can also be positioned in parallel or in free chosen directions. The growth rate has not been dependent on the crystallographic directions or planes and varied from 1 to 3 m=h. Again, like for LPE configuration, a decreasing of growth rate in time has been observed. Therefore, for obtaining more than 300 m thick GaN layer on one of the seed planes, the minimum 200 h crystallization process has had to be carried out.

10.4.4 Growth on HVPE Seeds in the c-Directions During the long period of crystallization under pressure .400 h/, the HVPE seeds have been grown up to 700 m per side. The growth in nonpolar directions has not been observed. The parasitic crystals have been grown on the seed’s edges. The growth of GaN “wings” from the edges of the seed, on the gallium side, has also been detected. These phenomena have been associated with the local (close to the seed edges) increasing of the supersaturation. The cross section analysis of the growing samples (see Fig. 10.9) has shown that the new grown material has been deposited in a uniform way, from the thickness point of view. The nitrogen side has been grown showing the presence of several hillocks (growth centers). On the contrary, the gallium side has been grown by macrosteps propagation. Cellular instabilities have been observed on this side. They have increased in the time of the crystallization run. Figure 10.10a represents the photograph of one of the seed used for the bulk growth. Figure 10.10b shows the same seed after 300 h of the crystallization process. The part of this crystal has been cut perpendicular to the c-direction and the pressure grown free-standing material has been obtained (see Fig. 10.10c). The Hall effect measurements on Ga side of the HP free standing crystal have shown that the carrier concentration in the material has been uniform- about 5  1019 cm3 . This result has been confirmed by the X-ray measurements of the c parameter in the crystal. The variation of this parameter has been 0.0003 Å (see Fig. 10.11a). This variation corresponds to the difference in the free carrier concentration of 1:4  1019 cm3 [38]. In turn, the FWHM has been of the order of 0:015ı for the reflection 004 at the beam of 0:2  1 mm2 (see Fig. 10.11b). The bowing of the crystal has been of the order of 10 m. An evident change in the FWHM of the HP-GaN coupled to the HVPE seed and the HP free standing GaN crystal should

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N side

HVPE seed

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{1120} cleaved

{1010} polished

Ga side

Fig. 10.9 Cross section of GaN crystal grown by HPS method on HVPE seed; grid 1 mm

Fig. 10.10 (a) Free standing HVPE crystal (700 m thick) used as a seed; (b) free standing HVPE crystal after HP crystallization process (1.5 mm thick); (c) 300 m thick hp-GaN crystal cut from the HVPE seed; grid 1 mm

be noted. The FWHM for the HP-GaN coupled to the seed has been of the order of 0:12ı for the reflection 002 and beam dimension of 3  10 mm2 . The FWHM of the free standing HP-GaN has been of the order of 0:025ı for the same reflection and beam dimension. It seems that for the free standing material, some stresses due to the difference in lattice constant between HVPE seed and pressure grown GaN crystal have been released. The etching process in KOH/NaOH eutectic at 450 ıC of the HP free standing GaN crystal has shown that the etch pits density on the gallium side has been 1  106 cm2 . The main disadvantage of this material has been the presence of graphite precipitates. Their density has locally attained 102 cm2 .

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a 11

5.18625 5.18628 5.18630 5.18633

Y [mm]

10 9

5.18636 5.18639

8

5.18641 5.18644

7

5.18647 5.18649 5.18652

6

5.18655 5.18658 5.18660 5.18663 5.18666

5 4

5.18668 5.18671

3

5.18674 5.18676

2

5.18679 5.18682 5.18685

1

5.18687 5.18690

0

b

1

12000

2

3

5 4 X [mm]

6

7

8

FWHM=0.015 deg. Beam: 0.2 mm x 1 mm

counts

8000

4000

0 36.40

36.47 omega, deg

36.54

Fig. 10.11 (a) Variation of the c parameter in the pressure grown free-standing crystal; the difference between green and red areas corresponds to 0.0003 Å; (b) X-ray rocking curves for this crystal

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10.4.5 Growth on HVPE Seeds in Nonpolar Directions The seeded growth in nonpolar directions under high nitrogen pressure was carried N out only in <1010> direction. The seed for this growth was prepared by HVPE crystallization of thick GaN layer (few millimeters) on the sapphirenGaN MOCVD template. The˚HVPE  free standing material was cut in such a way that small platelets N planes were obtained. The photograph of the seed is presented of GaN with 1010 in Fig. 10.12a. The crystallization was carried out at typical parameters already presented in this chapter (see Fig. 10.8a). After 100 h the seed grew about 150 m per side. The growth mode was unstable and it was very difficult to detect. Some regions with rectangle steps were observed. The parasitic deposition of small GaN crystals on the seed surfaces was watched. Thus, the crystal was polished on both sides. Since we were interested in free-standing material the sample was polished in such a manner that one of its side and the seed were removed. The X-ray measurement showed that the FWHM of the free standing crystal was 0:05ı for the reflection 100 and beam diameter of 3  10 mm2 . In the luminescence spectrum measured at 10 K the strong peak of the yellow luminescence was found. This suggested that the material was n-type. The Raman spectroscopy confirmed the good structural quality of the crystal and showed that the free electron concentration was higher than 5  1018 cm3 (the Raman spectrum is shown in the Fig. 10.12b). The Hall Effect measurements confirmed it. The free carrier concentration was of the order of 1:6  1019 cm3 (at mobility of 95 cm2 =Vs) thus lower by a factor three than the free carrier concentration in HP-GaN grown in polar direction. It ˚ suggests  that more N planes. oxygen can be built on the polar surfaces than on the nonpolar 1010 The crystal was etched in KOH/NaOH eutectic at 450 ı C in order to reveal the etch pits. Only one kind of pit (in sense of their size) was detected. The defect density on the 70% of the crystal surface was 3  106 cm2 , on 20% was 5  105 cm2 , on 10% was 5  107 cm2 . Figure 10.13 presents the surface of the crystal after the DSE procedure. The regions with various defect densities are marked.

10.4.6 Modeling of the Convective Transport in Gallium (Time Dependent Solutions) In comparison to the previous calculations presented in the Sect. 10.4.2, the stationary and time independent solutions have been resigned. The time-dependent set of equations has been used. The change of the equations has been associated with an increase of the size of the experimental setup and changing of the dimension of the calculated system from 2D to 3D. With an increase in the volume of gallium, the Reynolds number has increased too and the solutions have not been convergent in the stationary model. In a case of time-dependent solutions the convergence has been obtained. Thus, the flow in any moment of the simulation and the change in the mass transfer in time could be analyzed.

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˚  Fig. 10.12 (a) Seed; 101N 0 plane of HVPE-GaN grid – 1 mm; (b) Raman spectrum of the pressure grown free standing material; the crystal is also presented; the c-direction is marked; grid – 1 mm

Fig. 10.13 Pressure grown nonpolar GaN crystal after the DSE procedure with three regions of various defect densities

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The analysis of the convective transport in the gallium at three-dimensional approximation for the configuration presented in Fig. 10.8a has been carried out with the code Fluent supplied by Fluent Inc. All assumptions have been the same as for calculations presented in Sect. 10.4.2. Only the mesh has been changed and composed of 2  106 elements. In order to find time dependent solutions the modified set of the equations (10.2)–(10.6) have been solved:   r  uE D 0;   @Eu E

0 C uE.rEu/ D rp C r 2 uE  0 ˇT0 .T  T0 /g; @t   @T C uE  rT D r  .krT /:

0 cp @t

(10.9) (10.10) (10.11)

Since uE D 0 in the crucible and seeds, the following equation was taken for graphite and GaN crystals: @T D r  .ki rT /;

i cpi (10.12) @t where ki represents the thermal conductivity coefficient, cpi the specific heat capacity, ¡i the density, for the graphite and gallium nitride, respectively. The solution of equations presented above has allowed us to determine uE and T in any time of the simulation as the initial conditions uE D 0 has been taken. The simulations have been carried out to the moment of time when the solution for isotherm lines has stopped its evolution and has not been changed in time. The material properties of gallium, graphite, and GaN have been the same as those presented in Table 10.1. The heat capacity .cp / of gallium, graphite, and GaN has been taken as constant for the temperature range used. The thermal conductivity and viscosity of gallium for the given temperature have been determined according to (10.7)–(10.9) presented earlier. Figure 10.14a represents the convectional velocity vectors in the gallium at 2,270th second of the simulations for the temperature distribution presented in Fig. 10.8a. As seeds, GaN platelets with the dimensions 30  20  1 mm3 have been assumed. The hot gallium flows up through the graphite tube with the maximum velocity of 2.20 cm/s at the upper part of the crucible (red area in Fig. 10.14a). The main flow jets from the tube and hits the graphite closure of the crucible. Then, it flows to the sides and down. The flow close to the seeds is slower, up to 3 mm/s. Figure 10.14b shows the region just close to the seed. Between the wall of the crucible and the seed the gallium flows down but in the region between the tube and the seed it flows up. The correlation between direction of the gallium flow and morphology has not been found. The experimental configuration presented in Fig. 10.8a appears to be very interesting and important for further crystallization processes. Varying the tube diameter, thickness of the tube’s wall, or its length, the flow velocities and rolls positions can be changed for the given temperature distribution. Varying the radial temperature gradient (the temperature at the bottom of the crucible) the flow velocities can be changed for the given crucible configuration. For example at radial temperature

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Fig. 10.14 (a) Convectional velocity vectors in the gallium at the end of the simulations (in 2,270th second); simulation for the temperature distribution presented in Fig. 10.8a, (b) convectional velocity vectors in the gallium close to the seed

gradient of 5 K, the maximum convective velocity has only been 2 mm/s and the flow velocity close to the seed has been lower than 1 mm/s [37]. Finally, it should be possible to find an appropriate seed place in the solution, an appropriate place for relatively fast but stable growth of GaN. For that, however, an equation for the nitrogen concentration in liquid gallium should be added to the set of equations presented in this subchapter. Then, the supersaturation in any place of the solution might be determined.

10.5 Applications of Pressure Grown GaN Substrates: Blue Laser Diodes in TopGaN Ltd It is well known that GaN substrates for nitride based LDs technology should have high structural quality and low threading dislocation density. The substrates have to be relatively thick, for a possibility to miscut them appropriately for epitaxial growth and so that they are highly conducting which allows for preparation of stable and low-resistance ohmic contact. The pressure grown GaN crystals (HPGaN) almost fill these conditions. Spontaneous HP-GaN platelets have, as shown in Sect. 10.3.2, very good structural quality and very low defect density. They are also highly conductive. Unfortunately, they are relatively thin and miscuting them is rather impossible. For the use of the platelets as substrates for epitaxy by both MOCVD [39] and MBE [40] methods, the mechano-chemical polishing procedures have been used.

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Fig. 10.15 Structure of a laser diode grown on bulk GaN crystals in TopGaN Ltd

Figure 10.15 shows schematically a typical structure of a LD fabricated in TopGaN Ltd. The active layer of the device has been composed of one to five In0:1 Ga0:9 N quantum wells. The diode has been mounted p-side down on metalized diamond heat-spreaders. TopGaN LDs during CW operations have shown a threshold voltage of around 6 V@5 kA=cm2 at emission wavelength from 395 nm to 420 nm. The lifetime of the devices has achieved a few thousands hours. The LDs structures have been characterized by a density of dislocations at the level of around 105 cm2 . The origin of the dislocations has been the large mechanical strain existing in cladding and blocking AlGaN layers, high magnesium content in p-type part of the epi-structure, and segregation tendency of indium in InGaN quantum wells [41]. In spite of the relatively high defect density and graphite precipitates, pressure grown free standing HP-GaN crystals, described in Sect. 10.4.4, have also rendered good substrates for fabrication of the blue LDs. The defect density in these laser structures (made by MOCVD) has never exceeded 5  106 cm2 . During CW operation the lasers of the emission wavelength of 395 nm showed a threshold voltage of around 5:5 V@5 kA=cm2 . The lasers also show very satisfying spatial distribution of the emission wavelength with the variation 395 ˙ 2 nm. No doubt, it is due

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to good substrate properties - first of all applicable thickness allowing miscut of the substrate for the epitaxy. Today, at TopGaN and the Institute of High Pressure Physics laser technology level, it appears that the more important features of good substrates are their thickness and the possibility of miscutting them rather than their exclusively low defect density.

10.6 Summary and Perspectives of HPS Growth Method In this chapter the state of the art HPS growth method has been described. The spontaneous growth of platelets and their use as substrates for blue LDs by TopGaN Ltd. has been presented. The progress in seeded growth by HPS method has also been shown. The pathway from LPE to the crystallization on the seeds with convective flow of gallium under control has been demonstrated. The modeling of the convective transport in gallium has been used to better control the gallium flow during the crystallization process. The new class of HP-GaN crystals has been obtained. These crystals, grown by HPS method on the HVPE seeds, satisfy well the conditions required for the substrates for LDs. Today, their main disadvantage is the presence of graphite precipitates and defect density of the order of 106 cm2 . Despite these, they have been successfully used as substrates for the LDs. In the future, the growth on the HVPE-GaN crystals as seeds would be in obtaining HP-GaN crystals bigger than spontaneously grown platelets, uniform in shape, and in physical properties. Obviously, the spontaneous growth of the platelets is still being carried out, since these crystals seem to be very good seeds for crystallization with convective flow of gallium under control and/or for the HVPE growth (see Chap. 3 of this book) due to their extremely low dislocation density. Generally in the HPS method, the growth rate should be increased. Most probably, addition of appropriate impurities to the solution may be very helpful. As has been shown by Kawamura et al. [42] the LPE at sodium flux method with Li additive allows the obtaining of growth rate of the order of 20 m=h. Another approach is to crystallize ˚  in nonpolar and semipolar directions. The first tests of the growth on the N planes have shown that the growth rate has not been changed or increased. 1010 It seems that rather highly indexed planes of the seeds should be used. However for that, thicker HVPE free standing crystals are needed. Thus, the HVPE technique and crystallization on the sapphire/GaN MOCVD templates and the pressure grown crystals have to be developed fast. The growth with appropriate impurity in the solution and on appropriate seed (with highly indexed planes exposed to the solution) becomes the goal of our future work. Acknowledgements This chapter could be written thanks to the results obtained inter alia during realization of two scientific projects: Research Grant No R0203601 supported by the Polish Ministry of Science and Higher Education and Structural Founds Grant No WKP_1/1.4.3/2/2005/ 13/132/322/2007/U supported partially by the Polish Ministry of Science and Higher Education (34%) and the European Union (66%). The authors would like to thank Drs E. Litwin-Staszewska, J. Domagala, M. Krysko, H. Teisseyre, A. Khachapuridze, and G. Kamler for the Hall, X-ray, Raman, luminescence, and DSE measurements.

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Chapter 11

A Brief Review on the Na-Flux Method Toward Growth of Large-Size GaN Crystal Dirk Ehrentraut and Elke Meissner

Abstract The growth of bulk GaN by the Na-flux method is reviewed. The largest GaN crystal thus far produced is two inches in size with thickness of a few millimeters along (0001) axis. The historical development of the method, experimental conditions, and quality of the grown GaN crystals are summarized.

11.1 Introduction The sodium (Na)-flux method in its original version is one of the older approaches to grow GaN; however, a relatively new development among all the techniques toward the crystal growth of bulk GaN is the Na-flux method as a liquid phase epitaxy (LPE) variant. Firstly, Na was named as the flux component, the other part being gallium (Ga), and later a larger number of other group-I and -II elements like lithium (Li), potassium (K), calcium (Ca), strontium (Sr), barium (Ba) have been worked with and their effects on the crystal formation was studied by various research groups in the past. Major progress in terms of crystal size, quality, and scalability of the crystal growth system has been reported by the group from the Osaka University who employs the LPE technique [1]. The first growth of a 2-inch size GaN crystal has been reported in 2006 [2]. This chapter will therefore have its main focus on the achievements reported by the Osaka group, which by now is considered the leader in this particular technology, and will begin with a brief overview over the historical developments in the Na-flux technique.

11.2 Historical Development in Brief Yamane et al. have observed in the mid-1990s that the synthesis of ternary nitrides like Ba2 ZnN2 ; Sr2 ZnN2 , etc. from a Na flux sometimes yielded hexagonal GaN crystals as stable phase [3]. They consequently focused on the preparation of GaN 235

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crystals from the Na flux using sodium azide .Na3 N/ as Na source and were first to report the growth of hexagonal GaN crystals up to 2 mm in size in 1997 [4]. Sealed steel tubes were used at temperatures .T / ranging between 600 and 800 ı C, pressure .p/  11 MPa, and under nitrogen (N) excess. A high Na content is required to form GaN and it was reported that the formation of GaN requires less amount of Na when the temperatures rises. Moreover, the GaN size increased with decreasing Na content in the Ga–Na solution. This technology had been improved and soon Aoki and Yamane reported the growth of up to 3 mm large GaN crystals with an X-ray full-width half-maximum (FWHM) from the (0004) diffraction peak as small as 25 arcsec, synthesized by heating a Na–Ga melt placed in a boron nitride (BN) crucible at T D 750 ı C and p D 5 MPa of N2 for 200 h growth time [5]. Yields and morphology of GaN single crystals varied with the molar ratio Na=.Na C Ga/ in the starting composition. Electrical properties were measured: resistivity of 0:04 cm, carrier concentration of 1–2  1018 cm3 (n-type), and carrier mobility of 100 cm2 V1 s1 . Cathodoluminescence (CL) spectra at room temperature revealed a single peak around  D 362–365 nm. The crystal size was further increased to 3  5 mm2 platelets and bulky prismatic crystals of 1 mm along the c-axis [6]. The average growth rates were 14 and 2 m h1 along c direction for the platelet and prismatic crystals, respectively. Sodium was later speculated to act as catalyst for the dissociation of N2 molecule into N ions by donating electrons [7]. Consequently, a high content of Na would accelerate the introduction of N ions into the Na–Ga melt where Nax –.GaNy / complexes would be formed. This phenomenon has been further investigated by Iwahashi et al. and Kawamura et al. [8,12]. They concluded that N2 would be ionized at the gas–liquid interface of the Na flux at T > 900 K and successively entered into the Ga–Na flux. Pressure is also promoting the solubility of N radicals. The additives Ca and Li in the Ga–Na flux improve the quantity of dissolved N, obviously by the formation of a metastable phase in the flux, thus decreasing the N activity. Some of the larger crystals are shown in Fig. 11.1. A maximum size obtained for self-nucleated GaN crystals was around 6 mm along the c-axis. Using a selfnucleated GaN crystal as seed, a second generation of GaN was deposited at T D 760 ıC at low average growth rates around 0:2 m h1 for the nonpolar faces [9]. Li3 N was added to the flux to improve the solubility of N. Figure 11.2 compares the size of a seed and a GaN crystal grown on such seed. The growth rate for seeded growth was increased to about 2 m h1 though they are assumed to be higher because of initial dissolution of the seed crystal [10]. Growth temperature and pressure were T D 900 ıC and 0:8  p  7 MPa of N2 for 72 h and the Ga melt was heated in Na vapor, which forms the Na–Ga melt. The presently largest GaN crystal grown by the Na-flux method was produced by a group from the Osaka University [2, 12]. Figure 11.3 depicts the 2-inch size GaN crystal fabricated on a (0001) HVPE-GaN substrate. The first report by the Osaka group was published in 1998 [13]. In order to increase the crystal size as well as obtaining oriented crystals, the LPE technique employing a substrate was used successfully [14].

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Fig. 11.1 GaN prismatic and needle crystals prepared at 900 ı C and 9.5 MPa for 400 h. (Image courtesy of H. Yamane, Tohoku University, Japan.)

Fig. 11.2 Self-nucleated GaN seed and GaN crystal grown on the seed. (Image courtesy of H. Yamane, Tohoku University, Japan.)

It was recognized that a forced flow of the solution in the vicinity of the growing GaN crystal would lead to improve growth conditions, hence, leading to increase the growth rate. The flux-film-coated LPE method (FFC-LPE) was introduced in 2003 [15]. The heated growth chamber with the flux-containing crucible was mounted such that swinging could be applied. At a swing frequency of 1:5 min1 , the growth rate of 4 m h1 at T D 800 ıC was achieved although the nitrogen pressure was as low as p D 0:95 MPa. Nonpolar substrates are also in the focus and the fabrication of an a-plane GaN substrate has recently been reported [16]. Strontium (Sr) was added as the additive to

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Fig. 11.3 A 2-inch GaN crystal grown by the Na-flux LPE method. About 0.5 at% C had been used as additive to reduce the parasitic nucleation of GaN [12] (Reprinted with permission. c 2008, Elsevier Publishers.) 

gain the desired morphology. X-ray diffraction revealed the improved crystallinity, N reflection plane decreased from 1,152 to 236 arcsec i.e., the FWHM from the .1120/ for the GaN template (MOCVD GaN film on Al2 O3 ) and the LPE-grown crystal, respectively. Aside from the Na flux, there are some more attempts using alkaline and earth alkaline salts to obtain a reasonable solubility of GaN under moderate conditions. Feigelson and Henry have recently employed the system LiF  BaF2  Li3 N to grow GaN single crystals about 0.5 mm long [17]. The growth temperature was 800 ıC and pressure about 2.5 MPa. Song et al. used the mixture Li3 N C Ga [18] and Wang et al. made use of the mixture Li3 N C Ca C Ga [19] at temperature around 800 ı C and around 200 kPa N2 pressure. Free-standing GaN platelets of up to 4 mm in size were produced. As appealing as all the above mentioned low temperature and pressure approaches are, it must be however noted here that apart from the technology currently under development at the Osaka University, bulk GaN crystals of sufficient size by the Na-flux or related flux methods are yet to be demonstrated.

11.3 Experimental Conditions for the LPE Growth of GaN by the Na-Flux Method Experimental conditions are typically subject to modification in order to target a specific goal like doping, control of morphology, and the likes. Basically, the following procedure is employed according to [20]: Starting materials are Ga and Na,

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both placed in an alumina crucible. A GaN=Al2 O3 template would serve as seeding substrate and is placed at the bottom of the crucible prior to filling with the starting materials. A stainless steel tube resistant to high temperature and pressure is housing the crucible. Before mentioned steps are carried out in a glove box under argon (Ar) atmosphere to prevent the Na from the reaction with moisture from the air. Next, the steel tube is pressurized with N2 gas and heated in an electric furnace. GaN crystals are typically grown in the Ga–Na solution by maintaining the temperature (around 850 ıC) and pressure (around 5 MPa) over the entire growth time. After the growth has been completed, the tube is taken out from the furnace to remove the crucible, and the Na is dissolved in cold ethanol and water. Careful handling is requisite to prevent from strong reaction of the Na. Remaining Ga can be removed with ethanol and hot concentrated chloride acid.

11.4 Growth Mechanism and Dislocations 11.4.1 Effect of Flux Composition on Growth Stability and Crystal Morphology The effect of additives on solubility, growth stability, crystal morphology as well as alloying with GaN was studied to improve the GaN crystal growth and eventually to prepare new phases. Lithium increases the solubility of GaN probably through increased solubility of N [21]. Lithium-based fluxes like the Li3 N C Ga [18] and also Li3 N C Ca C Ga [19] were applied; however, it has been noted that excessive Ca in the flux hinders the formation of GaN. Bao et al. have recently reported the use of a flux composed of Ba3 N2 C Ga to grow small crystals of GaN [22]. As for the flux with Ca, higher concentration of Ba strongly reduces the formation of GaN and also affects the growth rates of the different crystal facets; consequently, the morphology of the crystal can be modified. Sekiguchi et al. have reported that using K instead of Na would yield cubic GaN crystals [23]. According to CL measurements, the cubic GaN crystals prepared and peaking around 3.2 eV were highly defective. Iwahashi et al. illustrated the change of the crystal morphology that is caused by Sr as an additive to the melt [16]: As the portion of Sr ˚ was increased from 0 to N 1.5 mol% with respect to Na, the formation of m-faces 1100 gained strength at ˚  N the cost of the pyramidal facets 1011 , Fig. 11.4. At 1.5 mol% Sr in the Na flux, self-nucleated GaN would only be bound by both basal faces (0001) and ˚ crystals  N and the 1100 N .0001/, faces. Further increase of the Sr content to 3.5% did not yield GaN crystals at all. Kawamura et al. stated on the effect carbon (C) would have as an additive [24]. The growth rate was significantly increased above 20 m h1 (even 30 m h1 was reported [24]) because of suppressing theparasitic nucleation and stabilizing the

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Fig. 11.4 Photograph of GaN produced on a-plane GaN templates in (a) the pure Na flux and c 2007, The Japan Society of (b) the Sr-added Na flux. (After [16]. Reprinted with permission.  Applied Physics)

formation of the nonpolar m-face. Already 1 at% C at T D 800 ıC would cause complete suppression of the parasitic nucleation. Moreover, seemingly C does not lead to strongly affect the C-level in the grown GaN. SIMS analysis revealed concentrations about 2  1017 cm3 which is the same amount as for those crystals grown under intentionally C-free conditions. The transition metal additives chromium (Cr), manganese (Mn), iron (Fe), cobalt (Co), and nickel (Ni) have been employed [25]. Chromium resulted in coprecipitation of CrN single crystals whereas Mn, Fe, Co, and Ni rather had an effect on the GaN crystal morphology in that growth in the c- direction was enhanced with Ni illustrating the largest effect. Only Mn could be incorporated into GaN at concentrations up to 0.35 at% as detected by ICPMS. Small crystals of the —-phase of Mn2 N could be prepared from an indium In–Na flux at 700 ıC [26].

11.4.2 Growth Mechanism and Effect on Dislocation Population Several factors do influence the defect formation and propagation during seeding and crystal growth, which was studied for the growth of GaN on (0001) MOCVDGaN substrates [27]. Namely, this is the concentration of dislocations at the initial growth stage (nucleation of GaN islands on the substrate), which can be as high as this of the substrate typically for MOCVD layers  109 cm2 . This is followed by the dislocation lines bending within a few micrometers above the interface of the substrate and the LPE grown GaN, which results in the low dislocation density of 103 cm2 [28] (Fig. 11.5).

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241 104~5 cm–2 LPE – GaN

106 cm–2

Interface MOCVD – GaN

108~9 cm–2

Fig. 11.5 Reduction of dislocations during the LPE growth of GaN on a MOCVD-GaN substrate c 2008, Elsevier Publishers.) by the Na-flux method. [12] (Reprinted with permission. 

11.4.3 Solubility and Growth Rate The Solubility of GaN is an essential parameter that would govern the crystal growth process. Kawamura et al. have investigated the solubility and reported that only 0.036 mol% of GaN can be dissolved in a flux of 27 mol% Ga and 73 mol% Na, while keeping the process running for 200 h at T D 973 K [12]. Adding Na to a Ga melt obviously improves the solubility of GaN and also allows applying lower pressure conditions. By comparison, Madar et al. reported a solubility of GaN in the pure Ga melt of only 0.008 mol% at T D 1; 473 K and p D 800 MPa [29] and Karpinski et al. have observed a rise in the solubility to about 0.16 mol% at T D 1; 773 K and p D 1; 600 MPa [30]. The solubility of GaN can be expressed in terms of system pressure and temperature. A minimum pressure is required above which crystal growth of GaN is commencing. Above T D 900 K, Na is clearly promoting the dissolution of N in the Ga–Na melt. As a result, about 0.12 at% N dissolves at T D 1; 073 K and p D 0:75 MPa in the flux of 27 mol% Ga and 73 mol% Na [9]. What is more, it was reported that the minimum pressure to grow GaN in the Ga–Na melt decreases until T D 1;023 K and after that increases. This phenomenon is explained in terms of the stronger rise of the solubility of GaN which exceeds that of N [9]. The growth rate depends on the thermal gradient between the substrate and the Na flux. Thermal convection inside the flux container is formed such way. Kawamura et al. recently reported a growth rate as large as 20 m h1 in the direction of h0001i [20]. They furthermore found that a dwell time of about 20 h from starting the experiment until first growth appears is needed, which is mainly caused by the slow saturation of the Na flux with nitrogen until reaching a level of supersaturation that initiates the growth process on the substrate. The control over parasitic nucleation is also a big challenge in the Na-flux method. It w as observed that they preferentially crystallize at the interface of the

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Table 11.1 Properties of GaN crystals grown by the Na-flux method Self-seeded GaN

GaN grown on substrate /

Impurities

a D 3:1903.3/ a D 3:1896.1/ a D‹ Å [18] Å [4] c D 5:1877 Å [5] c D 5:1864.6/ c D 5:1854.2/ Å [18] Å [4] / / 25 arcsec (0004) 40–60 arcsec (0002) [9] / / /

Resistivity

/

/

Carrier / concentration Mobility /

/

Lattice parameter

XRC FWHM

/

/

Na D 4:2  1014 cm3 [2] C D 2  1017 cm3 [24] /

0:04 cm (T D 110  350 K) [5] 1–2  1018 cm3 / (RT), n-type [5] 100 cm2 V1 s1 / (RT) [5]

liquid Na flux to the gas atmosphere, thus consuming a large fraction of the free N radicals needed to form the GaN on the substrate crystal. However, adding carbon (C) as additive would completely suppress the formation of parasitic GaN crystals and therefore support the growth of GaN on the substrate crystal [20].

11.5 Properties of GaN There is only very limited data available on the structural and physical properties of GaN produced by the Na-flux method. Table 11.1 has listed the lattice parameter a and c, XRC FWHM, concentration of the impurities Na and C, electrical resistivity, carrier concentration, and carrier mobility as far as available. Skromme at al. have investigated some optical properties of colorless prismatic GaN crystals grown by Yamane [31] by photoluminescence (PL), reflectance, and micro-Raman scattering. Raman scattering revealed an A1 .LO/ phonon mode at 739 cm1 indicating a free electron concentration around 2–3  1017 cm3 for 3–4 mm large platelets. Smaller platelets grown in a pyrolytic BN (pBN, purity > 99:9999%) crucible showed the A1 (LO) phonon mode at 733 cm1 , implying a free electron concentration in the mid 1016 cm3 . PL measurements disclosed a residual donor species with the binding energy of 33.6 meV, probably because of ON . Residual zinc (Zn) acceptors have also been evidenced peaking around 2.9 eV with a broader peak in the low-temperature .T D 1:7 eV/ spectrum. Other

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impurities evidenced were Si and Mg, both at low levels. The yellow band around 2.2 eV, and usually associated with gallium vacancies, was rarely observed. In the low-temperature .T D 1:7 eV/ PL measurement, the neutral donor-bound exciton .Do X/ emission was high in intensity with a small FWHM of only 2.2 meV. The authors concluded that the material was of excellent optical quality.

11.6 Industrialization Potential for the Na-Flux Method The Na-flux method has proven the possibility to grow GaN crystals under relatively mild conditions with respect to pressure and temperature, i.e., p  10 MPa and 600  T  900 ı C, respectively. The growth rate of up to 30 m h1 seems reasonable for industrial production purposes, particularly in the light of the intrinsically lower growth rates for growth techniques from the solution; refer also to Chaps. 7–10 and 12 in this book. The structural quality certainly can further be improved, especially in the view of single crystallinity and related grain boundaries for the wafer-sized area. A critical issue in terms of industrialization of the Na-flux technology, essentially for all technologies considered to produce GaN crystals for GaN wafer fabrication, is the feasibility for upscaling in order to achieve a high crystal throughput, which would positively affect the cost per epi-ready GaN substrate. The ammonothermal technique seems to be most promising in this regard. In fact, the cost for an epiready GaN substrate is an increasingly important issue considered by device makers as it covers the largest fraction in the cost structure for a mass-produced GaN-based device. For a vast sector of GaN devices, the cost for an epi-ready GaN substrate has to compete with that for the widely used sapphire-based templates that can achieve high-quality GaN devices through sophisticated template technology in combination with epitaxial lateral overgrowth (ELO).

References 1. Department of Electrical Engineering, Osaka University, Japan 2. F. Kawamura, H. Umeda, M. Morishita, M. Kawahara, M. Yoshimura, Y. Mori, T. Sasaki, Y. Kitaoka, Jpn. J. Appl. Phys. 45, L1136 (2006) 3. H. Yamane, F.J. DiSalvo, J. Solid State Chem. 119, 375 (1995) 4. H. Yamane, M. Shimada, S.J. Clarke, F.J. DiSalvo, Chem. Mater. 9, 413 (1997) 5. M. Aoki, H. Yamane, M. Shimada, T. Sekiguchi, T. Hanada, T. Yao, S. Sarayama, F.J. DiSalvo, J. Cryst. Growth 218, 7 (2000) 6. M. Aoki, H. Yamane, M. Shimada, S. Sarayama, F.J. DiSalvo, Cryst. Growth Des. 2, 119 (2001) 7. H. Yamane, D. Kinno, M. Shimada, T. Sekiguchi, F.J. DiSalvo, J. Mater. Sci. 35, 801 (2000) 8. T. Iwahashi, F. Kawamura, M. Morishita, Y. Kai, M. Yoshimura, T. Sasaki, J. Cryst. Growth 253, 1 (2003) 9. F. Kawamura, M. Morishita, K. Omae, M. Yoshimura, Y. Mori, T. Sasaki, J. Mater. Sci. Mater. Electron. 16, 29 (2005)

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10. M. Aoki, H. Yamane, M. Shimada, S. Sarayama, H. Iwata, F.J. DiSalvo, J. Cryst. Growth 266, 461 (2004) 11. T. Yamada, H. Yamane, Y. Yao, M. Yokoyama, T. Sekiguchi, Mater. Res. Bull. 44, 594 (2009) 12. T. Sasaki, Y. Mori, F. Kawamura, M. Yoshimura, Y. Kitaoka, J. Cryst. Growth 310, 1288 (2008) 13. Y. Mori, M. Yano, M. Okamoto, T. Sasaki, H. Yamane, Bull. Solid State Phys. Appl. 4, 198 (1998) (In Japanese) 14. M. Yano, M. Okamoto, Y.K. Yap, M. Yoshimura, Y. Mori, T. Sasaki, Jpn. J. Appl. Phys. 38, L1121 (1999) 15. F. Kawamura, M. Morishita, K. Omae, M. Yoshimura, Y. Mori, T. Sasaki, Jpn. J. Appl. Phys. Part II 42, L879 (2003) 16. T. Iwahashi, Y. Kitaoka, M. Kawahara, F. Kawamura, M. Yoshimura, Y. Mori, T. Sasaki, R. Armitage, H. Hirayama, Jpn. J. Appl. Phys. 46, L103 (2007) 17. B.N. Feigelson, R.L. Henry, J. Cryst. Growth 281, 5 (2005) 18. Y. Song, W. Wang, W. Yuan, X. Wu, X. Chen, J. Cryst. Growth 247, 275 (2003) 19. G. Wang, J. Jian, W. Yuan, X. Chen, Cryst. Growth Des. 6, 1157 (2006) 20. F. Kawamura, M. Morishita, M. Tanpo, M. Imade, M. Yoshimura, Y. Kitaoka, Y. Mori, T. Sasaki, J. Cryst. Growth 310, 3946 (2008) 21. M. Aoki, H. Yamane, M. Shimada, S. Sarayama, H. Iwata, F.J. DiSalvo, Jpn. J. Appl. Phys. 42, 7272 (2003) 22. H.Q. Bao, H. Li, G. Wang, B. Song, W.J. Wang, X.L. Chen, J. Cryst. Growth 310, 2955 (2008) 23. T. Sekiguchi, H. Yamane, M. Aoki, T. Araki, M. Shimada, Sci. Technol. Adv. Mater. 3, 91 (2002) 24. F. Kawamura, M. Tanpo, Y. Kitano, M. Imade, M. Yoshimura, Y. Kitaoka, Y. Mori, T. Sasaki, 5th International Workshop on Bulk Nitride Semiconductors (IWBNS-V), Itaparica, Salvador, Bahia, Brazil (2007) 25. M. Aoki, H. Yamane, M. Shimada, S. Sarayama, H. Iwata, F.J. DiSalvo, Jpn. J. Appl. Phys. 42, 5445 (2003) 26. M. Aoki, H. Yamane, M. Shimada, T. Kajiwara, Mater. Res. Bull. 39, 827 (2004) 27. F. Kawamura, H. Umeda, M. Kawahara, M. Yoshimura, Y. Mori, T. Sasaki, H. Okado, K. Arakawa, H. Mori, Jpn. J. Appl. Phys. 45, 2528 (2006) 28. F. Kawamura, M. Imade, M. Yoshimura, Y. Mori, Y. Kitaoka, T. Sasaki, Proc. SPIE, 7216, 72160B-1 (2009) 29. R. Madar, G. Jacob, J. Hallais, R. Fruchart, J. Cryst. Growth 31, 197 (1975) 30. J. Karpinski, J. Jun, S. Porowski, J. Cryst. Growth 66, 1 (1984) 31. B.J. Skromme, K. Palle, C.D. Poweleit, H. Yamane, M. Aoki, F.J. DiSalvo, J. Cryst. Growth 246, 299 (2002)

Chapter 12

Low Pressure Solution Growth of Gallium Nitride E. Meissner, S. Hussy, and J. Friedrich

12.1 Introduction Because of the extremely high decomposition pressure of GaN at its melting point [1], it is technically unfeasible to grow GaN from the melt by means of classical crystal growth techniques, like Czochralski, Bridgman, or Vertical Gradient Freeze, commonly used for the production of industrially important materials like silicon, GaAs, and many more. At present there are several promising techniques under development to produce thick GaN layers or bulk crystals, which are described in detail in other chapters of this book. Among the different crystal growth processes, the solution growth methods do have theoretical, as well as the already demonstrated practical potential for the production of GaN with the lowest possible dislocation density. The theoretical expectations on the solution growth are based on the fact that such a process would operate close to the thermodynamic equilibrium compared to any of the gas-phase methods and at much lower absolute temperatures. This expectation was successfully proven by Porowski et al. [2] with their high pressure solution growth method (HPSG) resulting in GaN material with an extraordinary low dislocation density of 10–100 cm2 . The HPSG method is performed under very high nitrogen pressure ranging up to 17 kbar and temperatures between 1,300 and 1;600 ıC [2], which indicates the obvious drawback of the method – the high pressure which needs to be applied. High temperatures are necessary in order to achieve a reasonable solubility of nitrogen in the gallium melt. The solubility of nitrogen in liquid gallium metal is very low [3] compared to that in classical solution growth processes. GaN can also be grown under much lower pressures and temperatures if it is not pure Ga-melt but gallium with the addition of a solvent. The solvent is required to enhance the reactivity and solubility of nitrogen in the gallium solution. Then, one could either still use pure nitrogen gas with a certain over pressure of up to, e.g., 50 bar to deliver nitrogen to the system or one can apply other alternative chemical species containing atomic nitrogen in order to get the process pressure down and maintain nitrogen solubility at a reasonable level. A certain variety of alkali or earth alkali metals, their respective nitrides, or mixtures thereof were already investigated by several researchers. The corresponding method originating from that approach 245

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is the so-called Na-flux technique, originally developed by Yamane et al. mainly between 1997 and 2000 [4, 5] and recently further investigated for the purpose of the design of a liquid phase epitaxy process (LPE) by Kawamura et al. [6]. The sodium flux method is based on the usage of a gallium–sodium melt made out of a mixture of the pure metals and NaN3 or through Na vapor generated in the system [7]. Other group one or group two elements from the periodic system were also tried as flux compositions, like lithium or lithium nitride [8, 9], lithium – potassium mixtures [10], or additions of certain amounts of calcium or carbon [11, 12]. The choice of the solvent is not easy, as it has to fulfill several criteria. In the ideal case, the solution additive should have a high solubility for nitrogen, but should not itself form a stable compound with nitrogen, which could compete with the formation of GaN under the given conditions, nor should it produce stable binary or ternary alloys, especially in the case of more than one element added to the solution. Finally, a handful of metallic elements from the periodic table remain, which could be promising candidates as additives for a GaN solution growth process, e.g., under reduced pressure. Systems with more than two elements other than gallium were not reported for the GaN crystal growth, since multicomponent systems are even more difficult to control in terms of solubility, thermodynamics, and reaction kinetics of possible phases formed. Logan and Thurmond [13] started already in 1972 with a very detailed investigation of the solution growth of GaN under a flowing gas atmosphere of ammonia and hydrogen, using a solution of Ga and Bi. They succeeded in growing epitaxial GaN layers on sapphire substrates. Elwell et al. [14], who had similar experimental conditions like Logan and Thurmond, used Bi and Sn as solvents and demonstrated again in 1984 that the growth of GaN from a solution is possible even under ambient pressure. Meissner et al. [15] and Hussy et al. [16] used other metal additives to the Ga melt, not applied so far by other research groups and managed to grow GaN under ammonia atmosphere from those solutions. The metals applied in their study do not have the same difficulties as sodium or potassium that have high vapor pressure. Their growth process was performed without any overpressure applied to the system. Table 12.1 shows a compilation of approaches, with the respective pressures, temperatures, and solvents, made for the solution growth of GaN (ordered after pressure). The sodium flux method still needs the application of certain overpressure and getting the pressure threshold down is unlikely, unless the chemical species delivering the nitrogen to the system is one other than the stable nitrogen molecule. Ammonia is the simplest choice to add the nitrogen into the growth process from the gas phase and adjust a condition close to the equilibrium curve of GaN, within the GaN stability field, thus preventing the GaN from decomposition during the process. Recently, over the last few years, a solution growth process for the growth of GaN under room pressure was developed and carefully investigated in detail at the Crystal Growth Department of the Fraunhofer Institute of Integrated Systems and Device Technology. The idea was to pick up the basic aims for the design of a crystal growth process for GaN, which could be of technical relevance because of simplicity, materials quality, scalability, and cost effectiveness. The result was the so-called low pressure solution growth method (LPSG), which is the subject of this chapter.

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Table 12.1 Data available in the literature for experimental conditions and solvents used for the crystal growth of Gallium nitride from a solution Solvents

Atmosphere

Temperature Pressure Reference, Year [bar] Œı C

NaN3 Na C NaN3 Na K Li Na C Ca Na C Li Na Li Li3 N Bi Bi, Sn Bi Ge Ge, Sn, Au, Ag Ge, Sn, Bi, Au, Ag and mixtures thereof

N2 N2 N2 N2 N2 N2 N2 N2 C NH3 NH3 N2 H2 C NH3 H2 C NH3 N2 C NH3 H2 C NH3 H2 C NH3 H2 C N2 C NH3

600–800 650–840 750 750 800 800 800 800 300–750 740–800 850–1,050 870–990 900 930–990 960 900–1,050

110 120 50 70 50 30 15 10–20 1 1–2 1 1 1 1 1 1

Yamane [17], 1997 Yamane [4], 2000 Aoki [5], 2000 Yamane [18], 2000 Morishita [19], 2003 Kawamura [10], 2002 Morishita [19], 2003 Iwahashi [20], 2003 Barry [8], 1998 Song [9], 2003 Logan [13], 1972 Elwell [14], 1984 Klemenz [21], 2000 Hussy [16], 2008; this chapter Hussy [22], 2008; this chapter Hussy [23], 2008

Most of the studies mentioned above lack a thorough investigation of the solubility of nitrogen in the chosen system and detailed considerations on the formation reaction of GaN from the particular system. Therefore, important information for the design of a technically relevant production process based on solution growth technique is missing for the crystal growers. The mechanism of action of the solvents in the solution is not sufficiently clarified. The solvents influence not only the solubility of nitrogen, but also the kinetics of nucleation and crystallization of GaN in the Ga-solution. In turn, very relevant questions necessary to be addressed like control over the nucleation density, growth rates, crystallization mechanism and others, which are related to the basic understanding of the influence of the solvents on the formation reaction of GaN as a function of the process parameters could not be answered so far with data available from the literature. The thermodynamic aspects of the “pure system,” namely the GaN formation from a Ga melt under nitrogen atmosphere and the solubility of nitrogen in pure liquid Ga, were, however, quite well investigated for the HPSG process [3, 24].

12.2 Technology of Solution Growth Under Ambient Pressure, the LPSG Method The LPSG technique is a relatively simple method for growing GaN, as it can be performed in easy horizontal or vertical furnace setups. Basically, the LPSG technique is very flexible. It can be modified in order to aim the growth of GaN bulk crystals, as

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well as it can be considered as a variant of multiwafer liquid phase epitaxy methods. Like with any other method, some important features must be addressed in order to gain a stable growth process and good quality material. A thorough study of this technique was performed during the last few years [15, 16, 25–27] and a detailed review of our work will be given in this chapter. The study answers at least some of the basic questions mentioned above. This chapter will also emphasize and discuss most recent achievements and the remaining challenges in LPSG GaN growth.

12.2.1 The Formation Reaction of GaN Under Ammonia Atmosphere The Fig. 12.1 shows the equilibrium pressures of N2 (upper region) and NH3 respectively over GaN as a function of temperature. The data were compiled from the literature and can be found in [28] and [29]. The stability curve was determined by thermal decomposition of GaN, the open and solid symbols represent the experimental data. The HPSG growth of GaN happens in the stability region of the stable nitride close to the equilibrium curve. The sodium flux technique works slightly away from the GaNGa C N2 phase boundary, but well situated within the GaN stability field, for the benefit of lower process pressure. Logan and Thurmond [13] and Elwell et al. [14] observed the decomposition or formation of GaN under ammonia atmosphere instead of pure nitrogen, in order to determine the boundary curve for the coexistence between GaN with NH3 . The LPSG growth of GaN and the LPE approaches from [13] and [21] were located in this respective lower region of the diagram in Fig. 12.1. Many efforts have been made in the last few years to control nucleation density, avoiding homogeneous nucleation in the solution and to establish a stable growth process by a careful choice of the solution composition, gas fluxes, process pressure, and temperature. Sun et al. [25] studied the formation reaction of GaN from a solution under ammonia atmosphere in order to retrieve the necessary data for the crystallization of GaN during the LPSG process. The work of Sun contains a systematic approach by means of high resolution thermogravimetry (TG); the details of the investigation can be found in the related publications [28, 30]. These considerations have the unquestionable advantage that an in situ technique was applied and that the measurements not only reflect a comparable situation to the LPSG growth process, but also to other solution growth techniques. The overall chemical reaction for the nitridation of gallium in an ammonia atmosphere can be simplified in the following expression: 2Ga C 2NH3 ! 2GaN C 3H2 During the TG experiments, the conversion of Ga from the solution to GaN crystals or epilayers was monitored as a weight signal as a function of reaction time, which compares to the ongoing crystal growth process. The formation of GaN from the

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Fig. 12.1 Equilibrium pressure of N2 (upper region) and NH3 (in H2 ) (lower region) over GaN as a function of temperature. The solid symbols denote experimentally found stable GaN and the open ones denote decomposed GaN. For the high and low pressure solution growth methods the calculated equilibrium (partial) pressures are indicated by dashed and solid lines, respectively

solution is expressed as conversion ratio ˛ which is plotted vs. the reaction time. ˛ is defined as mt ˛ D Ga ; m0Ga where m0Ga is the initial mass of Ga and mtGa the mass of gallium which has reacted to form GaN until the time t. As described in [29], the conversion ratio ˛ can be related to the total weight change measured during the TG experiments as a function of process time. Generally it is assumed that the formation of GaN from the Ga-solution occurs by a nucleation step followed by subsequent growth of a nucleus into a crystal as discussed, e.g., by [31]. For the solution growth process like LPSG, there are in fact several stages for the formation of GaN, first characterized by an initial reaction upon the introduction of NH3 into the system, followed by the so-called induction period for the formation of nuclei. The reaction accelerates when the nuclei start to grow and pass into a period of crystal growth with a constant formation rate .d˛=dt D const:/, which compares to the finding of a constant growth rate of GaN with time during the LPSG process if a stable growth condition is maintained (as discussed later in this chapter). The induction time tind is an important parameter, and although it is not a fundamental property for the system, it can be used to describe the nucleation rate J , which can be written as, 1 J DK tind with K being a constant.

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However, some assumptions are needed for this formulation, which can be found in a thorough study and interpretation of the induction period of crystallization processes in the paper from Söhnel and Mullin [32] and for the LPSG process in [29]. An empirical formula was expressed by He [33], which can be applied to describe the influence of temperature on the induction time and hence can give an idea about nucleation. Generally it is found that a high supersaturation leads to a high reaction rate and a short induction period, and in turn a high nucleation rate in the solution is expected. Thus, from the observation of the induction time, one can draw conclusions as to the supersaturation present in the growth system, since the nucleation rate J is proportional to the supersaturation S by the following formulation, where KN and n are constants: J D KN S n : Sun et al. [29] demonstrated how this relation can be used to retrieve the equilibrium partial pressure of ammonia for the LPSG growth process, by the observation of the length of the induction period as a function of ammonia partial pressure, p.NH3 /. It is essential to know the equilibrium ammonia partial pressure, since this parameter must be carefully controlled in order to achieve a stable growth condition, as will be explained later in this chapter in detail. The study of the induction period as a function of growth temperature generated valuable knowledge about the activation energy for homogeneous nucleation in the system. It was found that the activation energy for the reaction of Ga under ammonia remained the same if the ammonia partial pressure was changed [29]. If the activation energy for the reaction is not a function of p.NH3 /, the underlying reaction mechanism is the same for all p.NH3 / applied. However, the activation energy of the reaction was found to be different if the GaN was crystallized from a solution rather than from the pure Ga melt, which gives interesting implications as to the action of the solvents in the growth solution. The solvents seem to act kinetically as well as in terms of enhancement of solubility for nitrogen. In view of the growth rates, it would be important to define the rate limiting steps in the formation reaction of GaN. However, the determination of rate limiting steps is not easy. The soluble species is unknown and is expected to differ for any particular composition of a solution. Because the nitrogen as a component is supplied from the gas phase, the first interface that could define the reaction rates is the gas–liquid interface. Here the nitrogen compound must be dissociated and the atomic nitrogen must be dissolved in the melt or a soluble species must be formed. Krukowski et al. [34] examined the dissociation of molecular nitrogen on the surface of a gallium melt in the HPSG process by theoretical means and also Yamane et al. [35] described the dissociation of molecular nitrogen and the formation of a complex species which solves and contains the nitrogen. In the later case, sodium solvent plays an important role. The knowledge of the kind of soluble species would be a key issue in order to describe and understand the formation of GaN from a solution and to actively handle this parameter in view of the growth rates. But the soluble species remains subject to speculations or assumptions in theoretical considerations, as it is not easily accessible by any experimental technique. The approach made by Sun et al. [30] resulted in the cognition that the kinetics at the gas liquid interface plays

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an important role and represents a first trial of an experimental proof. The scenario of processes happening at the second rate limiting interface, the liquid–solid interface is not easily amenable by experimental means either; it consequently depends on the situation manifested in the solution “behind” the kinetics at the melt surface and may be even more difficult to investigate. Thus, the role of the solvents added to the solution growth is not fully clarified up to now, but it is definitely a combination of the enhancement of solubility for nitrogen in the Ga-melt and a kinetic influence on the key reactions at the respective interfaces.

12.2.2 Solubility of Nitrogen in Gallium-Metal Solutions The knowledge of the total concentration of nitrogen present in the Ga-melt is of great interest, in order to be able to quantitatively express a supersaturation state for the system and define the expected Ostwald–Miers region for stable growth. Moreover, if known, it would be possible to choose the most effective solvent with regard to nitrogen solubility. However, the access to the solubility by direct measurement is not easy, due to the very low values expected for the solubility of nitrogen in gallium or other metals. Thus, most of the solubility data were produced by ex situ analysis after quenching the sample [3,11,13,36] and the accuracy could suffer from desorption of N upon cooling the specimens. But, even at the highest available temperatures of 1;687 ıC, the solubility was not higher than 1 at%. In order to gain data for the solubility which would represent the condition in the low pressure solution growth under ammonia atmosphere and at room pressure, we made the approach to measure the solubility in-situ by means of thermogravimetry [28]. We used the same experimental setup as for the study of the reaction kinetics, as described above and the experimental procedure was very similar, but much more delicate in terms of artifacts and the ammono partial pressures are adapted to the equilibrium condition. Any kind of disturbing reaction, despite the dissolution of nitrogen to the melt, has to be avoided, like the formation of GaN in the solution, oxidation, and formation of GaN on the surface. But also the weight change from evaporation of Ga from the sample will interfere with the weight signal arising from the dissolution of nitrogen in the solution. All those parasitic reactions must be carefully avoided or at least minimized or quantitatively controlled such that a correction is possible. As the actual soluble species is subject to speculations a simple model for the reactions taking place in that system is formulated, assuming the dissolution of atomic nitrogen into the melt. This can be justified by the consideration of the weight signal resulting from the addition of nitrogen atoms only, which would be independent of the soluble species itself. According to Henry’s law for diluted systems and if the ammonia partial pressure is kept below the equilibrium partial pressure, one can treat the NH3 H2 equilibrium thermodynamically and the amount of dissolved nitrogen or solubility of nitrogen is expected to be a function of p.NH3 /. Figure 12.2 depicts the schematic illustration of the process of adsorption of nitrogen from an ammonia containing atmosphere to the solution in the LPSG process.

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NH3

process gas atmosphere

(1)

NH3,ad Nad + 3Had (2) (3) [N] (4) +Ga=GaN Fig. 12.2 Schematic illustration of the processes of absorption and desorption of nitrogen and the formation of GaN in a sequence of steps (numbers 1–4) (after [28])

The interaction of NH3 and the Ga solution includes the following steps according to the figure above: (1) NH3 adsorption to the free solution surface (2) Dissociation of the adsorbed NH3 molecules and dissolution of the atomic nitrogen into the melt (3) Transportation of the nitrogen atoms through the bulk solution (4) Formation of GaN from the solution if the solubility limit is reached A detailed discussion of the model may be found elsewhere [28]. Figure 12.3 shows as an example the graph of the corrected weight change as a function of time for a Ga–Sn solution under an ammonia–hydrogen atmosphere. The region ’ represents the deoxidation of the sample under ArH2 flow, followed by the dissolution of nitrogen of the ammonia containing atmosphere until equilibrium is attained (sector “) and the nitrogen desorption process in the region ” after the NH3 - flow was stopped. The curve must be corrected for the gallium evaporation taking place at all stages of the reaction. The solubility was calculated from the total weight gain in the equilibrium state. Table 12.2 combines exemplary solubility data from this study and data existing in the literature. It can be seen that a solubility that can be achieved with the LPSG method is comparable to those values reported for high pressures and high temperatures. This means that the solvent indeed enhances the solubility of nitrogen. In comparison to the modeled solubility of nitrogen in pure gallium under the same conditions, the solvents enhance the solubility by roughly three orders of magnitude. However, care has to be taken when comparing solubility data at different conditions as the gas solubility is a function of temperature and pressure. For many dilute solutions of gases in metal melts, the relation between lgXN (with XN D solubility mole fraction) and the inverse temperature is linear at a given pressure. This indicates that the changes in enthalpy and entropy of the dissolution of nitrogen in those systems is virtually independent of temperature. For  and the assumption that the solubility of N is very XN D 0:11 exp 12  28;745 T

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0.8 0.6 0.4

b

0.2

g

a

Dm (mg)

0.0 Ar+H2

–0.2 –0.4 –0.6

Ar+H2+NH3

Ar+H2

(1) Ga evaporation (1) Ga evaporation (2) Nitrogen incorporation (2) Deoxidization (3) Nitrogen transportation

(1) Ga evaporation (2) Desorption

–0.8 –1.0

900°C, Ga (28 at%)+Sn (72 at%) Ar=60ml/min; H2=40ml/min; NH3=0.075ml/min PNH =0.75 mbar

–1.2

3

–1.4 200

800

1400

2000

2600

3200

t (minutes) Fig. 12.3 Plot of weight change vs. time for a Ga–Sn solution. The curve is corrected by the subtraction of the Ga evaporation as well as by an offset (redrawn after [28]) Table 12.2 Comparison of nitrogen solubility data for different conditions present in GaN solution growth Study

T Œı C

P [bar]

Solution composition

Atmosphere

Solubility mole fraction N2 ; XN

Sun 2006 [28] Sun 2006 [28] Sun 2006 [28] Logan 1972 [13] Morishita 2005 [11] Madar 1975 [36] Grzegory 2002 [3]

900 930 950 1;150 900 1;200 1;520

1 1 1 1 42 8;000 11;000

Ga C Sn Ga C Ge Ga C Ge Ga Ga C Na Ga Ga

Ar C H2 C NH3 Ar C H2 C NH3 Ar C H2 C NH3 H2 C NH3 N2 N2 N2

6  104 8  104 8  104 3  105 104 1  103 3  103

small and therefore, Henry’s law is obeyed all gases can be treated as ideal. Using existing data in the literature and the solubility measured by the TG experiments, one can express the Gibbs free energy change as a function of temperature as shown in Fig. 12.4a and subsequently calculate the temperature dependence of the nitrogen solubility. The run of the curve in Fig. 12.4a shows that G for the dissolution of N into Ga is positive in a wide temperature range and implies that the nitrogen solubility will increase with temperature under NH3 atmosphere. Figure 12.4b depicts the respective comparison of the TG data gained by Sun, with the solubility data for pure gallium systems under high N2 pressure determined by Grzegory [3] and Logan [13] for pure gallium under ammonia atmosphere. For details of the calculation see Sun [28]. Overall, it could be experimentally shown by direct in situ measurements that the nitrogen solubility is indeed enhanced by the addition of suitable solvents. The

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b

1650 1500

200

0.01

150

1E-3

XN(mole fraction)

ΔG (kJ/mol)

a

100

ΔG1

50

0

–50 800

1400

T(K)

1600

1800

2000

1050

900

1E-5 Grzegory (Pure Ga+N2 at high pressure)

Nitrogen dissolution in Ga melt under NH3 atmosphere

1200

1200

1E-4

1E-6

1000

1350

1E-7 5.0

Logan (Pure Ga, NH3+H2 at 1 bar) Calculation for pure Ga + NH3 at 1bar This study (Ga+Solvent, NH3+H2+Ar, at 1 bar) Detection limit for TG measurements

5.5

6.0

6.5

7.0

7.5

8.0

8.5

1/T×10000(1/K)

Fig. 12.4 (left) Calculated Gibbs free energy change of nitrogen dissolution in a Ga melt under ammonia atmosphere as a function of temperature (right) Temperature dependent nitrogen solubility for different systems

nucleation as well as the growth rates will depend on temperature and ammonia partial pressure. Moreover, the solvent acts kinetically as was discussed in the previous chapter. The situation in the case of LPSG growth is therefore very much comparable to the high pressure solution growth but with much favorable process conditions like flowing gases and low temperatures. The TG method was also found to be a suitable way for studying the kinetics of the reactions involved and defining rate limiting steps as well as selecting the best solvents and best composition for the growth solution.

12.2.3 Growth Setup, Process, and Basic Challenges Different growth setups, vertical [21] as well as horizontal, [13, 14] were used by the respective working groups. Only Logan and Thurmond used a thermal gradient in the melt for most of experiments, claiming that it enhances the growth rate [13], while Elwell et al. found that “the temperature gradient transport is not a significant effect in crystal growth of GaN from the solution” [14]. Most of the growth runs reported in the literature have the substrates fixed or floating in the melt during the whole process [13, 14, 16], which has some pros (simple setup), but also some negative effects (like thermal etching or homogeneous nucleation, see later in this chapter). The furnaces discussed in the literature were resistively heated up to temperatures of 1;150 ıC and quartz glass was used as main furnace tubing material. Only small samples were processed in the furnaces of Logan et al. [13], Elwell et al. [14], or Klemenz et al. [21], whereas recent studies in our laboratory developed the method even further leading to growth furnaces capable of processing several 2 in. substrates in one run and demonstrating the first 3 in. GaN template grown from the solution [16]. However, one has to cope with some basic challenges with this method in order to attain stable growth conditions and a high quality material, which will be discussed in the following section.

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Formation of a Crust on the Melt Surface The formation of a crust of nitrides or oxynitrides on top of the melt can principally impede the reaction of ammonia with the solution and hence reduces or stops crystal growth. Logan and Thurmond [13] reported that the crust will start to form when process temperatures were set below 950 ı C. As a measure to prevent this, Logan [13] suggested to agitate the melt by mechanical vibration or a temperature gradient, whereas our recent studies indicated that this effect is more likely related to the purity of the gas phase. If oxygen and moisture is kept out of the system, it is possible to perform LPSG growth at temperatures as low as 800 ıC without the formation of a surface blocking layer. This also points out that the gas purity in general must be a major concern and the application of point-of-use purifiers for the process gases is advisable.

The Creeping Phenomenon The upwards flow of the Ga containing melts along the inner wall of the crucible is a more “striking feature.” This effect was already observed by Logan [13] who mentioned that the flow of the solution occurs if the ammonia partial pressure is set higher than the respective equilibrium partial pressure. Under such conditions, a thin GaN layer is probably formed on the inner crucible wall right above the melt surface, which is then wetted by the gallium solution. The thin layer in the vicinity of the new melt surface forms again, is subsequently wetted, and thus kicks off a fast progressing up flow of the melt along the crucible wall. The higher the ammonia partial pressure (supersaturation), the stronger is this effect. According to our studies, there are three ways to overcome or at least reduce the creeping problem. First, the supersaturation must be kept well within the Ostwald–Miers region, which will help to avoid parasitic formation of GaN and also the melt creeping phenomenon. Secondly, a short stop of the growth process after some hours for a certain time by reducing the supersaturation and/or the temperature can help to pin the up flowing melt at the position reached until that stage. Lastly the choice of the solvent has a strong influence on this inconvenient behavior of the melts. A solvent with a high vapor pressure compared to gallium, like sodium, potassium or others that are equally evaporating, will have a strong disadvantageous effect as the gallium concentration increases locally and therewith the supersaturation accelerating the creeping, whereas solvents with low vapor pressures will help to circumvent this phenomenon.

Materials in the Hot Zone and Dissociation of Ammonia Attention has to be turned also on the materials present in the interior of the furnace in order to avoid uncontrolled or unwanted dissociation of ammonia on other surfaces than the solution and corrosion of the furnace interior. But, this is not a too serious issue as long as no graphite containing parts are used inside the hot zone.

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The presence of graphite in the hot part of the furnace can be very deleterious to the growth process [23]. However, the situation can be different if the reaction gases are changed. The same debate holds concerning the choice of the crucible material. The most stable material in direct contact to the melt is pyrolytic boron nitride. Other materials like quartz glass, graphite, glassy carbon, silicon nitride, or aluminum nitride can be corroded by the melt and thus cause a pollution of the solution with unwanted impurities, beside the effect of uncontrolled dissociation of the ammonia as discussed above. Also the concentration of hydrogen in the carrier gas can have a noticeable influence on the thermal decomposition reaction of ammonia on the inner surfaces of the furnace [37] due to a passivation effect. This effect develops stronger with increasing hydrogen concentration in the atmosphere [23, 37]. A very detailed discussion on how to overcome or circumvent all this technological obstacles is given in Hussy [23]. In contrast to standard LPE techniques, the supersaturation in the LPSG process is not attained by cooling of the solution, but by applying a higher ammonia partial pressure than the equilibrium one [14, 16, 22]. In turn, the careful control of the partial pressure of the gases used in the process is a very important parameter. Besides that, the temperature and composition of the solution play the major role as described before. Regarding a complicated gas phase process, the handling of process parameters in LPSG is comparably simple and the most important points will be discussed in the following section of this chapter.

12.2.4 The Influence of Process Parameters on the Epitaxial and Parasitic Growth of GaN Ammonia Partial Pressure As described before in this chapter, the nitrogen concentration in the growth solution is a function of the ammonia partial pressure. Therefore, the equilibrium nitrogen concentration in the melt (no growth of GaN, but also no dissociation of GaN) corresponds to a certain ammonia partial pressure peq .T /. The value ppara .T / is the maximum ammonia partial pressure which corresponds to a maximum supersaturation of the solution before homogenous nucleation starts under the given conditions. The homogeneous nucleation is considered parasitical as long as the stable growth on a seed is wanted. The difference between this supersaturation concentration and the equilibrium concentration (D solubility) corresponds to the so-called Ostwald–Miers region. The effect of the ammonia partial pressures on the formation or dissociation of GaN in hydrogen atmosphere was investigated in a temperature range from 870 to 1;150 ıC using different approaches for either pure gallium [14, 38] or gallium–germanium solutions [16]. Limiting values of p.NH3 / for the dissociation of GaN .pdiss / and the formation of epitaxial, respectively parasitic GaN were found. Therewith it could be considered that any GaN that was grown on the seed

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Fig. 12.5 Limits of ammonia partial pressure in hydrogen over GaN(s)–Ga(l) for the formation of epitaxial and parasitic .pepi ; ppara / GaN respectively; the dissociation of GaN .pdiss / in the temperature range from 870 to 1;150 ı C. The calculated equilibrium ammonia partial pressure [38] is plotted as a solid line

was epitaxial GaN, .pepi / and any GaN that was nucleated in the solution elsewhere was parasitic GaN, .ppara /. In Fig. 12.5 the corresponding ammonia partial pressures .pdiss ; pepi ; ppara / are plotted as a function of temperature for the investigated temperature range. Thurmond et al. calculated the equilibrium ammonia partial pressure [38] which is increasing with temperature (see Fig. 12.5, solid line). The corresponding experimental results .pdiss ; pepi / show the same tendency [16, 38]; however, they can scatter widely especially above 1;000 ı C. For the formation of (parasitic) GaN crystallites in the solution, a higher ammonia partial pressure is necessary than for the growth of an epitaxial layer or for preventing the GaN from decomposition. This supersaturation is necessary to overcome a kinetic barrier. As the height of the kinetic barrier is temperature dependent itself (most likely decreasing with temperature) different temperature dependence for ppara compared to pepi or pdiss is anticipated and indeed the experimental results show this correlation. In Fig. 12.5 a region exists between pdiss and ppara and respective pepi and ppara , which corresponds to the classical Ostwald–Miers region, bordered by the solubility line (i.e., peq ) and the maximum supersaturation without homogeneous nucleation (super solubility line). A stable supersaturation exists within the Ostwald–Miers space. The corresponding sets of parameters (T and p.NH3 /) are favorable for the growth of seeded layers or crystals while the formation of parasitic GaN is suppressed.

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It was demonstrated that the supersaturation within the Ostwald–Miers region .pepi .T / < p.NH3 / < ppara .T // leads to a quite low epitaxial growth rate, nearly independent of the growth time [16]. This means that thicker GaN epitaxial layers can be principially grown just by increasing the process time, which will however result in quite long growth runs. In contrast, the application of a supersaturation above the upper limit of the Ostwald–Miers space .p.NH3 / > ppara .T // results in a strongly time dependent epitaxial growth rate. The growth rate is high at the beginning of the process and decreases drastically with process time [16]. The formation and growth of parasitical GaN crystallites which consume reasonable amounts of the nitrogen present in the solution is the obvious reason for this. A similar observation was made for LPE of GaN by the Na-flux method under high supersaturation [39].

Influence of the Carrier Gas The influence of the carrier gas is manifold. Most experimental studies published up to now used hydrogen as a carrier gas [13–16] whereas [21] used nitrogen. First of all hydrogen itself has a high solubility in the gallium solution [28]. It is several orders of magnitude higher than the solubility of nitrogen introduced by ammonia under the same conditions. The influence of the composition of the carrier gas on the equilibrium ammonia partial pressure was experimentally investigated in [23]. At a given temperature of 960 ıC pepi is 1.35 mbar for 75–100% H2 in the carrier gas, dropping to 0.45 mbar for 30% H2 . As no stable epitaxial growth could be obtained at 960 ı C for lower hydrogen concentrations than 30%, pepi could not be investigated experimentally since the creeping phenomenon becomes severe under this growth condition. We found that technological difficulties like the above mentioned up flow of the solution on the crucible wall show an increasing tendency to occur with lower fractions of hydrogen in the carrier gas. In parallel, it was found that it is more difficult to attain stable epitaxial growth when lower amounts of hydrogen are present in the process gas. For much lower temperatures .850 ıC/ Klemenz et al. [21] achieved stable epitaxial growth with pure nitrogen as carrier gas and an ammonia partial pressure as high as 250 mbar. However, there was no description of up flow of the solution under such process conditions and using the same conditions in our furnaces did not result in epitaxial growth of GaN on the seed. But, according to our investigations this could also display the finding that the decomposition of the ammonia is not independent of the flow rates of the gases, especially in the case of low total flow rates and high nitrogen content. On the other hand, in the case of pure hydrogen atmosphere, the epitaxial as well as the parasitic growth rates were found to be independent of gas flow rates [23]. This might also be a result of the surface passivation effect as described before. With low flow rates, the time the gases remain in the hot zone mainly determines the dissociation of ammonia. If pure hydrogen is applied as a carrier gas, the passivation effect is dominant and the time for passing the gases through the hot zone does not play a major role. The thermal dissociation is then independent of flow rates, especially in the case of a high total gas flow.

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Composition of the Solution The most crucial process parameter is the type and amount of the solvent added to gallium. As described before in this chapter, nitrogen solubility, as well as the kinetics of rate limiting steps, is strongly determined by the solvent and its concentration. Logan et al. [13] investigated the influence of the additive bismuth on the epitaxial GaN growth. They reported that as the ammonia partial pressure, necessary to initiate epitaxial growth, becomes higher, more bismuth is added to the solution. This is very pronounced for very high bismuth concentrations (>75 at%). But for solution additives with a high vapor pressure such as bismuth and lead, the exact values for pepi and ppara cannot be extracted easily from long-term experiments since the solution composition will change with time due to the evaporation loss of the components, resulting in nonstable growth conditions. Therefore, additives with a low vapor pressure should be favored, such as tin and germanium. Thus, instead of pepi and ppara , the ammonia partial pressure pgrowth , which leads to relatively stable epitaxial growth (up to 36 h growth time) was plotted against the initial solution composition, in the case of lead and bismuth as solvents. The results for the different solution additives, bismuth [13], germanium [23], tin, and lead are compiled in Fig. 12.6. The relationship between the concentration of the additives and the equilibrium ammonia partial pressure is qualitatively the same for all solvents shown in Fig. 12.6, whereas the effect of additives on the formation of parasitic GaN

ammonia partial pressure [bar]

0·1 pgrowth(NH3)

GaxBi1-x

1000 °C (Logan, 1972)

pgrowth(NH3)

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pgrowth(NH3)

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ppara(NH3)

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ppara(NH3)

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pepi(NH3)

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0·01

1E-3 0·0

0·2

0·4 0·6 gallium atom fraction xGa

0·8

1·0 (Ga)

   Fig. 12.6 Dependence of ammonia partial pressures .pepi ; ppara ; pgrowth / on the additive concen tration. For the solvents Bi and Pb the corresponding ammonia partial pressures .pgrowth / are given, which lead to a stable epitaxial growth in experiments with small process time (e.g., 20 h)

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crystallites is controversially discussed in the literature. Logan found that the addition of bismuth suppresses the formation of parasitic GaN, while in Elwells experiments additions of both bismuth and tin increased the nucleation rate of GaN [13, 14]. The thermogravimetric measurements performed by Sun can be taken for clarification [28]. The additives germanium and tin help to suppress a high nucleation rate for GaN as they extend the so-called induction period for low supersaturations (time interval until first detectable weight gain related to GaN formation). This was proven by our recent studies suggesting that the concentration of the germanium has an effect on the width of the Ostwald–Miers region [23]. It was found, that small additions of germanium (10–40 at% Ge) lead to the suppression of the formation of parasitic crystallites and thereby enhance the epitaxial growth, but no stable epitaxial growth could be established for high additive concentrations (>75 at% Ge). This can also be explained by the study of Sun [25] who found that the nucleation rate is a function of ammonia partial pressure and can be different for any individual solution composition. Changing the composition of the solution does not only mean different equilibrium conditions and different widths of the Ostwald–Miers region but also affects the total density of the solution and the wetting angle of GaN. It could be shown that the solution density strongly influences the morphology of the as-grown GaN layers [22] and the formation of macroscopic defects resulting from parasitically grown GaN crystallites. If the formation of parasitic GaN can not be completely eliminated or is accepted for the benefit of higher growth rates, the solution density is a key issue to control as otherwise crystallites formed in the solution will drop on the growth interface and cause macroscopic defects. This needs special consideration since the LPSG process is a bottom seeded technique because of the fact that the nitrogen is supplied from the gas phase.

Temperature–Time Program Another important factor is the temperature – time program of the process, especially the cooling procedure. If the seed is kept in the solution for the whole time until the furnace is cold, it must be taken care that the solution does not get undersaturated with respect to the equilibrium partial pressure. This can happen if the ammonia flow is reduced too strongly at an early stage. Then, the grown GaN surface will be etched non-uniformly and the as-grown surface morphology will be affected negatively. On the other hand, if the ammonia partial pressure is higher than the equilibrium one during cooling, the melt may excessively supersaturate leading to parasitical GaN formation in the volume of the melt and subsequent development of macroscopic defects. To avoid both of the two scenarios, the ammonia partial pressure must be carefully adjusted according to the equilibrium partial pressure during the cooling procedure.

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12.3 Evolution of Structure and Morphology of the GaN Layers 12.3.1 Importance of the Initial Growth Stage The growth of GaN on a GaN seed surface is considered homoepitxial. Therefore, it is expected, because supersaturation is low, that the growth should proceed by a layer by layer growth mechanism or by step flow along an off-cut direction, because the sapphire substrate with a MOCVD GaN layer applied as a seed, generally has an off orientation toward the c-direction of 0:25ı . However, it was observed [16] that a pronounced as-grown surface morphology is present after LPSG growth, characterized by hexagonal shaped features like terraces, macrosteps, or hillocks of various heights and densities. The roughness of the surface of the LPSG layers can vary strongly depending on the additives used for the solution growth process, the conditions of the growth process itself, and in turn, on the nucleation regime occurring in the very beginning of growth. Initial Growth of the LPSG Layer, Growth Stages The growth starts with the back etching of the seed, then 3D facetted islands nucleate. They are visible in Fig. 12.7 by faint contrast in secondary electron imaging. This contrast might develop from different doping levels of the material along the growth direction. It is not per se clear why the islands form. Respective experiments showed that their nucleation is suppressed if the seed applied in the growth process has a large off cut equal or larger than 1ı .

Fig. 12.7 SEM cross-section image of the interfaces between the seed layer and the LPSG GaN. The back etching region (grey line) is clearly visible. The nucleation islands developed on the seed surface are obvious from the darker contrast

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The islands proceed with lateral growth until coalescence of the GaN layer is attained. The distribution of the islands over the seed surface is non-uniform. Their density and size depend on the growth temperature and the solvents applied. In order to clarify why the islands can form in a homoepitaxial growth case, it is important to investigate the nature of the islands in chemical and structural respect. Nature of the Growth Islands Investigations by spectroscopic and microscopic means were performed in the region where the growth islands occur. These investigations should help to clarify why the growth, unintentionally, always starts with the formation of nucleation islands. For the luminescence of GaN, it is generally expected that variations in dopant concentration scale reasonably with the luminescence intensity of the material if the luminescence from the donor acceptor pair (DAP) recombination is considered. Cathodoluminescence images taken at the wavelength of the near band edge emission, around 364 nm at room temperature, clearly revealed the presence of the growth islands at the seed–layer interface with different contrast. A complementary secondary ion mass spectroscopy measurement was carried out, monitoring the impurities in the GaN layers from the growth surface through the seed-layer interface and into the sapphire substrate. TEM microscopy mirrored the microstructure situation in the same region. Figure 12.8 shows a compilation of the corresponding observations. The impurity content of 1019 cm3 for silicon or oxygen characterizes a highly n-type doped GaN in this region; thus it’s counterintuitive that the luminescence intensity of the highly doped, island containing, layer of the LPSG GaN is as low as the seed layer which is a MOCVD GaN with a carrier density in the order of 1016 cm3 , since the luminescence intensity of the DAP usually increases with increasing n-doping level. However, it is visible from the corresponding TEM image that within the same region after the seed-layer interface there is a high recombination activity of the dislocation lines of the threading dislocations penetrating from the seed which are most likely the reason for the low luminescence intensity of this region because of nonradiative recombination. Figure 12.8 clearly shows that most of the reactions between the dislocations take place in the first micrometer of the LPSG layer after its interface with the seed. The related diagram of the impurity contents measured by SIMS does correspond to this situation. The highest measured concentrations of impurity elements are located in the island region. Figure 12.9 confirms that the dislocations act as non-radiative recombination centers destroying the luminescence in those particular parts of the layer, which explains the finding that the luminescence intensity is low although the island region is highly n-doped GaN. Though, in the first micrometer after the growth interface, the coincidence of effects is numerous and does not allow concluding too rigidly why the nucleation of islands happens although the growth is carried out homoepitaxially. The high impurity level in the island region can have more than one reason, e.g., gettering of impurities by dislocations and incorporation of different elements on the particular crystallographic facets. The accumulation of impurities on the growth interface can

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Fig. 12.8 Correlation between the nucleation islands (SEM image, middle) and the abundance of dislocation lines and recombination action of dislocation in the first micrometer behind the seed–layer interface (TEM image, left) together with the depth distribution of impurity elements measured by SIMS analysis (right). Images are aligned towards each other along the interface Nucleation islands

Fig. 12.9 Panchromatic CL image of the interface region between the seed layer and the LPSG GaN. The nucleation islands are clearly distinguishable by darker contrast respectively lower CL intensity. The dislocations act as nonradiative recombination centers too and are visible as black lines (arrows)

however also result in a locally altered supersaturation with respect to the growing material and could cause a different nucleation regime although the general situation is homoepitaxial for GaN. On the other hand, the presence of the nucleation island is extremely beneficial in terms of reduction of the dislocation density. It was found in [41] that the dislocation density after the LPSG process was reduced by one to two orders of magnitude compared to the initial dislocation density of the seed as it is also obvious from the images above. The corresponding mechanism acting during the LPSG process is related to the nucleation of islands on the seed surface in the very beginning of the growth process. Therefore, it is very important to have sufficient control over the initial growth of the GaN in situ in the LPSG process, even though the nucleation situation is not fully understood up to now.

12.3.2 Macroscopic Defects in LPSG GaN Up to now most as-grown LPSG layers were reported not to be directly suitable for device production. This can be due to incomplete coalescence, existence of cracks,

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Fig. 12.10 SEM images of the most prominent surface defects on as-grown GaN grown from the solution. Macrosteps are characteristic for most of the surface area (a), with hexagonal shaped flat pyramids as exemplarily shown in (b). The circular shaped surface defects are illustrated in (c) whereas (d) depicts particle induced morphological disturbances

blistering of the layer, or off-oriented GaN particles incorporated into the layers during growth (e.g., [13, 21]). The layers grown on c-plane sapphire or c-plane gas-phase seeding layers by Logan et al. exhibit strong surface structures characterized by hexagonal features. The surface morphology of recent LPSG layers show mainly terraces with macrosteps (Fig. 12.10a), interrupted by a variety of possible larger defects, like surface depressions and circular shaped defects [16] depending on seeding-layer-orientation, misfit, and process parameters (Fig. 12.10b–d). Most critical are surface depressions which have diameters up to 100 m and sometimes penetrate the whole epitaxial layer down to the seed surface. Therefore, they cannot be removed by polishing. These depressions are similar to the ones described for LPE studies of GaAs [42]. Several possible triggers for the formation of theses defects were described, with one reason being particles present on the seed or existing in the solution. Indeed, it was shown by Hussy et al. that GaN crystallites (particles) present in solution can induce the surface depressions in the LPSG GaN [22]. Avoiding the formation of particles in the solution is therefore a prerequisite for the production of GaN material free of such surface defects. One way to prevent the formation of such surface depressions is to adjust the density of the growth solution [22]. In those solutions with a density higher than that of the detrimental particles, the particles will float on the surface of the solution and hence cannot get into contact with the seed or the growth surface of the crystal at the bottom of the crucible. Another typical macro defect found in GaN-LPSG layers are circular shaped defects (Fig. 12.10c). They can be correlated to gas bubbles sticking on the growth

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interface during the solution growth process [23]. However, the origin of those bubbles, which is among other things, correlated to the hydrogen content in the atmosphere is still under investigation. By adequate polishing technique most surface defects can be removed. Those layers as well as most recent as-grown layers were suitable for realization of LED test devices on the LPSG samples exhibiting an electroluminescence signal.

12.4 Properties of the LPSG GaN Material Many working groups made huge efforts to obtain GaN crystals with improved quality and structural perfection. Among all the methods tried, during the last few years, the solution growth is quite interesting as the crystallization of the material occurs under conditions which are close to thermodynamic equilibrium. Results of the high pressure solution growth demonstrated that very low dislocation density of some hundred dislocations per square centimeter can be attained [3]. Moreover, the growth situation under excess gallium can help to hinder the formation of point defects related to gallium vacancies as was discussed in [26]. The growth from a solution is however quite slow which is a great disadvantage for a technical process, but from many studies dedicated to the structural, optical, or electrical characterization of GaN, it is obvious that fast processes with simply high growth rates cannot deliver the desired materials quality [43], with the dislocation density of the materials being the most crucial issue. All fabrication methods rely on a full diameter seed in order to initiate growth of thick layers or quasi bulk GaN with usable diameters. In the majority of the cases, a seed with a GaN layer deposited by MOCVD is applied, which means that the initial dislocation density of the seed is in principle unacceptably high, compared to growth and seeding of classical compound semiconductor crystals. Measures have to be taken in order to reduce the dislocation density during the following growth process. The HVPE method reduces the dislocation density in the GaN boule as a function of thickness [44], whereas the common gas phase techniques, like MOCVD, take advantage from advanced measures like epitaxial lateral overgrowth (ELO) or related techniques [45]. The following section illustrates the situation in the LPSG material and describes how the dislocation density is reduced during the growth.

12.4.1 Structural Properties and Dislocation Density in LPSG GaN Material Dislocation Density of LPSG GaN Material Defect selective etching (DSE) was performed to evaluate the total dislocation density and the relation between the initial dislocation density (DD) and the one found

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in the LPSG material. Many investigations about DD in GaN were published over the last few years. However, as can be seen from the literature, GaN originating from different production methods can behave quite differently in the wet etching procedures, such that it was suspected that many values given for the dislocation density might not represent the real DD present in the crystal [46]. The DD for the LPSG GaN was determined with TEM as well as with wet etching in NaOH/KOH eutectic melt. The etching temperature and time were carefully investigated and calibrated by comparison to TEM results. The optimum conditions for revealing the true DD density in LPSG material were found to be 420 ı C and 4 min etching time. Concerning the etching temperature, this is significantly higher than for the etching of common MOCVD layers as e.g., reported by [46]. The DSE of LPSG GaN reveals the same picture as often reported in the literature. The dislocations cropping out on the surface of the crystal were decorated with three different sizes of etch pits which could be counted separately. However, care has to be taken for sufficient statistics, especially for the large etch pits which are the fewest. Figure 12.11 also gives an example of the etched surface of an LPSG GaN material. By TEM it was validated that the large etch pits do decorate the c-type screw dislocations and the medium pits correspond to the mixed type dislocations, whereas the pure edge dislocations are decorated by small pits only. This agrees with the expectations from the literature [46]. Those small etch pits are the most difficult ones to etch and visualize because of the fact that the activation energy for the formation of such a pit is highest among the three types of defects [46]. The average DD in the LPSG GaN grown on a MOCVD GaN seed layer is some 108 cm2 , but can be as low as 1  107 cm2 .

Reduction of Dislocation Density by Bending of Dislocation Lines It was observed that during LPSG growth as well as epitaxial growth by the sodium flux LPE growth, the dislocation density was reduced by one to two orders of magnitude compared to the dislocation density offered by the MOCVD-GaN seed [41]. This mechanism functions because of a dislocation recombination action in a narrow range behind the seed-LPSG GaN interface. The annihilation can happen as a

s

m

l

Fig. 12.11 As grown c-plane GaN, Ga-polar surfaces of the MOCVD seed layer used for subsequent solution growth and the LPSG GaN in direct comparison. The samples were etched at 420 ı C for 4 min. The errors indicate the small (s), medium (m), and large (l) etch pits

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VIIc V^c

Seedlayer

Substrate

Fig. 12.12 Schematic illustration of the dislocation reduction in a narrow region behind the seed surface. Nucleation of facetted islands (indicated by dashed lines) is the key for the induction of frequent interactions between bended dislocations. During further growth and lateral coalescence (indicated by grey colour) the dislocations can proceed towards each other and recombine or annihilate

result of the bending of the threading dislocations out of the c-growth direction. The bending enlarges the probability for a dislocation to meet another one and react or directly annihilate, in case of opposite Burgers vectors. The key issue is the presence of isolated, facetted nucleation islands in the very early growth state as described above and the tendency of the GaN material to form facets very easily, at least in the LPSG case as it was evaluated by [26]. This situation is illustrated by the schematic drawing in Fig. 12.12. The presence of crystal facets offers an energetic attraction to a dislocation to minimize its total line energy by changing its direction and proceed towards the facet surface as indicated in the figure. As the lateral growth velocity, v?c , is larger than the growth rate along the c-direction vIIc , the inclined dislocations can move towards each other upon the coalescence of the islands and interact if they get close enough [6, 41]. If the dislocation lines bend by large angles, a short distance after the seed surface is enough for a reduction of the dislocation density in the LPSG material. Bending by ninety degrees into the basal plane is most favorable. It will hinder the dislocations to propagate along the c-direction into the final active area of the layer, where the device is manufactured. But, a bending of dislocation lines by 90ı is not easily attained, because it requires the growth of steep side facets of the nucleation islands with particular crystallographic orientation. Only in that case the dislocation line can find its energy minimum and the driving force for the bending is high enough. Then the overall reduction of the number of dislocations is a function of the process time until island coalescence is attained, as the bending will only happen as long as the facets exist. Smaller bending angles of the dislocation lines would allow the coalescence time to be maintained for a longer time. Thus, in the LPSG growth also, the most beneficial combination are steep side facets of the nucleation islands provoking large bending angles. A detailed investigation and corresponding TEM investigations of the LPSG material by [41] showed, that dislocations behave differently, depending on their character with respect to their line and burgers vectors. The bending angle in relation to the c-direction is not only determined by the character of the dislocation, but also by the crystallographic index of the interacting

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facet. A corresponding theoretical model for the observed behavior of dislocations in LPSG GaN was developed, which is on the basis of considerations as to the total energy budget of a dislocation and the minimization of line energy as the main driving force for various types of dislocations in GaN and their interaction with crystal surfaces. Further details can be found in [41], where also a thorough validation of the theoretical model by experimental observations give a good proof for the assumptions made. As can be determined from DSE images exemplarily shown in Fig. 12.11 it is obvious that the main difference between the DD in the seed and the LPSG GaN, lies in the number of small etch pits. This means that the dislocation reduction mechanism as described above does mainly act on the edge or the edge component of mixed type dislocations, if no special action – with respect to the type of facets of the nucleation islands – is taken as described in [41]. The pure c-direction screw dislocations are hard to influence as the screw configuration is energetically already very favorable. In general the DD can be reduced by one to two orders of magnitude by that mechanism compared to the initial DD present in the seed.

Structural Quality of the Lattice The structural quality of the synthesized materials was not only characterized by means of TEM and defect selective etching, but additionally investigated by high resolution x-ray diffraction. Despite the fact that no reliable model exists which directly correlates the dislocation density to the width of a x-ray rocking curve, one can qualitatively state that the smaller the width of the rocking curve, the smaller is the number of structural disturbances in the material measured. This includes all factors contributing to the width of a x-ray reflex, like twist and tilt of crystal blocks, strain, dislocations, lateral coherence length, and so on. Thus, rocking curve measurement is generally used as a relatively fast monitor for the crystal quality and for comparison purposes between different materials. In order to get the best idea about the real materials quality besides the effects of other structural contributions to the line width, one has to perform a different rocking curve measurement, which helps to deconvolute the effect. Hence, two types of rocking curve were measured for the characterization of the structural quality of the LPSG-GaN. First we investigated the typical omega scan rocking curves with open detector geometry. The spot size is large in order to integrate over a larger area of the sample. In this geometry, structural imperfections like block tilts towards the growth direction, residual strain, and threading screw dislocations contribute to the line width and shape of the rocking curve. Secondly we measured omega 2 theta scans of the material with the same large beam spot but with a 3 bounce 220 Ge analyzer crystal on the diffracted beam side, which monitors the real materials properties itself despite the other factors like the mosaicity. Figure 12.13 contains two exemplary curves measured for the [002] direction. Typically the samples from the LPSG process exhibit a rocking curve width

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Fig. 12.13 Measurement of [002] rocking curves in two different diffraction geometries. (a) shows an example of a regular rocking curve measurement with an open detector, large beam spot and maximum acceptance angle for detection. Curve (b) exemplarily depicts a measurement with an analyzer on the diffracted beam side, monitoring the entire materials quality

around 200 ˙ 20 arcsec, measured with the open detector configuration as shown in Fig. 12.13a. This value is comparable to a high quality MOCVD GaN. No block tilts are visible from this curve, although the acceptance angle for detection of orientation differences is large. The entire materials quality of the LPSG GaN itself, blanking the other influences from the mosaicity, is expressed in the very narrow rocking curves as depicted in Fig. 12.13b. The width of those curves was found to be between 18 and 27 arcsec for all measured samples from standard LPSG processes, which was quite small and generally only found on material with high structural quality like GaN from high pressure solution growth or very recently from the ammonothermal method, although, principally the expectation holds that a process operated close to thermodynamic equilibrium should result in excellent structural properties of the material. Although the overall materials quality of the LPSG GaN is quite good and the dislocation reduction method does act usefully if the right conditions are chosen, the material shows a reduced DD compared to the seed layer, but it will not be possible to remove all the dislocations from the LPSG material that way. As soon as the dislocation density approaches a value on the order of 106 cm2 , the dislocation elimination by a recombination process is not effective any more as the interaction distances necessary for dislocation reactions are on the order of tens of nanometers, whereas the average distance of the dislocations at this DD is on the micrometer scale. However, it is a crucial fact for all existing growth processes for nitrides that the total number of dislocations in the material must be further reduced. But, by the first glance this can only be reached if either homogeneous nucleation is used or high quality seeds are applied. Here one could see a basic advantage of a solution growth process as free nucleation in the volume of the solution is theoretically possible, but it is a big challenge. So, one of the most urgent questions left to be solved now is how to generate a high quality seed for the GaN growth processes.

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12.4.2 Electrical Properties Conductive n-type or p-type GaN substrates are attractive for optoelectronic devices, as the conductivity of the substrate offers the possibility for back side contacts and a free path to the front side for out coupling of light, whereas semi insulating material is required for RF applications. In this sense, the solution growth has an advantage and a disadvantage. Typically solution grown material can suffer from incorporation of the solvent into the material produced. In those cases where the solvent can be used as a dopant this can be taken as an advantage. In the case of the LPSG growth of GaN, it is easy to use a solvent like germanium which will result in a highly n-type doped gallium nitride. The electrical properties of the LPSG GaN were determined by classical Hall measurements as well as by Fourier Transformed Infrared (FTIR) reflectivity measurements. In comparison to the Hall measurements the FTIR measurements can be performed without the necessity of contacts and give reliable results down to n D 5  1017 cm3 [27]. The carrier density of the LPSG GaN is typically around 4  1019 cm3 if an n-type solvent like germanium is applied which readily incorporates into the GaN structure. Variations among the processes or within the wafers produced by multi wafer processing are marginal only. Other solvents like bismuth lead to GaN with around one order of magnitude lower carrier density compared to the germanium doped material. So far, p-type doping has not been tried.

12.4.3 Impurities An average analysis of impurities in LPSG GaN, in the case of germanium containing solutions, is given in the Table 12.3. The values were determined by secondary ion mass spectroscopy and averaged over the entire layer thickness. Note, that this is not entirely true if one compares Fig. 12.8, where an accumulation of impurities close to the seed surface is obvious. It is clear from those measurements that the system can be kept quite clean and the major impurity is the solvent resp. the dopant. Despite that, Si and O do have a Table 12.3 Concentration of impurities in LPSG GaN as measured by secondary ion mass spectroscopy Impurity Germanium Oxygen Silicon Hydrogen Aluminum Boron Carbon

Average concentration in LPSG GaN, per cm3 9  1019 3  1018 1  1018 1  1017 3  1016 2  1016 1  1016

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noticeable abundance, whereas the other elements are quite low in concentration. No other impurities like alkali or earth alkali metals are present in this type of material because they do not find any use in the LPSG process. As discussed above, the impurities in the gallium nitride from the LPSG process mainly stem from the solvents added and therefore the gallium as well as the solvents should be of highest possible purity. However, an evaluation of the process concerning sources of impurities was necessary too. The oxygen contamination is most likely caused by the effect of hydrogen on the quartz glass ampoule. Boron can be mobilized from the boron nitride crucible. The process atmosphere containing ammonia and hydrogen is quite aggressive and it is important to choose appropriate materials in the interior of the furnace, especially in the hot zone. Graphite or oxide materials in the hot zone are no good choice as discussed in the Sect. 12.2.3 and are therefore not present in the process and the light carbon content obviously comes from another source.

12.5 Summary and Prospect The consequence of all the detailed studies discussed in this chapter is a deep understanding of the overall LPSG growth process as well as the, before more or less unclear, role of the solvents in terms of kinetics and solubility of nitrogen. Many of those findings are transferable to other solution growth methods which are probably very comparable among all the variants of solution growth. The intrinsic big capacity of a solution growth process is definitely the expectable materials quality from thermodynamic considerations, which were best demonstrated by the high pressure solution growth material. On the other hand, there are unwanted aspects like low supersaturations and in turn the very low growth rates leading to long process times. However, it was demonstrated by other groups like the sodium flux researchers that the growth rates can be higher by a factor of 10 or 20 compared to the actual state of the low pressure or high pressure solution growth. The materials quality which resulted from those Na-flux processes would need additional research efforts to be acceptable. The investigation and understanding of the action of the solvents as shown in this chapter can now help to perform a consequent screening of possible solution additives. With view to the enhancement of growth rates the real potential of this method can be tapped as much as possible by identifying the best solvents. But, all joint efforts did not clarify the experimentally maximal attainable growth rate for the solution growth methods up to now and more research efforts are necessary. It is anyway a mater of discussion whether bulk crystals are always needed for all application cases. The LPSG process can reduce the dislocation density remarkably and a model was built describing the underlying physical mechanism in a way that it can be used to give a forecast on the behavior of the dislocations. Because of the easiness of the whole LPSG setup, it is not difficult and very straight forward to up scale the process to larger diameters, which was successfully demonstrated. Seeing this, there seem to be no real obstacles for the growth on larger than 3 in. diameters,

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despite the availability of seeds. The method can also be considered as a variant of a liquid phase epitaxy process, which found wide application as a multi wafer route and can produce large numbers of wafers in parallel. For the multi wafer production of GaN from the LPSG process we could demonstrate that it is possible to process e.g., three wafers in parallel without any loss in quality and comparable electrical properties within experimental errors. However, evaluating and exploiting the full potential of the low pressure solution growth method do need further research work and the question how to produce high quality seeds remains to be solved for all the nitride crystal growth methods. Acknowledgements The authors would like to acknowledge the support of all people who have contributed to this work by various amounts over the years. Especially, the authors are grateful to Rainer Apelt, Carmen Maier, and Michael Lang for their unremitting and perfect technical support. G. Li Sun for his contribution in the frame of thermogravimetry. B. Birkmann for his work in the development of the LPSG process during the first period of the project. I. Knoke for the TEM image and L. Kirste, Fraunhofer IAF, for discussion and help with the rocking curve measurements. Thanks to B. Sieber, University of Lille for the panchromatic CL image and J. Weyher for fruitful discussion about the wet etching results. The authors would also like to thank Prof. G. Müller for continuous discussions and valuable suggestions over the years. This work was supported by the German ministry of education and science (BMBF) under contracts number 01-BM-158 and 01BM-580. The TEM work by I. Knoke, contributing to this work was supported by a Fraunhofer PhD scholarship.

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Part V

Characterization of GaN Crystals

Chapter 13

Optical Properties of GaN Substrates Shigefusa F. Chichibu

Abstract Optical properties of GaN substrates are surveyed to discuss the superiority and critical issues as a substrate for device fabrication. At first, optical properties of a few hundred-micrometer-thick free-standing GaN substrates grown heteroepitaxially on (0001) Al2 O3 substrates by metalorganic vapor phase epitaxy and halide vapor phase epitaxy using lateral overgrowth techniques are introduced, as they can also be used as a seed substrate for the ammonothermal overgrowth. Then, spatiallyresolved cathodoluminescence (CL) spectra of GaN films grown on free-standing GaN seeds via fluid transport by the ammonothermal method are correlated with the microstructure and growth polarity. The spectral lineshape of local CL was nearly position-independent for a 4-m-thick N-polar film exhibiting featureless morphology: the spectra exclusively exhibited a broad near-band-edge (NBE) free-carrier recombination emission with Burstein–Moss shift. Conversely, CL spectra at 100 K N and (1012) N facets with ridges origof a 5-m-thick Ga-polar film having (1011) inating from central pits exhibited an NBE peak at 3.444 eV and emission bands at 3.27, 2.92, and 2.22 eV, all of which showed rich intensity contrasts in the CL mapping images. The NBE peak intensity was remarkably enhanced at crests of the ridges, where the density of threading dislocations (TDs) having edge components was greatly reduced by the dislocation bending. The results encourage one to grow low TD density GaN wafers by slicing thick crystals grown by the ammonothermal method.

13.1 Introduction Optoelectronic devices such as bright blue, green, and white light-emitting diodes, laser diodes operated at 405 nm [1–4], high electron mobility transistors, and bipolar transistors have been realized using GaN and related (Al,In,Ga)N alloys grown heteroepitaxially on (0001) Al2 O3 substrates. However, demands are thrown for obtaining low threading dislocation (TD) density, large-area, and low-cost homoepitaxial substrates to realize high-efficiency, highly reliable, and low-cost devices.

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To comply with the requirements, the use of GaN wafer sliced from bulk GaN ingots is an ultimate solution. Several growth techniques using a Ga melt have been investigated to grow bulk GaN crystals [5–7]. However, as nitrogen has a very limited solubility in molten Ga, thin platelets of GaN were exclusively obtained. In addition, a bulk growth method must be expandable to an industrial mass-production scale to compete with existing GaN substrates prepared by metalorganic vapor phase epitaxy (MOVPE) [8] and halide vapor phase epitaxy (HVPE) on a patterned (111)A GaAs substrate [9] or on a GaN/(0001) Al2 O3 template whose surface is modified with TiN [10]. One of the promising techniques is an ammonothermal (AT) method, which is a solvothermal growth method that uses supercritical ammonia as a fluid medium, and is considered to have excellent scalability similar to the case for ’-quartz and ZnO hydrothermal [11] ones. Growth of AT-GaN has been investigated by several groups [12–18], and Hashimoto et al. [19–22] have reported on seeded growth of GaN on a free-standing (FS) GaN substrate prepared by HVPE [23] demonstrating the single crystalline growth [19] of large-area [20] .34 cm2 / and low TD density [21] (lower than 1106 cm2 for Ga-polar face) GaN films using basic mineralizers. Also, over 5-mm-long, 5mm-diameter hexagonal bulk GaN ingot has been grown recently [22]. Large-area (46-mm-diameter) seeded growth of GaN by AT method has also been reported by Ehrentraut et al. [24] using acidic mineralizers. Very recently, Dwilinski et al. have demonstrated 1-in.- [25] and 2-in.-diameter [26] GaN wafers grown by the AT method, both of which had nearly zero wafer bowing (curvature radius greater than 1 km). Structural properties of AT-GaN have been investigated by several groups. However, direct correlation between the microstructure and luminescence spectrum had not been reported before our work [27], although precise information obtainable from those analyses such as exciton energies versus strain, impurities, and carrier lifetime is needed to put AT-GaN in practical use as a homoepitaxial substrate. In this chapter, optical properties of GaN substrates are surveyed to discuss the superiority and critical issues as a substrate for device fabrication. At first, optical properties [28, 29] of a few hundred-micrometer-thick FS-GaN substrates grown heteroepitaxially on (0001) Al2 O3 substrates by MOVPE [8] and HVPE [30] using lateral epitaxial overgrowth (LEO) techniques are introduced, as they can also be used as a seed crystal for AT growth. Then, results of spatially-resolved cathodoluminescence (CL) measurements on Ga- and N-polar faces of GaN crystals grown on HVPE FS-GaN seeds via fluid transport by AT method are shown to correlate the local CL spectrum with the morphological feature, microstructure, and growth polarity [27]. We demonstrate the possibility to prepare low TD density GaN wafers by slicing thick crystals grown toward the Ga-polar direction by AT method, whose TD density is reduced by dislocation bending.

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13.2 Optical Properties of GaN Substrates Grown by Metalorganic Vapor Phase Epitaxy and Halide Vapor Phase Epitaxy 13.2.1 Photoreflectance Spectra of Excitonic Polaritons in MOVPE-LEO GaN Substrate As purely bulk, single crystals of GaN with sufficient area have not been available for a long time, the use of thick GaN films grown on heterogeneous substrates has been one of the alternative solutions to grow thick bulk GaN films. Nevertheless, one can pursue intrinsic material physics using such high quality GaN substrate with reduced TD and residual impurity density. Yamaguchi et al. [31, 32] have studied optical reflectance (OR) and photoluminescence (PL) spectra of thick GaN and GaN substrates prepared by HVPE-LEO to determine some valence band parameters. Concerning exciton–polariton [33] in GaN, Gil et al. [34] and Stepniewski et al. [35] have discussed exciton–polaritons in GaN=Al2 O3 [34] and homoepitaxial layers on bulk GaN crystals grown under very high-pressure N2 ambient [35]. Also, the authors have investigated low temperature OR and PL spectra of the GaN substrate grown by MOVPE-LEO [8] to show that the spectra were well described by a model exciton–polariton picture in which A, B, and C free excitons couple simultaneously to an electromagnetic wave [36]. Electromodulation spectroscopy such as electroreflectance (ER) yields a very sharp transition structure related to the third derivative of dielectric functions [37–39], and is a useful technique to obtain information on the bandgap. ER can be used to investigate exciton resonances in wide bandgap materials under the limit of low electric field [40] especially at low temperature. However, very little is known of the ER structure of FS-GaN. Here the spectra of photo-modulated electroreflectance, a contactless form of ER and is referred to as photoreflectance (PR), and PL of the high quality GaN substrate [8] are shown [29]. The GaN substrate [8] was grown by MOVPE-LEO with subsequent Al2 O3 removal. Details of the preparation procedures and structural properties are given in [8, 29, 36]. The lattice constants are a D 0:31898 ˙ 0:00002 nm and c D 0:51855 ˙ 0:00002 nm, and the crystal was close to strain-free. The residual electron density, n, at 300 K was too low to measure by the Hall effect measurement, i.e., n is estimated to be lower than 1014 cm3 , which is the limit of our equipment. Reduced nonradiative defect density was also characterized by time-resolved photoluminescence (TRPL) [28] and positron annihilation measurements [41–43]. The GaN substrate exhibited biexponential-type decay with decay constants 1 D 130 ps and 2 D 860 ps at 300 K, which will be shown later, while a standard 1-m-thick GaN=Al2 O3 generally exhibits a short decay of 2 shorter than 300 ps. OR spectra were measured at 10 K with incident angle smaller than 10ı . PR and PL spectra were measured as a function of temperature using the 325.0 nm line of a cw He–Cd laser as an excitor. The reflected white light from a Xe lamp and emitted light were dispersed by a f D 67-cm-focal-length grating monochromator, and phase

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0.4 0.3

3.47

3.48

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a

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Theory Experiment

b

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PR SIGNAL ΔR / R (arb. units)

REFLECTANCE

sensitive detection was carried out using a GaAs:Cs photomultiplier. The spectral resolution was ˙0:3 meV at a wavelength of 350 nm. The OR, PR, and PL spectra of the surface and the back of the substrate, namely Ga-polar and N-polar faces, were confirmed to be identical. Here we introduce the results on the Ga-polar face. Figure 13.1 shows (a) OR, (b) PR, and (c) PL spectra of MOVPE FS-GaN substrate measured at 10 K. In Fig. 13.1a, b, sharp reflection anomalies corresponding to the ground states of A, B, and C free excitons and the first excited states of A exciton .AnD2 / are found around 3.480, 3.485, 3.503, and 3.498 eV, respectively. The sample exhibited correspondingly sharp A, B, and AnD2 free excitonic PL peaks and a neutral donor bound exciton peak I2 . The full width at half maximum (FWHM) of I2 was as small as 0.7 meV. However, PL peaks corresponding to A and B free excitons were rather broad (2.6 meV for the A peak) and they cannot be fitted by single Lorentzian line-shape functions, suggesting the formation of excitonic polaritons. Indeed, emissions from lower and upper polariton branches of A exciton (LPBA and UPBA , respectively) and those of B exciton (LPBB and UPBB ) have been resolved using near-resonant excitation sources [36]. To interpret the OR spectrum, theoretical calculations have been carried out on the basis of a model exciton–

B

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I2

A,Bn=2,3..

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Fig. 13.1 (a) Reflectance OR, (b) PR, and (c) PL spectra of FS-GaN measured at 10 K. The transition structures labelled A, B, C, and AnD2 in the PR spectrum (b) are related to A, B, and C free excitons and the first excited states of A exciton, respectively. The PL peak I2 is due to recombination of excitons bound to a neutral donor. Theoretically fitted reflectance spectrum is shown by the solid line in (a). The calculated lifetimes  of LBPA , a branch combined UPBA and LPBB , and a branch combined UPBB and LPBC are shown in (d) where LPB’ .UPB’ / means the lower (upper) polariton branch associated with the ’ exciton. Dotted vertical lines through (a) to (d) show the transition energies obtained from the analysis of the PR spectrum (b). (After [29])

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Table 13.1 Values of free exciton energies at 10 K of the MOVPE-LEO FS-GaN substrate obtained from the fitting procedure of the OR spectrum. Transition energies and broadening parameters obtained from the analysis of the PR spectrum assuming the transitions as pure free excitons are also listed. (After [29]) Method

Parameter

Exciton/Transition A

B

C

Optical Exciton resonance 3:4791 ˙ 0:0002 3:4844 ˙ 0:0002 3:5027 ˙ 0:0002 reflectance energy (eV) (OR) Photoreflectance Transition 3:4803 ˙ 0:0003 3:4852 ˙ 0:0003 3:5033 ˙ 0:0003 (PR) energy (eV) Broadening parameter 1.0 1.5 1.1  (meV)

polariton picture in which A, B, and C free excitons couple simultaneously to an electromagnetic wave, where the effective mass anisotropy, the optical anisotropy, free exciton damping, and exciton “dead layer” are fully taken into account. Details of the calculation procedures are given in [36], in which the polariton lifetime  (Fig. 13.1d) has also been estimated. The best-fit curve of R D j.1  n/=.1 N C n/j N 2, where nN is the effective refractive index, is shown by a solid line in Fig. 13.1a. Satisfactory agreement was achieved, and the exciton parameters are summarized in Table 13.1. The PR spectrum shown in Fig. 13.1b was analyzed by the low-field ER lineshape function [37–39], 9 8 p = <X R Cj e ij .E  Eg;j C i j /mj ; (13.1) D Re ; : R j D1

where p is the number of the spectral function to be fitted, E is the photon energy, and Cj ; ™j ; Eg;j , and j are the amplitude, phase, energy, and broadening parameters of the j th feature, respectively. Here we tentatively assumed that the transitions are due to free exciton resonances .mj D 2/, which will be proved by the small  vales. The values obtained are also summarized in Table 13.1. Also, the transition energies are shown by dotted lines in Fig. 13.1a–d. Different from the PR spectra reported for unrelaxed GaN=Al2 O3 by Shan et al. [44] or Chichibu et al. [40], A of the FS-GaN was as small as 1.0 meV. From the fact that the PR transition energies agreed with the energies where the lifetime of the combined exciton–polariton branches approach their maxima, i.e., bottlenecks of the excitonic polaritons in terms of the polariton dispersion, the result means that PR monitors exciton– polaritons at low temperature. It is natural to admit that PR monitors exciton– polaritons rather than excitons or free carriers as the PR signal R=R is deduced as [37–39]

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R 2na D Re D Re Œ.˛  iˇ/" D a"1 C ˇ"2 ; R n."  "a /

(13.2)

where n2 D ©; na 2 D ©a ; na is the real refractive index, " D "1 C i"2 is the perturbation-induced change in the dielectric function ©, and ’ and “ are the Seraphin coefficients. Equation (13.2) means that the signal R=R is purely sensitive to ©. Thus, the PL peaks corresponding to A and B transitions shown in Fig. 13.1c are assigned to convolutions of upper and lower excitonic polaritons. The reason why the corresponding PL peaks are not resolved in this study but were resolved using the resonant excitation [36] in spite of smaller FWHM of I2 line here (0.7 meV in this study and 1.8 meV in [36]) may be due to the use of a He–Cd laser having high energy photons (3.81 eV), which produces excitons having large momentum. Figure 13.2a, b shows the PR and PL spectra, respectively, of the FS-GaN substrate as a function of temperature. The A and B excitonic transitions were well recognized up to 300 K, as is the case with GaN=Al2 O3 [40]. The PL spectra

a A

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PHOTON ENERGY (eV) Fig. 13.2 (a) PR and (b) PL spectra of FS-GaN as a function of temperature. Labels A, B, C, and AnD2 in the PR spectra denote the transitions related to respective excitons. (After [29])

Optical Properties of GaN Substrates

Fig. 13.3 Transition energies obtained from the PR spectra and PL peak energies of MOVPE FS-GaN as a function of temperature. (After [29])

283 3.52

PHOTON ENERGY (eV)

13

PR C PR A,Bn=2 PR B PR A

3.50 3.48 3.46 3.44

EA(T)=3.480-0.162/[exp(366/T)-1]

3.42

PL A,Bn=2 PL B PL A PL I2

3.40 3.38

MOVPE FS-GaN 0

50

100

150

200

250

300

TEMPERATURE (K)

were dominated by A and B free exciton emissions above 75 K. Surprisingly, recombination of AnD2 was recognized up to 200 K. Transition energies obtained from the PR spectra and PL peak energies are plotted as a function of temperature in Fig. 13.3. They showed the identical temperature variation. It has been shown [45] that in polar semiconductors that have optical phonon branches energetically close to the acoustic ones, fitting of temperature variation of a bandgap using the empirical equation assuming Bose–Einstein statistics considering Einstein phonons [46, 47] gave better results than using Varshni fitting [48] assuming Debye phonons. Therefore, EA .T / is fitted by [46, 47] EA .T / D EA .0/ 

0:162 .eV /; exp.‚E =T /  1

(13.3)

where free A-exciton peak energy at 0 K, EA .0/, is 3.480 eV and Einstein characteristic temperature (also named as Vina parameter) ‚E is 366 K. The value of ‚E nearly agrees with that obtained from GaN=Al2 O3 (349 K) [45], and this value corresponds to the energy maximum of the low-energy group of LO phonon branches near the Brillouin zone edge (K and M points) at 260 cm1 (370 K) [49].

13.2.2 TRPL of MOVPE-LEO GaN Substrate The sample used for the TRPL measurement was an 80-m-thick FS-GaN substrate that was fabricated by polishing away the Al2 O3 substrate after growing over 100-m-thick GaN on the LEO GaN base structure. For comparison, a 1-m-thick undoped GaN grown quasi-homoepitaxially on an 8-m-thick fully coalesced LEO GaN grown laterally over the SiO2 mask (LEO wing region) [28] and a 1-m-thick undoped GaN grown on an Al2 O3 substrate were prepared. For the preparation of the former two samples, LEO GaN was overgrown from a 2-m-thick GaN on Al2 O3 using a SiO2 pattern with 5-m-wide openings separated by 15-m-wide N > direction [28]. The edge-character TD denSiO2 stripes oriented in the < 1100 sity in the wing region was lower than 106 cm2 [50, 51]. We note that the wing

S.F. Chichibu

PL INTENSITY (arb. units)

284

300K

103 τ1= 130 ps

FS-GaN substrate τ2= 860 ps

102

101

τ1= 87 ps 1 μm GaN/Al2O3

0.0

0.5

LEO GaN (wing) τ2= 400 psec

τ2= 260 ps

1.0

1.5

2.0

TIME (ns) Fig. 13.4 TRPL signals measured at 300 K of the FS-GaN substrate, the wing region of an LEO GaN, and a GaN epilayer on Al2 O3 , which were grown by MOVPE. (After [28])

region of the LEO GaN suffered from a compressive biaxial strain up to 0.18% [52–54] after the coalescence. The TRPL of the NBE excitonic emission peak was excited by an approximately 150 fs pulse of a frequency-tripled, tuneable mode-locked Al2 O3 W Ti laser operating at 80 MHz repetition rate. The excitation energy and power at the sample were typically 4.64 eV and approximately 200 nJ cm2 per pulse, respectively. The signal was collected using a standard streak-camera acquisition technique. As the TRPL decay shapes could be fitted by a double-exponential function, fast and slow time constants (1 and 2 , respectively) were deduced. The FWHM value of the NBE PL peak decreased from 32 meV to 29 meV at 300 K by using LEO technique. From their peak energies and FWHM values, the PL peaks are assigned as being due to the recombinations of A- and B-free exciton (FE) emissions [40]. In general, PL lifetime at room temperature is limited by nonradiative processes because of thermal activation of nonradiative recombination centers (NRCs). Therefore, the recombination lifetimes of FE emissions shown in Fig. 13.4 may reflect the free (minority) carrier lifetime in the GaN crystals. From Fig. 13.4, it is noted that 1 was constantly 130 ps for the two low TD density samples, and was 87 ps for the most defective GaN epilayer. These results indicate that TDs may be related to NRCs [28, 42, 43]. However, nonradiative lifetime in GaN crystals has been shown to be mainly limited by the concentrations of point defect complexes involving Ga vacancies .VGa / [42, 43, 55]. Indeed, both 1 and 2 changed remarkably by changing the growth parameters or background carrier concentration, even if a standard GaN on Al2 O3 with intermediate TD density .5  108 cm2 / was analyzed. The pronounced long 2 of FS-GaN (860 ps) may indicate improved purity of the crystal.

13

Optical Properties of GaN Substrates

285

13.2.3 Low Temperature PL Spectra of HVPE-LEO GaN Substrate As described in Sect. 13.1, HVPE method is one of the most promising ways to grow thick GaN substrates, according to its high growth rate. Consequently, FS-GaN grown by HVPE can be used as a seed crystal for the boule growth by AT method. However, undoped GaN grown by HVPE in general shows n-type conductivity because of the residual impurities such as O and Si, which may give rise to the degradation of device performances as a result of the parasitic capacitance, current leakage, and the resultant cross-talk. Also, the wafer bowing problem originating from in-plain compressive strain in FS-GaN is unavoidable, as long as lattice-mismatched Al2 O3 is used as an original substrate. Therefore, growths of purely strain-free FS-GaN crystal must be carried out to improve the device performances. In this subsection, low temperature PL spectra of a FS-GaN substrate [30] prepared by HVPE using the facet-initiated epitaxial lateral overgrowth (FIELO) method [56] are shown to survey its optical quality. The sample measured was a 180-m-thick undoped FS-GaN grown by HVPE using the FIELO method. The FIELO-GaN was grown on a 2-m-thick GaN template prepared by MOVPE on 2-in. diam. (0001) Al2 O3 . Ammonia .NH3 / was used for an N source and GaCl formed from HCl flown on metallic Ga was used for a Ga precursor. The TD density was lower than 107 cm2 . Steady-state PL was excited by the 325.0 nm line of a cw He–Cd laser .6 W cm2 /, and was dispersed by a 25cm-focal-length grating monochromator. Phase-sensitive detection was carried out using a GaAs:Cs photomultiplier. PL spectra of the FIELO FS-GaN measured at 8 K and 300 K are shown in Fig. 13.5a, b, respectively. The NBE PL spectrum at 8 K is magnified in Fig. 13.5c. As shown in Fig. 13.5a, c, the spectra exhibited sharp PL peaks originating from

WAVELENGTH (nm) 380

1LO 2LO

(a) 8K FEA,B

(b) 300K YL

1.5

BL

2.0 2.5 3.0 PHOTON ENERGY (eV)

3.5

(c) 8K

350

TES I23.473 A 3.478 B 3.484 A,Bn=2,3… 3.502

PL INTENSITY (arb. units)

PL INTENSITY (arb. units)

FEA

WAVELENGTH (nm) 370 360

3.453

350

400

3.457 3.461

500

1LO

600

2LO

800

3.3 3.4 3.5 PHOTON ENERGY (eV)

Fig. 13.5 PL spectra of undoped very thick FIELO-GaN measured at (a) 8 K and (b) 300 K. (c) The near-band-edge PL spectrum of undoped FIELO-GaN at 8 K

286

S.F. Chichibu

free A- and B-excitons at 3.478 and 3.484 eV, respectively [29]. A peak originating from the excited states of free A- and B-excitons .A; BnD2;3:: / was also detected at around 3.502 eV. In addition, a sharp PL peak due to the recombination of excitons bound to a neutral donor .I2 / at 3.473 eV as well as two-electron satellites (TES) at around 3.453 eV [57] were observed. Remarkable emission peaks labeled 1LO and 2LO were observed in the lower energy side of the NBE peaks. They are assigned to LO phonon replicas of the main peaks, since their energy separation is close to the longitudinal optical (LO) phonon energy of 91 meV [49, 58]. In addition to the NBE peaks, broad luminescence bands in blue and yellow color region (BL and YL, respectively) at around 2.8 and 2.2 eV were recognized, especially at 300 K. Comparing the results of PL and positron annihilation measurements, the origin of YL band has been assigned to acceptor-type defect complexes composed of VGa and/or VGa  O [42,59–61]. The origin of BL band has been assigned to a donor-to-valence band transition, where the donor level has been proposed to be located at 0.58 eV below the conduction band minimum in undoped HVPE-GaN [62,63]. Another possible origin is a conduction band-to-acceptor transition like carbon. Anyhow, the intensities of BL and YL were more than two orders of magnitude lower than the NBE peak, indicating low concentrations of those point defects.

13.3 Effects of Growth Polar Direction on the Optical Properties of Seeded GaN Substrates Grown by Ammonothermal Method The samples investigated were a 5-m-thick Ga-polar and a 4-m-thick N-polar GaN films grown simultaneously on the front and back sides of a 400-m-thick (0001) FS-GaN seed, which was prepared by HVPE on a (0001) Al2 O3 followed by laser lift-off [23]. The TD density of the seed was approximately 107 –108 cm2 . The films were grown by AT method using a high-pressure autoclave [19–21] made of an Ni-based superalloy. Metallic Ga nutrient was placed in the low-temperature zone .520–530 ıC/ and the seed crystal was placed in the high-temperature zone .550–565 ıC/, because GaN has retrograde solubility at around 500 ıC in supercritical ammonia with basic NaNH2  NaI mineralizers [19] used in this study. The pressure was maintained at 150–200 MPa, and the growth was carried out for 72 h. Details of the growth have been described in [19, 20]. The surface morphology was observed with scanning electron microscopy (SEM) and atomic force microscopy (AFM). The microstructure was characterized with transmission electron microscopy (TEM) operated at 200 keV. TEM samples were N > axis by mechanical polishing prepared for observation nearly along the < 1010 C followed by Ar milling. For determination of polarity, convergent beam electron diffraction was simultaneously observed using the TEM specimen. Trace impurities were identified by secondary-ion mass spectrometry (SIMS). Wide-area .200 m¥/ PL spectra were recorded at 293 K using the 325.0 nm line of a cw He  Cd laser .38 W cm2 / as an excitor.

13

Optical Properties of GaN Substrates

Fig. 13.6 Plan-view and cross-sectional SEM images of Ga-polar (a, c) and N-polar (b, d) ammonothermal GaN films grown on a 400-m-thick HVPE FS-GaN seed wafer. (After [27])

287 Ga-polar

N-polar

a

b

RMS: 2.1 nm

RMS: 0.6 nm

10 µm (surface)

c

d

5 µm (cross section)

WAVELENGTH (nm) PL or CL INT. (arb. units)

Fig. 13.7 Wide-area (a) CL spectra at 100 K and (b) PL spectra at 293 K of AT-GaN films. For comparison, PL spectrum of the HVPE FS-GaN seed is shown in (b). (After [27])

600

500

(a) 100K CL 15kV, 1nA

400

C

D

Ga-polar (b) 293K PL He-Cd 325.0nm 12mW

340

A B

H

N-polar

HVPE GaN seed (x 1/10)

2.0

2.5 3.0 3.5 PHOTON ENERGY (eV)

Plan-view and cross-sectional SEM images of the 5-m-thick Ga-polar (4-mthick N-polar) AT-GaN films are shown in Fig. 13.6a, c (Fig. 13.6b, d), respectively. Under the particular growth conditions used in this study, average film thickness was comparable with each other. The Ga-polar surface was filled with grooves originating from linearly aligned inverted hexagonal pyramids having preferentially N and (1012) N facets, which have a pit at the center, whereas N-polar surface (1011) was nearly featureless. However, the local root-mean-square (RMS) roughness for the flat region of the Ga-polar surface (0.6 nm) was smaller than that for the Npolar surface (2.1 nm). Suspicious source of the grooves and their influence on the behavior of TDs and on the CL spectra will be discussed later. As shown in Fig. 13.7a, b, the N-polar film exclusively exhibited a broad NBE emission band both at 100 and 293 K. As the peak energy (3.440 eV) was higher by 30 meV than the A- and B-free exciton peak of strain-free GaN . 3:41 eV/ [29] and the FWHM value (231 meV) was so large at 293 K, the emission is assigned as being due to the free-carrier recombination with pronounced Burstein–Moss (BM) shift [64]. From the FWHM value at 100 K being 200 meV, the residual electron concentration is estimated to be approximately 1020 cm3 [65]. Culprit impurities were inspected by SIMS measurement to be Si and O. On the other hand, PL spectrum at 293 K of the Ga-polar film resembled that of the HVPE FS-GaN seed, though the broadening of the NBE peak was obvious and a weak band at 3 eV was observed, as shown in Fig. 13.7b. Remarkable difference

288

S.F. Chichibu

against the N-polar film was found at 100 K, as shown in Fig. 13.7a: the wide-area CL spectrum exhibited a predominant peak at 3.444 eV and a higher energy shoulder around 3.57 eV, and bands at 3.27 eV, 2.92 eV, and 2.22 eV. They are labelled A, H, B, C, and D, respectively. Among these, D peak is the well-known YL band, of which origin has been assigned to VGa and/or VGa  O complexes [42,59–61]. The A peak energy was lower by 30 meV than the free A-exciton [29] energy .EA / at 100 K but it approached EA at 293 K being 3.407 eV. The residual electron concentration of the sample is estimated from the FWHM value of A peak at 100 K being 94 meV to be of the order of 1019 cm3 [65], which is higher than the Mott density [40]. Therefore, the A peak is assigned to a free-to-bound (FB) transition at 100 K and a free carrier recombination at 293 K. Note that the excitonic contribution cannot be ruled out. The energy of H peak was approximately 3.58 eV and 3.51 eV at 100 and 293 K, respectively, and is assigned to a free carrier recombination with a BM shift. Apparently, the two NBE emissions A and H may not coexist in a homogeneous bulk film. As the film had odd-shaped morphology, the spectra in Fig. 13.7a, b are considered to be convolutions of position-specific local spectra.

13.4 Effects of Dislocation Bending on the Optical Properties of Seeded GaN Substrates Grown by Ammonothermal Method Spatially-integrated or resolved CL was excited by a cw electron beam with or without the beam scanning, and dispersed by a grating monochromator .f D 20 cm/ equipped with SEM. The acceleration voltage and probe current were 15 kV and 1 nA, respectively. Monochromatic CL intensity mapping images were taken under a fixed wavelength with the beam scanning at 100 K. Spot-excitation CL spectra at 100 K, the corresponding SEM image, and monochromatic CL intensity mapping images for A  D peaks of the Ga-polar film are summarized in Fig. 13.8. The local spectra numbered 1–6 were taken at respective spots in Fig. 13.8b. As shown, local spectra varied spot to spot. The overall CL intensity was weak at the spots 2, 3, and 6, as shown in Fig. 13.8c–f, in which white areas correspond to those emitting bright lights. Considering our experimental configuration, low light collection efficiency might be the cause for the spots 2 and 3. The reason for the low CL intensity of spot 6 is due to the presence of the central pit that corresponds to the apex of a TD, which initiates the facet growth. At the planar region 5, BL band (C band) intensity was strong and therefore the incorporation efficiency of a specific impurity responsible for C band is higher for the (0001) face. The most suspicious impurity is Mg or C, both of which were detected by SIMS measurement. The peak B at 3.27 eV was found at all spots 1–6. Although the peak energy was close to that of the NBE emission of zincblende GaN [or, basal plane stacking fault (SF) in wurtzite GaN], noticeable SF was not found, as shown in the cross-sectional TEM images in Fig. 13.9a, c. Therefore, the origin of B peak is assigned to an impurity such as Mg. The most interesting

13

Optical Properties of GaN Substrates

289

WAVELENGTH (nm) CL INTENSITY (arb. units)

600 500

Ga-polar GaN

400 340

15kV, 1nA, 100K wide-area C

D

(a) CL spectra B A H

1 2 3 4 5 6

6 4

2.0 2.5 3.0 3.5

1 mm

2

1

5

3

(b) SEM top view

PHOTON ENERGY (eV)

6 4

2 3

1

6 5

4

2 3

1

5

(c) CL map A (3.44 eV) (d) CL map B (3.27 eV)

6 4

1

2 3

6 5

4

1

2 3

5

(e) CL map C (2.92 eV) (f) CL map D (2.23 eV)

Fig. 13.8 (a) Wide-area and spot-excitation CL spectra at 100 K, (b) representative local SEM image, and (c)–(f) corresponding monochromatic CL intensity mapping images recorded for various photon energies at 100 K of the 5 m-thick Ga-polar AT GaN film. The spectra numbered 1–6 in (a) were taken at the numbered spots in (b)–(f). In (c)–(f), white areas correspond to those emitting bright lights. (After [27])

feature is that A peak intensity was remarkably enhanced at the spots 1, 2, and 3, where edge TD density was reduced by the dislocation bending owing to the facet growth [56], as shown in Fig. 13.9c. The average TD density was approximately 109 cm2 at the planar region and lower than 106 cm2 at the crests of the ridges (spot 1 for example). Therefore, wide-area low TD density region may be prepared by this mechanism when thicker film is continuously grown. Similar to the case for the FS-GaN substrate prepared by HVPE [9], a flat GaN wafer can then be prepared by slicing and polishing thick layers or boules. In contrast to the Ga-polar film, lineshape of the local CL spectra of N-polar film was nearly position-independent, as shown in Fig. 13.10a. The result may reflect the featureless morphology of the film. However, the NBE peak exhibited a noticeable intensity contrast in the CL mapping image, as shown in Fig. 13.10c. As the diameter of circular shaped bright (white in the figure) areas ranged between smaller than a few hundred nm and 1 m, they are considered to correspond to those between TDs.

290

S.F. Chichibu N-polar GaN

Ga-polar GaN

a

b

g=0002

g=0002

c

g=1120

1.3 mm

d

g=1120

1.1 mm

Fig. 13.9 Cross-sectional TEM images of (a,c) Ga-polar and (b,d) N-polar AT GaN films. The g vectors are < 0002 > for (a) and (b), < 1N 1N 20 > for (c), and < 112N 0 > for (d). (After [19, 27])

N-polar GaN CL INTENSITY (arb. units)

WAVELENGTH (nm) 600 500

1 mm

1

2

400 340

3

15kV, 1nA, 100K wide-area 1 2 3 4

2.0 2.5 3.0 3.5 PHOTON ENERGY (eV)

(a) CL spectra

4

(b) SEM top view 1

2 3

4

(c) CL map (3.50 eV)

Fig. 13.10 (a) Wide-area and spot-excitation CL spectra at 100 K, (b) representative local SEM image, and (c) monochromatic CL intensity mapping images recorded at 3.50 eV at 100 K of the 4 m-thick N-polar AT GaN film. The spectra numbered 1–4 in (a) were taken at the numbered spots in (b). In (c), white areas correspond to those emitting bright lights. (After [27])

Indeed, the TD density estimated from Fig. 13.9b, d is in the order of 109 cm2 , which correspond to the TD spacing of approximately 316 nm.

13.5 Summary In summary, optical properties of GaN substrates prepared by three methods were surveyed to discuss the superiority and critical issues of each method. Both the FSGaN substrates prepared by MOVPE using LEO technique and by HVPE using FIELO technique exhibited excellent PL properties, in terms of the observation of free exciton fine structures in their low temperature PL spectra. The intensities of

13

Optical Properties of GaN Substrates

291

characteristic YL and BL bands, which originate from point defect complexes and impurities, were more than two orders of magnitude lower than that of the NBE emission at room temperature for both cases. In addition, exciton–polariton features were found in the low temperature PR spectra of the MOVPE FS-GaN substrate, which exhibited a long PL lifetime of 860 ps at 300 K. The results indicate that they can also be used as a seed crystal for AT growth, except for the wafer bowing problem originating from in-plain compressive strain. Then, spot-excitation local CL spectra of Ga- and N-polar AT GaN films were correlated with submicrometer scale morphological features. The N-polar film contained high density residual electrons due presumably to the incorporation of O and Si, which gave rise to pronounced BM shift in the free carrier recombination emission band. On the other hand, the Ga-polar film exhibited rich spectral and intensity variations originating from facet growth. As a result of dislocation bending during the facet growth, TD density was greatly reduced at the crests of ridges. The results encourage one to grow low TD density GaN substrates by growing thick layers and/or boules by AT method and subsequent slicing. Acknowledgements The author would like to thank S. Nakamura, S.P. DenBaars, J.S. Speck, U.K. Mishra, A. Uedono, and T. Sota for having long-lived collaborations. He is also thankful to T. Hashimoto, S. Keller, K. Fujito, H. Marchand, M.S. Minsky, P.T. Fini, J.P. Ibbetson, and S.B. Fleischer of UCSB and Y. Ishihara and A. Usui of Furukawa Co. Ltd. for providing the samples. He is also grateful to T. Onuma and M. Kubota for optical measurements. This work was supported in part by Grant-in-Aids of CANTech, IMRAM, Tohoku Univ., NEDO project by METI and Scientific Research in Priority Areas No. 18069001 under MEXT, Japan.

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16. M.J. Callahan, B. Wang, L.O. Bouthillette, S.Q. Wang, J.W. Kolis, D.F. Bliss, Mater. Res. Soc. Symp. Proc. 798, Y2.10 (2004) 17. A. Yoshikawa, E. Ohshima, T. Fukuda, H. Tsuji, K. Ohshima, J. Cryst. Growth 260, 67 (2004) 18. T. Hashimoto, K. Fujito, B.A. Haskell, P.T. Fini, J.S. Speck, S. Nakamura, J. Cryst. Growth 275, e525 (2005) 19. T. Hashimoto, K. Fujito, F. Wu, B.A. Haskell, P.T. Fini, J.S. Speck, S. Nakamura, Mater. Res. Soc. Symp. Proc. 831, E2.8.1 (2005) 20. T. Hashimoto, K. Fujito, M. Saito, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 44, L1570 (2005) 21. T. Hashimoto, F. Wu, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 46, L525 (2007) 22. T. Hashimoto, F. Wu, J.S. Speck, S. Nakamura, Jpn. J. Appl. Phys. 46, L889 (2007) 23. B.A. Haskell, F. Wu, M.D. Craven, S. Matsuda, P.T. Fini, T. Fujii, K. Fujito, S.P. DenBaars, J.S. Speck, S. Nakamura, Appl. Phys. Lett. 83, 644 (2003) 24. D. Ehrentraut, Y. Kagamitani, T. Fukuda, F. Orito, S. Kawabata, K. Katano, S. Terada, J. Cryst. Growth 310, 3902 (2008) 25. R. Dwilinski, R. Doradzinski, J. Garczynski, L. P. Sierzputowski, A. Puchalski, Y. Kanbara, K. Yagi, H. Minakuchi, H. Hayashi, J. Cryst. Growth 310, 3911 (2008) 26. R. Dwilinski, R. Doradzinski, J. Garczynski, L. Sierzputowski, A. Puchalski, M. Rudzinski, M. Zajac, presented at International Workshop on Nitride Semiconductors 2008, Oct. 6–10, Montreux, Switzerland, No. We1-C1 (2008) 27. S.F. Chichibu, T. Onuma, T. Hashimoto, K. Fujito, F. Wu, J.S. Speck, S. Nakamura, Appl. Phys. Lett. 91, 251911 (2007) 28. S.F. Chichibu, H. Marchand, M.S. Minsky, S. Keller, P.T. Fini, J.P. Ibbetson, S.B. Fleischer, J.S. Speck, J.E. Bowers, E. Hu, U.K. Mishra, S.P. DenBaars, T. Deguchi, T. Sota, S. Nakamura, Appl. Phys. Lett. 74, 1460 (1999) 29. S.F. Chichibu, K. Torii, T. Deguchi, T. Sota, A. Setoguchi, H. Nakanishi, T. Azuhata, S. Nakamura, Appl. Phys. Lett. 76, 1576 (2000) 30. T. Kashiwagi, S. Sonoda, H. Yashiro, Y. Ishihara, A. Usui, Y. Akasaka, M. Hagiwara, Jpn. J. Appl. Phys. 46, 581 (2007) 31. A.A. Yamaguchi, Y. Mochizuki, H. Sunakawa, A. Usui, J. Appl. Phys. 83, 4542 (1998) 32. A.A. Yamaguchi, Y. Mochizuki, C. Sasaoka, A. Kimura, M. Nido, A. Usui, Mater. Res. Soc. Symp. Proc. 482, 893 (1998) 33. J.J. Hopfield, Phys. Rev. 112, 1555 (1958) 34. B. Gil, S. Clur, O. Briot, Solid State Commun. 104, 267 (1997) 35. P. Stepniewski, K.P. Korona, A. Wysmolek, J.M. Baranowski, K. Pakula, M. Potemski, I. Gregory, S. Porowski, Phys. Rev. B 56, 15151 (1997) 36. K. Torii, T. Deguchi, T. Sota, K. Suzuki, S. Chichibu, S. Nakamura, Phys. Rev. B 60, 4723 (1999) 37. M. Cardona, Modulation Spectroscopy, Solid State Physics. Suppl. 11, ed. by S. Seitz, D. Turnbull, H. Ehrenreich (Academic, New York, 1969) 38. D.E. Aspnes, Handbook on Semiconductors, Vol. 2, Chap. 4A, ed. by T.S. Moss (North-Holland, Amsterdam, 1980), p. 109 39. D.E. Aspnes, Surf. Sci. 37, 418 (1973) 40. S. Chichibu, T. Azuhata, T. Sota, S. Nakamura, J. Appl. Phys. 79, 2784 (1996) 41. A. Uedono, S.F. Chichibu, Z.Q. Chen, M. Sumiya, R. Suzuki, T. Ohdaira, T. Mikado, T. Mukai, S. Nakamura, J. Appl. Phys. 90, 181 (2001) 42. S.F. Chichibu, A. Uedono, T. Onuma, T. Sota, B.A. Haskell, S.P. DenBaars, J.S. Speck, S. Nakamura, Appl. Phys. Lett. 86, 021914 (2005) 43. S.F. Chichibu, A. Uedono, T. Onuma, B.A. Haskell, A. Chakraborty, T. Koyama, P.T. Fini, S. Keller, S.P. DenBaars, J.S. Speck, U.K. Mishra, S. Nakamura, S. Yamaguchi, S. Kamiyama, H. Amano, I. Akasaki, J. Han, T. Sota, Nat. Mater. 5, 810 (2006) 44. W. Shan, T.J. Shmidt, X.H. Yang, S.J. Hwang, J.J. Song, B. Goldenberg, Appl. Phys. Lett. 66, 985 (1995) 45. S. Chichibu, T. Mizutani, T. Shioda, H. Nakanishi, T. Deguchi, T. Azuhata, T. Sota, S. Nakamura, Appl. Phys. Lett. 70, 3440 (1997)

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46. G.D. Cody, T. Tiedje, B. Abeles, Y. Goldstein, Phys. Rev. Lett. 47, 1480 (1981) 47. L. Vina, S. Logothetidis, M. Cardona, Phys. Rev. B 30, 1979 (1984) 48. Y.P. Varshni, Physica 34, 149 (1967) 49. T. Azuhata, T. Matsunaga, K. Shimada, K. Yoshida, T. Sota, K. Suzuki, S. Nakamura, Physica B 219–220, 493 (1996) 50. H. Marchand, J. Ibbetson, P.T. Fini, P. Kozodoy, S. Keller, S.P. DenBaars, J.S. Speck, U.K. Mishra, Mater. Res. Soc. Internet J. Nitride Semicond. Res. 3, 3 (1998) 51. H. Marchand, X.H. Wu, J. Ibbetson, P.T. Fini, P. Kozodoy, S. Keller, J.S. Speck, S.P. DenBaars, U.K. Mishra, Appl. Phys. Lett. 73, 747 (1998) 52. S. Chichibu, A. Shikanai, T. Azuhata, T. Sota, A. Kuramata, K. Horino, S. Nakamura, Appl. Phys. Lett. 68, 3766 (1996) 53. A. Shikanai, T. Azuhata, T. Sota, S. Chichibu, A. Kuramata, K. Horino, S. Nakamura, J. Appl. Phys. 81, 417 (1997) 54. S. Chichibu, T. Azuhata, T. Sota, H. Amano, I. Akasaki, Appl. Phys. Lett. 70, 2085 (1997) 55. S.F. Chichibu, A. Uedono, T. Onuma, B.A. Haskell, A. Chakraborty, T. Koyama, P.T. Fini, S. Keller, S.P. DenBaars, J.S. Speck, U.K. Mishra, S. Nakamura, S. Yamaguchi, S. Kamiyama, H. Amano, I. Akasaki, J. Han, T. Sota, Philos. Mag. 87, 2019 (2007) 56. A. Usui, H. Sunakawa, A. Sakai, A.A. Yamaguchi, Jpn. J. Appl. Phys. 36, L899 (1997) 57. A. Wysmolek, M. Potemski, R. Stepniewski, J.M. Baranowski, D.C. Look, S.K. Lee, J.Y. Han, Phys. Status Solidi B 235, 36 (2003) 58. T. Azuhata, T. Sota, K. Suzuki, S. Nakamura, J. Phys. Condens. Matter 7, 129 (1995) 59. J. Neugebauer, C.G. Van de Walle, Phys. Rev. B 50, 8067 (1994) 60. J. Neugebauer, C.G. Van de Walle, Appl. Phys. Lett. 69, 503 (1996) 61. K. Saarinen, T. Laine, S. Kuisma, J. Nissilä, P. Hautojärvi, L. Dobrzynski, J. M. Baranowski, K. Pakula, R. Stepniewski, M. Wojdak, A. Wysmolek, T. Suski, M. Leszczynski, I. Grzegory, S. Porowski, Phys. Rev. Lett. 79, 3030 (1997) 62. P. Hacke, T. Detchprohm, K. Hiramatsu, N. Sawaki, K. Tadatomo, K. Miyake, J. Appl. Phys. 76, 304 (1994) 63. Z.-Q. Fang, D.C. Look, J. Jasinski, M. Benamara, Z. Liliental-Weber, R.J. Molnar, Appl. Phys. Lett. 78, 332 (2001) 64. E. Burstein, Phys. Rev. 93, 632 (1954) 65. E. Iliopoulos, D. Doppalapudi, H.M. Ng, T.D. Moustakas, Appl. Phys. Lett. 73, 375 (1998)

Chapter 14

Point Defects and Impurities in Bulk GaN Studied by Positron Annihilation Spectroscopy Filip Tuomisto

Abstract Positron annihilation spectroscopy is an experimental technique that allows for the selective detection of vacancy defects in semiconductors, providing a means to both identify and quantify them. This chapter first gives a short introduction to positron annihilation techniques, and then proceeds to present an overview of the positron studies of vacancy defects and impurities in bulk and quasi-bulk GaN. Both the in-grown and processing induced defects are addressed.

14.1 Introduction A variety of techniques can be applied to identify defects in semiconductors on the atomic scale. The main advantage of positron annihilation spectroscopy lies in its ability to selectively detect vacancy-type defects. The positron has two special properties that give this asset: it has a positive charge, and it annihilates with electrons. An energetic positron that has penetrated into a solid rapidly loses its energy and then lives a few hundred picoseconds in thermal equilibrium with the environment. During its thermal motion, the positron interacts with defects, which may lead to its getting trapped in a localized state. Hence, the final positron annihilation with an electron can happen from various states. The sensitivity of positrons to vacancy-type defects is rather easy to understand. A free positron in a crystal lattice feels strong repulsion from the positive ion cores, and an open-volume defect like a vacant lattice site is hence an attractive center where the positron gets trapped. The reduced electron density at the vacant site increases the positron lifetime, while the missing valence and core electrons cause substantial changes in the momentum distribution of the annihilating electrons. Two positron techniques have been efficiently used in defect studies in semiconductors, namely the positron lifetime and the Doppler broadening of the 511 keV line. These methods provide a straightforward tool for the identification of vacancy-type defects. In addition, the positron techniques are strongly supported by theory, as the annihilation characteristics can be calculated from first principles. A significant advantage of positron annihilation spectroscopy is that it can be applied to both bulk crystals and thin layers of any electrical conduction type. 295

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Vacancy defects in GaN can exist in both sublattices, those in the Ga sublattice (called Ga vacancies) being more likely to be observed by positrons because of their larger size and nonpositive charge states. Defect studies in GaN using positron annihilation spectroscopy have been carried out mostly by one research group, namely the one at TKK, since the late 1990s. The aim of this chapter is to give an overview of the results on vacancies and other point defects obtained with positron annihilation spectroscopy in bulk GaN, grown by either the high nitrogen pressure (HNP) method or by hydride vapor phase epitaxy (HVPE). The focus will be on the results obtained on the vacancy defects in bulk GaN; the principles of positron annihilation spectroscopy are described very shortly – for a full picture of positron annihilation in semiconductors, the reader is referred to earlier review articles [1–4], books, and book chapters [5–10], to the conference proceedings of the International Conference on Positron Annihilation (ICPA), and to other references therein. This chapter is organized as follows: The principles of positron annihilation spectroscopy are very briefly reviewed in Sect. 14.2. In Sect. 14.3, the results obtained on in-grown defects in HNP [11–17] and HVPE [16–29] GaN are reviewed. The positron results on artificial creation (e.g., by irradiation or implantation) and manipulation (e.g., by thermal annealing) of point defects in bulk GaN [30–35] are reviewed in Sect. 14.4. Finally, some concluding remarks are made in Sect. 14.5, summarizing the chapter.

14.2 Positron Annihilation Spectroscopy This section gives a brief overview of the principles of positron annihilation spectroscopy and of the experimental techniques. A thermalized positron in a crystal lattice behaves like a free electron or a hole. Analogously, positrons have shallow hydrogenic states at negative ions such as acceptor impurities. Furthermore, vacancies and other centers with open volume act as deep traps for positrons. These defects can be experimentally detected by measuring either the positron lifetime or the momentum density of the annihilating positron–electron pairs (Doppler broadening of the annihilation radiation). For the sake of clarity, we will concentrate on the measurements of those two quantities, as these methods are the most used in defect studies in semiconductors. Descriptions of other techniques can be found in, e.g., [9].

14.2.1 Positrons in Solids Positrons are easily obtained from radioactive .ˇ C / isotopes such as 22 Na, where the positron emission is accompanied by a 1.27 MeV photon. This photon is used in positron lifetime experiments as the time signal of the positron emission from the source. The stopping profile of positrons from the ˇ C emission is exponential

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[6, 36]: for the 22 Na source where the positron energy distribution extends to Emax D 0:54 MeV, the positron mean stopping depth is about 40 m in GaN. The positrons emitted directly from a radioactive source thus probe the bulk of a solid. Low-energy positrons are needed for studying thin overlayers and near-surface regions. Positrons from the ˇ C emission are first slowed down and thermalized in a moderator, consisting usually of a thin film placed in front of the positron source and made of a material (e.g., W) that has a negative affinity for positrons. Thermalized positrons close to the moderator surface are emitted into the vacuum with energy of the order of 1 eV, and a beam is formed using electric and magnetic fields. The positron beam can be accelerated to a tunable energy ranging from a few tens of eV to 100 keV (corresponding to mean implantation depths from a few nanometers to more than 10 m), giving the possibility to control the positron stopping depth in the sample. After implantation and rapid (taking a few picoseconds) thermalization, positrons in semiconductors behave like free carriers, i.e., the positron state is a Bloch state in a defect-free lattice (living for a few hundred picoseconds). Various positron states yield specific annihilation characteristics that can be experimentally observed in the positron lifetime and Doppler broadening experiments. The positron wave function can be calculated from a one-particle Schrödinger equation [3], where the positron potential consists of two parts: the electrostatic Coulomb potential and a term that takes into account the electron–positron correlation effects. Many practical schemes exist for solving the positron state ‰C from the Schrödinger equation [3, 37, 38]. A positron state can be experimentally characterized by measuring the positron lifetime and the momentum distribution of the annihilation radiation. These quantities can also be calculated once the corresponding electronic structure of the solid system is known. The positron annihilation rate , the inverse of the positron lifetime , is proportional to the overlap of the electron and positron densities: 1= D  D  r02 c

Z

d r j‰C .r/j2 n.r/ Œn.r/ ;

(14.1)

where r0 is the classical electron radius, c the velocity of light, n.r/ the electron density, and Œn the enhancement factor that accounts for the pileup of the electron density at the positron beyond the (average) density n.r/ [3]. The momentum distribution .p/ of the annihilation radiation is a nonlocal quantity and requires knowledge of all the electron wave functions ‰i overlapping with the positron. It can be written in the form ˇ2 ˇZ p ˇ  r0 c X ˇˇ i pr ‰C .r/‰i .r/ i .r/ˇˇ ;

.p/ D dre ˇ V

(14.2)

i

where V is the normalization volume and i .r/ may depend only on i or on r. The Doppler broadening experiment measures the longitudinal momentum distribution along the direction of the emitted 511 keV photons, defined here as the z-axis:

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Z1 Z1

.pL / D

dpx dpy .p/ :

(14.3)

1 1

It should be noted that the momentum distribution .p/ of the annihilation radiation is mainly that of the annihilating electrons “seen by the positron,” because the momentum of the thermalized positron is negligible.

14.2.2 Positrons at Defects Positron States at Defects In analogy to free carriers, the positron also has localized states at lattice imperfections. At vacancy-type defects where one or several ions are missing, the repulsion sensed by the positron is lowered and the positron feels these kinds of defects as potential wells. As a result, localized positron states at open-volume defects are formed. The positron ground state at a vacancy-type defect is generally deep, the binding energy is about 1 eV or more [3]. In a vacancy defect, the electron density is locally reduced. This is reflected in the positron lifetimes, which are longer than in the defect-free lattice. For example, the calculated lifetimes in the unrelaxed Ga and N vacancies are 55 and 5 ps longer than in the perfect lattice [13]. The longer positron lifetime at VGa is due to the larger open volume compared with that of VN . The positron lifetime measurement is thus a probe of vacancy defects in materials. Positron annihilation at a vacancytype defect leads also to changes in the momentum distribution .p/ probed by the Doppler broadening experiment. The momentum distribution arising from valence electron annihilation becomes narrower due to a lower electron density. In addition, the localized positron at a vacancy has a reduced overlap with ion cores leading to a considerable decrease in annihilation with high momentum core electrons. In the case of GaN, where the dominant contribution to the high momentum part of the distribution comes from the Ga 3d electrons, the changes in the momentum distribution are more pronounced when positrons are trapped at Ga vacancies. The comparison of the measured positron lifetimes and Doppler broadening spectra with the theoretically calculated data for specific defects provides a very efficient tool for identification of the observed vacancy defects. Several ab initio approaches have been studied in recent years (see [37] and the references therein). The agreement between theory and experiment is excellent in terms of differences or ratios between the data for defects and the perfect lattice. Also quantitative agreement has been obtained in many materials, but all the theoretical schemes applied so far seem to have problems with the treatment of the partially bound 3d electrons, either under- or overestimating their contribution to the positron annihilation data. For example, in the case of GaN, the calculated lifetimes in the perfect lattice range from about 130 to 160 ps [13, 26], depending on the scheme, while the

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experimentally determined bulk lifetime is 160 ps [11,17,26,34]. On the other hand, the calculated differences in the lifetimes and the ratios of the momentum distributions between the defect and bulk states agree very well with the experiments [26, 37, 39]. A negatively charged impurity atom or an intrinsic point defect can bind positrons at shallow states even if these defects do not contain open volume [40, 41]. Being a positive particle, the positron can be localized at the hydrogenic (Rydberg) state of the Coulomb field around a negatively charged center. The situation is analogous to the binding of an electron to a shallow donor atom. The positron binding energy is typically Eion D 30–120 meV, indicating that positrons are thermally excited from the Rydberg states at 100–300 K. The hydrogenic positron state around a negative ion has a typical extension of 10–100 Å and thus positrons probe the same electron density as in the defect-free lattice. As a consequence, the annihilation characteristics (positron lifetime, positron–electron momentum distribution) are not different from those in the lattice. Although the negative ions cannot be identified with the experimental parameters, information on their concentration can be obtained in the positron lifetime and Doppler broadening experiments when they compete with vacancies in positron trapping.

Positron Trapping at Defects The transition of a positron from a free Bloch state to a localized state at a defect is called positron trapping. The trapping is analogous to carrier capture. However, in order to be observed, it must be fast enough to compete with annihilation. The positron trapping rate D onto a defect D is proportional to the defect concentration cD through D D D  cD . The trapping coefficient D depends on the defect and the host lattice. As the positron binding energy at vacancies is typically > 1 eV, the thermal escape (detrapping) of positrons from the vacancies can be usually neglected. Because of the Coulomb repulsion, the trapping coefficient at positively charged vacancies is so small that the trapping does not occur during the short positron lifetime of a few hundred picoseconds [42]. Hence, the detection of vacancies or other defects in a positive charge state is impossible with positron annihilation techniques. The trapping coefficient at neutral vacancies is typically D  1014 –1015 s1 independently of temperature [42–44]. This value means that neutral vacancies are observed when their concentration is 1016 cm3 (i.e., about 100–200 ppb). The positron trapping coefficient at negative vacancies is typically D  1015 –1016 s1 at room temperature [42–44]. The sensitivity to detect negative vacancies is thus 1015 cm3 (i.e., about 10–20 ppb). The experimental fingerprint of a negative vacancy is the increase of D with decreasing temperature .D  T 1=2 / [42– 44], which allows to experimentally distinguish negative vacancy defects from neutral ones. The positron trapping coefficient ion at the hydrogenic states around negative ions is of the same order of magnitude as that at negative vacancies [41, 45].

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Furthermore, the trapping coefficient exhibits a similar T 1=2 temperature dependence. Unlike in the case of vacancy defects, the thermal escape of positrons from the negative ions plays a crucial role at usual experimental temperatures. The principle of detailed balance yields the following equation for the detrapping rate ıion from the hydrogenic state [3]:  ıion D ion

2 m kB T h2

3=2

  Eion exp  : kB T

(14.4)

Typically, ion concentrations above 1016 cm3 influence positron annihilation at low temperatures .T < 100 K/, but the ions are not observed at high temperatures .T > 300 K/, where the detrapping rate (14.4) is large. In practice, the positron annihilation data is analyzed in terms of kinetic rate equations describing the positron transitions between the free Bloch states and localized states at defects [8–10]. Very often, the experimental data show the presence of two defects, one of which is a vacancy and the other is a negative ion. In this case, the average lifetime ave (the center of mass of the lifetime spectrum), positron–electron momentum distribution .pL /, and shape parameters S and W of the Doppler broadened annihilation line (representing annihilations with low momentum valence electrons and high momentum core electrons, respectively) are given by P D B PB C ion Pion C V PV ;

(14.5)

Here P is the experimental parameter and the subscripts “B, ion, and V” denote the characteristic parameters of the bulk, negative ions, and vacancies, respectively. Equation (14.5) allows the experimental determination of the trapping rates V and ion related to the experimental trapping fractions through B D 1  ion  V ; V V D ; B C V C 1Cıionion=ion ion ion D .1 C ıion =ion / B C V C

(14.6) (14.7)

ion 1Cıion =ion

:

(14.8)

and consequently the defect concentrations can be obtained. Furthermore, these equations enable the combination of positron lifetime and Doppler broadening results and various correlations between ave ; .pL /; S , and W can be studied. At high temperatures, all positrons escape from the hydrogenic state of the negative ions and no annihilations take place at them. Then ion D 0 and ıion =ion >> 1, and hence the determination of the positron trapping rate and vacancy concentration is straightforward using (14.5–14.8) V D V cV D B

ave  B S  SB W  WB D B D B : V  ave SV  S WV  W

(14.9)

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It should be noted that in this case, ave ; S , and W depend linearly on one another. The linearity of experimental points in the .ave ; S /; .ave ; W /, and (S, W) plots provides thus evidence that positrons annihilate from two distinguishable states, indicating that they are trapped at only a single type of vacancy defect in the samples.

14.2.3 Experimental Techniques Positron lifetime spectroscopy is a powerful technique in defect studies, because the various positron states appear as different exponential decay components. The number of positron states, their annihilation rates, and relative intensities can be determined. In a positron lifetime measurement, one needs to detect the start and stop signals corresponding to the positron entrance and annihilation times in the sample, respectively. The 1.27 MeV photon that follows the positron emission from the 22 Na isotope is a suitable start signal. The 511 keV annihilation photon serves as the stop signal. The positron source is prepared by sealing about 10 Ci (about 105 –106 Bq) of radioactive isotope between two thin foils. The source is then sandwiched between two identical pieces (e.g., 5  5  0:5 mm3 ) of the sample material. This technique is standard for bulk crystal studies. Pulsed positron beams have been constructed for lifetime spectroscopy in thin layers [46, 47], but so far they have not been used much in defect studies. Typically, about 106 lifetime events are recorded in 1 h. The experimental spectrum represents the probability of positron annihilation at time t and it consists of exponential decay components: 

X dn.t/ Ii i ei t ; D dt

(14.10)

i

where n.t/ is the probability of the positron to be alive at time t. The decay constants i D 1=i are called annihilation rates and they are the inverses on the positron lifetimes i . Each positron lifetime has the intensity of Ii . In practice, the ideal spectrum of (14.10) is convoluted by a Gaussian resolution function which has a width of 200– 250 ps (full width at half maximum, FWHM). About 5–10% of positrons annihilate in the source material and proper “source corrections” must be made. Because of the finite time resolution, annihilations in the source materials, and random background, typically only 1–3 lifetime components can be resolved in the analysis of the experimental spectra. The separation of two lifetimes is successful only if the ratio œ1 =œ2 is  1.3–1.5. Figure 14.1 shows positron lifetime spectra recorded in two freestanding GaN samples grown by HVPE and the HNP method [17, 34]. Positrons enter the sample and thermalize at the time t D 0. The vertical axis of Fig. 14.1 gives the number of annihilations at a time interval of 25 ps. In the HVPE-GaN sample where the O concentration is low, the positron annihilation spectrum has a single lifetime

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105 n-type HNP GaN [O] ~ 1020 cm–3 τave = 190 ps

Number of counts

104

103 n-type HVPE GaN [O] ~ 1016 cm–3 τave = 160 ps

102

101 0.0

0.5

1.0 Time (ns)

1.5

2.0

Fig. 14.1 Examples of positron lifetime spectra in free-standing GaN crystals grown by two different methods with different concentrations of unintentional oxygen: hydride vapor phase epitaxy (HVPE) and the high nitrogen pressure (HNP) method. A constant background and annihilations in the source materials have been subtracted from the spectra, which consist of about 2106 recorded annihilation events. The solid lines are fits of the sum of exponential components convoluted with the resolution function of the spectrometer. The data in samples were recorded at 300 K. In the case of the HVPE sample, the annihilation spectrum has only a single component of 160 ˙ 1 ps. The spectrum in the HNP sample can be decomposed into two components of which the higher is 2 D 235 ˙ 5 ps

component of 160 ˙ 1 ps at 300 K corresponding to positron annihilations in the defect-free lattice. The spectrum measured in the HNP-GaN sample with high O content has two lifetime components, the longer of which .2 D 235 ps/ is due to positrons annihilating at Ga vacancies. The experimental results are often presented in terms of the average positron lifetime ave defined as Z1 ave D 0

 Z1  X dn D dt n.t/ D dt t  Ii  i : dt 0

(14.11)

i

The average lifetime is a statistically accurate parameter, because it is equal to the center-of-mass of the experimental lifetime spectrum. Hence, it can be correctly calculated from the intensity and lifetime values even if the decomposition represented only a good fit to the experimental data without any physical meaning. For example, the positron average lifetimes in the two spectra of Fig. 14.1 are 160 ps

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(HVPE-GaN) and 190 ps (HNP-GaN). The difference is very significant as changes in the average lifetime less than 1 ps can be reliably observed in the experiments. Doppler broadening spectroscopy is often applied especially in the low-energy positron beam experiments, where lifetime spectroscopy is usually very difficult because of the missing start signal. The motion of the annihilating electron–positron pair causes a Doppler shift in the annihilation radiation with E D cpL =2, where pL is the longitudinal momentum component of the pair in the direction of the annihilation photon emission. This causes the broadening of the 511 keV annihilation line. The shape of the 511 keV peak gives thus the one-dimensional momentum distribution .pL / of the annihilating electron–positron pairs. A Doppler shift of 1 keV corresponds to a momentum value of pL D 0:54 a.u. .3:91  103 m0 c/. The Doppler broadening can be experimentally measured using a Ge gamma ray detector with a good energy resolution. For measurements of bulk samples, the same source–sample sandwich is used as in the lifetime experiments. For layer studies, the positron beam hits the sample and the Doppler broadening is often monitored as a function of the beam energy. The typical resolution of a detector is around 1–1.5 keV at 500 keV. This is considerable compared to the total width of 2–3 keV of the annihilation peak meaning that the experimental lineshape is strongly influenced by the detector resolution. Therefore, various shape parameters are used to characterize the 511 keV line. The low electron-momentum parameter S is defined as the ratio of the counts in the central region (typically pL < 0:4 a.u.) of the annihilation line to the total number of the counts in the line. In the same way, the high electron-momentum parameter W is the fraction of the counts in the wing regions of the line (typically pL > 1:5 a.u.). A good way of choosing the integration windows for the S and W parameters is best seen in Fig. 14.2. The S parameter is integrated from the Gaussian part (given by the momentum distribution of the unbound or only weakly bound valence electrons) so that it includes roughly 50% of the total counts in the peak. The lower limit of the W parameter window is chosen so that the dominant contribution to that part of the spectrum comes from the exponential tails (linear in the semilog plot) of the core electron distributions. In order to have as good statistics as possible, the upper limit is set as high as reasonable from the data scatter point of view. The proper choice of the lower limit of the W parameter window depends on the studied material. Because of the above considerations S and W are sometimes called the valence and core annihilation parameters, respectively. The high-momentum part of the Doppler broadening spectrum arises from annihilations with core electrons that contain information on the chemical identity of the atoms. Thus, the detailed investigation of core electron annihilation can reveal the nature of the atoms in the regions where positrons annihilate. In order to study the high-momentum part in detail, the experimental background needs to be reduced. A second gamma ray detector can be placed opposite to the Ge detector and the only events that are accepted are those for which both 511 keV photons are detected [48, 49]. Depending on the type of the second detector, electron momenta even up to p  8 a.u. .60  103 m0 c/ can be measured with the coincidence detection of the Doppler broadening (see Fig. 14.2).

304

F. Tuomisto Electron momentum (10–3 m0c) 10 20

0

30

Momentum distribution (a.u.)

1 Gallium nitride GaN lattice Ga vacancy 0.1

0.01

0.001

S 0

W 1 2 3 Electron momentum (a.u.)

4

Fig. 14.2 The positron–electron momentum distribution measured in the GaN lattice and in an electron irradiation-induced Ga vacancy. Possible definitions of the integration areas for the S and W parameters are shown. It should be noted that the Doppler broadened spectrum is symmetrical: it has been folded around zero momentum in order to increase statistics

14.3 In-Grown Defects GaN crystals grown by HVPE and the HNP method, both doped and nominally undoped, have been the subject of numerous positron defect studies since the first identification of the native Ga vacancy in GaN [11]. These samples are characterized by various impurity profiles and growth temperatures, giving rise to different vacancy concentrations. Furthermore, the growth polarity has an important role in determining both the incorporation of impurities and the formation of extended and point defects in the material. In the following, an attempt is made to gather the pieces of understanding presented in [11–13, 15–20, 23–29], providing a consistent picture of the formation of vacancy defects in bulk (or bulk-like) GaN growth.

14.3.1 Defect Formation: Growth Methods and Doping Figure 14.3 presents the average positron lifetime ave measured as a function of temperature in six differently doped samples grown by HVPE and HNP. In all samples where ave is above the bulk lifetime B D 160 ps indicating that positrons annihilate as trapped at vacancy defects, the annihilation spectrum is two-componential with the higher component being 2 D 235 ˙ 10 ps (not shown here). This component is characteristic of the Ga vacancy [11, 17, 24], and hence Ga vacancies are

Point Defects and Impurities in Bulk GaN

Fig. 14.3 The average positron lifetime measured as function of temperature in selected HVPE-GaN and HNP-GaN samples

305 undoped HNP GaN, [O] = 4 × 1019 cm–3 Mg-doped HNP GaN, [Mg] = 6 × 1019 cm–3

Average positron lifetime τave (ps)

14

Mg-doped HNP GaN, [Mg] = 1 × 1020 cm–3 210 200 O-doped HVPE GaN, [O] = 5 × 1019 cm–3 O-doped HVPE GaN, [O] = 4 × 1018 cm–3

190

undoped HVPE GaN, [O] = 1 × 1016 cm–3 180 170 160 0

100 200 300 400 Measurement temperature (K)

500

present at concentrations above 1015 cm3 (the sensitivity limit of the temperaturedependent positron measurements) in these samples. The slight increase of the average positron lifetime with increasing temperature in the heavily Mg-doped HNP-GaN sample, where positrons annihilate from the delocalized state in the defect-free lattice, is due to the thermal expansion of the GaN lattice. In the HVPE-GaN samples, the average positron lifetime increases with decreasing measurement temperature. This is a fingerprint of negatively charged vacancies, the trapping coefficient of which increases at low temperatures (see Sect. 14.2). Hence, the Ga vacancies are negatively charged in n-type HVPE-GaN, irrespective of the O content. In the nominally undoped and less heavily Mg-doped HNP-GaN samples the average positron lifetime decreases with decreasing measurement temperature. This indicates that negative ions are abundant enough to compete with vacancies in positron trapping at low temperatures, where the escape rate from the shallow Rydberg states is negligible. The data also show that the vacancies are negatively charged, as in the case of neutral vacancies the decrease would be faster in the Mg-doped sample, and the average positron lifetime could not stabilize at a value higher than the bulk lifetime in the nominally undoped sample. It should be noted that the negative ions cannot be identified on the basis of the positron data alone. However, their concentrations can be determined and compared to acceptor impurity concentrations measured with secondary ion mass spectrometry (SIMS). In the above HNP-GaN samples, this comparison has allowed the identification of the negative ions as Mg impurities [12]. The temperature dependence of the average lifetime can be modeled with kinetic trapping equations introduced in Sect. 14.2. The positron trapping coefficients at negative Ga vacancies V and negative ions ion vary as T 1=2 as a function of

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Table 14.1 Defect concentrations (given in 1017 cm3 ) in selected bulk or quasi-bulk GaN samples grown with different methods, doping and polarities [12, 17, 24, 28]. The impurity data were obtained by SIMS ŒVGa 

[Mg]

cion

2 0.7 <0.01

10 600 1,000

30 600

7

10

30

4 200

<0.01 7

0.4 2

50

0.3 2 4 500

0.02 0.04 0.2 1.5

1

3

SamplenDefect

[O]

HNP GaN Ga side – undoped –Mg-doped 1 –Mg-doped 2 N side –undoped

400 1,200 900 800

Homo-epitaxial HVPE GaN – Ga-polar (undoped) – N-polar (undoped) Hetero-epitaxial HVPE GaN c-plane (Ga-polar): – undoped – O-doped 1 –O-doped 2 –O-doped 3 a-plane: –undoped

temperature [3, 8]. The positron escape rate from the ions can be expressed as ı.T / / ion T 3=2 exp.Eion =kB T /, where Eion is the positron binding energy at the Rydberg state of the ions (14.4). The fractions of annihilations at Ga vacancies V and at negative ions ion are given in (14.6–14.8) and they depend on the concentrations cV D V =V and cion D ion =ion of Ga vacancies and negative ions (14.9), respectively, as well as on the detrapping rate ıion .T / (14.4). We take the conventional value V D 3  1015 s1 for the positron trapping coefficient at 300 K [8, 9]. By inserting the annihilation fractions B ; ion , and V from (14.6–14.8) into the equation for the average lifetime ave D B B C ion ion C V V (14.5), the resulting formula can be fitted to the experimental data. The vacancy and negative ion concentrations obtained this way in HVPE and HNP samples with various impurity contents are shown in Table 14.1. The positron binding energy to the shallow states around the negative ions is obtained as 60 ˙ 10 meV. The positron data in Fig. 14.3 show further that in the HVPE-GaN samples the defects involving Ga vacancies are the dominant negatively charged acceptors, while the negative ions (Mg impurities) dominate in HNP-GaN. Also observable in Table 14.1, the Ga vacancy concentrations correlate with the O concentration in O-doped HVPE-GaN where the O content ranges from ŒO D 2  1017 cm3 to ŒO D 2  1020 cm3 . The fact that the Ga vacancy concentration decreases with Mg doping in HNP-GaN can be explained by the increase in formation energy of acceptor-like defects when the Fermi level is lowered toward the mid-gap (the

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heavily Mg-doped HNP-GaN samples are semi-insulating). On the other hand, the increase of the free carrier concentration with increasing O doping in the HVPE-GaN samples does not move the Fermi level significantly, and hence the thermodynamical formation energy is not enough to explain the increase in Ga vacancy concentration. A logical explanation is that the Ga vacancies are present as VGa  ON complexes, where the factor limiting the Ga vacancy concentration at lower O doping levels would be the availability of O impurities instead of the formation energy dictated by the carrier concentration through the Fermi level. In fact, recent detailed studies of the Doppler broadening show that the dominant vacancy defect, responsible for the positron lifetime of 235 ps, is a complex of Ga vacancy and oxygen [24]. The positron results presented in this section show that Ga vacancies act as dominant compensating centers in n-type GaN. It is worth considering the situation in p-type GaN, where the formation of Ga vacancies is energetically unfavorable due their acceptor nature. Hence, the natural question is whether N vacancies could compensate the doping. The detection of vacancy defects on the N sublattice with positrons is not evident as a result of the small open volume generated by the missing N atom. Nevertheless, evidence of the existence of N vacancies complexed with Mg .VN  MgGa / has been obtained with positrons in Mg-doped (p-type) GaN grown by metal-organic chemical vapor deposition (MOCVD) [50]. It should be noted that even though the vacancy concentrations are similar relative to the doping densities (a few percent at most) in both n- and p-type GaN, the vacancies are dominant compensating centers only in n-type GaN, while in p-type GaN other defects and impurities such as hydrogen seem to play the most important role.

14.3.2 Defects and Growth Polarity The main conclusion of the previous section states that the Ga vacancies are present in n-type GaN crystals (grown by either HNP or HVPE) as VGa  ON pairs if the O concentration is high enough. However, the presented experiments did not reveal the formation process of these defects. Fortunately, the wurtzite structure of GaN gives a possibility to study the effects related to the growth polarity (or direction) of the material. The comparison of both homo- and hetero-epitaxial HVPE, grown in either polar or nonpolar direction, with HNP shines light on the question of impurity incorporation and defect formation in GaN [17, 18, 27, 28]. Table 14.1 shows the comparison between two homo-epitaxial HVPE-GaN layers. The layers were 30–60 m thick and homogeneous across the thickness, and they were grown in the same run on the Ga and N faces of an essentially dislocationfree HNP-GaN substrate. The data obtained in these two layers can be compared to those obtained from the Ga and N faces of the substrate (also in Table 14.1). Interestingly, the Ga vacancy concentrations .7  1017 cm3 / coincide in the N -polar HVPE-GaN and the N side of the HNP-GaN samples, similarly as the impurity concentrations obtained from SIMS. On the other hand, the difference between the Ga polar HVPE-GaN .ŒVGa  < 1015 cm3 / and the Ga side of the HNP-GaN

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GaN lattice

W parameter

0.95 0.90 0.85 0.80 0.75

HVPE a-plane GaN HVPE c-plane GaN He-irradiated GaN 1.00

1.01

1.02 1.03 S parameter

VGa − ON

VGa 1.04

1.05

Ga vacancy concentration (cm–3)

1.00

a-plane HVPE GaN/sapphire c-plane HVPE GaN/sapphire

1019

1018

1017

1016 0

20 40 60 Distance from the interface (μm)

Fig. 14.4 The .S; W / plot showing the Doppler broadening parameters in c-plane and a-plane GaN. Parameters from irradiated GaN are presented for reference (left). Ga vacancy concentrations in a-plane and c-plane HVPE GaN as a function of the distance from the sapphire interface (right)

bulk crystal .ŒVGa  D 2  1017 cm3 / is substantial. These observations support the idea presented in [51] that the oxygen incorporation (and subsequent Ga vacancy formation) is stronger in the nonpolar directions, in which the N polar growth mainly proceeds. The difference between the polarities is larger in the HVPE GaN samples than in the HNP bulk GaN crystals. This can be explained by the lower temperature and pressure at HVPE growth, which reduce the oxygen diffusion, and by the presence of more oxygen in the high-pressure growth. In order to further study the role of the growth polarity on the defect incorporation in GaN, a-plane GaN layers (thicknesses 5–25 m) grown by HVPE on sapphire were measured with a variable-energy positron beam. The left panel of Fig. 14.4 shows the S and W parameters measured in the a-plane GaN layers. The data are compared to those obtained in c-plane HVPE GaN [18, 52] and He-irradiated GaN. In the latter, the data points fall on the line connecting the parameters of the GaN lattice and the isolated VGa [24], indicating that these are created in the He-irradiation. On the other hand, the data from both a-plane and c-plane GaN follow the line connecting the parameters of the VGa  ON pair [24] with those of the GaN lattice, showing that these are the dominant defects observed in the layers. The concentrations of the VGa  ON pairs are shown as a function of distance from the interface in the right panel of Fig. 14.4 for both the a-plane and c-plane samples. The difference between the polar and nonpolar HVPE-GaN layers is clear: the Ga vacancy concentration is constant in the a-plane HVPE-GaN, whereas it decreases with increasing distance from the sapphire interface in c-plane HVPE-GaN. SIMS results show that the O concentration is also constant in the a-plane HVPE-GaN layers, as is the density of extended defects (observed with cross-sectional transmission electron microscopy) [28]. These results give further support for the model based on growth–surface– dependent oxygen incorporation and subsequent Ga vacancy formation. In c-plane hetero-epitaxial Ga-polar HVPE GaN, the O concentration profile is determined by the dislocation profile likely due to diffusion from the sapphire substrate. On the other hand, in homo-epitaxial c-plane Ga-polar HVPE GaN, where the dislocation

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density is low and the amount of oxygen in the substrate is significantly lower, no vacancies are observed even in relatively thin layers, while in N -polar GaN both the Ga vacancy and O concentrations are high. Hence, as the growth modes are similar (an important part of the growth proceeds with nonpolar surfaces) in the N -polar and nonpolar GaN, it is natural that the O incorporation from the growth ambient is effective in both, giving rise to a high O concentration and subsequent Ga vacancy concentration, independently from the possible extended defects. It should be noted, however, that the higher vacancy concentration in hetero-epitaxial a-plane HVPEGaN than in c-plane HVPE-GaN with similar oxygen content is most likely related to the high extended defect density. The above results, together with the fact that isolated Ga vacancies are mobile at around 600 K [30, 31] and the theoretically estimated [53] formation energies that have been recently experimentally verified [32], the following model of Ga vacancy formation can be proposed. The Ga vacancies are formed as isolated during growth of n-type GaN. They migrate fast at the high growth temperatures that are typically above 1,300 K in case of HVPE and HNP crystals, and are stabilized (quenched) by donor impurities during cooling down. This is demonstrated by the fact that the VGa  ON concentrations are similar in materials grown by HVPE and the HNP method when the O concentrations are similar, in spite of the large difference of about 500 K in the growth temperatures. In fact, it has been shown that the VGa  ON pairs are stable up to about 1,300 K [30–32]. On the other hand, the concentration of Ga vacancies in Si doped n-type GaN is significantly lower than in O doped ntype GaN with similar free electron concentration [54] because of the lower binding energy of the VGa  SiGa pair originating from the larger distance between the individual acceptor (Ga vacancy) and donor (substitutional Si) defects, and screening by N atoms.

14.4 Defect Engineering In order to gain a broad understanding of the nature of point defects in a semiconductor material, it is often practical to produce them at concentrations much higher than those given by the thermodynamic equilibrium, e.g., by electron irradiation. It is also of interest to study the behavior of both the irradiation-induced and in-grown defects in thermal treatments. The irradiation by electrons can also be used to mimic the defects introduced in ion implantation, a typical processing step in semiconductor technology. Ion implantation itself can be used for defect studies as well, in which case the focus is on studying larger defect complexes such as vacancy clusters and their formation, or on the interactions between the implanted species and the intrinsic defects created in the implantation. In the following, a synthesis of the knowledge gained in electron-irradiation and thermal annealing studies is presented. Positron experiments in ion-implanted GaN have until now been rather preliminary [55, 56], and will not be discussed here.

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14.4.1 High Pressure Thermal Annealing Earlier attempts, where thermal annealings on as-grown bulk GaN samples were performed up to temperatures of about 1,300 K [30, 31], were not successful in modifying the in-grown point defect distributions. As can be seen in Fig. 14.5, temperatures above 1,400 K were needed to produce a change in the positron data [32]. In the experiments, a high nitrogen pressure of 10 kbar was used to prevent the GaN samples from dissociating. The samples were 270 m thick GaN films grown by HVPE at 1,350 K on a MOCVD-GaN template on sapphire. Four selected crack-free, self-separated samples from one wafer were annealed for 1 h at four different temperatures in the range of 1,423–1,723 K. Positron lifetime experiments were performed at temperatures 10–300 K for high sensitivity. The fast positrons used in positron lifetime experiments enter the GaN lattice to an average depth of 30 m with an exponential stopping profile, and thus they probe approximately onethird of the sample thickness. Hence, in order to study the distribution of the defects along the c axis, the positron lifetime in the samples was measured with both the Ga and N -polar sides facing the source (in the same way as in the case of the HNP-GaN crystals in Sect. 14.3.2). Figure 14.5 shows the VGa  ON concentrations obtained from the positron data in the samples, together with the O concentrations estimated from comparing SIMS and photoluminescence (PL) data [32, 57]. The O concentration near the N face of the HVPE GaN samples is relatively high because of the higher dislocation density in the heteroepitaxial interface region, even when a buffer layer is used. The Ga vacancies that are more easily formed in the region of high O content (high n-type conductivity) during growth, survive the cooling down by forming complexes with the O impurities. Hence, the decrease in the Ga vacancy concentration is due to the dissociation of the VGa  ON pairs. An activation energy .EA / fitted to the Ga

1016 As-grown

Ga vacancies, N face O impurities, N face Ga vacancies, Ga face O impurities, Ga face

1018

1017 1015

Estimated O concentration (cm–3)

VGa-ON concentration (cm–3)

1019

1400 1500 1600 1700 Annealing temperature (K)

Fig. 14.5 The VGa  ON and oxygen impurity concentrations determined from the N and Ga-polar sides of the HVPE GaN layers as a function of annealing temperature. The dashed lines are guides to the eye

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Point Defects and Impurities in Bulk GaN

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vacancy concentrations as a function of annealing temperature gives EA D 3:1.2/ eV, which can be interpreted as the sum EM CEB of the migration energy of VGa and the binding energy of the VGa  ON complex. Taking the migration energy EM D 1:5.2/ eV determined from electron irradiation data in [30], we obtain a binding energy of EB D 1:6.4/ eV, in good agreement with the value of EB D 1:8–2:1 eV predicted by theory [53, 58]. The increase of the Ga vacancy concentration after the annealing at 1,723 K measured near the Ga side of the sample indicates that thermally formed vacancies are present in the material. The flattening of the vacancy concentration and O impurity profiles suggests that the sample has reached thermal equilibrium. In fact, also the compressive in-plane stress and the differences between both a and c lattice constants of the two faces disappear in the annealing [57]. The present results show that the VGa  ON pairs become unstable at 1,523 K, and hence the Ga vacancy concentration observed after cooling down of the sample is determined by the equilibrium concentration at 1,423–1,523 K. This is in good agreement with the early experiments, where the dissociation temperature of the VGa  ON pairs has been estimated to be 1,300–1,500 K [30]. These results support the model presented in Sect. 14.3, i.e., the Ga vacancies are formed as isolated at high growth temperatures and the final concentration is determined by the ability of the Ga vacancies to diffuse and bind to O impurities. Finally, assuming that all the O donors are ionized, the increase in concentration on the Ga side shifts the Fermi level up toward the conduction band from EC  0:7 eV to EC  0:5 eV at 1,723 K, at which temperature the Ga vacancies are stabilized by the O impurities. This shift is sufficient to decrease the formation energy of the negative Ga vacancies enough for them to be formed at concentrations detectable by positrons. Taking a typical value for the formation entropy of S D .5–10/ kB , the formation energy of the isolated Ga vacancy can be estimated from the equilibrium concentration at 1,723 K to be E f D 2:5–3:2 eV at Fermi level position EC  EF D 0:5 eV, in perfect agreement with the theoretical results of E f D 2:5–3:5 eV [53, 58].

14.4.2 Electron Irradiation Experiments The low dislocation density and impurity concentrations in state-of-the-art freestanding HVPE-GaN make it possible to perform controlled studies of defect introduction by irradiation with, e.g., electrons. Figure 14.6 shows results obtained in HVPE-GaN samples irradiated with 0.45 and 2 MeV electrons to fluences of 2  1017 cm2 and 5  1017 cm2 , respectively, and subsequently annealed at temperatures above room temperature where the irradiation was performed [34, 35]. In the as-grown samples a low concentration, about 4  1015 cm3 of in-grown Ga vacancies (VGa  ON pairs) was found. The in-grown defect structure was reached after sufficiently high annealing temperature in both irradiations: 600 K in the case of 0.45 MeV electrons and 1,120 K in the case of 2 MeV electrons. Treatments at higher temperatures did not produce any further changes in the data. Another

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F. Tuomisto 2 MeV e--irr. GaN, annealed up to 440 K annealed at 625 K annealed at 775 K annealed at 1080 K

200

190

180

170

160

0.45 MeV e--irr. GaN, annealed up to 460 K as-grown, 0.45 MeV + 600 K, 2-MeV + 1220 K

0

100

200

300

400

Measurement temperature (K)

2-MeV irradiation Ga vacancies 2- neg. ions 1- neg. ions

1018 Defect concentration (cm–3)

Average positron lifetime (ps)

210

500

1017

1016

1015

0.45-MeV irradiation N vacancies 400

600

800

1000

Annealing temperature (K)

Fig. 14.6 The average positron lifetime measured as a function of temperature in electronirradiated and subsequently annealed HVPE-GaN samples (left). The defect concentrations extracted from the positron data as a function of annealing temperature (right)

common feature is that the irradiation-induced defects were stable up to about 450 K for the two electron energies. The 0.45 MeV electron energy was chosen in order to produce damage in the N sublattice only, as the theoretically predicted threshold for producing damage in the Ga sublattice is 0.53 MeV [59]. Indeed, even if the average positron lifetime increased in the irradiation (Fig. 14.6), no sign of Ga vacancies was observed. Instead, a second lifetime component (not shown) 2 D 190 ˙ 15 ps could be separated from the annihilation spectra at temperatures where the difference to the as-grown case was the largest [34]. In addition, as the Doppler broadening spectrum does not change in the 0.45 MeV irradiation, the defect producing this lifetime is likely to be the neutral N vacancy: the changes in the Doppler spectrum are expected to be minimal because of the similar overlap with the Ga 3d electrons both in the GaN lattice and in the N vacancy. By comparison to electrical and optical data [59] and theoretical predictions [60], it is likely that these N vacancies have formed neutral complexes with hydrogen that is abundant enough in HVPE-GaN. Figure 14.6 also shows the evolution of the N vacancy concentration with annealing temperature: an activation energy fit for annealing gives EA D 1:8 ˙ 0:1 eV. As the N vacancies are predicted to be much more stable [61], this result suggests that the disappearance of the N vacancies in thermal annealing is due to N interstitials (presumably also bound to hydrogen) becoming mobile above 450 K and filling the vacancies. As can be seen in Fig. 14.6, the positron data and their behavior both as a function of measurement temperature and annealing temperature are much more complex in the case of 2 MeV electron irradiation. The separation of lifetime components

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(giving 2 D 235 ˙ 10 ps) in the annihilation spectra shows that Ga vacancies are present after the irradiation and throughout the whole annealing process, and the temperature dependence of the average positron lifetime shows that they are in the negative charge state. A detailed analysis [34] of the average positron lifetime as a function of measurement temperature after annealings at different temperatures shows that two kinds of negative ions are present at different annealing stages. After irradiation, negative ions with high binding energy (and charge 2 ) are present up to 600 K, at which temperature 90% of them disappear and negative ions with low binding energy (charge 1 ) emerge. At the same time, half of the negatively charged Ga vacancies disappear. The 2 and 1 ions coexist up to about 700 K, above which the remaining 10% of the 2 ions disappear rapidly. The 1 ions slowly disappear with increasing annealing temperature starting at 700 K up to about 1,100 K, and the remaining 50% of the Ga vacancies begin annealing out of the samples above 800 K, also slowly disappearing up to about 1,100 K. The first annealing stage of the Ga vacancies at around 600 K is in good agreement with the earlier observation [30] and the theoretical prediction [61], and verifies the value of the isolated Ga vacancy migration barrier as EM D 1:8˙0:1 eV. The activation energies of 1.9 and 2.3 eV could be fitted to the annealing data and change of charge state of the 2 ions. The changes in the charge state of the negative ions are likely to be associated with two subsequent reconstructions of the defects themselves. This interpretation is based on the general behavior of the Fermi level in the samples: in originally n-type material, the Fermi level typically moves toward the conduction band when the temperature is decreased, and when compensating irradiation-induced defects are removed by annealing. Hence, if the negative ion defects changed their charge state because of the movement of the Fermi level, they would become more negative in both scenarios, whereas in our experiments, they become less negative (from 2 to 1 ). One possible explanation could be, e.g., the release of hydrogen from other defect complexes and the subsequent migration and trapping at these negative ions, making them less negative. However, as the structure of these defects remains unresolved, further experiments are needed in order to understand the observed behavior of the negative ions. No single activation energy could be fitted to the annealing data above 800 K, where the 1 ions anneal out together with the other half of the Ga vacancies, which indicates that complex migration processes are involved. As the isolated Ga vacancies are known to be mobile already at 600 K, the remaining Ga vacancies need to be stabilized by other defects. N vacancies that are also produced in the 2 MeV electron irradiation, although at a lower rate, and hydrogen, are possible candidates as it would be difficult to distinguish Ga vacancy complexes with either (or both) of them from the isolated Ga vacancies. In addition, a significant part of the Ga vacancies cannot be complexed with oxygen as the O concentration is too low in our samples, and the breakup of VGa  ON complexes is known to occur only above 1,500 K (see Sect. 14.3) Interestingly, recent results of ODEPR measurements [62] (also in highpurity HVPE GaN samples irradiated with 2.5 MeV electrons to a similar fluence as in the work presented here) were interpreted as isolated Ga vacancies surviving the annealing at 600 K and disappearing above 800 K. However, the effect of hydrogen

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and Ga interstitials, which are likely to be positively charged and thus invisible to positrons, should be considered in more detail, and further experiments are needed here as well to clarify the situation (note that the H and Ga interstitials will be partly charged and partly neutral, as the Fermi level is likely to be close to their energy levels).

14.5 Summary In general, positron annihilation spectroscopy gives microscopic information about vacancy defects in semiconductors in the concentration range 1015 –1019 cm3 . The positron lifetime is the fingerprint of the open volume associated with a defect, and it can be used to identify mono- and divacancies and larger vacancy clusters. The Doppler broadening of the annihilation radiation is a measure of the momentum distribution of the annihilating electrons. It can be used to identify the nature of the atoms surrounding the vacancy. Consequently, vacancies on different sublattices of a compound semiconductor can be distinguished and impurities associated with the vacancies can be identified. The charge state of a vacancy defect can be determined by the temperature dependence of the positron trapping coefficient. Positron localization into Rydberg states around negative centers yields information about ionic acceptors that have no open volume. In this chapter, the results obtained on vacancy defects and impurities in bulk or bulk-like GaN by applying positron annihilation spectroscopy are described and reviewed. The most important results include the identification of the Ga vacancy complexes as dominant intrinsic acceptors in n-type O-doped GaN and the increase in the understanding of physical processes leading to the formation of VGa  ON pairs in the high-temperature growth of GaN by HVPE or HNP. Also, many of the intrinsic properties of the point defects on both Ga and N sublattices have been addressed by performing electron irradiation and annealing experiments. Many of the results, especially those concerning the thermodynamic formation energies or migration barriers of vacancies, have been found to be in very good or even excellent agreement with theoretical predictions providing valuable benchmarks for theorists.

References 1. R.N. West, Adv. Phys. 22, 263 (1973) 2. P.J. Schultz, K.G. Lynn, Rev. Mod. Phys. 60, 701 (1988) 3. M.J. Puska, R.M. Nieminen, Rev. Mod. Phys. 66, 841 (1994) 4. P. Asoka-Kumar, K.G. Lynn, D.O. Welch, J. Appl. Phys. 76, 4935 (1994) 5. P. Hautojärvi (ed.), Topics in Current Physics, vol. 12, Positrons in Solids (Springer, Heidelberg, 1979) 6. W. Brandt, A. Dupasquier (eds.), Positron Solid-State Physics (North-Holland, Amsterdam, 1983)

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Index

Acceptor dopants Zn and Mg, 144 Acidic, 184 Acidic ammonothermal, 185, 190, 198, 199 Acidic mineralizers, 162, 184, 188, 198 Acidic regime, 139 Activation energies, 47, 250 Activation energy for homogeneous nucleation, 250 Additives, 236, 239, 242, 261 Additives germanium and tin, 260 (111)A GaAs, 278 AlGaN, 120, 125 AlGaN layers, 231 Alkali-metal amides, 139 Alkaline, 238 Alkyl gallium, 33 AlN, 120, 161, 165 Al2 O3 , 177, 183 Aluminum (Al), 166 Aluminum nitride, 256 Amide, 163 Ammonia, 125, 187, 190, 199 Ammonia partial pressure, 250, 256, 257, 260 Ammonium chlorogallate, 190 Ammonium halogenide, 139 Ammonium hexafluorogallate .NH4 /3 GaF6 , 190 Ammonoacidic, 162 Ammonobasic, 139, 162 Ammonolysis, 163, 190 Ammonothermal, 22, 138, 188–193, 195, 197–202 growth, 162, 185 method, 278 technique, 185 Annihilation fractions, 306 Annihilation radiation, 296 Annihilation rates, 301 Arsenic (As) desorption, 44, 46

Arsenides and phosphides of Ga and In, 34 As-grown surface morphology, 260 AsHn , 46 AsHn with n D 1–3, 47 Atomic force microscopy (AFM), 43, 99 Auger loss, 14 Autoclave, 162, 171, 172 Autoprotolysis, 190 Average positron lifetime, 302 Axial temperature gradient configurations, 212

Background donor impurities, 53 Baffle, 171, 217, 222 Ba3 N2 , 239 Band-edge absorption, 31 Band-to-band absorption, 157 Basal plane SFD, 106 Basal-plane stacking faults, 100, 102, 171 Basic ammonothermal, 190 Basic mineralizers, 162, 165, 184, 188, 199 Bending, 52, 219 Bending of dislocations, 130 Bi and Sn as solvents, 246 Biaxial strain, 13 Biexciton emission, 155 Bio-sensing, 9 Bismuth, 270 BL, 286 Blue/green/white LED, 161 Blue laser diodes, 161 Blue light emitting diodes, 161 Blue/near-UV laser diodes, 13 Blue/violet optical storage, 19 Blu-ray, 9 Blu-ray DVD, 162 Boiling point Tb of NH3 , 188 Boron, 271 Boron nitride crucible, 271

317

318 Bose–Einstein statistics, 283 Bowing, 184 Bowing-related broadening of XRD rocking curves, 107 Breakdown voltage, 20 Bridgman method, 61, 131, 138, 237, 245 Brillouin zone, 283 Broad luminescence bands in blue and yellow color region, 286 BSFs, 72 Buffer layers, 17, 38, 43, 44, 102 Bulk crystals, 295 Bulk GaN crystallization, 222 Bulk GaN substrates, 10 Bulk semi-insulating GaN substrates, 15 Burgers vector, 101, 267 Burstein–Moss (BM) shift, 287

Carbon (C), 176, 239, 242 concentration, 130 contamination, 130 incorporation, 126 Carrier concentration, 53, 90, 152, 172, 224, 227 Carrier density, 262, 270 Carrier injection, 14 Carrier lifetime, 278 Carrier localization, 15 Carrier mobilities, 15 Carrier recombination efficiency, 20 Catalytic reaction, 125 Cathodoluminescence, 262 Cellular growth, 214 Cellular instabilities, 224 Channel mobilities, 21 Chemico-mechanical polishing (CMP), 139, 195 Chemistry of the solution, 185 .CH3 /2 GaNH2 , 121, 125 .CH3 /2 GaNH2 W NH3 , 125 Œ.CH3 /3 Ga W NH2 x , 125 .CH3 /3 M W NH3 , 123 Cleaved surface, 45 Coalescence of islands, 267 Coincidence detection of the Doppler broadening, 303 Collision frequency, 125 Columnar growth mechanism, 191 Comineralizer, 139 Commercial printing, 9 Compensation of O donors, 56 Compressive biaxial strain, 284 Computational fluid dynamics, 125

Index Conductivity, 270 Contactless electroreflectance (CER), 156 Continuous wave (CW) operation, 10 Convectional flow, 222 Convectional flow velocity, 219 Convection flow, 172, 209, 212, 217, 223 Convection mechanism, 209 Convective transport, 219, 222, 229 Convergent beam electron diffraction (CBED), 171, 215 Coprecipitation, 240 Cracking, 79 Creeping phenomenon, 258 Critical breakdown electric field, 7 CrN, 240 Cross-sectional SEM, 99 Cross-sectional TEM, 171, 174 Crystal defect, 184 Crystal morphology, 239 Cubic GaN, 144, 193 Cubic symmetry, 32 Current leakage, 285 Curvature, 148 CVD, 122, 125 Czochralski method, 61, 138, 245 Dark currents, 16 Dark spot density, 150 Debye phonons, 283 Decomposition of nitrides, 138 Decomposition temperature, 15 Decomposition velocity, 125 Deep acceptor states, 55 Deep-level traps, 51 Defect density, 13, 172, 232 Defect generation, 42 Defect selective etching (DSE), 62, 64, 68, 70–72, 152, 216, 265, 268 Density, 229 of dislocations, 31, 231 of the solution, 260 Density-function theory (DFT), 32, 45 Detrapping, 299 Device efficiency, 14 Device lifetime, 10, 13 Device output power and reliability, 13 Device performance and reliability, 19 Dielectric constant ", 188 Dielectrics, 21 Differential interference contrast (DIC) microscopy, 150 Diffusion, 13 boundary layer, 191 coefficient, 191

Index Dihydride, 46 Diodes, 15 Dislocation density (DD), 10, 13, 22, 23, 53, 106, 119, 130, 150, 171, 201, 208, 216, 218, 232, 245, 265, 268, 271 Dislocation density for InGaN, 15 Dislocations, 13, 79, 171, 175, 196 bending, 278 lines, 13, 262 loops, 215 Dissociation of ammonia, 140, 256 of GaN, 256 of molecular nitrogen, 250 Domain-like structure, 22 Donor acceptor pair (DAP), 262 Donor acceptor transition, 152 Donor dopants, 144 Doping profiles, 13 Doppler broadening, 295 Doppler broadening spectroscopy, 303 Double crystal XRD (DCXRD), 51 Driving force for deposition, 34

Earth alkaline salts, 238 Edge-character TD density, 283 Edge-component dislocations, 102 Edge dislocations, 219, 266 Effective refractive index, 281 Efficacy, 20 Einstein phonons, 283 Electroluminescence, 265 Electron concentration, 172, 216, 227 Electron diffraction pattern, 41 Electron-dispersive X-ray analysis (EDX), 166 Electron irradiation, 309 Electron mobility, 7, 53, 90, 172 Electroreflectance (ER), 279 Energy barrier, 124 Energy consumption, 5 Energy cutoff value, 46 Energy efficiency, 5 Energy of formation HF , 189 Energy of solvation, 188 Enthalpy, 168 Entrapment of inclusions, 191 Entropy, 168 Entropy of solution, 169 Epitaxial growth, 260 Epitaxial growth front, 13 Epitaxial growth rate, 258 Epitaxial lateral overgrowth (ELO), 12, 13, 31, 62, 80, 243, 265

319 Epitaxial techniques (MOCVD and MBE), 208 Equilibrium concentration, 256 Equilibrium partial pressure of ammonia, 250 Equilibrium partial pressures, 32 Equilibrium pressures of N2 , 248 Etched back, 260 Etching, 150 Etch pit density (EPD), 151, 216, 219 Etch pits, 227 Exciton energies, 278 Exciton–polariton, 279

Facet formation, 130 Faceting, 99 Facets, 99 Facetted nucleation islands, 267 FeCl2 , 55 Fe-doped freestanding GaN, 32 Fermi level energy, 216 Ferrocene ŒCp2 Fe W .C5 H5 /2 Fe, 55 FIDAP 8.5, 219 Field effect transistors (FETs), 7, 20 FIELO, 285 Fill factor, 167 Finite element calculation, 219 First-principles pseudopotential approach, 46 Flat heterointerface, 13 Flow dynamic simulation, 122 Flow rate, 119, 127, 128, 191, 192 Flow reactor, 119 Flow speed, 120 Flow velocities, 229 Fluid transport, 170 Flux-film-coated LPE method (FFC-LPE), 237 Flux growth, 23 Focused ion beam (FIB), 174 Foreign substrates, 8 Fourier transformed infrared (FTIR) reflectivity, 270 Free carrier recombination, 287, 288 Free electron concentration, 219 Free exciton resonance, 277 free excitons, 280 Free-standing, 184 crystals, 232 GaN, 79 GaN wafers, 31 HVPE-GaN, 138, 150 wafers, 102 Free-to-bound (FB) transition, 288 FWHM, 216, 219, 224

320 G, 168 Ga, 163, 171 density, 220 inclusions, 215 metal, 33 thermal expansion coefficient, 221 vacancies .VGa /, 216, 242, 265, 284, 296 GaAs, 11, 31, 61, 245 GaAs-based HBTs, 18 GaAs-based laser diodes, 18 GaAs substrate, 8 GaCl, 31, 32, 63, 65, 132 GaCl3 , 32 GaGaNN2 equilibrium, 210 GaN2 GaN equilibrium conditions, 212 Gallium amide, 168 Gallium–germanium solutions, 256 Gallium imide, 167, 168 Gallium oxide Ga2 O3 , 195 Gallium–sodium melt, 246 GaN epitaxial growth, 11 FET, 15 laser diodes (LDs), 9 LED market, 5 MOSFET, 21 platelets, 208, 214 quantum wells, 69 RF FET, 21 Schottky diodes, 15 substrate, 8 synthesis, 210 wings, 224 ŒGaN4 , 126 GaN/AlGaN, 64, 69 GaN/(0001) Al2 O3 template, 278 GaN-based LED, 14 GaN-based solid solution, 161 GaN-based technologies, 61 Ga.NH2 /3 , 140 Ga2 .NH/3 , 140 ŒGa.NH3 /5 Cl2C , 140, 190, 191 Ga3C  3NH 2 complexes, 169 ŒGaN4 molecules, 120 400 GaN wafer, 11 Ga  O  Cl complex, 167 GaOx , 166 Ga-polar surfaces, 64, 73 Gas–liquid interface, 250 Gas source MBE (GSMBE), 37 Gate leakage current, 16 Generation gradient approximation (GGA), 46 Germanium, 259, 270 Germanium containing solutions, 270

Index Gettering of impurities by dislocations, 262 Gibbs free energy of formation, 163 Glassy carbon, 256 Grain boundaries, 183 Graphite baffles, 222 crucible, 211 Gravity vector, 221 Green LEDs, 15 Grown from solutions in liquid gallium at nitrogen gas pressure, 62 Growth, 23, 207 from melt, 193 hillocks, 217 polarity, 304, 308 stability, 239 Growth rate, 22, 23, 37, 39, 42, 119, 121, 122, 125, 127, 130, 190–193, 201, 207, 217, 232, 236, 239, 243, 250, 260, 265, 271 along the c-direction, 267 of GaN, 50

H2 , 32, 64 H , 168 Halide vapor phase epitaxy, 278 Hall effect measurements, 53, 224, 227, 270, 279 HCl, 32, 63, 132, 166 Heat capacity .cp /, 229 Helium, 32 Heteroepitaxial growth, 99 Heteroepitaxially grown GaN layers, 12 Heteroepitaxy, 137, 161 Heteroepitaxy of GaN, 31 Heterostructures, 102 Hexagonal symmetry, 32 High-brightness LED, 16 High current densities, 16 High definition DVD players and recorders, 3 High definition video projectors and displays, 9 High electron mobility transistor (HEMT), 20 High-flow-speed reactor, 127 High-frequency applications, 55 High nitrogen pressure (HNP), 138, 304 High power electronic (HPE), 6 High pressure reactors, 208 High-pressure solution (HPS), 185, 207–209, 211, 216, 222, 232 High pressure solution growth (HPSG), 62, 245 High resolution TEM, 39

Index High resolution thermogravimetry (TG), 248 High temperature solution growth, 207 Hillocks, 108, 224, 261 H2 =N2 ambient, 166 H3 N W .CH3 /3 Ga W NH3 , 123 HNO3 =HCl acid solution, 212 Homoepitaxy, 208 Homogeneous nucleation, 248, 254, 256, 269 Horizontal reactor, 120, 121 Hot-wall quartz reactor, 55 HP chambers, 211 H3 PO4 , 175 HP platelets, 64 HP-GaN, 216–218, 224, 225, 227, 230–232 HP plate-like seeds, 68 HP seeds, 69 HP substrates, 65 HP technique, 208 HVPE seeds, 224, 232 4H-SiC, 217 6H-SiC, 102 H2 SO4 W H2 O2 W H2 O, 38 HVPE-GaN, 217, 219 Hydride vapor phase epitaxy, 13, 22, 31, 61, 79, 98, 130, 131, 138, 161, 163, 168, 175, 180, 183, 185, 187, 189, 195, 197–199, 208, 224, 265, 304 free standing crystal, 219 GaN crystals, 65 seeds, 224, 232 Hydrogen, 103 Hydrogen chloride, 103 Hydrothermal, 165, 185, 190, 191, 200, 201, 278 crystallization, 138 growth, 22, 162 technique, 185

Impurities, 176, 177, 232, 256, 270, 278 Impurity content, 262 2-inches-diameter freestanding GaN, 53 Inert gas, 32 InGaN, 119, 161 LED, 131 quantum wells, 231 InGaN-based solar cells, 16 In0:1 Ga0:9 N quantum wells, 231 Initial dislocation density, 263, 265 InP, 11, 55 In-plain compressive strain, 285 In-plane rotation, 101 In situ gravimetric monitoring method, 48

321 Institute of High Pressure Physics PAS in Warsaw, Poland, 62 Institute of physics PAS, 77 Interface, 197 Intermolecular collision, 121, 123 Intermolecular reaction, 125 Internal quantum efficiency (IQE), 14 Intramolecular reaction, 123 Inversion domains, 108 Inverted hexagonal pyramids, 287 Ionic salts, 168 Ion implantation, 309 Island–island coalescence, 102 Isotherm lines, 229

Joule heating, 13

KCl, 163 KGa.NH2 /4 , 140 KI, 163 Kinetic barrier, 257 Kinetic trapping equations, 305 KNH2 , 139, 188 KOH/NaOH eutectic, 216, 225, 227

LiNH2 , 139, 140 Laminar flow, 122, 220 Laser diodes, 161, 230, 232 Laser lift-off technique, 20 Laser scattering, 119 Lateral epitaxial overgrowth (LEO), 105, 278 Lateral growth rates, 109 Lateral growth velocity, 267 Lattice and thermally matched substrates, 183 Lattice constant, 31 Lattice energy, 188 Lattice imperfections, 51 Lattice mismatch, 15, 79, 99 Lattice tilt, 23 Lattice-matched alloyed substrate, 185 Layer-by-layer growth mode, 15 Lead and bismuth as solvent, 259 LEO GaN, 283 Li additive, 232 LiCl, 163 Lifetime, 231 Lifetime components, 301 Lift-off, 61 Light extraction efficiency, 20 Li3 N C Ga, 239 Light-emitting diodes (LEDs), 4, 7, 119, 161

322

-LiAlO2 , 102 Li, Na, Mg, Al, Ni, Cr, Mn, 175 Li3N, 239 Liquid phase epitaxy (LPE), 201, 208, 216, 217, 219, 224, 232, 246 Liquid solution, 184 Liquid–solid interface, 251 Lithium aluminate, 103 Lithium nitride, 246 Lithium – potassium mixtures, 246 Low angle grain boundaries, 22 Low dislocation density, 265 Low pressure solution growth (LPSG), 23, 218, 246, 265 Low temperature phase of quartz, ’-SiO2 , 184 Low-angle grain boundaries, 147 Lower voids, 85 LPE at sodium flux method, 232 Luminous efficiency, 4

Macroscopic defects, 260, 263 Macrosteps, 214, 218, 261 Macrosteps propagation, 224 Manganine gauges, 211 Mass density, 92 Mass transport coefficient, 37 II–VI materials, 161 Maximum growth rate, 131 Maximum stable growth rate Vmax , 191 Maximum supersaturation of the solution, 256 Maximum velocity of the convectional flow, 222 M.CH3 /3 molecule, 123 Mechanical polishing, 224 Mechanical vibration, 255 Mechano-chemical polishing, 230 Melt growth, 201 Melting point, 188 Mesh, 220 Metal additives, 246 Metallic conductivity, 216 Metal nitride, 81 Metalorganic chemical vapor deposition (MOCVD), 23, 97, 152, 208, 230, 265 Metalorganic hydrogen chloride vapor phase epitaxy (MOHVPE), 33 Metalorganic vapor phase epitaxy (MOVPE), 32, 62, 119, 130, 131, 278 Metastable phase, 199, 236 Metastable zone, 223 Methane, 123 Mg, 15

Index MgAl2 O4 (spinel), 105 Mg-doped, 216 MgO, 37 Microdefects, 215 Microphotoluminescence (-PL), 151 Microprecipitates, 216 Micro-Raman scattering, 64, 70 Micro-Raman thermography, 17 Microstructure, 278 Migration, 120 Miller indices, 148 Mineralizers, 23, 139, 162, 187, 189, 190, 193, 195, 199, 201 Misfit, 264 Misfit dislocations, 41 Mixed-character dislocation, 175 Mixed mineralizer NH4 Cl.1x/ Xx , 190 Mixed-type dislocations, 219 Mn2 N, 240 Mobility, 216, 219 MOCVD GaN layer, 261 Modulation spectroscopy, 156 Molecular beam epitaxy (MBE), 15, 32, 97, 183, 231 Molecular diffusion, 127 Momentum density of the annihilating positron–electron pairs, 296 Monohydride, 46 Morphology of the as-grown GaN layers, 260 Mosaicity, 147 Mott density, 288 m-plane, 102 MQWs, 69 Multi wafer processing, 270 Multiwafer reactor, 128 Multizone cylindrical graphite furnaces, 211

NaCl, 163 Na-flux method, 138 Na-flux technique, 246 Na2 Ga.NH2 /5 , 140 NaGa.NH2 /4 , 163 NaI, 163 NaN3 , 246 NaNH2 , 140, 163, 165, 166, 170, 174, 188 Nano-net, 81 Nanoparticles, 121, 122, 125 Native bulk GaN substrate, 119 Native point defects, 216 NdGaO3 , 31 Nax –.GaNy /, 236 N-doped GaN, 262 Near-equilibrium growth conditions, 23

Index Negative ions, 296 Negative temperature coefficient, 142 Neutral donor bound exciton .D0 X/, 280 Neutral donor bound exciton recombination, 151 Neutral mineralizers, 162 Newtonian fluid, 220 NH, 46 NH2 , 46 NH 2 , 139, 162, 168, 193 NH3 , 31, 32, 46, 64, 121, 122, 131, 161, 188–191, 248 NHC 4 , 190 NH4 Br, 193 NH4 Cl, 139, 190, 193 NH4 Cl; NH4 I, 163 NH3 flow rate, 122, 131 .NH4 /3 GaCln , 190 .NH4 /3 GaF6 , 190 NH4 I, 140, 193 NHn , 46 Ni-based superalloy, 286 Ni  Cr based alloy, 163 Ni  Cr based superalloys, 163, 165, 171 NiTi layer, 57 Nitride-based laser diodes, 13 Nitrogen, 32 Nomarski micrograph, 175 Nonpolar, 162, 197, 214, 224, 227, 232, 236, 240 and semipolar GaN, 20 GaN, 21 planes, 97 substrates, 143 Nonradiative lifetime, 284 Nonradiative recombination, 13, 262 Nonradiative recombination at dislocations in InGaN, 15 Nonradiative recombination centers (NRCs), 284 N radicals, 236 n-type conductivity, 53 crystals, 72 GaN, 63 Nucleation rate, 249 Nucleation rate of GaN, 260 N vacancies, 307

Ohmic contact, 230 Oligomer formation, 125 ON , 242 On-resistance, 16

323 Open-volume defect, 295 Optical reflectance (OR), 279 Optical storage, 13 Optoelectronic device, 109 Organometallics, 122, 123, 127, 128 Oscillator strength, 20 Ostwald–Miers region, 223, 251, 255, 256 Overgrown regions (“wings”), 109 Oxygen, 176 Oxynitrides, 255

Parasitic capacitance, 285 Parasitic crystallites, 260 Parasitic GaN, 242, 257, 260 Parasitic growth rates, 258 Parasitic nucleation, 239, 240 Parasitic reaction, 120, 121, 125, 127 Partial pressures, 57 Passivation effect, 256 Patterned sapphire substrates (PSS), 12 pBN, 242 PEC, 64, 67 Pendeo- or cantilever-epitaxy, 12 Pentaaminechlorogallium (III) dichloride ŒGa.NH3 /5 ClCl2 , 190 Phonon replica, 152 Phosphors, 14, 20 Photolithographic processing, 109 Photoluminescence (PL), 51, 195, 199 Photoluminescence spectra, 216, 279 Photo-modulated electroreflectance, 279 Photoreflectance (PR), 156, 279 Photovoltaics, 16 Piezoelectric and spontaneous polarization, 20 Piezoelectric polarization, 14 Pinhole, 44 Planetary reactor, 128 Plasma-assisted MBE (PA-MBE), 69, 208 Plasma-enhanced chemical vapor deposition, 109 p–n junctions, 13 Point defects, 177, 265 Polariton dispersion, 277 Polariton lifetime, 281 Polarity, 42, 190, 215 Polar solvent, 184 Polar substrates, 143 Polytypes 4H, 113 6H, 113 Porous interlayer, 80 Positive solubility coefficient, 188

324 Positron annihilation, 216 measurements, 279 spectroscopy, 295 Positron beam, 297 Positron binding energy, 306 Positron escape rate, 306 Positron lifetime, 295 Positron lifetime spectroscopy, 301 Positron source, 301 Positron state, 297 Positron trapping, 299 coefficient, 299 rate, 299 Potassium amide, 139 Potential barrier, 211 Potential energy change, 47 Power output, 14 Precursor, 46, 187–190, 199 Pressure in the liquid metal, 221 p-type GaN, 270 PtRh6%–PtRh30% thermocouples, 211 p  T curve, 211 Pyrolytic boron nitride, 242, 256

Quadrupole Mass Spectrometer spectroscopy, 124 Quantum cascade laser (QCL), 16 Quantum-confined Stark effect (QCSE), 20 Quantum mechanical calculations, 210 Quantum well based LDs, 13 Quantum well intersubband transition based devices, 16 Quantum wells, 102 Quartz .’-SiO2 /, 23, 184, 278 Quartz glass, 256 ampoule, 271 production, 162

Radiative recombination, 151 Raman spectroscopy, 227 Rare earth elements, 139 Reaction kinetics, 251 Reaction rate, 250 Reactive ion etching, 224 Read-only LDs, 10 Recombinations of A- and B-free exciton, 284 Red-shifted emission, 20 Reduction of dislocation density, 263, 266 Resistance, 7 Resistive GaN:Mg, 63 Resistivity, 56, 90

Index Retrograde solubility, 141, 165, 167, 170, 171, 184, 188 Reverse-bias leakage current, 16 Reverse breakdown, 16 Reverse recover time, 16 Reynolds number, 219, 227 RF FETs, 16, 18 RF transistors, 15 RGB, 20 Room temperature thermal conductivity (RT-TC), 17 Root-mean-square (RMS) roughness, 99, 287 Roughness of surface, 261 r-plane sapphire substrates, 99 Rydberg states, 299

S, 168 Sapphire, 10, 61, 119, 130, 161, 163, 183, 217 Sapphire/GaN MOCVD templates, 217, 218, 227 Sapphire/GaN template, 219 Sapphire substrates, 31, 246, 261 Scalability, 143, 184, 201 Scanning electron microscopy (SEM), 45, 166, 171, 172 Schmidt number, 191 SC H2 O, 186, 192 Schottky diodes, 7 SC NH3 , 183 Screw dislocations, 16 Secondary ion mass spectrometry (SIMS), 56, 67, 175, 219, 262 Seeded crystallization, 213 Seeded growth, 216 Segregation, 231 Self-heating, 16 Self-separation, 81 Semi- and nonpolar, 201 Semi-insulating (SI), 53, 144, 270 Semipolar, 97, 162, 177, 214, 232 Seraphin coefficients, 282 Shallow hydrogenic states, 296 Sheet resistance, 15, 130 Shower head reactor, 120 Si, 31, 61 Si and O, 144 SiC, 37, 137, 162, 217, 219 Sidewall LEO (SLEO), 111 Silicon, 7, 11, 245 Silicon carbide, 161 Silicon nitride, 256 Single rotational domain, 108 SiO2 , 109

Index ’-SiO2 , 191 SiO2 layer, 52 SiO2 mask, 53 Small-volume autoclave, 166 Sodium, 139 Sodium flux LPE, 266 Solar cell photoresponse, 16 Solid solution, 184 Solid-state lighting (SSL), 4, 162 Solid-state power control devices, 6 Solubility, 22, 23, 141, 162, 165, 167, 170, 184, 185, 188, 189, 191, 201, 210, 236, 238, 239, 251, 252, 278 of AlN, 189 of nitrogen, 23, 245, 271 Solute, 22, 162, 184, 191 Solution, 23, 184, 185 Solution density, 260 Solution of Ga and Bi, 246 Solvent, 172, 184, 246 Solvothermal growth method, 22, 162, 165, 184, 185, 278 S parameter, 303 Spatially-resolved cathodoluminescence (CL), 278 Specific heat, 92 Specific heat capacity, 229 Specific heat capacity of gallium, 221 Spectroscopy, 9 Spontaneous crystallization, 214 Spontaneous growth, 232 Spontaneous nucleation, 224 Stabilization energy of complex formation, 123 Stable growth on a seed, 256 Stacking fault, 183 Stacking fault density (SFD), 104, 115 Standard free energy of solution, 168 Steady-state PL, 285 Stoney equation, 65 Strain, 184 Streamlines, 222 Stress relaxation, 41 Subgrains, 22 Sublimation method, 42 Supercell, 46 Supercooling, 214, 217 Supercritical (SC), 184 ammonia, 22, 138, 141, 278 ammonobasic, 165, 167 ammonobasic solution, 174 fluid, 165 NH3 , 162, 170 water, 22

325 Supersaturation, 165, 184, 191, 207, 209, 210, 214, 217, 224, 230, 250, 255, 271 Supersaturation concentration, 256 Surface kinetic factor, 218 Surface morphology, 50, 172 Surface passivation, 258 Surface polarities, 44 Surface roughness, 102

Terahertz applications, 16 Terraces, 218, 261 Terraces with macrosteps, 264 Testing and measurement applications, 9 Thermal agglomeration, 82 Thermal and lattice mismatch, 184 Thermal annealing, 41, 310 Thermal boundary layer, 119, 121 Thermal conductivity, 17, 91, 222 of device, 13 of gallium, 221 and viscosity of gallium, 229 Thermal conductivity coefficient, 221, 229 Thermal conductivity coefficient for graphite, 221 Thermal diffusivity, 92 Thermal dissociation, 258 Thermal escape, 299 Thermal etching, 254 Thermal expansion coefficients, 32, 79, 217 Thermal gradient, 120 Thermalized positron, 296 Thermally conductive substrate, 17 Thermal stresses, 80 Thermodynamical potential, 210 Thermodynamic equilibrium, 184, 245, 265 Thermogravimetry, 251 Thin layers, 295 Threading dislocation (TD), 15, 100, 101, 115, 183, 197, 199, 262, 267, 277 Threading dislocation density (TDD), 104, 108, 175, 230 Three-layer gas injection, 128 Three-layer gas-injection reactor, 127 Threshold voltage, 231 Time-resolved photoluminescence (TRPL), 279 TiN, 81, 278 Tin, 259 TMA, 121, 124 Total energy calculation, 46 Total lumen output, 20 Transition metal additives, 240

326

Index

Transmission electron microscopy (TEM), 64, 69, 71, 72, 100, 171, 174, 197, 216, 262 Trapping of carriers, 15 Tri-ethyl-gallium (TEG), 33 Trihydride, 46 Tri-methyl-aluminum, 121 Tri-methyl-gallium (TMG), 33, 121, 125, 128, 131 Twinning, 107 Two step growth technique, 119

Wafer bowing, 278, 285 Wafers, 184 Wet etching, 109 Wet etching in NaOH/KOH eutectic melt, 266 Wetting angle of GaN, 260 White light LEDs, 14 Wide bandgap, 7 Wireless communications, 16 W parameter, 303 Wurtzite GaN, 162 Wurtzite structure, 32, 214

Unstable growth, 215 Upper voids, 85 UV-blind avalanche photodetectors, 16

X-band, 15 X-ray, 163 measurements, 224, 227 photoelectron spectroscopy, 215 rocking curve, 171, 175, 180, 195, 198 X-ray diffraction (XRD), 64, 69, 100, 199, 216

Vacancy defects, 295, 296 Vapor phase, 119 Vapor phase epitaxy (VPE), 32, 138, 163 Vapor phase growth, 102 Variable-energy positron beam, 308 Varshni fitting, 283 V-defect formation, 15 V defects, 105 Velocity vector, 220 Vertical gradient freeze, 245 Vertical power FET, 20 VGa , 215, 286 VGa  O, 286 VGa  ON complexes, 307 V/III ratio, 99, 130, 132 Violet laser diodes, 53 Void-assisted separation (VAS), 79 Voids, 80, 115, 171, 175

Yellow band, 139 Yellow emission, 51 Yellow luminescence, 152, 216, 219 YL, 286

zb/wz, 163 Zero bias off-state, 20 Zinc (Zn) acceptors, 242 Zincblende, 162, 163 Zincblende structure, 32 Zincblende-type (cubic) GaN, 139 Zinc oxide (ZnO), 23, 185, 191 ZnO buffer layer, 31