Chemically Deposited Nanocrystalline Metal Oxide Thin Films: Synthesis, Characterizations, and Applications 303068461X, 9783030684617

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Fabian I. Ezema Chandrakant D. Lokhande Rajan Jose  Editors

Chemically Deposited Nanocrystalline Metal Oxide Thin Films Synthesis, Characterizations, and Applications

Chemically Deposited Nanocrystalline Metal Oxide Thin Films

Fabian I. Ezema  •  Chandrakant D. Lokhande  Rajan Jose Editors

Chemically Deposited Nanocrystalline Metal Oxide Thin Films Synthesis, Characterizations, and Applications

Editors Fabian I. Ezema Department of Physics & Astronomy University of Nigeria Nsukka, Nigeria

Chandrakant D. Lokhande D. Y. Patil Education Society (Deemed to be University) Kolhapur, India

Rajan Jose Department of Industrial Sciences & Technology Universiti Malaysia Pahang Gambang, Pahang, Malaysia

ISBN 978-3-030-68461-7    ISBN 978-3-030-68462-4 (eBook) https://doi.org/10.1007/978-3-030-68462-4 © The Editor(s) (if applicable) and The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 This work is subject to copyright. All rights are solely and exclusively licensed by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors, and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, expressed or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Switzerland AG The registered company address is: Gewerbestrasse 11, 6330 Cham, Switzerland

Foreword

The twenty-first century would observe a rapid technological furtherance to address challenges posed on resources and life sustainability due to the growth of a new billion population in the 10–15 years’ time. The fast-growing economy has dramatically increased the consumption of energy and other resources and often calls for alternative resources and solutions to sustain the growing population. Many countries have defined “sustainability” as the new normal owing to the doubling of waste in a decade, rapid depletion of traditional resources, and serious environmental pollution in the way we have been using the resources. Petroleum resources are depleting at a faster rate than ever, and the Earth is expected to run out of this resource soon. The pressing demand to introduce alternative energy sources to fossil fuels and minimize CO2 emissions generates considerable recent research interest in the development of renewables along with energy storage systems. Besides, it is necessary to avoid petroleum resources because of their adverse environmental impacts, such as water and air pollution, leading to global warming. Efficient and cost-effective technologies associated with energy efficiency, conversion, and storage are needed to solve the energy problems ravaging the world. The critical requirements for good energy system are affordability, availability as well as high energy and power densities; the recent advances in materials technology using the tools of nanotechnology gear metal oxides with appropriate property modification and value addition are one of the solutions. Holistic usage of metal oxides is a welcome development in nanotechnology because they belong to classes of functional materials; their practical distribution, natural abundance, and low cost enhance opportunities for application in diverse areas for the betterment of humankind.

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This book should be a welcome addition to the much-needed analytical studies on the application of metal oxides in photovoltaics, energy storage, and sensing, employing nanotechnological tools. I believe that the authors and editors of this book are responding positively to globalization and environmental concerns through promoting friendly energy sources, storage mechanisms, and sensing with essential approaches to appreciate the diversity and complexity of nanotechnology. University of Nigeria  Charles Arinzechukwu Igwe Nsukka, Nigeria

Preface

Metal oxides represent one of the important classes of functional materials in view of their abundance for their practical deployment and opportunities for value addition for continuous quality/performance improvement. The science and engineering literature on metal oxides is enormous with hundreds of thousands of research papers published, which can be retrieved through popular science and engineering databases such as Scopus and Web of Science. Nevertheless, new materials, processing techniques, structure–property and processing–property correlations are continuously reported; therefore, incessant efforts are required for consolidating these new knowledges. An increasing number of reviews are published continuously in journals; however, their diversity and scatter often defeat the purpose of having them in one place. This book is an attempt to consolidate important works synthesis, characterizations, and applications of chemically deposited nanocrystalline metal oxide thin films for various applications. We had our initial discussion on this topic during the International Conference on Advancements in Renewable Energy–2020 (ICARE 2020) during the second week of January 2020 at Swami Ramananda Teertha Marathwada University at Nanded, India. All the editors of the present book were invited speakers in the conference. We observed that such a book on electrochemical metal oxides for energy conversion, storage, and sensing could help the community in finding the relevant resources in one place. Consequently, book has been conceived to provide input on (1) chemically processed metal oxide films for energy conversion (photovoltaic), (2) chemically processed metal oxide films for energy storage (supercapacitors and batteries), (3) chemically processed metal oxide films for sensing, and (4) modeling and simulation. This book has 31 chapters, which are divided into five major parts. The first part has four chapters and is devoted to providing general information on the progress in solution-processed metal oxides, properties and applications of the electrochemically synthesized metal oxide thin films, structural and electronic properties of the various metal oxides, and the insights from first principle calculations on the properties of metal oxides. The second part contains six chapters, which elaborately reviews the photovoltaic applications of metal oxides. Reviews on the recent vii

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p­ rogress of the topic in general as well as the properties of metal oxides in favor of the photovoltaic action are discussed in the second part. Excellent reviews on the optically active metal oxides for photovoltaic applications and their doped analogues for dye-sensitized solar cells and perovskite solar cells are included in this part. The third part is devoted to photoelectrochemical cells, hydrogen production, and photocatalytic application of metal oxides elaborated in five chapters. The largest part in this book is that of energy storage detailed in the fourth part in 11 chapters. Various aspects of two intensively researched energy storage devices, viz. lithium ion batteries and supercapacitors, in addition to the general aspects of energy storage are detailed. Finally, the last part details the progress in electrochemical sensing applications of metal oxides in four chapters. We do hope that the classification and titling of these chapters are self-explanatory; and therefore, details of the chapters are not summarized here. We would like to thank the authors who have contributed to this book, which would not have been realized without their time and efforts. We do hope that the time and the efforts you have dedicated in preparing the chapters would be helpful in widening your visibility as a Science & Engineering research professional. The authors thank Professor Rajaram S. Mane of Swami Ramananda Teertha Marathwada University for bringing us on a common platform (ICARE2020), which made this book possible. It has been a pleasure working with the editorial team of Springer Nature, Charles Glaser, Michael Luby, and especially our project coordinator Ms. Olivia Ramya Chitranjan. Her timely help in proofreading and suggestions helped the book to be presented in a better way—thank you Olivia. Besides, as Editors of this book, we enjoyed working together on this project. And finally, you the reader, thanks in advance and we do hope that you have benefitted reading this book. Nsukka, Nigeria  Fabian I. Ezema Kolhapur, India   Chandrakant D. Lokhande Gambang, Pahang, Malaysia   Rajan Jose

Contents

1 Progress in Solution-Processed Mixed Oxides��������������������������������������    1 Swati N. Pusawale 2 Properties and Applications of the Electrochemically Synthesized Metal Oxide Thin Films ����������������������������������������������������   29 Abdellah Henni and Amina Karar 3 Structural and Electronic Properties of Various Useful Metal Oxides��������������������������������������������������������������������������������   49 Saima G. Sayyed, Annis A. Shaikh, Pankaj K. Bhujbal, Arif V. Shaikh, Habib M. Pathan, and Prafulla Kumar Jha 4 Properties of Metal Oxides: Insights from First Principles Calculations���������������������������������������������������������   85 Assa Aravindh Sasikala Devi and D. Murali 5 Recent Progress in Metal Oxide for Photovoltaic Application������������   99 Emmanuel O. Onah, Jude N. Udeh, Sabastine Ezugwu, Assumpta C. Nwanya, and Fabian I. Ezema 6 Structural and Electronic Properties of Metal Oxides and Their Applications in Solar Cells����������������������������������������������������  147 Agnes Chinecherem Nkele, Sabastine Ezugwu, Mutsumi Suguyima, and Fabian I. Ezema 7 Optically Active Metal Oxides for Photovoltaic Applications��������������  165 A. C. Lokhande, V. C. Lokhande, D. S. Dhawale, I. A. Qattan, Shashikant Patole, and Chandrakant D. Lokhande 8 Metal Oxides for Perovskite Solar Cells������������������������������������������������  197 V. C. Lokhande, C. H. Kim, A. C. Lokhande, Chandrakant D. Lokhande, and T. Ji

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9 Doped Metal Oxide Thin Films for Dye-­Sensitized Solar Cell and Other Non-Dye-­Loaded Photoelectrochemical (PEC) Solar Cell Applications����������������������������������������������������������������������������  235 M. D. Tyona 10 Doped Metal Oxide Thin Films for Enhanced Solar Energy Applications����������������������������������������������������������������������  261 Calister N. Eze, Raphael M. Obodo, Sabastine Ezugwu, and Fabian I. Ezema 11 Mixed Transition Metal Oxides for Photoelectrochemical Hydrogen Production������������������������������������������������������������������������������  279 Camillus Sunday Obayi and Paul Sunday Nnamchi 12 Plasmonic Metal Nanoparticles Decorated ZnO Nanostructures for Photoelectrochemical (PEC) Applications ��������������������������������������  293 Mangesh A. Desai and Shrikrishna D. Sartale 13 Oxygen-Deficient Metal Oxide Nanostructures for Photocatalytic Activities��������������������������������������������������������������������  329 Rahul B. Pujari and Dong-Weon Lee 14 Oxygen-Deficient Iron Oxide Nanostructures for Photocatalytic Activities��������������������������������������������������������������������  355 Sanjana S. Bhosale and Arpita P. Tiwari 15 Properties of Titanium Dioxide-Based Nanostructures on Transparent Glass Substrates for Water Splitting and Photocatalytic Application ��������������������������������������������������������������  389 Crispin Munyelele Mbulanga, Chinedu Christian Ahia, and Johannes Reinhardt Botha 16 Mixed Transition Metal Oxides for Energy Applications��������������������  405 Ravindra N. Bulakhe, Anuradha B. Bhalerao, and Insik In 17 Nanosheet-Derived Porous Materials and Coatings for Energy Storage Applications������������������������������������������������������������  431 Shirin P. Kulkarni, Akash S. Patil, Vinod V. Patil, Umakant M. Patil, and Jayavant L. Gunjakar 18 Role of Carbon Derivatives in Enhancing Metal Oxide Performances as Electrodes for Energy Storage Devices ��������������������  469 Raphael M. Obodo, Assumpta C. Nwanya, Innocent S. Ike, Ishaq Ahmad, and Fabian I. Ezema 19 Hydrothermal Synthesis of Metal Oxide Composite Cathode Materials for High Energy Application����������������������������������  489 Moses Kigozi, Blessing N. Ezealigo, Azikiwe Peter Onwualu, and Nelson Y. Dzade

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20 Metal Oxide Composite Cathode Material for High Energy Density Batteries ��������������������������������������������������������  509 Jin Kiong Ling and Rajan Jose 21 Chemically Processed Transition Metal Oxides for Post-Lithium-Ion Battery Applications��������������������������������������������  531 Amol Bhairuba Ikhe 22 Nanostructured Metal Oxide-Based Electrode Materials for Ultracapacitors����������������������������������������������������������������������������������  561 Chukwujekwu Augustine Okaro, Onyeka Stanislaus Okwundu, Philips Chidubem Tagbo, Cyril Oluchukwu Ugwuoke, Sabastine Ezugwu, and Fabian I. Ezema 23 Nanoporous Metal Oxides for Supercapacitor Applications����������������  601 Ved Prakash Joshi, Nitish Kumar, and Rahul R. Salunkhe 24 Nanoporous Transition Metal Oxide-Based Electrodes for Supercapacitor Application��������������������������������������������������������������  623 U. M. Patil, V. V. Patil, A. S. Patil, S. J. Marje, J. L. Gunjakar, and C. D. Lokhande 25 Hybrid Nanocomposite Metal Oxide Materials for Supercapacitor Application��������������������������������������������������������������  673 Vaishak Sunil and Rajan Jose 26 Liquid Phase Deposition of Nanostructured Materials for Supercapacitor Applications ������������������������������������������������������������  725 Shreelekha N. Khatavkar and Shrikrishna D. Sartale 27 Chemically Processed Metal Oxides for Sensing Application: Heterojunction Room Temperature LPG Sensor���������������������������������  765 Bidhan Pandit and Babasaheb R. Sankapal 28 Chemically Synthesized Novel Materials for Gas-Sensing Applications Based on Metal Oxide Nanostructure������������������������������  807 David C. Iwueke, Raphael M. Obodo, Chinedu Iroegbu, Ishaq Ahmad, and Fabian I. Ezema 29 Low-Temperature Processed Metal Oxides and Ion-Exchanging Surfaces as pH Sensor ����������������������������������������������������������������������������  821 Cyril Oluchukwu Ugwuoke, Philips Chidubem Tagbo, Onyeka Stanislaus Okwundu, Chukwujekwu Augustine Okaro, Sabastine Ezugwu, and Fabian I. Ezema

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30 Performance Evaluation of P-Type Semiconducting Metal Oxide-Based Gas Sensors ������������������������������������������������������������  863 Raphael M. Obodo, Sylvester M. Mbam, Ishaq Ahmad, and Fabian I. Ezema 31 Development of InSb Nanostructures on GaSb Substrate by Metal-Organic Chemical Vapour Deposition: Design Considerations and Characterization����������������������������������������  879 Chinedu Christian Ahia, Crispin Munyelele Mbulanga, Edson L. Meyer, and Johannes Reinhardt Botha Index������������������������������������������������������������������������������������������������������������������  903

About the Editors and Contributors

About the Editors Fabian  I.  Ezema  is a Professor at the University of Nigeria, Nsukka. He obtained a PhD and MSc in Physics and Astronomy from the University of Nigeria, Nsukka, and a BSc from the then Anambra State University of Science and Technology, Enugu. His researches are on several areas of Materials Science: synthesis and characterizations of particles and thin film materials through chemical routes with emphasis on energy applications. He is interested in materials for thin film solar cells fabrication and nanoparticle synthesis. For the last 15 years, he has been working on the nano/submicron-sized materials for energy conversion and storage (cathodes, anodes, supercapacitors, thin film solar cells, DSSC, etc.), including novel methods of synthesis, characterization, and evaluation of the electrochemical and optical properties. He was a CV Raman Fellow at the Shivaji University, Kolhapur, India, in 2011 and MIF Fellow at Tokyo University of Science, Japan, in 2013. He is a visiting professor to NRSEMT (MATECSS UNESCO Chair) in Varennes, Quebec, Canada, and iThemba Labs, South Africa. He is also a Fellow of UNESCO-­UNISA South Africa Chair in Nanosciences and Nanotechnology (U2ACN2). He was awarded University of South Africa (UNISA) visiting researcher program (VRP). Fabian is a senate representative to National Center for Energy Research and Development (NCERD), University of Nigeria, and due to his commitment to research and international visibility was appointed Pioneer Acting Director, International Office of the University of Nigeria after the approval of the office by the University Senate. He has served as Coordinator, Natural Science Unit, School of General Studies, UNN.  He is the Coordinator of Nano Research Group, University of Nigeria, Nsukka. He is the pioneer Dean, Faculty of Natural and Applied Sciences, Coal City University Enugu, Nigeria. Fabian has published over 180 high impact papers xiii

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in various national and international journals and gave over 50 talks at various conferences, workshops, and seminars. His h-index is 24, i10 index of 88 with over 2409 citations as in google scholar. Chandrakant  D.  Lokhande  is presently working as a Dean and Research Director at D Y Patil Education Society (Deemed to be University), Kolhapur, India. He has been working on several areas of thin film technology, ranging from chemical synthesis of thin films to their applications in solar cells, gas sensors, and supercapacitors. Moreover, he made a great contribution in designing several prototype devices such as supercapacitors and heterojunction-based room temperature gas sensors. He received his PhD from Shivaji University, Kolhapur, in 1984, without viva voce examination as his thesis was adjudged as “Excellent.” Later, in 1987, he joined as assistant professor in Physics and became professor and Head at Shivaji University, Kolhapur, immediately after accomplishing his first postdoctoral stay at the Weizmann Institute of Science, Israel. He has won many awards and received many honors. He was appointed as Fellow of Institute of Physics, London, in 1990; was visiting scientist in the Indo-Polish CEP scheme in 1991; was INSA Visiting Fellow in 1993; is the first recipient from Shivaji University of the prestigious Alexander von Humboldt Fellowship, Germany, in 1996 and Brain Pool fellowship of South Korea in 2003; was participant in Noble Laureates Meeting, Lindau, Germany, in 2001; was visiting professor at Hanyang University, South Korea, in 2006; was awarded a Rajya Shishak Purshakar, Government of Maharashtra State in 2009, and Best Teacher Award from Shivaji University in 2010. He is presently an editorial board member of Electrochemical Energy Technology, De Gruyter; the fellow, Maharashtra Academy of Sciences from 2012; an expert member, and a distinguished visiting professor in polymer chemistry, Institute of Chemical Technology, Mumbai, from 2012. He is the author of more than 600 papers in international journals with “h” index 90 and more than 29,000 citations, edited 11 books, filed more than 45 patents, and directed more than 60 PhD theses. Recently, he has been listed at first position in top 2% scientists in the subject of Applied Physics in India by a Stanford University Survey. Rajan Jose  supervises the Nanostructured Renewable Energy Materials Laboratory in the Universiti Malaysia Pahang (UMP) and is the Associate Editor-in-Chief of the Springer Nature journal Materials Circular Economy. He has investigated nanostructured perovskite ceramics for microwave and superconducting electronics during his doctoral research at the Council of Scientific and Industrial Research (CSIR), Trivandrum, India, and has received PhD degree in the year 2002. He has contributed to the science and engineering of diverse

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range of materials including metals and alloys, luminescent quantum dots for biological and energy applications, glass and glass ceramics for quantum electronics, and electrochemical materials for energy conversion and storage. He was employed as a scientist at the Indira Gandhi Centre for Atomic Research (India), AIST (Japan), Toyota Technological Institute (Japan), and the National University of Singapore (Singapore) before joining UMP.  He has published over 220 papers in Web of Science (Thomson Reuters) indexed journals which are cited nearly 11,000 times with an h-index of 55. He holds 25 patents nationally and internationally. He has been listed as the top 2% Materials Scientists by the Stanford University. He has supervised 6 postdoctoral, 23 doctoral, and 10 master’s researchers. His current research interests include circular economy, data science, renewable materials and devices; most of his research is on the structure–property relationship in materials for a desired device functionality.

Contributors

Chinedu Christian Ahia  Institute of Technology, University of Fort Hare, Alice, South Africa Ishaq Ahmad  National Center for Physics, Islamabad, Pakistan NPU-NCP Joint International Research Center on Advanced Nanomaterials and Defects Engineering, Northwestern Polytechnical University, Xi’an, China Nanosciences African Network (NANOAFNET), iThemba LABS-National Research Foundation, Somerset West, Western Cape Province, South Africa UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa (UNISA), Pretoria, South Africa Assa Aravindh Sasikala Devi  Nano and Molecular Systems Research Unit, Oulu, Finland Anuradha  B.  Bhalerao  Applied Science Department, K.  K. Wagh Institute of Engineering Education and Research, Nasik, India Sanjana S. Bhosale  College of Chemical Engineering, Fuzhou University, Fuzhou, China Pankaj  K.  Bhujbal  Department of Physics, Savitribai Phule Pune University, Pune, India Johannes Reinhardt Botha  Department of Physics, Nelson Mandela University, Port Elizabeth, South Africa Ravindra  N.  Bulakhe  Department of Polymer Science and Engineering, Korea National University of Transportation, Chungju, South Korea Mangesh  A.  Desai  Thin Films and Nanomaterials Laboratory, Department of Physics, Savitribai Phule Pune University, Pune, India D. S. Dhawale  Centre for Interdisciplinary Research, D. Y. Patil Education Society (Deemed to be University), Kolhapur, India Nelson Y. Dzade  School of Chemistry, Cardiff University, Cardiff, UK

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Contributors

O.  Onah  Emmanuel  Crystal Growth Laboratory, Physics and Astronomy Department, University of Nigeria, Nsukka, Nigeria Calister N. Eze  Department of Physics, Federal University of Technology, Minna, Niger State, Nigeria Blessing  N.  Ezealigo  Dipartimento di Ingegeneria Meccanica, Chimica e dei Materiali, Universita Degli Studi di Cagliari, Cagliari, Italy Fabian  I.  Ezema  Crystal Growth Laboratory, Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu State, Nigeria Nanosciences African Network (NANOAFNET), iThemba LABS-National Research Foundation, Somerset West, Western Cape Province, South Africa UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa (UNISA), Pretoria, South Africa Department of Electrical Engineering, Faculty of Science and Technology, Tokyo University of Science, Yamazaki, Noda, Japan Science and Engineering Unit, Nigerian Young Researchers Academy, Onitsha, Anambra State, Nigeria Sabastine Ezugwu  Department of Physics and Astronomy, University of Western Ontario, London, ON, Canada Jayavant L. Gunjakar  Centre for Interdisciplinary Research, D. Y. Patil Education Society (Deemed to be University), Kolhapur, India Abdellah  Henni  Laboratory, Dynamic Interactions and Reactivity of Systems, Kasdi Merbah University, Ouargla, Algeria Innocent  S.  Ike  African Centre of Excellence in Future Energies and Electrochemical Systems (ACE-FUELS), Federal University of Technology, Owerri, Nigeria Department of Chemical Engineering, Federal University of Technology, Owerri, Nigeria Amol Bhairuba Ikhe  Department of Printed Electronics Engineering, Suncheon National University, Chonnam, Republic of Korea Insik  In  Department of Polymer Science and Engineering, Korea National University of Transportation, Chungju, South Korea Department of IT Convergence (Brain Korea PLUS 21), Korea National University of Transportation, Chungju, South Korea Chinedu  Iroegbu  Department of Physics, Federal University of Technology, Owerri, Imo State, Nigeria David  C.  Iwueke  Department of Physics, Federal University of Technology, Owerri, Imo State, Nigeria

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Prafulla Kumar Jha  Faculty of Science, The Maharaja Sayajiraos University of Baroda, Vadodara, India T.  Ji  Department of Electronics and Computer Engineering, Chonnam National University, Gwangju, South Korea Rajan  Jose  Nanostructured Renewable Energy Materials Laboratory, Faculty of Industrial Sciences & Technology, Universiti Malaysia Pahang, Kuantan, Pahang, Malaysia Ved Prakash Joshi  Materials Research Laboratory, Department of Physics, Indian Institute of Technology Jammu, Jammu, J&K, India N.  Udeh  Jude  Crystal Growth Laboratory, Physics and Astronomy Department, University of Nigeria, Nsukka, Nigeria Amina Karar  Laboratory, Dynamic Interactions and Reactivity of Systems, Kasdi Merbah University, Ouargla, Algeria Shreelekha N. Khatavkar  Thin Films and Nanomaterials Laboratory, Department of Physics, Savitribai Phule Pune University, Pune, India Moses  Kigozi  Department of Materials Science and Engineering, African University of Science and Technology, Abuja, Nigeria Department of Chemistry, Faculty of Science and Education, Busitema University, Tororo, Uganda C.  H.  Kim  Department of Electronics and Computer Engineering, Chonnam National University, Gwangju, South Korea Shirin  P.  Kulkarni  Centre for Interdisciplinary Research, D.  Y. Patil Education Society (Deemed to be University), Kolhapur, India Nitish  Kumar  Materials Research Laboratory, Department of Physics, Indian Institute of Technology Jammu, Jammu, J&K, India Dong-Weon Lee  MEMS and Nanotechnology Laboratory, School of Mechanical System Engineering, Chonnam National University, Gwangju, Republic of Korea Jin Kiong Ling  Nanostructured Renewable Energy Material Laboratory, Faculty of Industrial Sciences & Technology, Universiti Malaysia Pahang, Pahang, Malaysia A. C. Lokhande  Applied Quantum Materials Laboratory (AQML), Department of Physics, Khalifa University of Science and Technology, Abu Dhabi, United Arab Emirates Chandrakant  D.  Lokhande  Centre for Interdisciplinary Research, D.  Y. Patil Education Society (Deemed to be University), Kolhapur, India V. C. Lokhande  Department of Electronics and Computer Engineering, Chonnam National University, Gwangju, South Korea

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Contributors

S.  J.  Marje  Centre for Interdisciplinary Research, D.Y.  Patil Education Society (Deemed to be University), Kasaba Bawada, Kolhapur, Maharashtra, India Sylvester M. Mbam  Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu State, Nigeria Department of Physics, Allama Iqbal Open University, Islamabad, Pakistan Crispin Munyelele Mbulanga  Department of Physics, Nelson Mandela University, Port Elizabeth, South Africa Edson  L.  Meyer  Institute of Technology, University of Fort Hare, Alice, South Africa D. Murali  Faculty of Physics, IIITDM, Kurnool, India Agnes Chinecherem Nkele  Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu, Nigeria Paul Sunday Nnamchi  Department of Metallurgical and Materials Engineering, University of Nigeria, Nsukka, Nigeria Assumpta  C.  Nwanya  Crystal Growth Laboratory, Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu State, Nigeria Nanosciences African Network (NANOAFNET), iThemba LABS-National Research Foundation, Somerset West, Western Cape Province, South Africa UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa (UNISA), Pretoria, South Africa Camillus Sunday Obayi  Department of Metallurgical and Materials Engineering, University of Nigeria, Nsukka, Nigeria Raphael M. Obodo  Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu State, Nigeria National Center for Physics, Islamabad, Pakistan NPU-NCP Joint International Research Center on Advanced Nanomaterials and Defects Engineering, Northwestern Polytechnical University, Xi’an, China Chukwujekwu Augustine Okaro  Science and Engineering Unit, Nigerian Young Researchers Academy, Onitsha, Anambra State, Nigeria Onyeka  Stanislaus  Okwundu  Science and Engineering Unit, Nigerian Young Researchers Academy, Onitsha, Anambra State, Nigeria Azikiwe  Peter  Onwualu  Department of Materials Science and Engineering, African University of Science and Technology, Abuja, Nigeria Bidhan  Pandit  Institut Charles Gerhardt Montpellier (ICGM), Université de Montpellier, Montpellier, France

Contributors

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Habib M. Pathan  Department of Physics, Savitribai Phule Pune University, Pune, India Akash S. Patil  Centre for Interdisciplinary Research, D. Y. Patil Education Society (Deemed to be University), Kolhapur, India Umakant  M.  Patil  Centre for Interdisciplinary Research, D.  Y. Patil Education Society (Deemed to be University), Kolhapur, India Vinod V. Patil  Centre for Interdisciplinary Research, D. Y. Patil Education Society (Deemed to be University), Kolhapur, India Shashikant Patole  Applied Quantum Materials Laboratory (AQML), Department of Physics, Khalifa University of Science and Technology, Abu Dhabi, United Arab Emirates Rahul B. Pujari  MEMS and Nanotechnology Laboratory, School of Mechanical System Engineering, Chonnam National University, Gwangju, Republic of Korea Swati  N.  Pusawale  Department of Sciences and Humanities, Rajarambapu Institute of Technology, Rajaramnagar, Maharashtra, India I.  A.  Qattan  Applied Quantum Materials Laboratory (AQML), Department of Physics, Khalifa University of Science and Technology, Abu Dhabi, United Arab Emirates Rahul R. Salunkhe  Materials Research Laboratory, Department of Physics, Indian Institute of Technology Jammu, Jammu, J&K, India Babasaheb R. Sankapal  Department of Physics, Visvesvaraya National Institute of Technology, Nagpur, Maharashtra, India Shrikrishna D. Sartale  Thin Films and Nanomaterials Laboratory, Department of Physics, Savitribai Phule Pune University, Pune, India Saima  G.  Sayyed  Department of Electronic Science, L.V.H Arts, Science and Commerce College Panchavati, Nashik, India Annis A. Shaikh  Department of Physics, Savitribai Phule Pune University, Pune, India Arif  V.  Shaikh  Department of Electronic Science, L.V.H Arts, Science and Commerce College Panchavati, Nashik, India Suryadatta Education Foundation’s, Suryadatta International of Cyber Security (SIICS), Pune, India Mutsumi  Suguyima  Department of Electrical Engineering, Faculty of Science and Technology, Tokyo University of Science, Yamazaki, Noda, Japan Vaishak  Sunil  Nanostructured Renewable Energy Materials Laboratory, Faculty of Industrial Sciences & Technology, Universiti Malaysia Pahang, Kuantan, Pahang, Malaysia

xxii

Contributors

Philips  Chidubem  Tagbo  Science and Engineering Unit, Nigerian Young Researchers Academy, Onitsha, Anambra State, Nigeria Arpita  P.  Tiwari  D.  Y. Patil Education Society (Institution Deemed to be University), Kolhapur, Maharashtra, India M.  D.  Tyona  Department of Physics, Benue State University, Makurdi, Benue State, Nigeria Cyril  Oluchukwu  Ugwuoke  Science and Engineering Unit, Nigerian Young Researchers Academy, Onitsha, Anambra State, Nigeria Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu State, Nigeria

Chapter 1

Progress in Solution-Processed Mixed Oxides Swati N. Pusawale

Abbreviations CBD CV EDS EGFET ESR FTIR FTO HRTEM ITO NCM SAED SCE SEM SILAR SS TEM XPS XRD ZNACO

Chemical bath deposition Cyclic voltammetry Energy dispersive X-ray spectroscopy Extended-gate field effect transistor Electron spin resonance Fourier transform infrared spectroscopy Fluorine-doped tin oxide High-resolution transmission electron microscopy Indium-doped tin oxide Nickel-cobalt-manganese Selected area electron diffraction Saturated calomel electrode Scanning electron microscope Successive ionic layer adsorption and reaction Stainless steel Transmission electron microscope X-ray photoelectron spectroscopy X-ray diffraction Zinc-nickel-aluminium-cobalt oxide

S. N. Pusawale (*) Department of Sciences and Humanities, Rajarambapu Institute of Technology, Rajaramnagar, Maharashtra, India © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_1

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1.1  Introduction Metal oxides are having a wide range of applications in different fields such as gas sensors, supercapacitors, catalysis, thin-film transistor, etc., due to their excellent electrical, physical, chemical, and optical properties. Mixed oxides are the materials having two or types of metal cations, and depending on the metal cations, they are classified as binary, ternary, and quaternary oxides. The mixed oxides are materials with two or more than two metal ions linked with oxygen in proportions that can vary or defined with firm stoichiometry [1]. When oxides are mixed, its results change in the electronic structure of the whole material. This includes the modifications in the bulk and surface properties, also properties like bulk electronic structure, bandgap energy, Fermi energy level, transport, etc. [2]. Mixing different metal oxides is a useful way to obtain specific materials with required properties, as the interaction of materials in this system occurs at the atomic level, due to which they have the advantage of chemical properties that can be totally different and promising when compared with single oxide material [1]. The mixing can lead to impressive catalytic and thermal stability and can generate new surface-active sites to participate in the reaction [3]. Mixed oxides are gaining considerable attention due to their excellent properties, for various applications such as catalysis, supercapacitor, etc. [4–10]. Synthesis of metal oxide can be possible using either physical or chemical methods. But compared to physical methods, solution-processed methods are more useful, as they offer control over the morphology with improved performance. Some of the popular solution-processed methods are hydrothermal, solvothermal, electrodeposition, chemical bath deposition, electrophoretic deposition, sol-gel, precipitation, successive ionic layer adsorption and reaction (SILAR), etc. Solution-processed methods have the potential advantages of (1) control over the stoichiometry of the material, (2) low-temperature deposition, (3) formation of homogenous materials, (4) formation of complex shapes, (5) useful for composite material synthesis, and (6) low cost of equipment [11]. Every solution-processed method has its own advantages and disadvantages, but they offer a synthesis of different types of materials easily so one needs to select among the different methods, which one will be beneficial for synthesis of required material, depending upon desired properties and application. In this chapter, we have described the synthesis approach for mixed oxides by different solution-­ processed methods. Among the referred literature, it was observed that the sol-gel method was majority used for deposition mixed oxides, but apart from this, mixed oxide formation using other chemical methods such as electrodeposition, chemical bath deposition, SILAR, precipitation, etc. are also described in this chapter. At the last, the literature review is given in Table 1.1, which depicts the variation in the morphology of the mixed oxide with the deposition method and its application.

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Table 1.1  Review of mixed oxides Sr. No. Mixed oxide 1. Manganese-­ nickel 2. Na3BiO4-­ Bi2O3 3. Ce-Co

Synthesis method Morphology Electrodeposition Iceland-like grains

Application Supercapacitor

Reference [14]

Electrodeposition Vertically aligned nanoplates Electrodeposition Nanoparticles

pH sensor

[15]

Counter electrode in electrochromic devices Supercapacitor

[16]

[17]

Supercapacitor

[18]

Charge storage devices Supercapacitor Supercapacitor

[20]

Supercapacitor

[23]

–

[25]

Photocatalysis

[25]

4.

Mn-Co

5.

7. 8.

MnO2-NiO on Sb-doped SnO2 Manganese-­ vanadium SnO2-RuO2 Cu-Co

9.

Co-Cu

SILAR

10.

CuO-NiO

Precipitation

11.

ZnO-SnO2

Precipitation

12.

Co3O4–ZnO

Precipitation

13.

ZnO–NiO

Precipitation

14.

ZnO-Fe2O3

Sol-gel

15. 16. 17.

WO3-TiO2 Cr2O3–TiO2 IrO2-TiO2

Sol-gel Sol-gel Sol-gel

18.

Sn-Zn

Sol-gel

19.

TiO2-MgO

Sol-gel

6.

Electrodeposition Hexagonal nanoplates Electrodeposition Nanowires

Electrodeposition Uniform and dense SILAR SILAR

Fibrous and porous Spherical grains interspersed with nanoporous agglomerates of short nanorods Short spherical granules embedded between nanorod agglomerates Agglomerated spherical particles Agglomerated spherical particles Agglomerated nanoparticles Agglomerated nanoparticles Aggregates of crystals – – Agglomerated particles Irregular sized nanoparticles Dumbell shape particles

[22] [23]

Probable in sensing, [27] optoelectronics, and photocatalysis Catalyst [28] Photocatalysis

[31]

Photocatalysis Photocatalysis –

[32] [33] [34]

DSSC

[35]

Probable in photocatalysis and DSSC

[37]

(continued)

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S. N. Pusawale

Table 1.1 (continued) Sr. No. Mixed oxide 20. TiO2-AgO

21.

Ni-Co-Mn

22.

Zn-Ni-Al-Co

Synthesis method Morphology Sol-gel Interconnected particles forming porous structure CBD Nanosheets with porous structure CBD Nanosheets

Application Gas sensor

Reference [38]

Supercapacitor

[42]

Supercapacitor

[43]

1.2  S  olution-Processed Methods for Synthesis of Mixed Oxide 1.2.1  Electrodeposition Electrodeposition is an easy and cost-effective method for the deposition of material in thin-film form. Typically, the synthesis involves two metal electrodes dipped into an electrolyte solution; the deposition of the material occurred on the cathode, which is also called as working electrode after applying the external field across the electrodes. The main condition for this method is that the deposition is possible only with conducting substrates. A wide variety of films can be obtained by selecting different parameters of electrodeposition. Preparation of thin films using the electrodeposition method has some advantages such as: 1. The method is useful for the deposition of structurally and compositionally modulated alloys and compounds. 2. Low-temperature deposition method. 3. Useful for deposition of material even on complex shapes. 4. Higher deposition rates. 5. More control over the deposition process as reactions involved in the method occur close to the equilibrium. The method is useful for the deposition of a variety of materials such as metals, oxides, magnetic materials, supercapacitive materials, and chalcogenides [12, 13]. Some of the mixed oxides deposited by the electrodeposition method are described below: Tahmasebi et al. used potentiodynamic anodic electrodeposition for the synthesis of thin films of manganese-nickel oxide. The effect of deposition scan rate and Ni: Mn molar ratio in the bath on the capacitive properties of mixed oxide was studied. The films showed the effect of the deposition scan rate on its topography and morphology. At 50 or 200 mV s−1 scan rate, the large island-like grains were observed in topography, and at the scan rate of 600 mV s−1 changes in the surface, topography was observed with size reduction of surface features and formation of particle-like morphology. The CV curves of the pure MnO2 and MnO2 with 10 and 17 wt% Ni in the 1 M Na2SO4 electrolyte in the potential range from 0 to 1 V vs. ­Ag/

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Fig. 1.1  The CV curves of MnO2 (MO), MnO2 with 10 wt% Ni (MNO2), and MnO2 with 17 wt% Ni (MNO10). (Reprinted with permission from [14], copyright 2016 Elsevier)

AgCl electrode at the scan rate of 20  mV  s−1 is shown in Fig.  1.1. The specific capacitance of the films increased with an increase in deposition scan rate from 50 to 600 mV s−1, but further increase in deposition scan rate to 800 mV s−1 decreased the specific capacitance value. This decrease in the capacitance was attributed to the decrease of the Ni content in the mixed oxide at the deposition scan rate of 800 mV s−1. The increase in Ni: Mn molar ratio in the bath increased the Ni content in the mixed oxide initially rapidly and then slowly after Ni: Mn ratio reached 4. The Ni content in the mixed oxide affects the morphology and porosity of the mixed oxide. The mixed oxide of MnO2 with 10 wt% Ni showed the high specific capacitance of 250  F  g−1 at 100  mV  s−1 with 122% capacitance retention for 10,000 cycles [14]. Na3BiO4-Bi2O3 films were deposited at room temperature on the ITO substrate. The XRD pattern confirmed the polycrystalline nature of the films. Diffraction peaks corresponding to Na3BiO4 and Bi2O3 were observed with monoclinic structure observed for both oxides. Figure  1.2a shows the SEM image of the sample revealing the formation of vertically aligned nanoplates, the average thickness of 90 nm was observed from the high magnification image shown in Fig. 1.2b. The semiquantitative concentration analysis confirmed 20% of Ni3BO4 and 80% of Bi2O3. The mixed oxide electrodes were utilized as pH sensors as sensing film for EGFET. A sensitivity of 49.63 mV/pH has been observed with good linearity [15]. Cathodic electrodeposition was carried out for the deposition of Ce-Co mixed oxide films on ITO substrates. The electrodeposited films for equimolar ratio showed no diffraction peaks confirming the formation of amorphous phase; however, the formation of nanocrystalline CeO2 was confirmed from the TEM studies. Further, the deposition was carried out for different temperatures as 30, 50, and 70 °C for an equimolar ratio of Ce: Co. The effect of deposition temperature on the morphology of the mixed film was observed and is shown in Fig.  1.3; with an

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Fig. 1.2  SEM image of Na3BiO4-Bi2O3 mixed oxide nanostructure (a) at bar scale 50μm, (b) showing average thickness of individual nanowall (bar scale 3μm). (Reprinted with permission from [15], copyright 2020 Elsevier)

Fig. 1.3  TEM images of Ce-Co mixed oxide with equimolar ratio at different deposition temperatures (a) 30, (b) 50, and (c) 70 °C. (in the inset corresponding SAED pattern). (Reprinted with permission from [16], copyright 2003, Elsevier)

increase in deposition temperature, the crystallinity of the mixed oxide increased (inset of Fig. 1.3) and nanoparticles grew larger. The IR studies showed the presence of peak corresponding to Ce-O-Co vibration in the films deposited from the bath temperature of 30 °C. The CV studies reflected larger current densities in anodic and cathodic regions for the films with an equimolar ratio of Ce: Co and deposited at a bath temperature of 30  °C.  The films showed higher transmittance of about 0.9 in the visible region, making it useful as a transparent counter electrode in electrochromic devices [16]. Mn-Co mixed oxide electrode materials were synthesized on SS substrate for different volume ratios as 50:50, 60:40, and 70:30. The XRD studies confirmed the poor crystallinity of the mixed oxide with the presence of weak peaks for MnO2 and Co3O4. The SEM studies revealed the presence of hexagonal plates interconnected

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with each other forming flower-like morphology as shown in Fig. 1.4a–i. The effect of the volume ratio on the thickness of the hexagonal-plated was observed with the least thickness of 22 nm for a 60:40 volume ratio as observed in Fig 1.4g–l. The films were studied for electrochemical properties. The CV curves for the Mn-Co oxide between the potential range of −0.8 to 0.6 V vs. SCE in 1 M NaOH for different volume ratios are shown in the Fig. 1.5a–c, and Fig. 1.5d shows the combined CVs of the Mn-Co mixed oxide for all the volume ratio at the scan rate of 5 mV s−1. The presence of the redox peaks in all CVs confirmed the pseudocapacitive behavior of the electrode, and the peaks were present even at a higher scan rate of 100 mV s−1 showed the better rate capability. In Fig. 1.5d, the A1, A2 and C1, C2 represent anodic and cathodic peaks, respectively. The film with 60:40 volume ratio of Mn:Co showed excellent supercapacitive response with specific capacitance and energy of 679 g−1 and 97 Wh kg−1, respectively [17]. Jiang et al. [18] coated Sb-doped SnO2 with electrodeposited MnO2-NiO mixed oxide layer. SnO2 doped with Sb deposited on the FTO substrate via the vapor-­ liquid-­solid (VLS) technique. The nanowires of SnO2 were grown on the substrate surface, which did not affect even with the change in the doping concentration of Sb, but it affects the electrical resistance of the nanowires. SnO2 nanowires doped

Fig. 1.4  SEM image (a, b, c) of 50:50, (d, e, f) of 60:40, (g, h, i) of 70:30 volume ratio, and cross-­ sectional SEM of (j) 50:50, (k) 60:40, (l) 70:30 of Mn-Co mixed oxide. (Reprinted with permission [17], copyright 2018, Elsevier)

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Fig. 1.5  CV curves of Mn-Co mixed oxide at different scan rates (5, 10, 20, 50, and 100 mV s−1) for volume ration (a) 50:50, (b) 60:40, (c) 70:30, and (d) CV curves of Mn-Co mixed oxide for all volume ratios at the scan rate of 5  mV  s−1. (Reprinted with permission [17], copyright 2018, Elsevier)

Fig. 1.6 (a) SEM image of mixed MnO2-NiO on Sb-doped SnO2 and (b) EDS analysis of mixed oxide. (Reprinted with permission from [18], copyright 2020, Elsevier)

with 5% Sb were utilized for deposition of MnO2-NiO mixed oxide. The SEM image of the mixed oxide and EDS analysis is shown in Fig. 1.6a, b. The diameter of SnO2 nanowires was increased from 110 to 250  nm after it was coated with MnO2-NiO mixed oxide. The presence of Mn and Ni confirmed from EDS analysis as shown in Fig. 1.6b. The addition of NiO in MnO2 reduced the reduction of Mn4+ in an acidic electrolyte, which is necessary for the faradaic reactions to occur. Figure 1.7 shows the CV curves of MnO2-NiO mixed oxide with various NiO content within the potential range of 0–1 V vs. Ag/AgCl in 0.5 M Na2SO4 electrolyte at

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Fig. 1.7  The CV curves of MnO2-NiO mixed oxide with various NiO content. (Reprinted with permission from [18], copyright 2020, Elsevier)

the scan rate of 10 mV s−1. The enhancement in the current response after the incorporation of NiO to MnO2 can be easily observed from the curves with the maximum current response for the mixed oxide with 20% of NiO. The mixed oxide films with 20 wt% of NiO showed a specific capacitance of 304 F g−1, which is 110% larger compared with pure MnO2 (146 F g−1). Thin films of manganese-vanadium mixed oxide were deposited on a platinum substrate and thereafter heat-treated at various temperatures from 25 to 400  °C under vacuum by Nakayama and group. From ESR and XPS studies, it was observed that only in the presence of vanadium, the reductive formation of Mn2+ occurs at 300 °C. This was due to the electron transfer from vanadium ions (V4+) to neighboring Mn sites. The voltammetric response of the deposited mixed oxide in borate solution increased with increasing the number of potentials, and the steady-state current was larger than that of pure manganese oxide [19, 20]. The morphology of the mixed oxide transferred from dense to the porous structure after cycling, and it showed better charge-discharge performance useful in charge storage devices.

1.2.2  Successive Ionic Layer Adsorption and Reaction (SILAR) Nicolau discovered the SILAR method in 1985 [21]. SILAR is a very simple and inexpensive method and useful to deposit a wide variety of materials. The deposition occurs directly on the substrate surface in thin-film form. The method is mainly based on the adsorption and reaction of the ions from the solutions and rinsing between every immersion with water to avoid homogeneous precipitation in the solution. The first step in the method is adsorption, where metal ions get adsorbed

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on the substrate surface immersed in the solution; in the second step, the substrate then immersed in a water solution to remove loosely bounded ions. The third step is a reaction, where substrate with some adsorbed cations immersed in an anionic solution, where the reaction occurs and a solid substance formed on the surface of the substrate, after that again the substrate will be immersed in the water solution, where loosely adsorbed material gets removed from the substrate surface. This mechanism of adsorption, reactions, and rinsing in water is repeated multiple times to achieve the appropriate film thickness. SnO2-RuO2 mixed films were deposited by this method by Pusawale and group. The SnO2-RuO2 mixed oxide was prepared by varying the deposition cycles of SnO2 and RuO2 using SILAR method. The SnO2-RuO2 mixed thin films were deposited on SS substrates. The schematic deposition process involves five beakers system shown in Fig. 1.8. It was observed that after incorporation of RuO2 in SnO2, there is a change in the structure and morphology. The films with major SnO2 content were nanocrystalline in nature, however, with the increase in RuO2 in the deposition nature shifted to amorphous form. The morphology also shifted to dense and compact from porous

Fig. 1.8  Schematic of SnO2-RuO2 film deposition using SILAR, a) immersion of substrate in SnCl2 solution, b) immersion of substrate in 1% H2O2 solution, c) rinsing the substrate in the double distilled water, d) immersion of substrate in the RuCl3 solution, and e) immersion of the hot water bath  ref. [22]

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Fig. 1.9  The charge-discharge curves of SnO2-RuO2 mixed oxide for different cycle ratios of SnO2: RUO2 as (a) 3:2, (b) 1:1, (c) 1:3, and (d) 2:3 at the current density of 1 mA cm−2 ref. [22]

morphology after RuO2 deposition cycles were increased. The charge-discharge curves of the mixed oxide with varying deposition cycle ratios of SnO2-RuO2 within the potential range from −0.2 to 0.6 V vs. SCE, in 0.5 M H2SO4 at the scan rate of 5  mV  s−1, are shown in Fig.  1.9. The mixed oxide electrode with 1:3 deposition cycle ratio showed good charge-discharge behavior compared with other cycle ratios. An increase in the specific capacitance value of SnO2 from 4 to 180 F g−1 was observed after combining with RuO2 [22]. Cu-Co and Co-Cu mixed oxide films deposited on ITO and SS substrates by Nwanya and group. The XRD studies confirmed the presence of CuO and Co3O4 peaks. Figures 1.10 and 1.11 show the SEM images of Cu-Co and Co-Cu mixed oxides at different magnifications. Cu-Co mixed oxide film showed spherical grains interspersed with nanoporous agglomerates of short nanorods, whereas Co-Cu mixed oxide film showed the formation of nanoporous short spherical granules from the nucleation center embedded between nanorod agglomerates on the surface. The CV curves of Cu-Co and Co-Cu oxide on ITO and SS substrates in the potential range from −0.2 to +1.2 V vs. Ag/AgCl in 0.5 M Na2SO4 electrolyte are shown in Fig. 1.11. In both cases, the mixed oxide deposited on the ITO substrate showed a better current response than the SS substrate. The Co-Cu mixed oxide showed less specific capacitance of about 385  F  g−1 compared with Cu-Co mixed oxide. The Cu-Co mixed oxide showed a specific capacitance of 919  F  g−1 with a specific energy of 28.78 Wh kg−1 and specific power of 51.8 Wh kg−1. The synergistic effects between the two oxides resulted in better rate capability and durability [23].

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Fig. 1.10  SEM of the deposited films on ITO substrate (a and b) Cu-Co and (c and d) Co-Cu mixed oxide at two different magnifications of 50 and 100 K. (Reprinted with permission from [23], copyright 2017, Elsevier)

Fig. 1.11  CV curves of (a) Cu-Co mixed oxide and (b) Co-Cu mixed oxide on ITO and SS substrates. (Reprinted with permission from [23], copyright 2017, Elsevier)

1.2.3  Precipitation Method Chemical precipitation is a simple method for the deposition of large-scale powder samples. “In this method, precipitation occurs when the concentration of one solid is above the solubility limit in the host solid, due to rapid quenching or ion implantation, and the temperature is high enough that diffusion can lead to segregation into precipitates” [24]. Typically, this method is useful to synthesize large-scale

1  Progress in Solution-Processed Mixed Oxides

13

materials either in macro or nano size. This method occurs at low temperatures; however, as it occurs at a very fast rate, control over morphology is not possible. Some of the mixed oxides deposited by chemical precipitation are described below: CuO-NiO mixed oxide systems were prepared by this method. The precursor solutions used were 0.2 M copper and nickel nitrates and 1.0 M oxalic acid. The mixed oxide was prepared with varying ratios of NiO from 1 to 50  wt%, with a calcination temperature of 400 and 600 °C. It was observed that due to the addition of NiO in CuO, the crystallinity of oxide increased, resulting in increase in the particle size. With the addition of NiO in pure CuO, the surface area decreased continuously. The synergistic effects between NiO and CuO decreased the surface excess oxygen of CuO [25]. Nanocrystalline ZnO-SnO2 mixed metal oxide powder was prepared using zinc acetate, stannic chloride, and ammonia as precursor materials by Chenari and group [26]. The white precipitates formed dried at 100  °C for 3  h and then calcined at 500 °C. The reaction mechanism for the formation of ZnO-SnO2 is proposed as follows:

SnCl 4 ·5H 2 O + 4 NH 4 OH → Sn ( OH )4 ↓ +4 NH 4 ( ac )



Zn ( CH 3 COO )2 ·2H 2 O + 2 NH 4 OH → Zn ( OH )2 ↓ +2 NH 4 ( ac )

(1.1)



(1.2)



After calcination of precipitated powder at 500 °C, the formation of ZnO-SnO2 mixed oxides occurred as:

Zn ( OH )2 + Sn ( OH )4 → ZnOSnO 4 + 4H 2 O ( g ) ↑

(1.3)



The XRD studies confirmed the formation of hexagonal and tetragonal structure for ZnO and SnO2, respectively. The average crystallite size was found to be 100.6 nm for ZnO and 43.1 nm for SnO2. SEM studies showed formation of agglomerated spherical shape particles. It was observed that the photocatalytic activity of mix ZnO-SnO2 was higher and stable for photo corrosion than pure ZnO. Co3O4-ZnO mixed oxide prepared by surfactant-free mechanism in less time by Sharma and Ghose [27]. The prepared powder was calcined at 350 and 500 °C in air and then used for further studies. The proposed reaction mechanism is described as follows: 6CoC2 O 4 ⋅ 2H 2 O ( aq ) + 6 Zn ( CH 3 COO )2 ⋅ 2H 2 O ( aq. ) + 24 NH 4 OH ↓ 80° C, 2 h 3Zn ( OH )2 + 3Zn ( CH 3 COO )2 ( NH 3 )2 + 3Co ( OH )2 + CO3 O 4

+ 6 ( NH 2 )2 C2 O 4 + 6CH 3 COONH 4 + 32H 2 O + H 2 ( g )

(1.4)

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After calcination,

Zn ( CH 3 COO )2 ) ( NH 3 )2 + [ O ] → ZnO + C2 H 6 ( g ) + 2CO2 ( g ) + 2NH 3 ( g ) (1.5)



Zn ( OH )2 → ZnO + H 2 O



3Co ( OH )2 → Co3 O 4 + 2H 2 O + H 2 ( g )

(1.6)



(1.7)



As observed from the XRD studies, the as-prepared powders showed traces of hydroxides of Zn and Co. Peaks corresponding to the cubic spinel structure of Co3O4 and hexagonal ZnO were observed only for calcined powders. The oxide nanoparticles calcined at 350 °C showed a higher specific area of 72.4 m2 g−1. The mixed oxide showed weak ferromagnetic behavior at low temperatures with a low coercive field of 175 Oe and a remanent magnetization of 0.0068 emu g−1. ZnO-NiO mixed metal oxide powder was prepared using a zinc and nickel acetate precursor. The reaction mechanism for the formation of mixed oxide is described as follows: Zn ( CH 3 COO )2 ⋅ 2H 2 O + Ni ( CH 3 COO )2 ⋅ 4H 2 O + 4 NH 4 OH

→ Zn ( OH )2 + Ni ( OH )2 + 4CH 3 COONH 4 + 6H 2 O Zn ( OH )2 → ZnO + H 2 O



(1.8) (1.9)



After Calcination

ZnO + Ni ( OH )2 → ZnO + NiO + H 2 O



(1.10)

The XRD patterns confirmed the formation of nanocrystalline ZnO-NiO mixed metal oxide after calcination at 350 °C. The effect of calcination temperature on the crystallite size of the oxide was observed. The crystallite size of ZnO was increased from 6.8 to 27 nm and for NiO, it was increased from 8.2 to 11.2 nm when calcination temperature was increased from 350 to 500 °C. Figure 1.12a shows the TEM and (b) the SAED pattern of ZnO-NiO mixed oxide, confirming its polycrystalline nature and particle size in the range of 8.5–10.3 nm. The mixed oxide showed ferromagnetic behavior at low temperatures [28].

1.2.4  Sol-Gel The sol-gel process is a popular method in solution-processed methods for the preparation of colloidal nanoparticles. In this method, first formation of sol occurs, which is nothing but an agglomeration of microparticles in solution under well-­ controlled environment, after this the prepared sol channels with each other to form

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Fig. 1.12 (a) TEM image and (b) corresponding SAED pattern of ZnO-NiO mixed oxide. (Reprinted with permission from [28], copyright 2016, Elsevier)

gel [24]. The method requires precursors such as metal ion precursors like alkoxides and alkoxysilanes. The commonly used precursors are tetramethoxysilane (TMOS) and tetraethoxysilanes (TEOS). The method is useful in the manufacturing ceramic nanomaterials, hydrophobic coatings, optical and optical and refractory ceramic fiber production, nanoscale powders, and injectable nanocomposites such as plasminogen activator entrapment in alumina [29]. The method is also proved as best for the preparation of excellent quality metal oxide nanoparticles as well as metal oxide composites. This method has excellent control over the texture and surface properties of the materials. “The method can be described in five key steps: hydrolysis, polycondensation, aging, drying, and thermal decomposition” [30]. The flowchart of different steps in sol-gel method is shown in Fig. 1.13. Some of the mixed oxides deposited by the sol-gel method are described below: The ZnO-Fe2O3 mixed oxide was prepared by using inorganic salts as precursors at pH 7 and 9. The precursor solutions involved 0.09 M zinc acetate and 1 wt% of ferric chloride hexahydrate and ammonia. The prepared powder was calcined at 350 °C for 3 h. The SEM images of the sample at different pH are shown in Fig. 1.9. Clearly one can see the difference in the morphology of mixed oxide powder prepared at different pH. The different hydrolysis rate of FeCl3 at pH 7 and 9 resulted in lowering the bandgap of ZnO from 2.84 to 2.66 eV. The powder obtained at neutral pH showed secondary particles with of linear and irregular shape having particle size more than 100 nm as shown in Fig. 1.14a. Whereas for the powder obtained at pH 9 showed the average diameter size of 100 nm with spherical particles forming compact agglomeration as confirmed from the SEM image as shown in Fig. 1.14b. In terms of the photocatalytic activity, the mixed oxide prepared with neutral pH

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Fig. 1.13  Flowchart of different processes in sol-gel method

Fig. 1.14  SEM images of calcined ZnO-Fe2O3 (a) at pH 7 and (b) at pH 9. (Reprinted with permission from [31], copyright 2007, Springer Nature)

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showed better performance than the mixed oxide prepared at pH 9 with a rate constant of 9.7 × 10−3 and a time of 71 min [31]. WO3–TiO2 mixed oxides were synthesized using an inorganic salt Na2WO4 and an organic alkoxide W(OC2H5)6 for tungsten precursor, with different W/Ti ratios. From XRD studies, for it was observed that the pure TiO2 was in the rutile phase with a 10% anatase phase, however, with the incorporation of tungsten in 1–5 mol% range, only anatase phase was observed even at high temperatures. As the wt% of the tungsten in TiO2 increased, the surface area also increased from 6.7 m2 g−1 (pure TiO2) to 52.9 m2 g−1 (5 mol%). The increase in photocatalytic activity was observed after WO3 combined with TiO2 [32]. Cr2O3–TiO2 mixed oxides with different compositions were prepared and studied for their photocatalytic properties by Jung and group [33]. In the XRD studies, it was observed that the TiO2 transformation from the anatase to rutile phase at a higher temperature in mixed oxide is dependent on Cr content. For Cr content of 10 mol%, anatase phase was maintained up to 600 °C, which means Cr acts as a structure stabilizer. In UV-visible spectroscopy, it was observed that the visible light absorption increased with the increase of Cr content. The graph of photocurrent response under visible light versus Cr content is shown in Fig. 1.15. The maximum photocurrent was observed with mixed oxide having 10 mol% Cr, however, with an increase in Cr content decrease in photocurrent, hence photocatalytic activity was observed. When large amounts of Cr were added, phase segregation was observed from XRD analysis, which resulted in a decrease of photocatalytic activity by trapping charge carriers at the boundaries. The IrO2–TiO2 mixed oxides were prepared by acid-catalyzed hydrolysis of an iridium solution by Osman and group. For the preparation of IrO2–TiO2 mixed oxide, the precursor involved were iridium alkoxides generated from Na2IrCl6

Fig. 1.15  Graph of photocurrent under visible light illumination as function of Cr content in Cr2O3-TiO2 mixed oxide after calcined at 600 °C. (Reprinted with permission from [33], copyright 2011, Elsevier)

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dissolved in ethanol and mixed with Ti(OEt)4. These were hydrolyzed by a water-­ ethanol solution containing an acid catalyst (HNO3) with a ratio of H2O:M:H+  =  5:1:0.1. It was observed that the gelation becomes difficult with an increase in Ir content. The dried gels contained Ir-rich nanoparticles dispersed in an amorphous TiO2. These nanoparticles may not be completely hydrolyzed to IrO2. After calcination, the amorphous phases crystallize to crystalline anatase [34]. Saeidi and group for dye-sensitized solar cells [35] prepared mixed tin and zinc oxide nanoparticles. The purpose of combining TiO2 with SnO2 was to increase the stability of TiO2 for UV illumination. SnO2-ZnO sample with five different weight ratios such as 1:0, 2:1, 1:1, 1:2, and 0:1 was prepared. The XRD studies confirmed the formation of tetragonal SnO2 and hexagonal ZnO with an average crystallite size of 20 and 25 nm, respectively. From optical studies, it was observed that the bandgap of ZnO increased from 3.19 to 3.66 eV due to an increase in SnO2 content. The SnO2-ZnO mixed oxide prepared with a 1:2 molar ratio exhibited the best electron lifetime and charge transport dynamics for DSSC. The TiO2-MgO mixed oxide nanoparticles with varying MgO content were prepared by Bayal and Jeevanandam. The preparation method for MgO was described elsewhere [36]. The XRD studies confirmed that in the presence of small Mg in mixed oxide, the transformation of anatase to rutile phase was not possible even at a high temperature about 900 °C. Figure 1.16 shows the SEM image of the sample calcined at 700 °C, which confirmed the presence of dumbell-shaped particles for Ti:Mg molar ratio of 1:2 and 1:1, the other images showed the spherical and agglomerated particles for Ti:Mg molar ratio of 1:0.5, 1:0.25, and 1:0.1. The presence of two different particles confirmed the presence of rutile and anatase phase together. The SEM images of the oxide calcined at 900 °C are shown in Fig.  1.17, which show the formation of large, irregular, and agglomerated particles mostly for all molar ratios. A blue shift in the bandgap was observed for TiO2 with the incorporation of MgO.  The presence of MgO in TiO2-MgO nanoparticles suppressed the crystal growth of TiO2, leading to an increase in the bandgap value, which was attributed to the formation of MgTiO3 [37]. Nanocrystalline TiO2-AgO mixed oxide in thin film and powder form was prepared at a low temperature of 573 K. From FTIR studies, it was observed that the phase composition of the mixed oxide depends upon the annealing temperature. The crystallite size of about 4 nm was observed for the mixed oxide annealed at 773 K and increased to 10  nm after annealed at 973  K.  The films were prepared using powders of mixed oxide under optimized conditions and utilized for gas sensing applications. The mixed oxide showed a fast, stable, and reproducible response toward carbon monoxide gas at the low-operating temperature of 200  °C.  The response magnitude of 34.6 and response time of 30 s were achieved for this sensor toward 400 ppm of carbon monooxide [38].

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Fig. 1.16  FE-SEM images of TiO2-MgO nanoparticles at (a) 1:2, (b) 1:1, (c) 1:0.5, (d) 1:0.25, and (e) 1:0.1 Ti:Mg molar ratio calcined at 900 °C. (Reprinted with permission from [37], copyright 2014, Elsevier)

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Fig. 1.17  FE-SEM images of TiO2-MgO nanoparticles at (a) 1:2, (b) 1:1, (c) 1:0.5, (d) 1:0.25, and (e) 1:0.1 Ti:Mg molar ratio calcined at 900 0C . (Reprinted with permission from [37], copyright 2014, Elsevier)

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1.2.5  Chemical Bath Deposition (CBD) Nagayama et al. first used the chemical bath deposition (CBD) method while preparing SiO2 films on silicon wafers [39]. The CBD method is based on the immersion of the substrate in precursor solution either with or without temperature, the chemical reactions in the solution give rise to formation of desired oxide or hydroxide precipitates, which get deposited on the substrate surface [24].The method has been used for many years for the deposition of sulfide, selenide, and nonoxides [40]. The CBD method is a promising solution-processed method; as like other solution-­ processed, it too occurs at relatively low temperature and low cost of deposition but apart from that, the one major benefit of this method is that it is highly suitable for large-scale deposition. The method requires a strong chemical oxidant or reducing agent to drive reactions to take place. There is little work published on the use of this method for deposition of mixed oxides, some of it is presented here: Nickel-cobalt-manganese (NCM) ternary oxide nanosheets were prepared on nickel foam. The precursor solutions involved were manganese chloride, nickel chloride, cobalt chloride, and ammonia. The reaction mechanism is proposed as follows [41]: 2+



 M ( H 2 O )6 − x ( NH 3 ) x 

+ 2OH − → M ( OH )2 + ( 6 − x ) H 2 O + xNH 3 (1.11)



2 M ( OH )2 + S2 O8 2 − → 2 MOOH + 2SO 4 2 − + 2H + ( M = Ni,Co,Mn ) (1.12)

The formation of a highly porous structure of ternary NCM oxide was confirmed from SEM studies shown in Fig. 1.18a. The NCM nanosheets grown with higher density can be seen in through higher magnification SEM image shown in Fig. 1.18b. The nanosheets were interconnected with each other resulting in the formation of a highly porous structure as shown in Fig.  1.18c; the porous structure is also confirmed from TEM image shown in Fig.  1.18d. The HRTEM image is shown in Fig. 1.18e shows the formation of larger mesopores dispersed over the surface of NCM oxide nanosheets with an interplanar spacing of 0.24 nm, which is in agreement with the (311) lattice planes of NCM oxide. The polycrystalline nature of the films was confirmed from the SAED pattern shown in Fig.  1.18f. The prepared electrode was studied for the supercapacitor application. Figure 1.19 shows the CV curves of the Ni foam, cobalt oxide, and NCM oxide within the potential range of 0–0.6 V vs. SCE in 1 M KOH electrolyte at the scan rate of 10 mV s−1. As observed from the CV response, the Ni foam contributes very less to the specific capacitance of the NCM oxide; hence its contribution to specific capacitance was neglected. The NCM oxide showed a higher current response compared to cobalt oxide, resulting into an energy density of 36.4 Wh kg−1 at a power density of 320 Wh kg−1 [42]. Zhang and group deposited zinc (Z), nickel (N), aluminum (Al), and cobalt (Co) (ZNACO) oxide nanosheets on Ni foam. The bath contains urea, zinc nitrate

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Fig. 1.18  SEM image at different magnifications (a) 200μm, (b) 2μm, and (c) 200 nm, respectively, (d) TEM image, (e) HRTEM image (f) corresponding SAED pattern of NCM oxide nanosheets. (Reprinted with permission from [42], copyright 2016, Elsevier)

hexahydrate, nickel (II) nitrate hexahydrate, aluminum nitrate monohydrate, and

Fig. 1.19  The CV curves of Ni foam, cobalt oxide, and NCM mixed oxide at the scan rate of 10 mV s−1. (Reprinted with permission from [42], copyright 2016, Elsevier)

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cobalt (II) nitrate hexahydrate with a total volume equal to 400 mL. The bath then was kept in the oven at a constant temperature of 90 °C for 4 h. The schematic illustration of ultrathin ZNACO nanosheets on Ni foam is shown in Fig. 1.20a. The XRD pattern, SEM image, and surface area analysis are shown in Fig. 1.20b–g. The XRD pattern of the as-prepared ZNACO nanosheets confirmed the formation of cobaltbased layered double hydroxides shown in Fig.  1.20f. The SEM images in Fig.  1.20b–e show the formation of ZNACO nanosheets covered all over the Ni foam surface forming flower-like morphology. The double hydroxides nanosheets calcined in presence of argon gas to transform it into hierarchically porous structures of ZNACO nanosheet with the release of CO2 and H2O gases as per the reaction below [43].

Fig. 1.20 (a) Schematic illustration of the synthetic process of ZNACO nanosheets grown on Ni foam and SEM images of ZNACO nanosheets on Ni foam at different magnifications (b) 5μm (inset at 50μm), (c) 50μm, (d) 3μm, (e) 500 nm, (f) XRD pattern, and (g) N2 adsorption isotherm of as deposited ZNACO nanosheets (inset pore size distribution). (Reprinted with permission from [43], copyright 2016, Elsevier)

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Fig. 1.21 (a) The schematic illustration of hybrid supercapacitor configuration of ZNACO and activated carbon electrode and (b) CV curves of hybrid supercapacitor at different scan rates. (Reprinted with permission from [43], copyright 2016, Elsevier)

 2 − 3z  Co32 −  nH 2 O +   O2 3z / 2  2  (1.13) ∆ 3z → Zn 3 x Ni3 y Al3 z Co3 (1 − x − y − z ) O 4 + ( n + 3 ) H 2 O ↑ + CO2 ↑ 2

 Zn 3 x Ni3 y Al3 z Co3 (1 − x − y − z )( OH )6 

3Z +

The N2 adsorption-desorption measurements with an inset image of pore distribution are shown in Fig.  1.20g. The surface area observed was 83.5  m2  g−1 with average pore size distribution fixed at approximately 5.7 nm. The ZNACO sheets utilized as an electrode for supercapacitor showed a specific capacity of 839 C g−1 at 1.0 A g−1 and rate capability of 688 C g−1 at 20 A g−1. A hybrid supercapacitor comprising ZNACO nanosheets as a positive electrode and activated carbon as a negative electrode was fabricated, the schematic illustration of it is represented in Fig. 1.21a. The CV curves of hybrid supercapacitor at different scan rates in the potential range from 0 to 1.5 V in 2 M KOH are shown in Fig. 1.21b, with the presence of two redox peaks at 0.8 and 1.5 V. The hybrid supercapacitor delivered high energy of 72.4 Wh kg−1 at a power density of 533Wh kg−1 with higher cyclic stability, making it a promising material for supercapacitors. Table 1.1 gives a review of mixed oxide discussed in this chapter.

1.3  Conclusions Mixed oxides undoubtedly are dominant materials for catalysis applications, but now they are also utilized in other applications such as sensors, DSSC’s, and supercapacitors owing to their promising properties. The synergistic effects occurring when two or more than two metal oxides mixed in the material are giving rise to different structural, morphological, and electrical properties compared with a pure

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oxide material. Compared to physical methods, solution-processed methods for the synthesis of mixed oxides are gaining a lot of attention due to their simplicity, less cost, and low temperature deposition. In the solution-processed methods, sol-gel and precipitation methods are widely used for deposition of mixed oxides, but these methods do not provide any control over the morphology of the oxide, resulting in limiting applications. The methods like electrodeposition, SILAR, and CBD, in this case, are beneficial as they provide greater advantage of controlling the morphology. Mixed oxides in the form of nanosheets, nanowires, nanoplates were formed with these methods, which in turn provide more porous structure, larger surface area, and pore volume, which is useful in different applications, and these are performance deciding characteristics for supercapacitor application. The synthesis of supercapacitor electrode based on mixed oxide material showed impressive electrochemical properties compared to the pure oxides in terms of specific capacitance, specific energy, charge-discharge, and cyclic stability. We believe that mixed oxide materials will bring exciting results in the future in different research areas, especially in supercapacitors.

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Chapter 2

Properties and Applications of the Electrochemically Synthesized Metal Oxide Thin Films Abdellah Henni and Amina Karar

2.1  Introduction Modern society has changed with the various innovative materials designed by the materials science and engineering community. Modem technology requires nanomaterials and thin films for different applications. Nanostructured materials or thin films are experiencing growing interest in vast technological fields because of their physicochemical properties, which are often more interesting than those of massive materials [1]. Among these nanostructures, metal oxide receives a lot of attention because of their various applications, thanks to the existence of their dual property, electrical conductivity, and transparency in the visible, making them ideal candidates for applications in fuel cells, supercapacitors, sensors, catalysis, solar cells, or electrochromic windows (Fig. 2.1) [2–6]. Metal oxide semiconductors are an interesting class of materials, which are extremely economical and can be produced on the scale necessary to meet widespread demand. Several parameters and factors influence the efficiency of metal oxides thin films. For this, several researches have been conducted for the improvement of the thin films according to their application. This metal oxide has been prepared by a variety of deposition techniques, such as sputtering [7], pulsed laser deposition [8], chemical vapor deposition [9–11], spray pyrolysis [12–15], hydrothermal [16–18], and the sol-gel process [19–22], which were extensively studied. For these industrial applications, the preparation of optically transparent films in a large area and the improvement of their photo-functional properties would be crucial tasks. So far, thin A. Henni (*) · A. Karar Laboratory, Dynamic Interactions and Reactivity of Systems, Kasdi Merbah University, Ouargla, Algeria e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_2

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Fig. 2.1  Applications of metal oxides as thin films

films of transparent oxide are classically deposited from the vapor phase. However, these methods have high costs, and the preparation of films in a large area is technically difficult. Currently, we find in the literature a significant amount of research carried out on these metal oxides, which have been able to improve their optical and electrical properties. The most famous metal oxides are the oxides of indium, copper, titanium, tin, zinc, and nickel. Usually, the oxides are doped with a metal. This chapter discusses the electrochemically synthesized metal oxide thin films with emphasis on recent applications related to thin films.

2.2  Electrochemical Synthesis The electrochemical synthesis method for the elaboration of metal oxide as thin films is particularly necessary for the researchers working on the thin films. However, a bibliographical synthesis on this method in the development of

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nanostructures electrode has been neglected in most articles. Different oxide synthesis modes exist the galvanostatic mode (constant current deposition; two electrodes), the galvanodynamic mode (current pulses), the potentiostatic mode (constant potential deposition; three-electrode assembly), and the potentiodynamic mode (cyclic voltammetry or potential pulses). Generally, the electrodeposition procedure is performed in a three-electrode electrochemical (Fig. 2.2). Electrodeposition was carried out using a galvanostat-­ potentiostat. In this method, very low voltage or current is used to produce a thin film on conductive substrates. During electrodeposition, the morphology and film thickness can be controlled by the operating parameters: applied potential (E), current density (i), deposition time (t), temperature (T), and concentrations (C). Recently, several thin films based on metal oxides such as NiO, ZrO, CeO2, Cu2O, TiO2, SnO2, and ZnO using aqueous solutions have been made with electrochemical methods. In some cases, the deposition leads to direct formation of crystalline thin films under near-room temperatures. Among technological advantages generally found for the electrodeposition are (1) economy and simplicity, (2) film thickness control, uniformity, and large-area deposition, and (3) good adhesion, high quality, and the possibility of using substrates of nonuniform and complex shape [23–25].

Fig. 2.2  Overview showing the electrodeposition method

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2.3  Electrodeposition of Metal Oxide as Thin Films A thin film of metal oxide is a solid material deposited on a substrate that has oxygen anions chemically bonded to one or more metal cations. This work is devoted to the most used metal oxides: ZnO, Cu2O, and NiO.

2.3.1  Zinc Oxide (ZnO) Zinc oxide (ZnO) is a semiconductor that is gaining more and more attention in many areas [26–28] due to its ability to form a wide variety of nanostructured forms such as nanorods, nanowires, nanofibers, nanobelts, nanospheres, nanocombs, and nanoflowers. ZnO nanowires (NWs) and nanorod (NRs) are the two forms most used in solar cells, thanks to their large contact surfaces [29, 30], and it is thermally and chemically very stable under hydrogen plasma processes used in photovoltaic cells production [31, 32]. Generally, in this deposition method, molecular oxygen (O2) [33–36], nitrate (NO3−) [37–40], or peroxide hydrogen (H2O2) [41–44] are used as precursors of hydroxide (OH−). The fundamental reaction characterizing the electrodeposition of ZnO is based on the reduction of the precursor of hydroxide ions. The OH−-generated and Zn2+ ions chemically form Zn(OH)2. Zn(OH)2 is then spontaneously dehydrated to form the ZnO film according to Eq. (2.1).

Zn 2   2OH   Zn  OH 2  ZnO  H 2 O

(2.1)

The effect of the main parameters (Zn2+ precursor concentration, electrodeposition time, support electrolyte concentration, and electroplating temperature) on ZnO growth was studied (Fig. 2.3), and the influence of those different parameters on the growth mechanism, density, thickness, and shape was discussed [45]. Henni et al. [46] have studied the effect of potential on the ZnO nanostructure electrodeposited onto ITO. They showed that the increase of potential affects the morphological characteristics of nanorods, including density, height, and diameter. Elias et al. [47] have investigated the effect of charge densities on the features of the ZnO nanostructures. Zheng et al. [48] used the NO3− reduction method to deposit ZnO nanowires in AAM-type matrices. After the deposition of ZnO in the pores, they dissolve the matrices and obtain ZnO nanowires. Subsequently, several studies were made with this method to deposit ZnO nanowires without the matrices. Xu et al. [49] used this same method, but adding several types of amines and inorganic salts to the deposition solution to alter the morphology of the deposit. They obtained different morphologies including ZnO nanowires. The use of O2 as a precursor makes it possible to deposit arrays of ZnO nanowires without matrices. This method was developed by the Lincot group [50]. It

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Fig. 2.3  Illustration of the electrochemical growth of ZnO NRs when (a) the OH− generation rate is higher than the Zn2+ diffusion rate and (b) the Zn2+ diffusion rate and the OH− generation rate is of the same order [51]. Top-SEM images of ZnO electrodeposited at different concentrations of Zn(NO3)2: (c) 5 × 10−2 mM, (d) 5 × 10−1 mM, and (e) 5 mM [51]

consists in saturating the solution with O2 in the presence of ZnCl2 and KCl. Once the solution is saturated with O2, a potential of −1.0 V vs. SCE is applied to reduce the O2 to OH− at the surface of an electrode (SnO2: F). The formed OH− is adsorbed on the electrode and react with Zn2+ to form a compact layer or nanowires of ZnO. In recent years, thin films of doped ZnO have been widely studied and many published papers have shown that the properties of ZnO can be improved by the change of dopant and its concentration. 2.3.1.1  Applications of ZnO The unique properties of ZnO have sparked great interest in industrial applications. Several studies show that ZnO nanowire arrays have better electronic properties than porous TiO2 layers [52]. Hendry et al. [53] determined the electron mobility of nanocrystalline TiO2 layers and found much lower values (10−2 cm2 V−1 s−1) than those estimated for ZnO nanowires by Konenkamp et  al. [54] (23  cm2  V−1  s−1). Recently, several groups have replaced the layers of TiO2 nanoparticles in dye-­ sensitized solar cells (DSSCs) with ZnO nanowires [55–57]. Conversion yields of ~2.5% were obtained by Law et al. [58] for this kind of cells. Other groups have used the same principle of the DSSCs based on ZnO nanowire arrays, replacing the

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dye with different types of absorber polymers with high hole mobility [59, 60] or with quantum dots [61]. On the other hand, Konenkamp et al. [62] proposed the concept of an extremely thin absorber (ETA) solar cell as a completely inorganic and solid variant. In 2005, the group of Lévy-Clément et  al. [63] showed the experimental validity of the concept of ETA cells using the ZnO/CdSe/CuSCN heterostructure based on ZnO nanowire arrays. Lévy-Clément et  al. [63] have proposed a heterostructure based on networks of ZnO nanowires obtained by the electrochemical reduction of O2 on an FTO substrate. The yield obtained (2.3%) showed the experimental validity of the concept of ETA type cell. For an understanding of the wettability of nanorod ZnO surfaces, Ghannam et al. [64] synthesized ZnO NRs of different sizes using an electrochemical route. The results change in the chemical and physical texturing of the ZnO surface causes a change in the contact angle, mainly due to a difference in the penetration of the drop of liquid. Alternating voltage electrochemical dispersion (AVED) has been successfully used to design a ZnO-NRs/RGO composite for Na-ion battery as an eco-friendly and easy electrochemical method [65]. As-prepared ZnO/RGO shows ultralong cycle life with ~92% capacity retention after 1000 cycles at 500 mA g−1. ZnO nanowires and nanorods are attractive materials for gas sensors due to their ultrahigh surface/volume ratio. Electrochemically deposited ZnO NRs layers for gas detection applications have been reported by Chetna et al. [66]. This route may offer new avenues fort more efficient and low-cost sensor with improved gas detection performance. Co-doped ZnO/rGO composite was synthesized method coinduced via a novel coinduced electrochemical method in aqueous solution at room temperature [67]. They observed that the cobalt-doped ZnO/rGO improved the photocatalytic activity and stability under light irradiation. Henni et  al. [68] studied the ZnO/graphene composite synthesized by a one-step electrochemical deposition approach for efficient photocatalytic degradation of organic pollutant. The results obtained show that the ZnO/rGO nanocomposite is a promising material for the degradation of organic contaminants such as the methylene blue dye.

2.3.2  Copper Oxide (Cu2O) Copper oxide thin layers can be produced using a wide variety of techniques, namely chemical bath deposition [69], hydrothermal [70], laser ablation [71], sputtering [72], and electrochemical deposition [73]. It is obvious that the time used for electroplating has a remarkable effect structure and morphology of the layers. Zhai et  al. [74] obtained thin Cu2O layers of good crystallinity and different morphological shapes: cubes, tetrahedron, truncated cubes, and truncated octahedron; these latter shapes were electrodeposited at different deposition times (1–4 h), respectively (Fig. 2.4). The optical properties of electrodeposited Cu2O thin films have shown that the optical gap changes from 2.21 to 2.43 eV depending on the electrodeposition time.

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Fig. 2.4  Top views of the SEM images and illustration of the Cu2O shape evolution process on ITO–PET substrate with different electrodeposition time: (a) 1 h, (b) 2 h, (c) 3 h, and (d) 4 h [74]

During electrodeposition, the surface condition of the substrates is an important data in understanding some properties such as the morphology, roughness, and average size of thin film crystallites. The choice of substrate has an essential role in controlling the growth of Cu2O. Liu et al. [75] developed thin layers of Cu2O in an alkaline medium containing lactic acid, copper sulfate, and NaOH. Different substrates were used: ITO, Si, and Au. Observations by SEM show that the deposits obtained on ITO have grains in the form of pyramids with a size of ∼2 μm. These pyramids grow in orientation (111) on the substrate. Also, the deposits deposited on the Si substrate exhibit pyramids with four faces developed according to orientation (111). A remarkable decrease in grain size was observed (~100 nm) when compared to this obtained on ITO. Films developed on the Au substrate exhibit an orientation of (100) with good crystallinity. These results could be explained by the difference between the mismatch of the substrate lattice and the deposit (parametric mismatch).

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El-mezayyen et al. [76] studied the effect of substrate conductivity on the morphology of Cu2O. These authors developed Cu2O nanostructures electrochemically using three different substrates: ITO, FTO, and (ITO/PET). Different Cu2O morphologies such as ferns, dendrites, as well as dense films were formed on ITO, FTO, and ITO/PET substrates, respectively. The Mott-Schottky curves showed that all elaborated films are p-type semiconductors with charge carrier density varying between 1.4 × 1018 and 1.2 × 1019 cm−3. The crystallinity of the layers is one of the most important factors concerning the functional properties of Cu2O. The pH of the electrolyte is a parameter that has a direct influence on the size of crystallites. The effect of the pH of the electrolyte has been studied by Golden et  al. [77]. They deposited films of p-type Cu2O by the reduction of copper (II) lactate in 0.4 M copper sulfate with 3 M lactic acid, and they concluded that the surface texture of Cu2O thin films is directly affected by the pH of the bath. At pH = 12, the preferential orientation of the Cu2O is (111), while at pH  =  9, the preferential orientation changes to (100). When the pH 14  nm), the rutile phase becomes more stable than the anatase phase, also it has high growth rate compared to anatase. This phase is almost stable at all temperature and at pressure up to 60 kbar; as a result, thermodynamically, TiO2 (II) becomes much favorable form [29]. The phase anatase TiO2 also has a tetragonal structure with six atoms per unit cell with large distortion in TiO6 octahedron as compared to rutile phase [30]. At temperature 0 °K, anatase TiO2 is more stable compared to the rutile TiO2, however, there is small (∼2–10 kJ/mol) energy difference between both the phases [31]. Whereas, brookite TiO2 has orthorhombic crystal structure, with eight formula units for each unit cell, which is formed by edge-sharing of TiO6 octahedron [22]. Abrahams and Bernstein [32] have carried out the experiment via single-crystals X-ray diffraction, where rutile structural parameters were determined to have a higher degree of accuracy. Various other investigations for single crystal and powder form have been done via X-ray and neutron methods to check the shift in the positional parameters of oxygen atoms between the two techniques because of polarization of the O atom. As a result, only a small shift difference was observed in the oxygen positional parameter between X-ray and neutron experiments [33]. The system’s ground state can be achieved by minimizing the total amount of energy for the P42/mnm structure with lattice constant of a, c, and u, which is performed by two-step process [14]. The unit cell volume, that is, V  =  ca2 is specified by minimizing the total amount of energy in regards with c/a and u through the calculated stresses and forces as a support. The step is then repeated for different volumes close to the proposed equilibrium structure. In this fashion, structural parameters and cohesive energies are obtained. The structural parameters in terms of the reduced volume, that is, V/Vo (Vo is hypothetical volume of unit cell at

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Fig. 3.2  Ideal structure parameters for TiO2 rutile structure predicted by ab-initio total energy minimization for different reducing volumes: (a) a (b) c, and (c) u. (Reproduced from ref. [14] with permission, American Physical Society, Copyright, 2020)

Fig. 3.3  Energy-volume relationship based on Vinet et al. [34] along with the cohesive energy Ecoh and the unit cell volume per TiO2 molecule. (Reproduced from ref. [14] with permission, American Physical Society, Copyright, 2020)

equilibrium condition) as present in the Fig. 3.2. Whereas, Fig. 3.3 illustrates the corresponding energy-volume relationship based on Vinet et al. [34] with the Ecoh, that is, cohesive energy and the unit cell volume per TiO2 molecule given. The energies of the isolated pseudo-atoms were found to be −89.88 eV and −429.59 for Ti and O, respectively. These energy values contain spin-polarization correction [35] of 0.68 and 1.41 eV for Ti and O by applying identical spin-polarized exchange correlation potentials [36, 37]. The ground-state structure is formed by minimization of total system energy with reference to volume. This is commonly performed

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through integrated equation of state (EOS) by substituting the theoretical values of energy and volume in it.

3.2.3  Electronic Structure The TiO2 electronic structure has been experimentally examined by various characterization techniques such as X-ray-induced Auger-electron spectroscopy [35], X-ray photoemission spectroscopy (XPS) [38], Auger-electron spectroscopy [39], X-ray emission (XES) [40, 41], absorption spectroscopy (XAS) [42, 43], electron energy-loss spectroscopy (EELS) [44, 45], ultraviolet photoelectron spectroscopy (UPS) [46], and resonant photoelectron spectroscopy (RUPS) [47]. For TiO2, calculations of the self-consistent ab initio electronic structure are more accurately available. Within the LDA, the calculations of self-consistent total energy could be performed by using a pseudopotential plane-wave formalism [48, 49], also it has been performed by using Hartree-Fock pseudopotential calculation [50]. Recently, these calculation are performed by LMTO, that is, linear muffin-tin orbital method [51], in order to illustrate the near edge structure in core level spectroscopies. TiO2 is a semiconductor material with energy bandgap of 3.02, 3.2, and 2.96 eV for the rutile, anatase, and brookite phases, respectively [52]. The conduction band of TiO2 is 3d orbitals of titanium, while valence band contains the 3d orbitals of titanium with 2p orbitals of oxygen hybridized [53]. It was found that the rutile and anatase phases of TiO2 can improve the absorption rate of visible light in comparison with pure phases [54]. The phase anatase TiO2 is known to be active photocatalytic component appropriate to chemical properties, charge-carrier dynamics, and photocatalytic activity. As compared to rutile, it has essential surface-band bending results in higher potential with deeper region [55]. Hence, the surface hole controlled by synthesis conditions, precursors, impurities, oxygen vacancies, and the initial particle size of the anatase phase. It was observed that the higher photocatalytic activity can be achieved by using high surface area and high crystallinity [56].

3.3  Anatase TiO2 Nanocrystals Nanocrystals of anatase TiO2 have attracted significant attention because of their numerous applications such as photovoltaics, fuel cells, and photocatalysis [7–9]. Because of its high catalytic performance, anatase nanocrystals have recently drawn considerable attention to study the characterization of TiO2 nanocrystals [57–60]. It is expected that the high performance of TiO2 nanocrystals would come from the amplified surface effect; though to explain the behavior of TiO2, nanocrystals are insufficient by simple estimation of the extended surfaces. The characterization of TiO2 nanocrystals is typically challenging task in both numerical simulations and experiments. Previously, to disclose the excess amount of electrons present in the reduced-anatase nanocrystals, many techniques have been used such as XPS and electron paramagnetic resonance. The more effective technique, that is, STM for

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characterization of extended surfaces, may not be used in nanocrystals, which makes it difficult to understand a single nanocrystal at the atomic level. The experimental results shows, the performance of TiO2 nanocrystals could be significantly enhanced through the reduction technique. Generally, reduction of the TiO2 surface leads toward the higher surface reactivity [61, 62]. This change was referred as the formation of oxygen vacancies and Ti interstitials, defects, and resultant-excess electrons. Here, by using DFT, the properties of defects produced in the reduced-anatase TiO2 nanocrystals and introduced-excess electrons by defects have been studied. The research focuses on small reduced-anatase nanocrystals with {101} plan as a nanocrystal model and shows the properties of spatial defects and introduce of excess electrons by defects. To clearly explore the properties of defects, nanocrystals reduction effects, oxygen vacancies (Vo), and Ti interstitials (Tiint) are examined together with the introduced excess electrons. Firstly, the stability of defects with various position is calculated, which shows the existence of Vo at both surface sites as well as subsurface, while Tiint is stable at subsurface. Next, the introduced-excess electrons through defects are studied, which shows they can spatially occupy various Ti sites though they localize to form polarons. The differences and similarities of nanocrystals defects are compared with those of the extended surfaces [63]. The spin-polarized calculations of TiO2 were carried out by using VASP, that is, vienna ab-initio simulation package carried out with the plane-wave basis sets, and the scheme of projector-augmented wave was used [64, 65]. Where the results were visualized by using VESTA [66]. Previously, many investigation have been used on TiO2 bulk and surface plane-wave basis sets [67, 68]; utilizing a plane-wave basis set provides the comparison between the results of this analysis with earlier work. Since for the valence electrons, 2s and 2p states and 3s, 3p, 3d, and 4s states were taken into consideration for O and Ti, respectively. In the previous studies, the anatase nanocrystals with facets of {101} were examined, where the shape is predicted by the Wulff construction [69, 70]. In addition to the complexities of handling nanocrystals, the excess electrons’ complex behavior must be addressed when addressing defects. Earlier investigation of TiO2 extended surfaces, it is concluded that the defects may lead to introduction ofexcess electrons and due to these excess electrons results in distortion of local lattice and forming polarons [71, 72]. The proper explanation of both electronic and atomic structures is needed for the polaronic nature of excess electrons. In handling the excess electrons, the problem appears since local and semi-local functionals in the DFT approach failed to explain the self-interaction of electrons in TiO2, as a result, it lead to deficient localization behaviors of excess electrons. This drawback can be improved by using GGA + U (GGA with Hubbard U corrections) or by using other hybrid functional. Firstly, in nanocrystals, the defects stability of 105 and 495 atoms were determined. For the exchange and correlation functional, GGA demonstrated by Perdew et al. [73] was employed. Secondly, the electronic property was investigated by hybrid functional, HSE06 [74, 75].

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3.3.1  E  lectronic Properties of Reduced TiO2 Nanocrystals and Stability of Defects Focussing on TiO2 nanocrystals with {101} plane as a model, several simulations of different oxygen vacancies and Ti interstitials positions were performed. Where VO showed equivalent stability on the subsurface and the surface sites, while Tiint was more stable at the subsurface site. For nanocrystals, surface VO can be stabilized when compared with the extended surface. Depending on their localization behavior, introduced-excess electrons through VO and Tiint were occupied deep and shallow states. Moreover, in some structures, the valence band tailing off up to 0.6 eV was noted, which used to provide the excess electrons’ aggregation at the surface of nanocrystal TiO2. These excess electrons’ behaviors are reported with the help of experiment, which indicates narrow bandgap of the anatase nanocrystals, resulting from valence-band tailing and in-gap state. This study demonstrates the band structure tuning possibility through defects and used to enhance the understanding of reduced-anatase nanocrystals. Within the surface sites, Tiint is more stable between the twofold-coordinated oxygen atoms for the calculations. This finding is similar to Cheng and Selloni’s extended surface analysis. From this analysis, they proposed that the diffusion of Tiint takes place at subsurface regions. Similar conclusion was presented by Cheng and Selloni [76]. The analogy of the extended surface and nanocrystals suggests that Tiint may diffuse from the surface to deeper areas.

3.4  T  iO2 Nanocluster and Dye–Nanocluster Systems: Photovoltaic or Photocatalytic Applications The properties (physical and chemical) of nanocrystals TiO2 are largely determined by its electronic structure, which are controlled through the shape, size, surface properties, and structure of the nanoparticles (NPs). Some electronic properties play an vital role in photovoltaic and photocatalytic applications. For example, the alignment of the energy level between TiO2’s valence or conductive band-edge, with respect to ground and the excited state of absorption of molecule or electrolyte’s redox level, determines whether or not a process can occur. Cluster approach [77] is normally used to analyze the interaction between TiO2 surface and individual dye, while a periodic approach should be used for the study of interfaces of periodic crystalline materials, such as perovskite, on TiO2. Despite major computational studies were stated so far for TiO2 clusters, studies can still not address the certain problems or particular complex systems because of cluster sizes used in the study are not appropriate to address the issue. To find out TiO2 nanoclusters with small size, the issues are related to cluster size, which should be chosen in correspondence with the adsorption of molecule size. While choosing the suitable cluster, the rigidity of the molecule’s backbone plays an important role or not it needs to be addressed in the study. To deal with these issues, five various computational studies of

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TinO2n + 2H4 clusters (where n = 14, 24, 34, 44, and 54) were reported, to realize the effects of size on geometrical optimization and transfer of charges. To investigate the structure, anchoring the optical properties, electronic spectra, and the charge injection mechanisms from the adsorbed molecule to such TiO2 clusters. Study based on adsorption of molecules is simple and commonly known methods like penicillin V [78], coumarin-based dye, C343 [79], in addition to new molecules, that is, OMCD1 [80].

3.4.1  Methods and Materials The titanium oxide clusters reported [81], irrespective of size, were firstly shorten from the experimental structure of anatase to model the surface of (101). As shown in Fig. 3.4, there are four atoms of hydrogen that have been utilized to resolve the dangling bonds of O atoms bounded by two atoms of Ti at the edge of Ti14O30H4 cluster or by threefold-coordinated Ti atoms at the edge of Ti54O110, Ti44O90H4, Ti34O70H4, and H4Ti24O50H4 clusters, to ensure the existence of charge neutrality under-coordinated Ti atoms. The molecules were investigated, OMCD1, C343, and penicillin V, both adsorbed and isolated on the cluster, using DFT approach for minimization of energy using the split valence 3-21G(d) basis set with polarization functions [83, 84] and hybrid B3LYP exchange-correlation functional [85, 86]. In order to find out the electronic state density, the single-point calculations were

Fig. 3.4  Optimized structures of nanocluster (left to right) Ti14O30H4, Ti24O50H4, Ti34O70H4, Ti44O90H4, and Ti54O110H4, for anatase TiO2 (101) plane. Ti atoms represented by gray, O atoms represented by red and H atoms represented by light gray color. (Reproduced from ref. [82] with permission, Nanomaterials, Multidisciplinary Digital Publishing Institute (MPDI Copyright, 2020)

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carried out via hybrid B3LYP exchange-correlation functional along with the effective core potentials (ECPs) for atoms (Ti, P, S, and I), also double-quality basis functions by using LANL2DZ [87] and attributed to aqueous solvent effects by using conductor-like polarizable continuum model (C-PCM) [88]. Based on the United Atom for Hartree–Fock procedure [89], the cavity was created from center of spheres on strong nuclei, which can be used in C-PCM calculation. By using time-dependent DFT (TD-DFT), the 50 lowest singlet-to-singlet transition of electrons were calculated. The calculations were carried out by using GAUSSIAN09 quantum chemistry package [90], whereas, the UV–visible spectrum on various system components and the projection of the DOS were achieved by GaussSum.

3.4.2  S  tructural and Electronic Properties of TiO2 Nanocluster and Dye–Nanocluster Systems The deformation of the lattice parameters can be examined with the average value of the Ti-O distance to be taken into consideration and its equivalent standard deviation for the different clusters and to compare it with the experimental values for the bulk. The average value of Ti-O distance is always small as compared to

Fig. 3.5  The DOS of system consist of molecule C343 adsorb on the Ti14O30H4 cluster through (a) two or (b) three O-Ti bonds, and on the Ti24O50H4 cluster (c). States density of the TinO2n + 2H4 clusters where n = 14, 24, represented by dotted curves. Calculated by DFT at the level of B3LYP/ LANL2DZ in a water solvent. (Reproduced from ref. [82] with permission, Nanomaterials, Multidisciplinary Digital Publishing Institute (MPDI Copyright, 2020)

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Fig. 3.6  Isodensity surfaces (0.03 e/bohr3) of molecular orbitals of the systems with C343 molecule adsorbs on the Ti14O30H4, cluster through two or three (left or middle column) O-Ti bonds, and on the Ti24O50H4 cluster (right column), DFT calculations performed at the level of B3LYP/ LANL2DZ in water solvent. LUMOs represented in first row and HOMOs represented in second row. Ti atoms in gray; O atoms in red; and H atoms in light gray color. (Reproduced from ref. [82] with permission, Nanomaterials, Multidisciplinary Digital Publishing Institute (MPDI) Copyright, 2020)

experimental observed value, that is, 1.950 Å, only effective for the bulk oxide, for equatorial and axial bonds, it has two different values of Ti-O distance. The average distance value of Ti-O increases as the size of cluster increases, smaller the clusters, stronger the distortion. The distortion in lattice can also be examined by the Ti-O distances distribution, in comparison with bulk. Furthermore, the Ti-O distances’ range increased with the size of cluster. For the smaller clusters, larger the surface-­ to-­bulk ratio, and lower the distribution of distances. In Ref [78], authors have examined states’ density for the three-model systems and compared with the typical graph for the bare TinO2n + 2H4 clusters where n = 14, 24 as shown in Fig. 3.5. The molecule C343 contributes more effectively in the valence band and specially in the bandgap in order to find out the density of states, at which complex system HOMO is placed. The data of state-curves density is presented in Fig. 3.5 and is finalized by the density of electron shown in Fig. 3.6 for the molecular orbitals. Throughout in all three mentioned systems, the HOMO energy level has a major contribution from the molecule absorption, because of the organic dye’s conjugated character, the molecular orbital having a nature. In compared with HOMO level, the LUMOs having major offering from the oxide, and the nature of the molecular orbital can be calculated by d-atomic orbitals of Ti.

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3.5  Photoexcited TiO2 Nanoparticles The properties of the photo-excited carriers are important to the photocatalysts and photovoltaic performance of TiO2. When photoexcitation occurs by UV light absorption, electrons present in TiO2 are excited from the valence band state to conduction band states, whereas, holes are formed in the valence band. Charge carriers, that is, electrons and holes, may react with adsorbed molecules after diffusion at the surface and being locked at defect location, then again move to a collecting electrode through TiO2 crystal. The maximum quantum efficiency can be achieved through recombination of electron for the photovoltaic and photocatalytic processes by separating of the charge and preventing from further recombination. In TiO2, the charge transport characterization plays significant role in models development and to control the reduce rates of recombination. For this reference, the nanocrystalline anatase TiO2 is commonly suggested as compared to rutile TiO2 due to its high electron mobility [91] and photocatalytic activity [92], which allows long diffusion pathways for photo-induced carriers as well as improves the quantum efficiency. The investigation of TiO2 from polaron states to NPs becomes more interesting, owing to the important role of TiO2 nanocrystalline in the application of photocatalysis and photovoltaics. The structure and excitons energy of anatase TiO2 NPs with single electron and hole polarons can be characterized by using hybrid DFT and TD-DFT calculations. This study deals with the impact of the HFexc in case of electron, hole, and exciton polaron states from the comparison of hybrid DFT functional results with pure GGA functional result by using various percentages of HFexc. The modeling study of TiO2 NPs, a stoichiometric cluster of TiO2 reveals the majority surface of (101), which is to be considered. Entire measurements were totally on the basis of DFT and used three hybrid functionals viz. B3LYP, B3LYP, and CAM-B3LYP [93, 94] and GGA along with BLYP functionals [85]. In addition

Fig. 3.7  Unrelaxed, S0: B3LYP-performed ground state, T1: photo-excited state of geometrical arrangement of (TiO2)38 cluster. Ti atoms represented in gray color, and O atoms represented in pink color. (Reproduced from ref. [96] with permission, American Chemical Society (ACS))

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Fig. 3.8  Difference in the value of Ti-O distances (Å) in the photo-excited state and ground state for the (TiO2)38 cluster, determined by using ground-state DFT (SCF, top of the panel) and TD-DFT (TD, the bottom of the panel) with different functional. (Reproduced from ref. [96] with permission, American Chemical Society (ACS))

to this, the calculations TD-DFT under different functionals, that is, B3LYP and CAM-B3LYP have been performed to determine the adiabatic, that is, vertically exciting energies of the excited state S1 (lower singlet state) and T1 (triplet state) at the optimal geometrical ground state. Where BLYP integrated with gradient-­ corrected replaced function of Becke [85] with the LYP correlation function [95].

3.5.1  Structural Properties Molecular structure of TiO2 with geometric parameters is shown in Fig.  3.7. A geometry reorganization of both photo-excited states, that is, the SCF T1 (e− – h+) and the single D1 (e−/h+) along with estimated bond distance differences of Ti-O is 0.2  Å. GGABLYP functional shows less reorganization, for particularly reduced (TiO2)38− cluster, where the modifications in the value of Ti-O distances never exceed to 0.07 Å with regards to state S0. In 3LYP-15 functional, the determined value of Ti-O bond distance differences found to be under 0.1 Å and 0.2 Å for the same cluster, that is, (TiO2)38− and the cation (TiO2)38+ along with photo-excited (e− − h+) clusters, respectively. For the T1(e− − h+) photo-excited state, functional CAM-B3LYP level showed the Ti-O bond distance differences of 0.2 Å, which is same as achieved with B3LYP functional. Furthermore, the differences of the Ti-O bond distance among ground state (S0 and S1) and excited states (T1) are measured by TD-DFT. The study shows that, in the TD-B3LYP the geometrical relaxations of excited states, S1 and T1 are identically equal to the TD-CAM-B3LYP, whereas, the relaxation of state S1 with reference to S0 in TD-B3LYP is slightly large as compared to TD-CAM-­B3LYP. Further,

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Fig. 3.9  MO energy level diagram of (TiO2)38 cluster, performed by different theories in DMF solvent; (a) BLYP, (b) B3LYP-15, (c) B3LYP, (d) CAM-B3LYP. H and L in black color shows HOMO and LUMO, SO in red shows SOMOs, and SU in green shows SUMOs. (Reproduced from ref. [96] with permission, American Chemical Society (ACS))

small difference occurs when compared with SCF T1 and TD-DFT S1-excited state calculated by CAM-B3LYP.  Because of these differences, the estimation of selftrapping and adiabatic excitation energy can be affected. The results show that, when photoexcitation occurs, the level of structural reorganization depends upon the electronic structure method as shown in Fig. 3.8.

3.5.2  Electronic Properties Figure 3.9 shows the TiO2 NP molecular orbital energy level diagram, which indicated after excitation, the charge carriers (electrons and holes) are locked into energy bands within the bandgap, and location of these energy bands is decided by the technique used for the calculation purpose. In photo-excited state of T1 (e− − h+), higher energy HOMO of the electron located at the unoccupied bottom states with ΔESO-L of 0.07 eV for BLYP and its energies are 4.88, 1.89, and 0.97 eV below the LUMO for the hybrid functionals of CAM-B3LYP, B3LYP, and B3LYP-15, respectively. Likewise, lower energy SUMO of holes placed at 0.56 eV over the occupied states for BLYP, whereas, its bandgap 6.13, 3.02, and 2.26 eV above the HOMO, for the hybrid functionals of CAM-B3LYP, B3LYP, and B3LYP-15, respectively. Figure 3.10 represents the electron-spin densities of NP TiO2 with photo-excited. The BLYP functional provides the spin density based on three sites of Ti and three O centers for T1 state. The hybrid B3LYP-15 functional provides nearly similar spin localize on the O centers, but for the Ti centers, it consists of strong localization on site Ti36 vs. Ti6. For both the level, that is, B3LYP, CAM-­B3LYP, the state T1 exhibits strong localized of electrons Ti3 and holes at O51. Furthermore, the BLYP spin density is bounded by Ti atoms with three centers for an excess electron, while functional B3LYP-15 locates the spin density only on center of Ti36 with small share on Ti6, whereas, functional B3LYP locates the excess electron particularly on the Ti3 center. Excess holes are moderately delocated on O

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Fig. 3.10  Spin densities of electron calculated by different theory in DMF solvent for single extra hole (right), the photo-excited T1 state (left), single extra electron (middle), and in (TiO2)38 cluster. (Reproduced from ref. [96] with permission, American Chemical Society (ACS))

atoms with three centers, that is, O94, O113, and O49 at the B3LYP-15 and BLYP levels that exhibits a strong contribution to the O94 atom. As indicated for the T1 state, the excess hole is placed in quite different location such as the O51 atom with a

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slightly shared from O70 and O79 atoms at the functional B3LYP level. The spin density of the photo-excited cluster nearly equal to the addition of the spin densities of the single excess hole and electron, the spin density localization degree increases with increment in the HFexc fraction. It is also found that, for all levels of theories, the excess hole preferred to be localized under-coordinated O2c instead of fully coordinated O3c centers. Whereas, excess electron preferred to be localized on fully coordinated Ti6c center for Ti51 with B3LYP and CAM-B3LYP and for other levels Ti36 sites, instead of on under-coordinated Ti4c sites. This may be based on: geometrical optimization of the nanoparticles from the bulk structure, tetragonal configuration strongly rearranged from an octahedral imposed by Ti4c centers, with a considerable reducing the Ti-O bond distances (0.1–0.2 Å at the B3LYP level), while the Ti6c centers holding to octahedral configuration only, with slight change in the Ti-O bonds distance.

3.6  Indium Oxide (In2O3) Indium oxide (In2O3) material plays a significant role in microelectronic device such as solar cells [97], flat-panel display sensors [98], transparent conductors [99], and architectural glasses. The In2O3 is the well-known TCO, that is, transparent conducting oxide, because it is customized as semiconductor or conductor or an insulator. TCOs show unique properties like either it can be conductors or semiconductors, lower electrical resistivity (80%), and higher infrared reflectance with a wider bandgap [100]. The performance and efficiency of the material totally based on their optical and electrical properties. TCOs have been studied for numerous applications, still its electronic structures is key challenge in terms of the character and size of the bandgap [101], band structure details [102], and the possibility of accumulation layer of charges on the surface of In2O3 [103]. The LDA approach is assumed as the class of Ceperley and Alder (CP) used to explain the exchange-correlation potential energy functional [104]. The SIESTA is a computational program, which stands for the Spanish initiative for electronic simulations with thousands of atoms, used to determine ground-state configurations implemented by numerical atomic basis approximation [105]. Here, basis sets, that is, double-zeta plus polarization (DZP), were used for simulation purpose [104]. The valence electrons have been addressed by norm-conserving pseudopotentials [106], where basis sets of I and O and the pseudopotentials were drawn from SIESTA code. The configurations of an atomic and core are I: [Kr] 4d105s25p1 and O: [He] 2s2 2p4. The three configurations of In2O3, that is, trigonal, cubic, and orthorhombic all were used for simulation. The lattice of size of (9.36, 8.83, 7.65), with k-points (4, 4, 4) and a mesh cutoff 450 Ry, was achieved in cubic structure as a result of their corresponding optimizations, similarly for orthorhombic and trigonal structure, the values are (3.36, 6.02, 13.60), (6, 6, 6), and 350 Ry and (6.18, 5.16, 4.82), (7, 7, 7), and 400 Ry, respectively. In conclusion, optimization of all the

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structures with respect to their k-points, optimized values of lattice constants, and cutoff until the required forces on each atom to be smaller than 0.005 eV/Å.

3.6.1  Structural and Electronic Properties The In2O3 crystallized into three configurations, that is, trigonal, cubic, and orthorhombic structures with space groups of: Pnma(62), Ia3(206), R3c(167). In cubic structure, the length of bonds between the atoms are In–In 3.63 Å, O–O 3.58 Å, and In–O 2.25 Å, while in orthorhombic and trigonal are In–In 3.67 Å, O–O 3.05 Å, and In–O 2.57 Å, and In–In 3.66 Å, O–O 1.42 Å, and In–O 2.59 Å, respectively. There exist a slight difference for the bond length of atoms In-In and In–O and a large variation in bond length of O–O. The structures size slowly decreased due to reduction of total number of atoms from cubic: 40, orthorhombic: 20 to trigonal: 10. The phase metastable rhombohedral structure of In2O3 has hexagonal unit cells with lattice parameters a = b = 5.476 Å and c = 14.51 Å, which contains six formula units per hexagonal cell. The In2O3 bcc-structure with 10.117 Å lattice constant contains 16 formula units per cubic unit cell [107]. It contains 80 atoms per unit cell, which is originated from supercell (2 × 2 × 2) of fluorite structure. To maintain an ordered structure, from fluorite structure 25% of oxygen atoms are removed. It includes two unequal indium atoms, which are placed at two different lattice positions. Where 48 identical oxygen atoms fill the edge positions, 8 indium atoms occupy In-b site, and 24 ones occupy In-d sites [108]. The projected cubic structure density states, that is, 0 eV shows the Fermi energy level, which is lower to the valence band and upper to the conduction band. Through valence band to Fermi energy level, the lower and broad peak of 5p-orbital states of indium (In) atoms and the higher and broad peak of 2p-orbital electronic states of oxygen (O) atoms, vice versa, were observed in the conduction band, where in the valence band, there is p–p hybridization, and in the conduction band, there is s–p hybridization. Approximately, in orthorhombic structure, similar type of orbital states was observed besides lower bandgap at higher levels of the conduction band. Whereas in trigonal structure, it was observed that there is a higher peaks of 5p-orbital electronic state of In at Fermi energy level within conduction band range, and few higher peaks of 2p-orbital electronic state of O atoms are found in the higher valence band region, while in the lower region of valence band, 2s-orbital electrons of In were found. In lower region of valence band, s–p hybridization is present and p–p hybridization in higher valence band region, whereas in conduction band region, there is p–p hybridization in all structures. Orthorhombic and cubic structures show zero bandgap, that is, no forbidden energy gaps, therefore, they are known as conductors. But in trigonal structure, it was found that there were two forbidden energy gaps, which are close to Fermi energy level with lower bandwidth of 1.0 eV and upper bandwidth of 0.5 eV, therefore, exhibits a semiconductor nature. In 1954, Rupprecht [109] have synthesized In2O3 polycrystalline films by deposition indium metal on a quartz substrate at high temperatures in air. He found the

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optical absorption at 3.55  eV, whereas, Weiher and Ley [110] observed a weak absorption peak at 2.6 eV in single crystalline In2O3 plates prepared via vapor transport technique. The bandgap of In2O3 was observed to be 1 eV confirmed through XPS analysis, optical experiments, and ARPES [111]. For bcc-In2O3, the direct bandgap of 3.1 eV was confirmed by Fuchs and Bechstedt [112] using DFT and many-body perturbation theory (MBPT). The difference between the experimental bandgap and theoretical value is due to life time, excitonic effects, and phonon broadening [113].

3.7  Tin(IV) Oxide (SnO2) In the optoelectronic and related applications, the wide-bandgap semiconductor SnO2 has garnered considerable interest. It has a crystal structure of tetragonal for rutile type [114]. SnO2 is a multifunctional semiconductor oxide, which can be used in many areas, for example, catalysis, chemical sensors, solar cells, lithium-ion batteries, etc. [115]. In addition to this, SnO2 is a promising candidate for photodetectors and ultraviolet light-emitting diodes (LED) [116] with 3.6  eV bandgap. The white powder of SnO2 having the 7.0096  g/cm3 of density with 2000  °C highest melting point [117]. SnO2 part of the space group of P42/ mnmorD144h (SG136) for rutile crystalline forms within ambient conditions [118]. In this structure, Sn atoms are located at the Wyckoff 2(a) sites (0,0,0) and (1/2,1/2,1/2), whereas, oxygen atoms are located at the Wyckoff 4(f) sites ± (u  +  1/2;1/2  −  u,1/2) and ±  (u,u,0). The arrangement of this rutile-type crystal structure is marked as hcp array of the anion, for which cations occupy one-half of octahedral holes. Electronic and structural parameters that represent the behavior of SnO2 attracted attention to focus on examine these parameters in terms of theory and experiments. There are several methods implemented to investigate this property. The calculations

Fig. 3.11  Unit cell of SnO2 (Sn atoms represented in gray and O atoms represented in red color). (Reproduced from ref. [122] with permission, Springer, Copyright, 2020)

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were evaluated by using eight different potentials: LDA, GGA along with Perdew-­ Burke-­ Ernzerhof (PBE) functional [119], HSE03, LDA(HSE03), GGA(HSE03), HSE06, LDA(HSE06), and GGA(HSE06) adopted in the Cambridge sequential total energy package code (CASTEP) code [120]. Ultrasoft pseudopotentials for Sn and O atoms represented the ionic cores. The oxygen 2s2, 2p4 electrons with tin 5s2, 5p2 electrons were referred to as a part of the valence states. To optimize the structure of SnO2, the Brillouin-zone integration was carried out through the 5 × 5 × 8 and 3 × 3 × 4 for simple and hybrid potentials mesh sizes by using Monkhorst-pack technique.

3.7.1  Structural and Electronic Properties In SnO2, every anion is bonded by three cations, which are enclosed with six anion atoms [121]. Figure  3.11 represents the SnO2-simulated primitive unit cell. The calculations were carried out by using different potentials to provide accurate and reliable optimized structures along with ultrasoft pseudopotential methods. The structure is characterized by using volume and two lattice constant of 4.737 and 3.186 Å with angles α, β, and γ to be 90°, which are experimentally achieved by the ref. [123]. Various functionals including HSE03, HSE06, GGA, GGA(HSE03), GGA(HSE06), LDA, LDA(HSE03), and LDA(HSE06) were brought to reduce total optimize energy and optimal SnO2 primitive unit cell. The best value of the bandgap was found to be 3.603 eV as shown in Fig. 3.12b. The cutoff energy of 420 eV, this value is the best suited to the experimental value [124] as presented in Fig. 3.13b. The SnO2 energy band structure for all abovementioned fuctionals are shown in Fig.  3.14. According to experimental result, SnO2 is a material which is having direct bandgap as shown in Fig. 3.14 [125]. The Fermi energy level is indicated by zero on the axis of energy, while valence band is marked as an occupied state at below to Fermi energy level, and the conduction band represents as an unoccupied state at above to Fermi energy. The G-point of the Brillouin zone is defining location as the top of the valence band and bottom of the conduction. The LDA and GGA used to calculated bandgap were found to be far underestimated in comparison with the experimental and other functionals. The bandgap was found to be 3.459 eV for HSE03 method with the 280 eV optimal cutoff energy. Although, it was found to be 3.494  eV for HSE06, LDA (HSE06), GGA (HSE06), LDA (HSE03), and GGA (HSE03), which is closed to the reported values in comparison with GGA and LDA functionals. This is may be due to LDA and GGA are ground-state theory and the energy gap referred as the property of excited state [125]. Among all the functional mentioned here, the only LDA (HSE06) functional shows close to ideal value that found by experimental data.

Fig. 3.12  Lattice parameters, bandgap, cell volume, and their (%) errors by LDA (HSE06) functional observed values. (Reproduced from ref. [122] with permission, Springer, Copyright, 2020)

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Fig. 3.13  Lattice parameters, bandgap, cell volume, and their (%) errors by GGA (HSE03) functional observed values. (Reproduced from ref. [122] with permission, Springer, Copyright, 2020)

3  Structural and Electronic Properties of Various Useful Metal Oxides 69

Fig. 3.14 SnO2 structure for different functional: (a) LDA, (b) GGA, (c) HSE03, (d) HSE06, (e) LDA (HSE03), (f) LDA (HSE06), (g) GGA (HSE03), and (h) GGA (HSE06). (Reproduced from ref. [122] with permission, Springer, Copyright, 2020)

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3.8  Zinc Oxide (ZnO) Special attention is given to the zinc oxide clusters, which are very remarkable in their interesting kind of chemical and physical properties, such as amphoteric chemical properties, semiconducting properties with a wider bandgap, piezoelectric properties, anisotropic crystalline structure, biocompatibility, and higher exciton energy [126, 127]. There are various investigation on ZnO in order to understand the structure, formation process, behavior, and properties of nanoparticles [128, 129]. Nanostructure ZnO thin films become promising candidates in the application of switches, UV emitters, ultrathin displays [130, 131], and gas sensors [132]. Dmytruk et  al. [133] and Zhao et  al. [134] have proposed that the core-cage structure for (ZnO)34 is most stable, while Wang et al. [135] have showed that the structure with hollow cage produced by (ZnO)3 hexagons and (ZnO)2 squares is stable. It was observed that, in case of (ZnO)60, vigorously proposed sodalite motif, whereas nested cage structure was found to be more stable by refs. [133, 134]. Here, the study structural and electronic properties of nanoclusters (ZnO)n with n is 34, 60 are done by DFT approach. To study the properties of ZnO structures with nanoscale, such as nanowires and nanotubes, the ab-initio calculations within DFT were successfully carried out [124]. For structural models, the geometrical optimization (determining the ion equilibrium coordinates, where the system’s full electronic energy is minimum) was carried out, which is estimated by using delocalized internal coordinates effective algorithm [136]. The GGA in a parameterization of Perdew-Burke-Ernzerhof functional was used for formulation the exchange and correlation energy of the electronic subsystem [73]. As it is well known that the calculation results of GGA approach is in the underestimation of the numerical values of the binding energy. The approach LDA is an alternative approach for exchange-correlation interaction, which leads to overestimation in the values of binding energy in comparison with the experimental results. If the calculated results state the cluster model is stable, then it is possible to say that the real system will also be stable by using GGA study. The electronic functions of electrons are categorized based on atomic orbitals, containing d-orbitals. The core electrons can be described by using potential-­ concerning relativistic corrections. The integration in the first Brillouin zone was conducted in the Monkhorst-pack k-point set [137].

3.8.1  Structural and Electronic Properties The structure of ZnO possesses one-dimension (1D), two-dimension (2D), and three-dimensional (3D). Structure 1D exists in the different shape such as nanorods, nanohelices, nanoneedles, nanowires, nanotubes, nanorings and springs, nanobelts, nanoribbons, and comb-like structure [138], whereas, nanoplate and nanosheet are the examples of 2D structures of zinc oxide [139]. The 3D nanostructure of ZnO

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includes flower, snowflakes, coniferous three-dimensional ordered macroporous (3DOM) structures, etc. [140]. The crystal structures of ZnO are wurtzite (B4 type), zinc blende (B3 type), and rocksalt (B1 type). The wurtzite structure shared by ZnO is thermodynamically stable at ambient conditions. This structure has hexagonal form, related to the space group of P63mc, which is characterized into two sublattices, that is, Zn2+ and O2− interconnected to each other in such a way that every Zn ion is surrounded by tetrahedral of O2− ions and vice versa [141, 142] The zinc-blende structure of ZnO may be stabilized by growth on cubic substrates, while rocksalt structure can be obtained comparatively at high pressure. In both the structures, that is, zinc blende and wurtzite, the bond distance of four closest neighborhood atoms and next 12 closest neighbors is same because of the tetrahedral coordination. The only difference between these two structures is the arrangement of atoms in the closed-pack planes. The wurtzite ZnO structure contains triangular arrangement of alternating bi-atomic closed-pack (0001) planes along with (0001) direction, whereas zinc-blende ZnO structure consists of organized triangular atoms in the closed-pack (111) planes in the (111) direction. The crystal structure of ZnO reveals the crystallographic polarity, which shows the direction of bonds, that is, closed-packed planes (0001) in the wurtzite structure, and (111) planes in rocksalt −− as well as zinc-blende structures are contradict from  000 1  and (1 1 1 ) planes, respectively. The polar Zn-terminated (0001) faces, which is used to have various chemical and physical properties, and O-terminated (000.1) faces (along c-axis) possess a lightly changed in electronic structure [143]. The properties of ZnO such as spontaneous polarization, piezoelectricity, crystal growth, plasticity, defect generation, etching etc. depends upon the polarity. If the bonds along with c direction are from Zn ions to O ions, then the polarity is known as Zn polarity, and when the bonds are along with c direction, that is, from O ions to Zn ions, then polarity is said to be O polarity [144]. Several isomers examined to find out most stable structure for magic-clusters (ZnO)34 and (ZnO)60. Among them, cage and hollow fullerene-­ like structures follows the six isolated quadrangles rule. The (ZnO)34 nanoclusters found to be most stable structure due to fullerene-like hollow structure as it satisfies the rule of six isolated quadrangles. For the (ZnO)60 nanoclusters, various types of isomers, along with hollow and sodalite-like structure consisting of nanoclusters (ZnO)12, were examined. The most promising structure was found to be sodalite-­ type structure, which consists of seven (ZnO)12 clusters with shared quadrangle edges. Both the structures having LUMO level and HOMO for (ZnO)60 and (ZnO)34, respectively [145]. In the year 1969, many experimental efforts have been done on wurtzite structure of ZnO to prove the bulk electronic structure predicted by Rössler. The experimental data on the electrons energy levels in ZnO via X-ray-induced photoemission spectroscopy was reported in ref. [144]. From the experiment, it is concluded that the definite position of the Zn 3d level in ZnO can be measured, and the difference between true and measured energy value depends upon angular momentum. Powell et  al. [146] have located the Zn 3d level below valence band maximum, that is, at 7.5 eV via UV-photoemission estimation on ZnO hexagonal cleaved in vacuum, which is 3 eV less than the prediction by Rössler’s band calculation, whereas, Zn 3d level at 8.5  eV and 8.81  eV was reported by using X-ray photoemission [147]. Another

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method tight-binding was also employed to find out the position of the 3d level of Zn in ZnO [143], however, the exact position of the Zn 3d states could not be predicted. Girard et al. [148] have observed two surface states at (0001) surface by using offnormal ARPES spectra. First state was observed at 7.5 eV, which was predicted by theoretical calculation and the second one at 4.5 eV below Fermi level, which was not predicted by theory. The electronic band structure of various phases of ZnO also studied [149] Jaffe et al. [150] uses GGA form (both the LDA and the PBE96) with optimized Gaussian basis sets to measure the electronic structure of ZnO in zinc blende, wurtzite, and rocksalt crystal structures. The peaks was observed at above the upper valence band and little decrease in the height with shifted down in energy, the Zn 3d-derived and O 2s peaks are little broader with shifted up in energy, and the separation shows in the 3d states of Zn, when wurtzite structure of ZnO is compressed. While in rocksalt structure at pT1, that is, transition pressure, remarkable changes were observed, the maximum peak near the valence band is significantly reduced in height, and the Zn 3d peaks become thinner and drop lightly in energy. The partial states densities from the sharing of various orbital components for nanocluster (ZnO)34 with conduction and valence band are represented [145]. Where each cluster’s valence band within the −7.0 and −4.0 eV consists mostly from 3d and 2p states of Zn and O, respectively. While the valence band within −4.0 and 0 eV consists of 3d and 2p states of Zn and O, respectively, also 3p and 3s states of Zn for small scale. Whereas the conduction band within 1 and 5 eV contains mostly from 4s states of Zn and 2p and 2s states of O [145]. Similarly, partial states densities from the contributions of various orbital components for (ZnO)60 nanocluster conduction and valence band are also represented, where each cluster shows same states of Zn and O as of (ZnO)34 nanocluster for valence and conduction bands [145].

3.9  C  opper (I) Oxide (Cu2O), Copper (II) Oxide (CuO), and Copper Dioxide (CuO2) Nanoclusters The copper combines with the oxygen to form different oxides such as cuprous oxide (Cu2O), cupric oxide (CuO), and copper dioxide (CuO2), also their combination in several phases like Cu4O3. Copper (I) oxide (Cu2O), also known as cuprous oxide, is a red powder with p-type semiconductor material; having promising application in the agriculture, and in many fungicides as an ingredient to save it from the fungal diseases. It can also be used in ceramic glazes, in antifouling paints, as catalyst, photocatalytic, and photovoltaic applications [151, 152]. Copper (II) oxide (CuO), also called as cupric oxide, which can be used as pigment by adding in to clay glazes [153] and used as abrasive for polishing optical components and lenses [154], while copper dioxide (CuO2) can be used in superconducting lattice [155, 156]. The thin film or nanoparticles of CuO, CuO2, and Cu2O synthesis via different techniques such as thermal evaporation, RF sputtering, chemical vapor deposition, laser ablation, chemical bath deposition, spray pyrolysis, electrodeposition, and sol-gel method [157, 158]. The surface morphology of nanoflakes, nanorods, and nanotube

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could be synthesized by appropriately controlling mechanisms [159, 160]. The formation of nanoclusters can be done by capping the material within surfactant or by synthesizing material in the plane-oriented substrates. By using NWChem package, the possible structures of Cu2O, CuO, and CuO2 clusters were optimized completely [161]. This package is used to simulate the complicated chemical and biological structures in a large-scale calculation. The Hartree-Fock calculation utilized by conventional methods, performed with Fock build. In DFT calculation, the computation and energy convergence done by algorithm (Kohn-Sham matrix element) of NWChem. The dynamic correlation in DFT is used to find out the motion of the discrete electrons, while electron correlation is addressed through

Fig. 3.15 Optimized structures of CuO, CuO2, and Cu2O clusters. (a) CuO-1. (b) CuO2-1. (c) Cu2O-1. (d) CuO-2. (e) CuO2-2. (f) Cu2O-2. (g) CuO-3. (h) CuO2-3. (i) Cu2O-3. (Reproduced from ref. [165] with permission, material Science, Kaunas univ. of technology, Copyright, 2020)

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exchange energy, which appears from the antisymmetry of quantum mechanical wave function. DFT calculation method includes the pseudopotential approximation, which is replacement of effects of the bound electrons in the atom of the cluster that modifies Schrödinger equation’s potential terms. DFT is utilized by hybrid B3LYP exchange-correlation with 6-31G basis set for CuO, CuO2, and Cu2O clusters [162– 164]. The atomic number of oxygen and copper are 8 and 29, respectively; 6-31G is appropriate basis set to improve the nanoclusters.

3.9.1  Structural and Electronic Properties Different structures of copper oxide (CuO, CuO2, and Cu2O) clusters are shown in Fig. 3.15. The CuO-1 structure shows the Cu and O atoms alternatively attached, which is having hexagonal pattern, CuO2-1 structure has hexagonal structure with two atoms of Cu and four atoms of O, while in Cu2O-1 also has hexagonal structure with two atoms of O attached to four atoms of Cu. In case of CuO-2, CuO2-2, and Cu2O-2, they exhibit two hexagonal layers with total 12 atoms. In case of CuO-3, CuO2-3, and Cu2O-3, they exhibit hexagonal structure, which is connected to the neighborhood cluster and forms a bee’s hive-like structure. Among the CuO-1, CuO2-1, and Cu2O-1 structure, Cu2O-1 has the energy of −6711.371 Hatrees, which is found to be stable as compared to other three structures. Maximum dipole moment was found to be 2.335 Debye with energy of −3590.941 Hatrees for CuO-1 because of asymmetricity of the charges in the both atoms (Cu and O), whereas least dipole moment was found to be 0.603 Debye for Cu2O-1, due to symmetry of the charge for all the three clusters. Similarly Cu2O cluster is more stable than the other structures, which has energy of −13,420.90 Hartrees. The maximum dipole moment was found to be 3.126 Debye for CuO2 structure and least dipole moment of 0.505 Debye for CuO-2 for all the cluster. Whereas, the energy increased to −13,567.948 Hartrees for Cu2O-3, which is stable, and CuO2-3 is least stable. The maximum dipole moment was found to be 3.2664 Debye for CuO-3. And low value of dipole moment found to be 0.0045 and 0.0003 Debye for CuO2-3 and Cu2O-3, respectively. From the DOS spectrum [165], it was observed that for CuO-1, the equal atoms of Cu and O are in the cluster, which results in overlapping the orbitals. In CuO-1 and CuO-3, the α and β orbitals are observed, whereas in the CuO-2, no alpha and beta orbitals are seen, as the atoms of Cu and O with double layer balance the spin of the electrons. The DOS spectrum provides understanding of the charges density in the occupied and virtual orbitals. DOS spectra and HOMO-LUMO gap of CuO, CuO2, and Cu2O clusters are shown in Fig. 3.16a–i. Figure 3.17a–c shows electron affinities (EA) of CuO, CuO2, and Cu2O clusters, respectively. The higher value of EA indicates the changes in the energy as soon as an electron gets added into cluster. Among the all clusters, Cu2O is found to have a higher value of EA because of maximum number of Cu atoms in the clusters. Cu2O clusters are promising candidate for the application in chemical sensor because of

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Fig. 3.16  DOS Spectra and gap of HOMO-LUMO level for CuO, CuO2, and Cu2O clusters. (a) CuO-1. (b) CuO2-1. (c) Cu2O-1. (d) CuO-2. (e) CuO2-2. (f) Cu2O-2. (g) CuO-3. (h) CuO2-3. (i) Cu2O-3. (Reproduced from ref. [165] with permission, material Science, Kaunas univ. of technology, Copyright, 2020)

5.4

4.0

EA

IP

3.8

5.2

3.6

5.0

3.4

4.8

3.2

4.6

3.0

4.4

2.8

4.2

2.6

4.0

5.0

4.5

3.0

CuO2 Cluster

Ionizaton Potential (eV)

2.6

4.0

IP

2.4

3.5

2.2 EA

3.0

Cu2O

(c)

2.8

2.0

2.4 CuO

3.2

CuO-2

CuO2-2 Cluster

Cu2O-2

5.0

5.5

4.5 5.0

4.5

4.0 IP

3.5 3.0

4.0

EA

2.5

Electron Affinity (eV)

5.6

(b)

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Electron Affinity (eV)

4.2 Electron Affinity (eV)

4.4

5.8 Ionization Potential (eV)

(a) 6.0

Ionization Potential (eV)

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2.0

3.5 CuO-3

CuO2-3

Cu2O-3

Cluster

Fig. 3.17  Variation of IP and EA of: (a) CuO-1, CuO2-1, and Cu2O-1 clusters; (b) CuO-2, CuO2-2, and Cu2O-2 clusters; (c) CuO-3, CuO2-3, and Cu2O-3 clusters. (Reproduced from ref. [165] with permission, material Science, Kaunas univ. of technology, Copyright, 2020)

rapid change in energy [166, 167]. The EA of various nanoclusters also deal with the presented findings of CunOn for n = (1–8) clusters [168] and Cu2Ox for x = (1–4) [169]. The ionization potential (IP) describes the energy required to separate the electron from the cluster. Among all the clusters, higher value of IP was found to be 5.83, 5.04, and 5.22 eV for CuO-1, Cu2O-1, and Cu2O-3, respectively.

3.10  Conclusion A comprehensive study on structural and electronic properties of various metal oxides such as titanium dioxide (TiO2), indium oxide (In2O3), tin (IV) oxide (SnO2), zinc oxide (ZnO), copper oxide (Cu2O, CuO), and copper dioxide (CuO2) has been presented here. Metal oxides have been widely studied material over a decade, for better understanding of the properties and to develop thin film for the wider range applications. These applications highly depend upon the morphology and physical properties of the material chosen. In short, this chapter provides the details of

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different techniques, functionals, and theories used, with a view to understand the structural and electronic properties of various metal oxides.

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Chapter 4

Properties of Metal Oxides: Insights from First Principles Calculations Assa Aravindh Sasikala Devi and D. Murali

4.1  Introduction Metal oxides are important class of materials, owing to the complex nature of metal-­ oxygen bonding, and have drawn research interest in recent years, owing to their application in dielectrics, magnetism, ferroelectricity, superconductivity, light emission, optical spectroscopy, etc. to name a few [1, 2]. Mostly, metal oxides possess wide bandgap and exist as conductors, semiconductors and insulators, and hence, they are used for applications in sensors, data storage, catalysis, energy, optical and displays. To understand the metal oxide device properties in detail, it is necessary to estimate the structural details, dielectric properties, optical characteristics, and luminescence. In this scenario, ab initio methods, such as density functional theory (DFT), have been very successful in modeling and predicting the properties of metal oxides. In first principles methods, the atomic number and structural parameters are taken into consideration and are devoid of empirical parameters [3, 4]. The geometry, bonding energies, dipole moments as well as electronic properties of metal oxides are calculated successfully using DFT. Different oxide materials, such as (La,Ba)2CuO4, YBa2Cu3O7 − δ, Bi2SrCu2O8, and HgBa2Ca2Cu3O8 + δ, are under investigation since long time [1, 2]. DFT calculations could provide important insights into the complexity of structure, shape of Fermi surfaces, etc. Further, DFT calculations could reveal the nonoxide like electronic structure of iron-pnictide superconductors, which was useful in understanding the superconducting properties [5, 6]. A. A. Sasikala Devi (*) Nano and Molecular Systems Research Unit, University of Oulu, Oulu, Finland e-mail: [email protected] D. Murali Faculty of Physics, IIITDM, Kurnool, India © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_4

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Most of the oxides are magnetic and present interesting electronic properties. Some of the prominent oxide materials are lanthanide codoped oxide Y2O3 with body-centered cubic structure, YBa2Cu3O7−x (YBCO), KNbO3, cobalt oxide (Co3O4), manganese oxide (MnO2) as well as nickel oxide (NiO) [7]. MnO2, Co3O4 and NiO are commonly used materials to enhance the electrode action of lithium-­ ion batteries [8, 9]. Some examples of underestimation of bandgaps are that of manganese oxide and nickel oxide, wherein standard DFT approaches, such as local-density approximation (LDA) and generalized gradient approximation (GGA) reported a bandgap that is much lower than the experimental value. Furthermore, metallic and magnetic properties are also reported wrongly, and well-known examples include cobalt oxide and iron oxide that were incorrectly described as metallic and ferromagnetic [10]. When it comes to mixed oxides of Mo and V, that were successfully employed for gas phase catalysis [11] also, standard DFT is not sufficient to describe the electronic properties. Figure 4.1 shows a comparison of density of states calculated for the oxides MnO, NiO, CoO, and FeO, where the change in electronic properties is visibly different with different functional [12]. The bandgaps of these materials also were

Fig. 4.1  Comparison of total density of states (DOS) in various DFT and beyond DFT methods. The panels labelled a, b, c and d indicate DOS of MnO, NiO, FeO and CoO respectively. (Reproduced from ref. [12])

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Fig. 4.2 Comparison of band structures and spectral functions computed in B3LYP and eDMFT. The panels labelled a, b, c and d represent band structures of MnO, NiO, FeO and CoO respectively. (Reproduced from ref. [12])

clearly improved [12]. Visible differences in band structure were also noticed with the use of different functional, as shown in Fig. 4.2. These days, tremendous advancement has taken place in metal oxide research with the emergence of machine learning and evolutionary algorithms that can predict accurate crystal structures [13]. For example, the CRYSTAL software, wherein the algorithm-based USPEX code is implemented (see Fig.  4.3) incorporating hybrid functional, could predict accurate crystal structure of complex transition metal oxides [13]. Figure 4.4 shows the energetically most stable structure predicted using USPEX algorithms [13]. To exploit the metal oxides to their full potential for application purposes, the effect of surfaces, interfaces, intrinsic and extrinsic defects must be included. In this scenario, theoretical modeling of the physical phenomena at the interface of oxide surfaces poses many challenges. The description of electronic properties of the chemically diverse semiconductors-electrolyte surface using a single, nonempirical quantum mechanical approach is a complex issue [14]. Another problem is to deal with the multiscale time and energy processes like light absorption. This will be accompanied by generation of charge carriers in the semiconductor, and these phenomena will occur in the femto seconds time frame and the typical energies are

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Fig. 4.3  The lowest energy unit cell of NiO predicted using USPEX code. (From ref. 13)

Fig. 4.4  The crystal structure prediction workflow of USPEX code

in electron volts. Further, the charge transfer at the solid-liquid interface, owing to the atomic motion at room temperature that occurs at the milli electron volt energy scales and time scales of pico to nano seconds. To study such large heterogeneous systems, advanced computational techniques methods, such as many-body perturbation theory, are employed, which often comes at a higher computational cost. Hence, a number of methods are often employed, starting from the very basic

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DFT methods and ranging to perturbation methods, such as GW approximations [15] or hybrid functional, which itself encompasses various flavors and approximations. The hybrid functional is less computationally demanding than full many-body approaches. A large number of potentially active catalyst materials are identified, such as MoVWO and MoVNb(Sb,Te)O systems, which have been successfully used for selective oxidation purposes. The DFT  +  U approach has proven to be a viable approach, considering the computational cost in eliminating the self-interaction correction [15]. In this method, an empirical correction is introduced that penalizes the electron localization arising out of the self-interaction correction. Some examples in which DFT + U has been successfully produced the desired results are the Pt/CeO2 system, where in the catalytic action occurs at the water-gas shift reaction [16]. Another instance is the oxidation catalysis involving propene activation [17]. These catalytic reactions involve formation and filling of the O vacancy defect sites that often contain localized electrons. However, DFT is an exact theory, the approximations considered to describe the exchange and correlation functional can make vast changes in calculating properties of these materials. The Kohn-Sham DFT, including functional to describe the exchange and correlation interactions, is the standard workhorse of computational catalysis involving solids. These simple functional, namely LDA and GGA, are successful in predicting and reproducing the properties of most materials, such as magnetism, high-pressure phases, Fermi surfaces, etc. However, in the case of strongly correlated systems such as oxides, simple LDA and GGA fail to reproduce experimental results. These functional often have the tendency to generate artifacts as a consequence of the erroneous description of localized electron configurations [18]. This phenomenon is due to the infamous self-interaction error (SIE) that exists in most local and semi-­ local formulations of the exchange–correlation (XC) functional [19]. These spurious interactions will get reflected in misinterpretation of electronic structures and inaccurate energetics, underestimation of highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) gaps [20] and inappropriately quenched spin densities. Some instances are that of the ground state of Mott insulators and underlying physical phenomena associated with the electron correlations responsible for these ground states. DFT often underestimates the bandgap and fails to predict if the metal is a metal or insulator. Nevertheless, if all the metallic centers of a material are in the highest oxidation state, devoid of localized orbitals of d and f electrons, standard LDA and GGA would be sufficient to reproduce the accurate electronic structures. Moreover, the excited-state properties are also often beyond the scope of DFT. In recent years, several methods like the Hubbard parameter included in DFT (DFT + U), hybrid functionals, meta-­ GGAs, GW, and DFT-embedded dynamical mean-field theory (eDMFT) have been developed to explain the electronic structure of correlated materials. Nevertheless, the lack of clarity about the accuracy of these methods in describing different strongly correlated oxide materials still persists. While attempting to estimate the accuracy of first principles calculations in reproducing experimental data, it is often unclear about what experimental data could be used as a reference. It can also

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happen in such a way that certain properties calculated experimentally are not directly related to the values calculated theoretically. Two important properties often pointed out are excitonic characteristics and the electron–phonon coupling. However, standard electronic structure calculations often don’t take these effects into consideration. Conversely, fundamental quantities like the optical gap are experimentally determined by extrapolation of measured data assuming specific theoretical models. This indicates that depending on the property under consideration, appropriate theoretical tools must be selected that gives reasonable agreement with experimental results [21].

4.2  An Example System: BaTiO3 Though recent years showed increased interest in solar energy harvesting, based on semiconductor materials, further advancements were often hindered by the problem of reduced open circuit voltage (VOC) in comparison to the bandgap [22]. In the search of new materials with enhanced efficiency and photovoltaic (PV) properties, ferroelectric oxides are identified for their exemplary PV effect and voltages higher than the bandgap [23]. Especially, the ABO3 perovskites find applications in sensors, solid oxide fuel cells, heterogeneous catalysis, etc. to name a few. Here, A represents a monovalent or divalent cation, whereas B corresponds to a transition metal with penta or tetravalency. The perfect perovskite crystal structure should be a cubic lattice, and the structure of a unit cell is presented in Fig. 4.5.

Fig. 4.5  The cubic phase of BaTiO3. The green, red and blue atoms represent Ba, O and Ti, respectively

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However, in reality, this cubic symmetry (paraelectric phase) will be distorted to a ferroelectric structure with reduced symmetry at lowered temperature. Hence, ferroelectric perovskites with noncentrosymmetric structure exhibit larger compared to the bandgap, attract enormous attention for prospective photovoltaic applications. One such candidate is LiNbO3 single crystal, which when doped with Fe showed giant PV response, and this superior PV feature is broadly attributed to ballistic and shift current phenomena [24]. Prominent ferroelectric materials under investigation are (Pb,La)(Zr,Ti)O3, BaTiO3 (BTO), PbTiO3, (K,Ba)(Nb,Ni)O3, Bi(Fe)O3, Pb(Ti)O3, Pb(Zr)O3, and KBiFe2O5 mainly due to their above bandgap PV, which used to be a hindrance for using semiconductors for PV applications. This observed PV is attributed to the interband carrier transitions and, therefore, depends on the degree of delocalization of valence band and conduction band states, determined by the strength of covalent bonding [25]. The PV response will be enhanced when the orbital states constituting the conduction band edge align with the polarization direction, as there will be large shift current. For example, (K,Ba)(Nb,Ni)O3  −  δ, with rhombohedral geometry, in which the orientation of polarization happens to be directed on the body diagonal, showed degenerate t2g orbitals at conduction band edge, resulting in enhanced shift current response. However, for the tetragonal geometry, where the polarization is oriented toward the z direction, the xy orbital constitutes the conduction band edge, resulting in reduced shift currents [26]. Among the ferroelectric ABO3 perovskite family, the tetragonal BTO, which was discovered in 1941 [27], exhibits low PV response owing to the d orbital of Ti, which forms the conduction band edge states [24]. For BTO, which is a standard 2:4:2 type perovskite material, the transformation from the paraelectric cubic to ferroelectric phases with lower symmetry occurs at 130 °C. When the temperatures are reduced, two phase transitions will take place accompanied with displacements and macroscopic strain occurring along the ferroelectric orientation. Apart from ferroelectricity, BTO also possesses outstanding properties, such as high-voltage tunability, pyroelectricity, dielectric permittivity and piezoelectricity [28, 29]. DFT studies could successfully reproduce most of the experimentally observed properties of BTO. Chen et al. had calculated the crystalline lattice parameter of 3.99 Å and a dielectric constant of 5.74 for the BTO in cubic crystal structure, and phonon frequencies that are in good agreement with experiments [30]. They also found that Ti displacements could induce ferroelectric phase transition in BTO, and during this phase transition, Ba atoms gain and Ti lose charge, whereas the charge reduction of one O atom is compensated by the charge gain of the other. Their calculations also found that the Ba-O bond possess covalent nature, and the covalent interactions between Ti 3d atomic orbitals and O 2p atomic orbitals play crucial role in stabilizing the ferroelectric, tetragonal phase of BTO. BTO can also be utilized as a semiconductor having positive temperature coefficient of resistivity with sufficient doping, as BTO samples with lowered resistivity are suitable for applications in conducting powders. Hence, research has been conducted focusing on the conducting properties of BTO with various dopant elements like Dy, Y, Er, Ho, Sm, Zn and Ag [31, 32]. Certain codoping pairs such as La and Ag have been found to be favorable for lowering the resistivity significantly [33–35].

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Fig. 4.6  The right panel shows the partial density of states of the T and O phases of BST7 in the vicinity of CBM. The middle panel shows the band dispersion of T, and the right panel shows the band extrema and along z-direction for the O phase. (From ref. [38])

It is seen that the PV properties of BTO are shown to be tuned by compositional change induced by doping, and subsequent symmetry breaking. Bulk photovoltaic effect (BPVE) has been shown in BTO single crystalline samples with higher power conversion efficiency than the Shockley-Queisser limit [36], owing to the shift current phenomena. As an after effect, interband transition of charge carriers occurs in noncentrosymmetric systems, followed by momentum shift of real-space vectors, giving rise to large PV response [37]. DFT studies have played an effective role in explaining these behaviors [38] by analyzing the consequences of symmetry breaking on electronic properties and subsequent change in PV response. Orthorhombic (O) as well as tetragonal (T) structures of BTO were studied, and the density of states calculations proved that the O phase possesses mixed character of dxy and dxz/yz orbitals, and well-separated dxy and dxz/yz states are shown in the T phase (Fig. 4.6). Change in the crystal field and subsequent enhanced hybridization among the dxy and dxz/yz orbitals are observed in O phase, leading to a combination of characteristics of these orbitals in the CBM. These characteristics observed in electronic structure could successfully explain that the O phase showed superior PV response than the T phase. The PV properties of the material are also influenced by the motion of the photo-­ generated charge carriers, as increased mobilities lead to suppression of charge carrier recombination, increasing the PV response. Therefore, an important property deciding the ferroelectric properties of materials is the effective mass in the direction of flow of current. By employing DFT simulations, the direction-dependent effective mass of the material could be calculated, which is inversely related to motion of charge carriers. Since systems with smaller effective masses possess more mobilities, they exhibit enhanced PV responses. In addition to this, an

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anisotropy occurs for the effective masses owing to the direction-dependent band dispersion in the reciprocal space. The right panel of Fig. 4.6 shows the effective mass calculated for the Z direction, which also happens to be the direction of flow of current and is the reason behind the superior PV performance of BTO. The large PV response of BTO is attributed to the breaking of symmetry, and subsequent transformation of structure from T to rhombohedral and O geometries is modeled using DFT studies successfully [39] that showed that this symmetry breaking transformation reduces effective masses of electrons and holes. This causes delocalization and reorientation of conduction charges along the axis of polarization, thus enhancing the PV performance. Figure  4.6 shows the band dispersion of T, O and rhombohedral phases centred on CBM and VBM along the Z-direction in the k-space. It is seen that the dispersion for T phase is very flat at the CBM. Whereas, the O and rhombohedral phases show large dispersion, indicating the smaller effective masses for these two phases. Another lead-free ferroelectric material that possesses large piezocoefficient and subsequent energy storage capacity is 0.5Ba(Zr0.2Ti0.8) O3 – 0.5(Ba0.7Ca0.3)TiO3 (BZT-BCT). DFT has been successfully used to calculate the effective mass of BZT-BCT [39]. The results indicate that the rhombohedral and O phases show improved PV response than the T phase. The smaller effective mass along polarization direction for these two phases indicated that the states in the vicinity of the Fermi level are less localized, thus reducing the charge recombination rate, resulting in high-PV response. Different doping species and concentrations have been introduced to the BTO system to see the influence of the dopant in affecting the band structure and further possibility of increasing the PV response. This is possible by the effect of the dopants in delocalizing the CBE, formed by the orbitals, which lie along the orientation of polarization. Two favourable elements at the A-position are Bi- and Pb-forming covalent bonding with the p orbitals of O, and thus subsequently delocalizing the states surrounding the Fermi energy [40]. Even though due to the toxicity of Pb, Bi is a more preferred dopant when Bi3+ is used as dopant at the Ba2+ position; another co-dopant with single valency is required to compensate the charge. In this scenario, Li is a desirable candidate, as its ionic nature will not influence the electronic structure significantly. BaTiO3 substituted with Bi and Li at equiatomic proportions at minimal doping concentration showed highest value of photo-generated field ever reported in any bulk polycrystalline noncentrosymmetric system [41]. Further, Bi and Li when doped together to BaTiO3 have shown that increase in lattice polarization accompanied with delocalization of the conduction band occurs, leading to giant PV response. BTO with proper codoping can be used as semiconductor with positive temperature coefficient of resistivity. Since standard DFT methods often fail to describe the strong electronic correlations, arising due to the d electrons, spin-­ polarized DFT calculations, including the Hubbard U parameter (DFT  +  U) conducted on tetragonal BTO crystal, doped together with La and Ag, including intrinsic points defects, such as O vacancy, showed superior electronic and electrical properties [42]. In this work, the rotationally invariant GGA + U formalism is used

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[43], and for Ti 3d electrons, U value of 3.5 eV is used [44]. While for 4d electrons of Ag [45] and 5d electrons of La [46], a U value of 6.0 eV is used. The inclusion of Hubbard parameters substantially increased the accurate prediction of electronic structure of this material, and it has shown that there exists the possibility of a reduction in bandgap with doping, apart from the defect complex consisting of La atom and oxygen vacancy. It turned out that the inclusion of O vacancy proved to be useful for the narrowing of bandgap, indicating the possibility of application in photocatalysis. This study also predicted that in the absence of intrinsic oxygen vacancy in the material, Ag doping alone induces p-type conductivity. This finding indicated that the observed n-type conductivity in experiments [47] with Ag doping could be arising out of the presence of intrinsic oxygen vacancies, which are unavoidable defects. They also suggested codoping of La with another element other than Ag if substantial reduction in resistivity is needed for application purposes. Furthermore, codoped BTO containing intrinsic oxygen vacancy showed conductivities about 6% higher than purely La-doped BTO and about 15% more with only Ag dopant. Therefore, this study emphasized the need to use codoping as an effective strategy to manipulate the conducting properties of BTO. Apart from the bulk BTO crystal, thin films were also investigated using DFT methods, revealing significant changes in properties with change in dimension. DFT-B3LYP simulations of BTO polar surfaces with (111) orientation, performed using the CRYSTAL code, have shown that covalent character of the Ti-O bond is increasing considerably at the surfaces, which eventually will play a significant role in the surface structure stabilities [48]. Therefore, it might be possible to tune the properties by changing the dimensionality. A DFT study involving the GGA has calculated the fundamental properties of BaTiO3 clusters and showed that all growth modes under investigation were found to consist of the similar structural units, comparable to the much studied (TiO2)2 clusters. Interestingly, the clusters containing one and two structural units of BaTiO3 have adopted equivalent geometries to the ground-state structure of (TiO2)n clusters with space groups Cs and D2h [49]. There were attempts to overcome the shortcomings of DFT by incorporating higher-level approaches, such as isotropic shell model employing the PBEsol functional, particularly aimed to reproduce the crystalline properties of solids [50]. Though this model could reproduce the ground-state properties and finite-­ temperature experimental properties, the temperature scale of the phase transitions came out to be compressed.

4.3  Summary To conclude, metal oxides are an important class of materials for many applications. Though there has been vast pool of experimental information available regarding metal oxides, insights gained from combining basic physics principles, advanced computational methods and artificial intelligence are their properties can be tuned to achieve better performance. DFT has been successful in predicting the structural

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parameters, choice of proper doping elements to tune the conducting properties, charge transfer, bonding nature and parameters, etc. Nevertheless, it is important to understand that, though first principles-based DFT has been successful in predicting and modeling properties of metal oxides to a great extent, the performance depends on the choice of functional. Dielectric-dependent hybrid functionals are known to accurately describe the defects in bulk phases, while the electronic structure of oxides is an ever challenging problem. The most important challenge to sort out is to identify a proper method with choice of parameters, hybrid functional or many-­ body method, which can describe the electronic structure accurately, to match with experimental findings.

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Chapter 5

Recent Progress in Metal Oxide for Photovoltaic Application Emmanuel O. Onah, Jude N. Udeh, Sabastine Ezugwu, Assumpta C. Nwanya, and Fabian I. Ezema

5.1  Introduction There is increase in global economics and human population. These have necessitated the great demand for energy supply, both now and in years to come. The shortage in fossil fuel reserve and impending global warming craves for immediate changes in world energy generation and utilization process. In this regard, sustainable energy resources, which are eco-friendly, cheap, clean, and easily practicable, serve as the bases for renewable energy system [1–3]. Solar energy is regarded as the practical source through which this energy can be generated, stored, and used at convenient time [4, 5]. The outstanding electrical, morphological, optical, electrochemical, magnetic, and catalytic properties of nanomaterials have been reported. Copper oxide, nickel oxide, titanium dioxide, gallium oxide, iron oxide, tin oxide, etc. are among the notable oxides used for these applications. Some of these metal oxides exist in different shapes and morphology. The different forms that have been reported E. O. Onah · J. N. Udeh Crystal Growth Laboratory, Physics and Astronomy Department, University of Nigeria, Nsukka, Nigeria S. Ezugwu Department of Physics and Astronomy, University of Western Ontario, London, ON, Canada A. C. Nwanya · F. I. Ezema (*) Crystal Growth Laboratory, Physics and Astronomy Department, University of Nigeria, Nsukka, Nigeria Nanosciences African Network (NANOAFNET), iThemba LABS-National Research Foundation, Somerset West, Western Cape Province, South Africa UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa (UNISA), Pretoria, South Africa e-mail: [email protected]; [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_5

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include nanowires, nanorods, stars, triangular, spherical, etc. These metal oxide nanomaterials have contributed immensely in several applications such as batteries, biosensors, chemical sensors, ceramics, solar cells, gas sensors, photocatalysis, and supercapacitors [6, 7]. However, their application as optically active materials in photovoltaic systems is of paramount important. This is because the optically active metal oxides harvest solar energy from the sun through photovoltaic process [2, 8]. The harvested energy can be stored for later use or be used immediately, as it is converted to electric current through viable processes. Pure metal oxides have been employed in trapping solar energy radiation. However, different modifications of the pure metal oxides for better efficient end product have been put into practices. This modification changes the structural, electrical, and morphological properties of the conventional metal oxides [9, 10]. Doping of metal oxides improves the conductivity and efficiency via reduction of the wide bandgap energy of the material and subsequent improvements of the absorbance of the materials. On the other hand, incorporating carbon derivatives into the sites of metal oxides is also important, as this increases the surface area of the material, thereby improving the efficiency of the cells [9, 11]. The metal oxides used for solar cell application exists in different forms. They can be organic or inorganic metal oxides depending on the constituents of the composing materials. The p-type and n-type metal oxide semiconductors are the common terms while discussing conventional solar cells for photovoltaic systems. The nature of the impurity infused into the system determines the functionability of these cells. Forming composites of different metal oxides also enhances the efficiency of the cells, since it involves the combination of two or more materials with different features [5, 12]. Some of these devices have low stability like perovskite solar cells (PSCs) and dye-sensitized solar cells fabricated with some metal oxides as active layers. Efforts have been made toward attaining high power conversion efficiency (PCE) and improving the stability of PSCs in order to facilitate its commercialization process. This solution-based cell lacks adequate crystalline nature, and thus possesses numerous grain boundaries and defects, which destabilize the device performances. When exposed to heat, moisture, light, or oxygen, the compositing materials degrade. As a result, inorganic ions are used to replace the organic contents that are unstable through compositional engineering, or the material can be crystallized by incorporating some functional additives. Moreover, another active way of ensuring lengthy stability of the PSCs is to employ interfacial engineering [13–15]. This chapter will delve into elaborate discussion of the metal oxides as active materials for photovoltaic applications.

5.2  Solar Cells for Photovoltaic Applications A solar cell, also known as photovoltaic cell, is a chemical or physical occurrence, where electrical device through a photovoltaic effect converts energy of the sun directly into electricity. The electrical features of a solar cell such as resistance, current, or voltage alters when the device is exposed to light [5, 10]. Aggregates of

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solar cells combined effectively in appropriate manner produces a solar module, which can be regarded as solar panels when properly arranged. The maximum open-­ circuit voltage (Voc) that can be obtained in a single-conventional silicon solar cell is 0.6 V approximately [16]. Artificial light or natural sunlight radiation can be used for the photovoltaic process. Photocells can be used to measure the intensity of light and to detect different paths of the electromagnetic radiation as photo detector. Three basic requirements in solar cell (photovoltaic) operations are: • The light absorption and electron-pair generation. • Charge carrier separation. • The extraction of the separate carriers to external circuit. Voltage and currents in solar cells are accumulated when solar cells are connected in series or parallel. However, possible loss in power and damages in solar cells can occur when the cells are connected and part of the cells staying under shade, thereby leading to reverse bias current. The surface of the well-connected solar cells (modules) is always sealed with transparent surfaces so as to allow solar radiation to pass toward the cells without diversion. The transparent sheets of glass prevent the module from rain, mists, fog, dust, and other atmospheric particle capable of striking the surface of the solar cells. Solar cells operate by absorbing energy from the sun when the radiation hits the surface of a photosensitive material, for example, silicon-doped semiconductor material. Absorption of solar radiation will adjust the state of the material so that the electrons in the material are excited. This excitation will either heat up the material and then return to its ground state or traverse through to an electrode, where it can easily be used for external work as electricity when the cells are connected in arrays or stored for later usage with the help of inverter. One layer of such silicon is usually doped with boron while the other layer is doped with phosphorous, and this produces different polarity from the former. The p-n junction type of silicon is considered the most commonly used solar cells. Other types of solar cell that have almost the same operation methods include perovskite solar cell, organic solar cells, quantum-dots solar cells, and dye-sensitized solar cells. These materials are optically active, and they allow light to pass through them. Conducting polymers, indium tin oxide, or nanowires are basically used as high-electrical conductive materials and high-­ transmitting materials in thin film of this type. Solar cell PCE is a measure of incident power transformed into direct electricity. The general efficiency of the solar cell is the value obtained when one multiplies the thermodynamic efficiency, the efficiency used during separation of charge carrier, conductive efficiency, and reflectance efficiency. These efficiencies are dependent on the temperature coefficients, voltage efficiency curve, and permissible shadow angle. The solar cell fill factor is a major determinant of solar cell efficiency. It is a measure of maximum power that can be obtained from a cell to the product of short circuit current and open circuit voltage. The maximum value of this parameter is unity, while the minimum value is zero. Thus, the closer the fill factor of a solar cell to one, the more efficient the solar cell, provided other operational parameters are

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working optimally. Cells with high fill factor losses less current, and they are characterized by having high shunt resistance, low series resistance, and low recombination. Theoretically, silicon cells have approached approximated 33.16% power efficiency for single p-n junction crystalline silicon [17]. However, estimated 25.6% PCE has been obtained in practical using silicon [18]. Solar cells are designed to absorb solar radiation reaching the earth, while some are designed for space use. A light absorbing material made of single layer can be used as solar cell and it is called single junction, while multijunction materials are those light absorbing materials with multiple configuration for easier absorption of light and separation of charges. Solar cells can be classified into three generation cells, namely: • First generation. • Second generation. • Third generation. The first-generation solar cells are made of crystals of silicon, and they occupy greatest percentage of photovoltaic market, since they show high conversion efficiency when compared to others. They can be referred to as wafer-based, traditional, or conventional solar cells. They are mainly made of monocrystalline and polycrystalline silicon. The monocrystalline type (mono-Si) produces better efficient cells, but this type of cell is very expensive. It has octagonal shape with small pattern of white diamond display on the assembled panels. Polycrystalline silicon is the most familiar type of silicon solar cells for photovoltaic application. They are less expensive when compared to monocrystalline silicon cells but on the contrary, their efficiency is relatively low. Second-generation solar cells are mostly amorphous silicon, CIGS, and CdTe. They are basically solar cells made into thin film. They are utilized in small-scale power station, small standalone photovoltaic systems. Their efficiency is lower than the conventional silicon solar cells. The third-generation solar cells are mostly recent emerging photovoltaic technologies that have little or no commercial recognition as a result of their limited PCE. They are made of thin film technologies, and they are still undergoing vigorous research process till date to ascertain their prospective attributes in terms of efficiency. Inorganic and organic or organometallic compounds are used to fabricate this type of solar cells. The absorbing materials for this type of cells are unstable, depriving the prospect of commercial application. However, research is going on to improve the technology, which center on using low-cost materials to fabricate high efficient cells. Conventionally, when solar energy impinges on the surface of a semiconductor material, electrons are migrated to the conduction band of the crystal structure from the valence band of the material. The outermost part of the crystallite material contains the electrons, and they are referred to as valence electron. The smallest amount of energy required to liberate electron from the shell is called the bandgap energy EG. Figure  5.1 shows the photon absorption process in a semiconductor, while Eq. (5.1) reveals the energy. The excited electron with Eph = hv migrates to the final energy level of the conduction band from the initial energy when energy is absorbed, and this causes the creation of hole at the valence band.

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Fig. 5.1  Absorption of photon by semiconductor material with bandgap Eg



= E E= Ei – Ef = hv ph

(5.1)

5.3  Solar Cell Output Parameters The characteristic solar cell is known with the following parameters:

5.3.1  Short-Circuit Current (Isc) The short-circuit current, Isc, is one of the major characteristic of a solar cell, and the value is taken when a solar cell is short circuited [19]. The short-circuit current is the current obtained when there is no load in the connection or when there are zero loads. The density of light flux incident on a solar cell establishes the value of the short-circuit current, and this density is determined by the electromagnetic spectrum reaching the cell. The short-circuit current, Isc, is conveyed when the voltage is zero [19]. The zero voltage can be obtained with wire of least possible resistance when the negative and positive terminal is connected with a wire.

5.3.2  Open-Circuit Voltage (Voc) The Voc is the voltage obtained when there is no load on the circuit or when the circuit is not delivering current to an external load. That is the maximum voltage a solar cell can possibly deliver. Its alteration is always in the direction of natural logarithm of solar irradiance Mertens (2014).

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5.3.3  Fill Factor (FF) This is a parameter that works with Voc and short-circuit current of a solar cell to facilitate the maximum power delivered by a solar cell. It is a point where the solar cell obtains its maximum power. Mathematically, the ratio of the product of current at the maximum point and voltage at the maximum point to the product of Voc and short-circuit current denotes the fill factor. It can also be described as the ratio of the maximum power delivered to the product of short-circuit current and open-circuit voltage. It is represented by Eq. (5.2).



FF 

I max  Vmax Pmax  I sc  Voc I sc  Voc

(5.2)

The fill factor parameter varies between 0 and 1. The closer the value of the fill factor is to one (1), the better and more likely the material will have better efficiency. Silicon devices from laboratory have demonstrated fill factors closer to one. Both commercial and laboratory solar cells have obtained maximum fill factors of 0.83 and 0.85 in that order.

5.3.4  Solar Cell Efficiency The solar cell efficiency is the utmost evident consideration used in illustrating or weighing the performance of solar. The performance of one solar cell can be different from another by its efficiency. Efficiency of a solar cell is the ratio of its output energy to the input energy. Many factures such as nature of the material, temperature, solar irradiant intensity, and the solar spectrum contributes immensely to the efficiency of the solar cell. Thus, during the performance evaluation of two or more solar cells, the standard conditions for solar cell testing should be considered. Air mass of 1.5 and 25 °C temperature are the main factors considered for testing of terrestrial solar cells. However, an air mass of aero is used while testing a solar cell in space. In order to compare the performance of one solar cell to another, there must be a control of conditions under which solar cell efficiency is determined. The efficiency of solar cells is known through the fraction of incident power that is transformed to electricity. The parameters can be illustrated with Eqs. (5.3) and (5.4). Pmax = Voc I sc FF





Pout Pm FF  I sc  Voc   Pin Pin EA

(5.3) (5.4)

where, FF is the fill factor, η denotes the efficiency, Isc, is the short-circuit current, the open-circuit voltage is represented by Voc, while Pin is the intensity of solar radiation.

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5.4  Oxides A compound containing one or more atom of oxygen is known as oxide. Oxides can be classified into simple and complex oxides. A simple oxide is one that is attached to one other element, while a complex oxide is attached to two or more element to form a compound. Due to the oxidation of individual elements found in earth crust, earth crust in entirety is composed of oxides. In a nutshell, we can say that oxides are earth abundant and stable and naturally occurring compounds. The tendency of the oxygen atom to form bonds by attracting electrons (electronegativity) made oxides to be naturally abundant. Thus, this element can chemically bond firmly with almost all other elements to form oxides. The progress in human technology cannot be fully explained without mentioning oxides. Ancient times during the iron and bronze ages can be explained currently with oxygen splitting advances. The Neolithic era made jewelry using gold or platinum, which are oxide-resistant. At the onset of twentieth century, the emergent of electricity and industrialization triggered the quest for the electrical features of material oxides by scientists and philosophers. Conductors, semiconductors, superconductors, and insulators contain oxides of different states [20]. Oxide materials are employed in electronic radios, solid state transistors, computers, and calculators as optoelectronic component. SiO2 specifically has been used as a technology in metal oxide semiconductor, which was a big breakthrough in technology. Recently, oxides are infused into different devices made of solid-state electronics. It is heavily employed in the fabrication of devices that converts solar energy into direct current electricity [21, 22]. Due to the low cost of oxide materials, they can be used in solar cells as semiconductors, insulators, and conductors. In photovoltaic technology, the basic structure of the cells is composed of two carrier-extraction layer sandwiched with light-harvesting layer leading to conversion of the harvested light to electricity [23]. The oxide materials can be optically transparent, which is used in the fabrication of transport layers, conductive electrodes, and photon absorber.

5.5  M  ethods of Synthesizing Metal Oxides for Photovoltaic Application Metal oxides have been extensively applied in photovoltaic and other areas. The synthesis and preparation methods are vast as listed and briefly explained. The synthesis methods have attracted huge interest as a result of the novel features of metal oxides. These synthesis and deposition techniques help to produce highly efficient materials from electronics to composites. Some of the methods such as chemical processes, chemical vapor deposition (CVD), vapor deposition (PVD), etc. are briefly discussed for familiarity. Table 5.1 shows examples of metal oxides synthesized using different methods and their related electrical and photoelectrical properties. Generally, the cost and strict guiding method of most of the equipment

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Table 5.1  PCE properties of some metal oxides synthesized using different methods Metal oxides TiO2 ZnO CuGaO2/spiro-OMeTAD NiOx TiO2 SnO2: GQDs Perovskite/SnO2 TiO2/NiO rGO/TiO2 TiO2 and ZnO

PCE (%) 7.59 ± 0.08–8.79 ± 0.15 11.86–15.92 16.94–18.55 14.53 and 17.60 0.13–0.48 15.84–20.30 4.5–18.8 2.30 18 10.2–17.0

Synthesis method Coprecipitation Hydrothermal Hydrothermal Coprecipitation Doctor blade Spin coating Sol-gel and hydrothermal Spray pyrolysis Spin coating Hydrothermal

Ref [24] [25] [26] [27] [2] [28] [29] [30] [31] [32]

for the synthesis of metal oxides might be a disadvantage. Some of the machines might not be good to produce massive nanoparticles. On the other hand, the techniques have been helpful in terms of the tenability and variable temperature-­ synthesizing environment, which has helped to improve the desired chemical and physical properties of the metal oxide materials.

5.5.1  Hydrothermal/Solvothermal Approach A hydrothermal method involves stirring of the mixed solution for homogeneity and transfer of the solution into a closed chamber, where it will be heated under high temperature above the boiling point of the solvent and high pressure [33]. The solvent helps to synthesize the metal oxide materials while looking into the pH of the solution. Basically, the hydrothermal temperature ranges from 150 to 250 °C. The concentration of the chemical constituents determines the size or thickness of the deposited samples when hydrothermal method is used [34].

5.5.2  Thermal Evaporation One of the simplest and oldest techniques used in deposition of metal oxides is thermal evaporation [35]. A furnace of high temperature is required to vaporize the materials to aid the deposition with thicknesses ranging from angstrom to microns [36]. With this technique, the electrical properties, thickness, grain structure, uniformity, adhesion strength, optical properties, and stress of the nanomaterials could be controlled [37]. To enhance the smoothness of the operation, substrates are needed for precise uniformity controls, while auxiliary is also needed for vacuum pressure [38].

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5.5.3  Sputtering Deposition In sputtering method of deposition, ion beam from inert gas is used to vaporize a solid material. This method is used mostly to deposit oxides of metals. Urban et al. [39] used magnetron sputtering techniques to form different metal oxides, and this method has been employed by many researchers to deposit nanostructured films. In this method, the ions and ion complexes interact energetically to grow films. The sputtering power and time determine the thickness of the deposited materials [39].

5.5.4  Coprecipitation This is a synthetic bottom-up method for the fabrication of metal oxide materials. In this technique, temperature, surface tension, concentration of solution, viscosity, stirring speed, and pH control the properties of the synthesized or deposited material. Coprecipitation production of metal oxide is a synthetic method that uses bottom-up approach, in which the features of the materials are controlled by the temperature, viscosity, pH, stirring speed, and surface tension.

5.5.5  Physical Vapor Deposition This also involves high-temperature evaporation of the constituent depositing materials beyond its melting point. The substrate is placed in such a way as to allow the evaporated sample in a closed vacuum chamber to be channeled along its path for deposition.

5.5.6  Chemical Vapor Deposition In this method, a solid deposited material is formed after the reactions of the chemicals, which involves passing of gaseous precursors toward a heated substrate. The exhaust chamber releases byproducts during the chemical reaction process. 200–1200 °C range of temperature and different sources of energy like electrons, lasers, photons, hot filaments, plasmas, and ions are used during the process. Structure, nucleation, and crystallinity strongly determine the physical features of the films deposited [40].

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5.5.7  Sol-Gel Approach In this synthesis technique, a high temperature is needed to fabricate the metal oxide by the use of colloidal dispersion or metal alkoxides with the sol and gel intermediate route formation. Sol-gel can be achieved by polymerization of molecules or by dispersion of particles in a liquid to form a solid. It involves mixing of many components at atomic level. The homogenous mixture is attained by stirring while heating and stirring until gel is formed to ensure uniformity [41]. In some cases, a complexing agent is used to enhance fast reactions. Li et al. [24] studied the optical and photoelectrical properties of TiO2 by forming nanostructural dual metal of the material while incorporating Au of different concentration into the active layer. The mechanisms used where analyzed theoretically, and the highest photoelectrical efficiency was approximately 8.79%, which occurred as a result of the improved fill factor and short-circuit current density. Hydrothermal techniques have been employed by controlling the morphology and sizes of metal oxide nanoparticles. Controlled size nanocrystals of ZnO that are highly homogenous were synthesized by Zhang et al., [25] and the nonagglomerated ultrafine particles were synthesized by using different precursors. The device performances were varied as a result of the tunable sizes of the nanocrystallites to produce the films [25]. The SEM images of ZnO synthesized using different solvents are shown in Fig. 5.2a–c. The morphology revealed different sizes of nanoparticles ranging from 25 to 50 nm. Figure 5.2d shows the XRD patterns, and it reveals the narrow widths and peak intensity of the wurtzite-structured ZnO, confirming valid agreement with the SEM micrograph and the photoelectrical efficiency and in the range of 11.86–15.92%. The quest to obtain clean, affordable, sustainable, and highly efficient cells necessitated the study of TiO2/NiO heterojunction metal oxide solar cells using spray pyrolysis method by Ukoba, Inambao, and Eloka-Eboka [30]. From the optoelectrical features, the cell revealed Voc of 350 mV, Isc of 16.8 mA, and 2.30% photo conversion efficiency at 0.39 fill factor. Reduced graphene oxide (rGO) and mesoporous TiO2 were used as a nanocomposite to fabricate electron transport layers of perovskite solar cells, in which the GO was employed to reduce the interfacial resistance and also increase the surface area of the electrode material. In this study, Han et al. [31] used rGO to improve the charge collection efficiency, and maximum efficiency of 18% was obtained for the 0.4% vol of rGO. Recently, Lu et al. [32] improved the efficiency of PSCs by employing interface engineering techniques. In comparison, the pristine TiO2 and ZnO produced respective photo conversion efficiency of 13.2 and 10.2%, while the TiCl4–TiO2 and TiCl4–ZnO-­ based devices produced respective photon conversion efficiency of 16.5 and 17.0%, which are improved efficiency when compared to the pristine samples. Zhu et al. [29] used high pressure and temperature environment to hydrothermally synthesize SnO2 nanoparticles. They observed a very high crystalline SnO2 that is very homogenous with 5–10 nm grain sizes. This SnO2 was found to be more crystallite

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Fig. 5.2  SEM images ZnO synthesized with (a) methanol, (b) ethanol, (c) n-propanol; and (d) the XRD pattern of ZnO synthesized [25]

than the one prepared using sol-gel technique, and the electron transport was faster, thereby reducing recombination effect [29]. Figure 5.2a revealed that the SnO2 synthesized with hydrothermal technique is highly crystalline when compared to the one synthesized using sol-gel method. Furthermore, Fig. 5.2b revealed homogenous and uniform nature of the SnO2 SEM micrograph, while the TEM micrograph in Fig. 5.2c confirms the particle size. The elevated crystallinity of the SnO2 was finally affirmed by the high-resolution TEM morphology. The hydrothermal and sol-gel fabricated cells produced 4.5 and 8.7% PCE, accordingly. The perovskite/SnO2 interface for the sol-gel fabricated cell had deficient charge transfer. However, the reverse and forward scan revealed 18.8 and 18.4% PCE, which is much enhanced. The topmost photovoltaic efficiency of 18.8%, 77% fill factor and external quantum efficiency (EQE) of 80% was recorded for the hydrothermally deposited samples as shown in Fig. 5.3a, b, respectively. Figure 5.4c, d revealed 18.5% steady state PCE at 0.92 V constant bias near the maximum power and 17.4% average efficiency from a histogram of 40 samples,

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Fig. 5.3 (a) J-V curve, (b) percentage EQE, (c) stabilized photocurrent density, and (d) efficiency histogram of perovskite/SnO2 synthesized using hydrothermal method [29]

respectively. Zhang et al. [26] investigated the use of CuGaO2 and spiro-OMaTED as hole-transporting material for perovskite solar cells. They used hydrothermal method for the synthesis of the materials, and the CuGaO2 showed high transmittance and high bandgap and hole mobility. The photoelectrical parameters reveal that the CuGaO2 exhibited 18.51% PCE and 0.77 fill factor, which is relatively higher than the one obtained from spiro-OMeTAD with 17.14% PCE and 0.74 fill factors, and this result reveals that the CuGaO2 had better charge-transport ability than its counterpart. To confirm this measured efficiency, the output powers for the devices were evaluated at constant bias of 0.94 and 0.91 V. The spiro-OMeTADbased device and the CuGaO2-based device produced 16.94 and 18.55% maximum PCE. This inconsistency might be as a result of the recombination problem occurring between the hole transport layer and the perovskite. Zhang et al. [27] proposed highly efficient perovskite solar cell that is highly tunable using coprecipitation technique to form the NiOx hole tansport layer. The maximum PCE for the flexible and rigid NiOx-based cells is 14.53 and 17.60% with 0.705 and 0.784 fill factors, respectively.

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Fig. 5.4 (a) The XRD pattern of SnO2, (b) SEM, (c) TEM, and (d) high-resolution TEM of SnO2 synthesized using hydrothermal method [29]

5.6  Organic Metal Oxide for Photovoltaic Application Organic metal oxides for photovoltaic systems are materials made of carbon derivatives with characteristics of semiconductors. In this system, organic molecules are loosely bonded by van der Waals force while p-bonds bind the atoms of the metal oxide semiconductors. Exceptional features such as low point of sublimation, light weight, and flexibility are some of the attributes of the bonding structures harnessed by the organic metal oxide compounds unlike the inorganic metal oxides that have no such features due to their huge covalent structure. However, the structural bands of organic metal oxides are similar to that of inorganic metal oxides at a large scale. While the conduction band always has space to accept electrons, the valence band is always occupied with electrons for the inorganic metal oxide materials. In analogy, the organic metal oxide materials have the lowest unoccupied molecular orbital

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(LUMO), which represents the corresponding conduction band in inorganic metal oxides, while the highest occupied molecular orbital (HOMO) of the organic metal oxide functions as the valence band in the inorganic metal oxide semiconductor materials. The conjugated p-electrons of the antibonding and bonding hybridization denote the LUMO and HOMO process in organic metal oxides semiconductors [42, 43]. P-conjugated networks form organic molecules in the organic metal oxide semiconductors. Three strong links are structured on the adjoining atoms by sp2-­ hybridized carbon atoms and bonds [43]. The development of fragile p-bonds on delocalized shades of electrons is structured by the residual p-orbitals of the carbon atoms. Thus, the conjugated organic metal oxide semiconductors forms a structural quasi-one-dimensional bond, while dissimilar bond configurations are found on the p-bond system as a result of the overlap in the electron-wave function of adjacent atoms. The antibonding and bonding states correspond to varying levels of energy for variable states of p-bonds. The diverse states of hybridized p-bonds found in HOMO and LUMO organic metal oxides lead to varying energy levels in the organic metal oxide semiconductors. When solar radiation hits the surface of the semiconductor/organic interface, the organic components absorb the light and get excited. Thus, an excited electron migrates from the HOMO to the LUMO of an organic metal oxide semiconductor if the absorbed energy is high enough. The organic molecule itself gets excited during this period. The electron at the LUMO part of the organic material migrates to the metal oxide semiconductor, where it delivers the electron for external work. The organic metal oxide materials have a different mechanism for carrier transport when compared to the inorganic metal oxides counterparts. Here, the energy blockades surrounded by the disarrayed structural-conjugated polymer are overcome by thermal agitation of carriers in the organic metal oxides [44]. Contrarily, the process of charge transport is characterized by migration of free carriers in the conduction and valence band of an inorganic metal oxide semiconductor. The organic metal oxide materials always hopping transport mechanism of the charge carriers, and this has resulted to the drastic decrease in the mobility of charge carriers in comparison to the inorganic metal oxides. Small organic molecules can achieve hole mobility of approximated 1.5  ×  10−3  m2  V−1  s−1 [45, 46], while an approximated 4.5  ×  10−3  m2  V−1  s−1 is recorded for conventional silicon [47]. In the same way, small organic molecules can achieve electron mobility of approximated 1 × 10−5 m2 V−1 s−1 [48, 49], while an approximated very high 0.1 m2 V−1 s−1 is recorded for conventional silicon. In photovoltaic applications, the mobilities for electrons and holes range from 10−7 to 10−8  m2  V−1  s−1 when blended to form a film [50]. When compared to inorganic metal oxides, the mobility of organic metal oxide is low, and this has been observed as major disadvantage, which should be overcome by proposing different devices of credible organic metal oxide.

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5.6.1  G  eneration of Exciton in Metal Oxides for Photovoltaic Application Electron in the organic metal oxide is stimulated to LUMO from HOMO region when photon is absorbed by the organic metal oxide (Fig. 5.5) in the same way an excited electron reaches the conduction band from the valence band in inorganic metal oxide semiconductor upon solar energy absorption. The hole and electron pair exhibit strong coulombic attraction between them because of the localized hole- and electron-wave function and meager dielectric constant [51, 52]. This result to hole-­electron pair bound. The exciton has 0.1–1.4 eV binding energy [52], in contrast to inorganic metal oxide semiconductors with very low as contrasted to a lot low connecting energy of a little meV. Thus, the relative possibility of generating free-­charge carriers is higher in inorganic metal oxides after light absorption due to the easier dissociation of holeelectron pairs from thermal energy, whereas excitons having strong bounds are continuously being generated by the inorganic metal oxide semiconductors. Organic metal oxides have absorption coefficient that is relatively close to 105 cm−1 [53]. The active layers of metal oxides have few 100 nm ranges of thickness, which is moderate to absorb enough quantity of photons so as to reveal significant features of solar cells [54]. Different light trapping methods such as lens concentrators [55, 56], folded cells [54], and gratings [57, 58] are employed to enhance the ability of these materials to absorb energy even when their natural absorption capacity is very weak. The small range of absorption and large bandgap has been major concern to organic materials since this has resulted to low efficiency in the absorption over wide range of electromagnetic spectrum. Approximately 77% of the energy change in the LUMO–HOMO variation leads to 1.1  eV when solar light is absorbed. However, bandgap energy of 1.9  eV has been produced by organic silicon cells, while many organic materials have energy bandgap of 2 eV approximately. Low-­ bandgap materials can be used to enhance the absorption efficiency of materials so that greater percentage of energy can be absorbed within the wide range of electromagnetic spectrum.

Fig. 5.5  Photon absorption and excitation of electrons

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5.6.2  Exciton Diffusion and Dissociation in Metal Oxides After generation of electron-hole pair that is bound to each other, it is pertinent to device a means of separating them to free charges. Two organic metal oxide materials having aligned band levels can be used to produce solar cells when heterojunction is formed between the two materials [59]. This is the basic design of organic solar cells using organic metal oxide materials. The adjacent arranged materials (X and Y) are aligned as shown in Fig. 5.6 so as to dissociate the exciton. The bound electron-hole pair’s potential difference must be larger than the energy difference in the HOMO of X material and LUMO of Y material. The generated exciton in sample X moves to the heterojunction. Energetically, the exciton transfers electron to LUMO Y since there is low difference potential in potential between HOMO X and LUMO Y.  Thus, a hole is created in HOMO X when the exciton transfers electron from HOMO Y. Sequel to the process through which the charges are transferred, material X is called donor material, while material Y is called acceptor material. The recombination effect of the excitons occurs within 1  ns timescale [60]. Contrarily, it takes a shorter time of about 45  fs for the charge transfer to occur [61]. The dissociation of the excitons will efficiently take place within the heterojunction as a result of these timescales. A coulombic bound, a geminate pair, is formed from the electron-hole pair when dissociation occurs, and this requires internal field to be separated. Exciton diffusion length is the distance, which an exciton can diffuse or traverse before recombination occurs, and it is few tens of nanometers for organic metal oxides [62, 63]. When excitons generated are longer than this length, from the heterojunction then recombination occurs before it reaches the heterojunction, thereby leading to the reduction in the efficiency of the excitation dissociation. For this reason, the demarcating phase existing between acceptor and donor is made to be inside the length of exciton dissociation by ensuring that the active layers of the metal oxide are very thin. Conversely, when the

Fig. 5.6  Band alignment for acceptor and donor heterojunction material

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active layer of the material is very thin, it leads to severe exchange of small absorption efficiency. Hence, the exciton dissociation requires surface area that is large enough while exciton dissociation efficiency requires sufficient phase disconnection. An active novel device like active layer of nanostructured material and bulk heterojunction devices can be fabricated.

5.6.3  Carrier Transport in Metal Oxide Semiconductors The generated geminate pairs obtained from exciton is meant to traverse to electrodes where it is attracted. Diffusion and drift currents propel the electrons and holes to the cathode and anode, respectively [64]. The solar cell’s potential gradient is attributed to the drift current from the movement of carriers, which is in turn determined by the type electrode material for the solar cell. A cathode and anode having low and elevated work function, respectively, are utilized, and the alteration in the work function leads to the development of in-built electric field in the solar cell’s Voc. The drift current as well as the interior electric field alters on application of exterior bias. For charge collection, the solar cell’s summed electric field drifts the carriers in the direction of individual electrodes. The diffusion current is another means of transporting carriers. It is a process where carriers are transported following the gradient of carrier concentration in a solar cell. There exists higher concentration of holes and electrons at the heterojunction, the solar cell generates germinate pairs at its hetero-junction. This enables diffusion of carriers away from this area of higher concentration to enforce diffusion current. This current leads when the interior electric field is altered by the applied bias to zero point approximately. The active-layer portability remains the setback of carrier transport. In order to allow carriers to migrate to electrodes at a faster rate, the active layer should be moderately thin because the mobility of electrons and holes in organic metal oxides is usually low. The difference in mobility is a big factor that is highly considered since the characteristic charge transport is governed by it. More than ten factor differences in mobility lead to limited current in space charge [65]. Limited current in space charge generally occurs when a carrier like the electron having better mobility in organic metal oxides is more efficiently transferred to the cathode [66]. Effect of space charge is thereby created, as cathode accumulates electrons due to the higher migration of electrons to the cathode when compared to the movement of holes to the anode. The features of charge transport are then altered at the active layer leading to the creation of upper limit for solar cell’s current output. Therefore, a system striking a balance in the mobility of electrons and holes at the active layer is necessary for efficient carrier transport (Fig. 5.7).

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Fig. 5.7  Charge transport in acceptor and donor heterojunction material

5.6.4  Extraction of Charges at the Electrodes The charge carriers are moved toward electrodes from the active layer after the charge carrier has been transported to the boundary of the electrodes and active layer. There is need to lessen the potential barrier at the boundary separating the electrodes and active layers in order to ensure magnified efficiency during the extraction of charges. For this reason, the LUMO acceptor should have a commensurable work function to that of the cathode, while the HOMO which acts as the donor should have a commensurable work function to that of the anode. This happens at the ohmic contact, where the variation in the HOMO donor and LUMO acceptor associates absolutely with the Voc [67]. Ohmic contact is not formed when the material’s cathode and anode have work function that is farther apart from the LUMO acceptor or HOMO donor. The behavior of the carrier extraction in this situation is overseen by metal-insulator-metal mold [68]. Different type of material electrodes can be used to enhance the work function to match the electrodes. Aside matching the electrodes by altering the materials, the boundary between the HOMOand LUMO- active layer and electrodes inter layers. Before a typical LiF’s cathode that is very thin forms an ohmic contact, it evaporates on the active layer [69, 70]. Materials such as ZnO and TiO2 that is processed in solution form have been used to improve the collection of electron effectively [71]. To form interlayer ohmic contact at the anode side, WO3 or MoO3 from oxides of transition metals oxides were used [72]. The use of transition metal oxides showed that at work function that is even low metals like Al evaporates to form anodes on top an overturned arrangement at the interlayer. Aside altering the electrode’s work function, increase in the surface area and roughness of the electrode materials can make a greater area available for efficient collection of charges [73].

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5.7  Inorganic Metal Oxide for Photovoltaic Applications Currently, about 99% of inorganic-based oxides are trending in the field of photovoltaic applications as one of the energy storage devices. The inorganic oxides for photovoltaic applications are very attractive because of their nontoxicity, chemical stability, and large quantity of oxides of metals that potentially permit the productions below the ambient temperature conditions. It has been reported recently in many research works that the demand and growth for photovoltaic modules has improved tremendously for some years now, and in 2011, power greater than 25 GW has been installed, which leads to a growing installed capacity of about 70 GW, and they are also serving as an apparent conducting oxides [74, 75]. The most widely used transparent conducting oxides for extraction of thin films or nanoparticles for solar cells applications in the recent time are aluminum-doped zinc oxide (AZO), fluorine-doped tin oxide (FTO), indium tin oxide (ITO), and many others [75]. Inorganic metal oxides for solar cells have been used widely in the recent technology for production of many electronic components such as touch screen, smart phones, flat panel displays, and other devices. Profound research has shown in the field of thin-film transistors that transistors based on wide bandgap metal oxides have displayed good optical transparency that has added a good advantage to the functionality of electronic components [76]. The classification of metal oxide for photovoltaic is shown in Fig. 5.8. Solar cells that are mainly made of inorganic lead hybrid perovskites, for instance, CH3NHPbI3, have attained great and extraordinary progress in the past few years [77]. Oxides of inorganic or inorganic oxide are dual compound that have at least oxygen as an element, and the oxides are grouped into acidic or basic according to their acid-base features in both theoretical and practical applications. A compound of an oxide, especially in inorganic metal oxide, must contain or include at least a lone of oxygen atom or oxygen molecules and one or other element in its chemical combination or chemical formula. It is important to state that simple oxides are called binary oxide, but they encompass supplementary elements or compound,

Fig. 5.8  Classification of metal oxide for photovoltaic

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while that of complex oxide is a compound that includes oxygen and as a minimum of two or more elements or compound [78]. Inorganic oxides can be synthesized through many chemical routes, such as: hydrothermal method, coprecipitation method, chemical bath method, electrodeposition method, reaction of oxygen with other compounds at a very high temperature, chemical precipitation, direct heating of elements with oxygen, etc.

5.7.1  C  ontributions of Various Inorganic Metal Oxides for the Development of Photovoltaic Cells Inorganic metal oxides that range from ZnO, TiO2, Al2O3, NiO, CuI, MoO3, CuSCN, CuO, VOx, Cu2O, and others have been broadly used as perovskite supporting scaffolds, photovoltaic cells employing p-type semiconductors, and charge transport in photovoltaic cells. They have shown excellent and convincing characteristics such as well chemical stability, applicable valence band energy level, and soaring mobility, and they are highly transparent within visible region of light. The advantages of using inorganic nanostructured materials are in a such way that they replace conventional inorganic materials in many applications demanding nanostructures, which includes solar cells, high-mechanical strength areas, catalysis, hydrogen storage batteries, fuel cells, biosensors, and many others [79]. Various forms of inorganic metal oxides for solar cells applications have been synthesized and reported in many research papers. For example, titanium dioxide metal (TiO2) has shown enormous efforts for the alteration of the broad bandgap semiconductor to improve the visible light performance in solar cells applications [80]. Furthermore, nitrogen-doped titanium dioxide (TiO2) confirmed superior reactivity over titanium dioxide (TiO2) under visible light because many dopants were substitutionally introduced into titanium dioxide (TiO2) lattice, which include metallic and nonmetallic, and demonstrated to be essential bandgap narrowing and photocatalytic activity [81]. Zhang and coworkers [82] worked on the chemical doping agent of cobaltocene (known as CoCp2). It demonstrated strong electron ability as a stabilizer and has good electrical property that is sound enough for photovoltaic applications.

5.7.2  E  fficiency of Inorganic Photovoltaic Solar Cells Made from Metal Oxides Photovoltaic solar cells based on inorganic metal oxides have contributed significantly in the area of power electronics and in production of reliable electricity that can serve for years without applying any maintenance to it. It has contributed significantly in charging our home electronics and our mobile devices such as smartphones, power banks, cell phone chargers, powered solar calculators, etc. The recent

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technology is showing that photovoltaic cells based on inorganic metal oxides are promising in the future applications of the smallest practical electric motor for a car that requires about 45 KW of energy source at 45% efficiency.

5.7.3  H  ybrid Metal Oxides as Active Materials for Photovoltaic Application Active layers of hybrid can be generated in three major ways. Deposition of a polymer film on a surface of oxide material, nanocomposite creation with metal oxides and polymer to form films, and intercalation of organic matters into ranges of nanostructured oxides are the three major ways leading to the formation of hybrid active materials. The basic scientific questions regarding the interaction of metal oxides and polymers are answered by the use of bilayer as a simple approach. It responds to the quests of Voc origination in hybrid material devices. Initially, there was a believe that the Voc is a major determinant of the material work function [83], but it has been observed recently that among some factors, the positions of active material’s band contribute immensely to the value of Voc of the materials. Fundamentally, the oxides’ conduction band edge (Ec) and the polymer’s HOMO decide utmost probable Voc in hybrid appliance. The higher boundary to the Voc is altered when the energy levels with reverence to the vacuum level alter. The length which the exciton traverses during diffusion has been discovered to be low in organic materials [84– 86]. Recombination occurs when much of the generated excitons diffuse through a polymer layer thicker, thereby leading to low dissociation at the interface. Increasing the surface area of the interface between polymer and oxides is achieved by using nanostructured materials to enhance dissociation. Parameters such as current density (Jsc) and efficiency are highly enhanced as a result of this. Techniques with low energy can be used to produce nanostructured hybrid materials [87]. Uniform arrays of nanorods that are perpendicularly aligned to the substrates have been devised. However, this technique lacks accurate power over the measurement and directions of the prepared arrays. Thus, patterning arrays of large area with perfect uniformity and well-aligned vertical nanostructures of suitable metal oxides have served as the solution [87, 88]. On the other hand, codissolution of polymer and nanoparticles in a pool of organic solvent before depositing a well-blended film from the mixture is another way of fabrication, a hybrid active layer for photon absorption. As an advantage, greater exciton diffusion is observed in this approach due to the huge surface area in comparison to intercalated nanostructured materials. The combined morphology leads to reduced charge transport on the films but enhances generation of charge carrier just as solar cells of organic-made having bulk heterojunction. Charge trapping is necessitated if there is no optimization of the nanoparticles morphology, and this retards the efficiency of charge collection. Thus, an organic molecule such as trioctylphosphine oxide or oleic acid is used as a capping agent on the metal oxide nanoparticles after the nanomaterial synthesis. The addition of the organic

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molecules on the surface helps to prevent insolubility and precipitation of the solution, thereby leading to uniform mixture formation. The oleic acid and trioctylphosphine oxide have to be removed to acquire a moderate surface since they act as barriers between metal oxides and polymer materials [89]. Pyridine has been used to replace the oleic acid organic molecules [89]. Placing of this active layer between the conducting polymer and metal oxide material reduces shunt pathways and improves the efficiency of the cell. The charge transfer dynamics of the pyridine at the interface influences the Jsc, especially due to high area of oxide/polymer surface, and more excitons are found to dissociate highly to facilitate large increase in PCE.  The processing history of the metal oxides determines the performance of hybrid devices. Growing different seed layers from solutions containing different chemicals and annealing the films at different temperatures alter the properties of the materials [90] treatment with heat enhance the performances of the metal oxide cells since the heat removes some absorbed molecules on the surface of the oxide. This leads to creation of moderate contact with the polymer. 5.7.3.1  Hybrid Perovskite Solar Cells In perovskite solar cells, the active layer makes use of materials structured in perovskite form. It is a form of hybrid organic-inorganic incorporated with tin or lead halide that is mostly processed in solution. In 2009, the maximum efficiency that was obtained from this cell was not up to 5%. However, the efficiency of the cell has increased to above 25% in a decade. This makes perovskite cells to have advance growth in terms of technology, and researchers have found the study very welcoming. The materials for the PSCs are relatively cheap when compared to conventional silicon cells. Thus, the cell is very interesting for commercial purposes. Unfortunately, the instability of the PSCs has retarded the prospects of its commercialization. However, researchers are working round the clock to enhance the stability of perovskite solar cells. Furthermore, cesium lead halide (CsPbX3  =  I, Br, Cl) perovskite, as one of inorganic nanocrystals have come out as current possible materials that is suitable for photovoltaic cells, is an excellent candidate for display and lighting technology [91]. Begum and coworkers [92] demonstrated that CsPbI3 and CsPb (Br/I)3 showed good optical and structural properties, for example, superior emission quantum and phase stability that yields up to 94.2%. In addition, zinc oxide (ZnO) has also proven to be a good candidate for photovoltaic application. Zinc oxide (ZnO) has shown good performance in solar cell heterostructures, as it serves as charge selective and electron acceptor in solar cells. Researchers have fabricated hybrid solar cells that comprise ZnO/ITO nanowire (NW) SQ2/P3HT/Ag. Their work shows that zinc oxide is broadly used in metal oxide semiconductor for photovoltaic application, and they also highlighted that the charge generation of metal oxide based on photovoltaic cells was improved considerately.

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5.7.3.2  Dye-Sensitized Solar Cells (DSSCs) Graetzel and O’Regan remain the pioneers of the DSSC, where they initiated the novel research area. In recent years, the idea of hybrid metal-oxide polymer photovoltaic is driven by the inspiration of the combination advantages of inorganic and organic solar cells [93, 94]. Photovoltaic cells employing organic and inorganic hybrid perovskite as absorber have shown considerable interest in the area of solar cells, with an extraordinary increase in energy/power conversion productivity from 3.9 to 25.3% in the recent years. They obtained approximated average PCE of 7.9% [95]. Due to the success they achieve, many researchers have done extensive work to improve the PCE of DSSC, especially on the photoelectrode materials by employing n-type metal oxide materials such as SnO2, SrTiO3, ZnO, TiO2, and many other composites. For suitable band site comparative to the sensitizer, wide bandgap greater than 3  eV is necessary for the production of DSSC.  For the fabrication of DSSC, metal oxide semiconductors are utilized to absorb solar radiation within the near-infrared and visible region because of their wide bandgap. The dye which acts as the photosensitizer is responsible for light absorption within the ultraviolet section of light spectrum. To further enhance the light absorption ability and PCE of the DSSC, a metal oxide semiconductor nanoporous material with lofty surface area is needed. The composition of DSSC is naturally abundant, simply synthesized, and cheap. The metal oxide materials can combine with solution during the synthesis for easy fabrication [96, 97]. TiO2 has been the most used metal oxide semiconductor for the fabrication of DSSC among many other wide bandgap materials. This is because of its unique features, since it had recorded the most promising results in terms of PCE.  TiO2 exists in different forms, but brookite, rutile, and anatase have been the most researched crystallite forms of TiO2. TiO2 is highly stable chemically, it is abundantly available, and it is not toxic. The rutile and anatase phase of TiO2 are symmetrically tetragonal in terms of the structure, and the oxygen atoms are coordinated in sixfold Ti4+. The location of the oxygen atom differentiates two phases of rutile and anatase. The anatase form is characterized as having shorter average distance between the Ti4+ sites when compared to the rutile. This makes the rutile to be more stable thermodynamically. Within 700–1000 °C temperature range, anatase is transformed to rutile phase, and the transforming temperature is determined by the purity of the material and its crystallite size [98, 99]. Due to the shift in the negative direction of the anatase’s conduction band, the anatase phase has more energy bandgap of 3.2 eV compared to 3.0 eV of the rutile TiO2. The TiO2 has fairly ionic and fairly covalent bond, and it has insulating crystals stoichiometrically [97, 100]. During some synthesis process, much quantity of trapped states is stimulated as a result of vacancies in oxygen. Alteration in conductivity by different sizes can occur when oxygen vacancies are structured in opposite direction at la lofty temperature or diminished pressure. The vacancies of oxygen lead to the development of Ti3+ causing a doped negative crystal known as n-type. The photo ability of this type of metal oxide semiconductor does not degrade unlike other semiconductors when excited. SnO2 is more stable than TiO2 in terms of light degradation due to its lofty work function and bandgap. Unfortunately, one

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of the major setbacks in the experimental results is the low efficiency attained from the mesoporous TiO2, which cannot be matched to its theoretical counterparts as a result of recorded 0.1 cm2 V−1 s−1 low-electron mobility. A typical standard thin film having thickness ranges of 2–15  μm with 20–30  nm sized particles used for the deposition of the films. A photo-scattering coating, which assists in trapping light, is a fabricated structure with double-layer having particle sizes of 200–300 nm with thickness of 2–4  μm [97, 101]. Dye containing ruthenium complexes is used to sensitize the metal oxide semiconductor material used for DSSC. The metal oxide semiconductor is deposited on a glass substrate, and this acts as the photoelectrode material while another conducting glass substrate is coated with counter electrode material. Between these two electrode materials is sandwiched with redox electrolyte, and these three layers provide adequate connection for external circuit. This set up functions like photosynthesis route in plant. The dye molecule got excited upon solar energy absorption, and an electron moves from the dye to the metal-oxide semiconductor, which acts as the conduction band of the TiO2. This makes the dye to be oxidized while the metal oxide TiO2 is reduced. After doing external work, the electron returns to the dye molecule through the back contact of the counter electrode, and this makes the state of the dye molecule to be restored to ground state [102]. Figure 5.9 shows the working mechanism of a DSSC. The oxidized dye with conduction band electron recombination occurred during the regeneration of the sensitizer by iodide/triiodide electrolyte. The charge carriers are moved to the external circuit when the oxidized redox species diffuses electron into the metal oxide semiconductor nanocrystallite film, which then moves through the back contact of

Fig. 5.9  Working mechanism of DSSC

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the conducting substrate. The cycle is completed when the electrolyte is reduced after the returning electron from external load enters the redox chamber [100, 101]. It is of paramount important to ensure that the rate of recombination or electron loss is far much lesser than the rate at which electrons are injected into the crystallite TiO2. Most notable recombination process is the loss of electron from the conduction band of the metal oxides directly into the electrolyte without delivering current to the external circuit. The process by which an excited electron moves from the LUMO part of the dye molecule to the photoelectrode’s conduction band is known as interfacial electron transfer, and this has k2 rate constant [97]. To a great extent, the energetic strength of the boundary linking metal oxide semiconductor, dye and electrolyte, and the electron density in metal oxide semiconductor determines the electron-transfer kinetics of the interfacial surfaces. This happens within picoseconds duration, and for many metal oxides and sensitizers, 10–12  s−1 have been recorded [97, 102]. The efficacy of electron injection at the layers of the mesoporous metal-oxide semiconductors to the conducting surface is determined by electron-diffusion coefficient (De) and lifetime of the electron. Thus, DSSC PCE should be produced from metal oxide materials whose particle sizes are greater than the Bohr radius of their exciton. For DSSCs to have enhanced performance, some factors must be accomplished. The metal oxide semiconductor photoelectrode must have large surface area and, naturally, mesoporous so as to encourage adsorption of dyes on this surface. The metal oxide semiconductor should have large amount of density of states than the dye’s molecular orbital to facilitate faster electron injection to the metal-oxide semiconductors through the dye. Electron trapping and detrapping in the electronic band of the surface atomic space emanates from the metal-oxide semiconductors as a result of insufficient transport of charges at the metal-oxide semiconductors. In addition, the great capacity of atoms on the surface of the nanostructured metal oxides results to high-trap density. Removal of the traps and its quantification might enhance the DSSCs’s photoelectrical efficiency 11–15% current record [98, 99, 103]. One of the major challenges of using the TiO2 mesoporous material is the reduction in the efficiency of DSSC due to small electron mobility of about 0.1 cm2 V−1 s−1, and this makes the efficiency to be lower than the theoretically obtained value [104]. The electron recombination with the electrolyte occurs when the thickness of the photoelectrode is very much compared to the length of diffusion and length of transition. On the other hand, the nanostructure of SnO2 is a very lofty-transparent conducting oxide used for electronics because of its lofty bandgap of approximately 3.6 eV and μn between (10 and 125 cm2 V−1 s−1) [105]. Unfortunately, TiO2 has a higher energy in the conduction band than the SnO2 [102], and this makes the DSSC fabricated using SnO2 to have open-circuit voltage. The use of wide bandgap metal oxide materials by preparing composite materials and core/shell electrode has improved the Voc up to 600  mV recently [106]. Modification of the surfaces of electrodes, the composition of the electrolyte, and combination of SnO2 with other materials has also enhanced the PCE of the material [107].

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5.8  Active Metal Oxide Roles in Photovoltaic Cells These active materials play variable roles in organic and inorganic-based photovoltaic technologies. The light absorption is performed by the active materials. The polymer absorbs photon, and the interfaces are responsible for the dissociation of the excitons. Generations of free carriers are not possible without the interfaceexciton dissociation of the oxides. Electrons are transported through the oxide, polymer provides the base for hole transportation, and they are all collected at the electrodes. Some of the parts of the active layers are enumerated.

5.8.1  Transparent Electrodes Due to recombination of charge carriers generated, the carriers traverse some hundred nanometers, and this leads to low-relative conductivity of the organic semiconductors. As the incident photon reaches active layer, charges are collected at the entire devices within the transparent electrodes. Suitable used materials for this purpose are wide-bandgap semiconductors. This is because their electrical properties can be tuned, and they also possess excellent optical features.

5.8.2  Charge-Blocking Layers Current leakage or dark current in organic photovoltaic system has contributed to its reduced efficiency. The performance of the photovoltaic device is drastically diminished as a result of the fact that the generated charges recombine. The unified nature of organic and hybrid systems contributes particularly to this problem. Thus, metal-oxide-blocking layers can be introduced on either side of the energetic layers to reduce the problem of current layer. One layer favors the transportation of electron and the other favors the transportation of holes as a result of the alteration in the position of the energy levels. With this, dark current is reduced, and the conductivity is improved.

5.8.3  Charge Collectors Relatively, some metal oxides possess elevated bulk-electron mobilities in orders greater than the magnitude of some organic materials. In conventional devices made from organic photovoltaic materials, these metal oxide materials operate as charge collectors. Hybrid- or organic-blended cells undergo intercalation during the formation of oxide nanostructures, where charge photogenerated carriers are conducted to the electrodes due to the high-mobility nanomaterials that extend into the active layer for conduction.

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5.8.4  Optical Spacers When light travels through multiple thin films such as hole/electron-blocking layer, active layer, glass slide, transparent electrode among others with variable optical features, and the solar radiation traverses throughout the device films with different optical features, the interface of the materials can encounter reflections. The light reflects back at the interface of the layers when it arrives at the electrodes. The refractive index and thickness of the films determine the intensity of the light and the wavelength in the devices’ regions as a result of destructive and constructive interference. The highest intensity of the spectrum is absorbed at the active region through the help of optical spacers.

5.8.5  Intermediate Layers in Tandem Cells Up-to-date, multiple layers of metal oxide and other materials having different energy bandgap have been stacked together to obtain the maximum efficiency in inorganic devices. An intermediate layer made of metal oxides of few nanometers must be incorporated to achieve this standard in devices made of organic and hybrid materials. One of the reasons for this is because the bottom cell is protected by this layer, thereby preventing it from being dissolved when subsequent cells are deposited. On the other hand, it facilitates cell contact and also allows recombination of unwanted electrons and holes from different layers.

5.8.6  Stability Enhancers It is observed that in the presence of oxygen and light, the semiconductor materials joined with polymers are unstable generally. Thus, in order to assuage this, oxygen-­ deficient oxides are used to improve in the stability of the organic layers.

5.9  R  eview of Some Metal Oxide Materials Used for Photovoltaic Application Onah et al. [2] fabricated DSSC using dyes from plants and flowers and TiO2 with the aid of doctor blade’s technique, and the photoelectrodes were annealed at 450  °C.  They analyzed the structural features of the nanoparticle using XRD spectrophotometer. The XRD pattern of the photoelectrode is shown in Fig. 5.10. It shows an anatase phase of TiO2 having high peaks at 2θ values of 25.28° and 48.05° with hkl index of (101) and (200), respectively. The results obtained here matched

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Fig. 5.10  XRD pattern of titanium dioxide (TiO2) annealed at 450 °C [2]

with JCPDS card number 00-021-1272, and other lower peaks at different reflection planes were also observed [108]. They also obtained 26.56 nm and 0.0042 mean crystallite size and lattice strain from the lattice parameters of the photoelectrodes. The surface characteristics of the deposited TiO2 nanoparticles were analyzed with SEM as revealed in Fig. 5.11a–c with increasing magnification from (a–c). The SEM micrograph shown in Fig. 5.12a–c revealed nanoparticles of 100–250 nm that are randomly positioned in granular-like shapes. They noted that the particles had strong binding capacity, which is as a result of the large surface area provided by the particle connections. On the other hand, the deposited samples were evidenced from the EDX spectroscopy, where the associating elements and its distribution were seen. They observed the presence of oxygen and titanium as the major reacting constituent elements in the presence of some minor impurities. The distribution and the percentage weight were shown in the image by the peaks in Fig. 5.11d [109]. The FTIR spectra were used to analyze the functional groups at the region bands for the extracted dyes as shown in Fig. 5.12a–e. The dyes had great peaks of absorption mainly characterized with O–H, C–H alkane, aromatic –C=C–, –O–C=O–, and carbonyl groups, which are good features of dye extracts from flowering plants. These dyes from anthocyanin pigments facilitated the light-absorption properties of the comonents [110]. C≡C triple bonds from alkynes, C–O–C esters and epoxides, and –O–N=O from nitro compounds as well as nitriles having viable peaks at 2632.24 cm−1, 2987.00 cm−1, 3221.49 cm−1, and 3681.98 cm−1 were also detected. Stretching vibration from O–H intermolecular bonds was recorded for the green leaves extracts. The absorption spectra of the dyes and dyed photoelectrodes were studied with the help of spectrophotometer. Figure 5.13a–f reveals the degree of incident solar energy absorption of the films and the dyed electrodes. It was observed that the

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Fig. 5.11 (a–c) SEM images for TiO2 films annealed at 450  °C with increasing magnification from (d) EDX spectra of TiO2 film [2]

Hibiscus rosa-sinensis flower displayed better absorption properties than the others as shown in Fig. 5.13a. Its absorption occurred within 714 nm wavelength of the electromagnetic spectrum. The dyes generally had higher absorbance than the photoelectrode. However, the spectrum for the dyed photoelectrodes occurred within 900 nm. This implies that the wide absorption range of the TiO2 photoelectrode has indirectly affected that of the dye components as shown in the figures. Moreover, the dyes showed higher absorbance than the TiO2 and the TiO2-dyed photoelectrodes. From Fig. 5.13a–f, it is evident that while the dyed photoelectrode absorbs toward the ultraviolet part of the spectrum, the dye absorbs within the visible region of light spectrum. The reduced absorption of the dyed photoelectrode might be due to the presence of the titanium metal oxide [111]. Figure 5.14 shows the bandgap energy of the film as represented in Eq. (5.5) [112]:

 h 

1

n

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where β is a constant representing band-tailing parameter, Eg denotes the energy bandgap, and n is the power factor mode. Plotting (αhυ)2 against hυ gives a distinct straight curve whose intercept is extrapolated along hυ at α = 0 point. 3.2 eV energy bandgap was obtained from the curve, which agrees with literature [113]. Furthermore, Fig. 5.15a–e shows the bandgap energy of the stained photoelectrodes calculated using Eq. 5.5.

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Fig. 5.12  FTIR spectra of (a) Hibiscus rosa-sinensis, (b) Helianthus annuus, (c) Luffa cylindrica, (d) Mimosa pudica, and (e) Lonchocarpus cyanescens dye extracts [2]

The dyed TiO2 film stained with Hibiscus rosa-sinensis dye extract revealed 2.15 eV energy bandgap, and this is followed by 2.7 eV obtained from TiO2 film stained with Helianthus annuus dye extract. 2.99  eV was obtained from Luffa cylindrica stained TiO2 films, while the films stained with Mimosa pudica and Lonchocarpus cyanescens dye extracts revealed energy bandgap of 2.99 and 2.37 eV, respectively. The decrease in energy bandgap as the photoelectrodes are sensitized with the various dyes extracts shows that the dye extracts have lower energy bandgap than the titania films. It also confirms that the dyes have better absorbance than the TiO2 photoelectrode. In other words, this would further improve the photo conversion ability of the TiO2 photoelectrode as the sensitizers would improve the transport system of the cell. The photoelectrical features of the assembled cells are shown in Fig. 5.16a–e, and among the five cells fabricated with different dye extracts as sensitizer, the one with Hibiscus rosa-sinensis flower dye extract gave the best result, followed by the

Fig. 5.13  Absorption spectra for (a) pristine TiO2 film and (b–f) the dyes and dyed TiO2 films for the various dyes [2]

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Fig. 5.14  Energy bandgap of TiO2 photoelectrode [2]

2.5

(αhυ)2 (eV/cm)2 X1014

TiO2

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Fig. 5.15 (a–e) Energy bandgap of the dye-sensitized photoelectrodes for the various dyes [2]

dye extracts of Helianthus annuus and Luffa cylindrica with PCE of 0.48, 0.35, and 0.29%, in that order [114]. Furthermore, the cells fabricated with Lonchocarpus cyanescens and Mimosa pudica revealed 0.17 and 0.13% PCE.  The dye extracts from the flowering plants had better efficiency due to their good absorbance and better adsorption ability to the photoelectrodes. However, the relatively poor photo conversion ability of the Lonchocarpus cyanescens- and Mimosa pudica-fabricated cells was because of their meager anchoring abilities [115]. Nkele [116] used chemical bath deposition technique to synthesize nickel oxide films as a hole transport layer in perovskite solar cells. The molar concentration of

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Fig. 5.16  I–V plots for the fabricated cells using dye extracts from (a) Hibiscus rosa-sinensis, (b) Helianthus annuus, (c) Luffa cylindrica, (d) Lonchocarpus cyanescens, and (e) Mimosa pudica [2]

the nickel oxides they used was different from the ones in literature, and the films were annealed at lower temperatures. Furthermore, they deposited methylammonium lead iodide (CH3NH3PbI3) as an initially synthesized perovskite layer using spin coating technique on the deposited nickel oxide layer [117]. They further compared the properties of the nickel oxide coated with CH3NH3PbI3 and the pristine nickel oxide for the optical, morphological, structural, and the electrochemical abilities using corresponding equipment [117]. The interfacial recombination and charge transport attributes were also analyzed for the nickel oxide and the perovskite. Figure 5.17 shows the SEM micrograph of the nickel oxide deposited at different time, and this morphology helps in getting the information of the material both in nanoscale and in bulk quantity [117]. Using different deposited time, the SEM of the synthesized nickel oxide films deposited for 10 min showed brick-like structure that converted to clusters of nest-like crystals as the deposition time increases. High

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Fig. 5.17  SEM images of the NiO films deposited at different times. (Adapted from ref. [117])

Fig. 5.18  SEM images of MALI spin-coated on NiO films. (Adapted from ref. [117])

performance of the films and light absorption properties are favored by the big brick-like structure obtained as deposition time increases [118]. The electrochemical features were favored also by the porous nature of the film surface due to the method employed for the deposition [119]. Furthermore, large surface interaction areas are provided by the porosity of the nickel oxide films, and this leads to high solar energy transmission through the film and subsequent charge transport enhancement. On the other hand, Fig. 5.18 shows the SEM micrograph of the spin-coated methylammonium lead iodide (MALI) on NiO films, which reduced gradually in size as the time of deposition increases. The hole transport ability was efficiently

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Fig. 5.19  Absorbance spectra for the (a) NiO and (b) MALI spin-coated on NiO. (Adapted from ref. [117])

evidenced due to the high porous nature of the clustered films. There exist spikes on the surface of the substrates from the MALI/NiO film deposited for 50 min. These spikes on the NiO film surface show the direction of the MALI growth. Spectrophotometer was used to analyze the optical properties of the films. Figure 5.19 revealed the optical properties of the NiO and NiO/MALI films and the films’ absorbance occurred between 200 and 1100 nm wavelengths of the electromagnetic spectrum. Maximum absorption was recorded at 300  nm, which is the ultraviolet part of electromagnetic spectrum. However, reduced absorption was observed toward the visible and then to the near-infrared region where it stabilizes at a constant value. The uncontrolled deposition thickness as the deposition time increases contributed to the variation in the absorbance [120, 121]. It is also seen from the graph that the thickness of the films contributed to the increase in the absorbance due to the increase in deposition time. Figure 5.20a, b revealed the bandgap energy of the NiO and NiO/MALI. The figure revealed 2.9–3.3 eV energy bandgaps. It shows systematic decrease in the bandgap energy as the time of deposition increases for both NiO- and NiO/MALI-­ deposited films. From Fig. 5.20b, it is obvious that the addition of perovskite on the NiO film caused minor reduction of the bandgap energy of the films [122, 123]. This would favor the conductivity of the perovskite film, since the reduction in bandgap energy closes the gap between the valence band and conduction band, thereby promoting easier transmission or migration of charge carriers. The electrochemical impedance of the cells fabricated using NiO and NiO/MALI was used to study the recombination effect and interfacial charge transport of these devices. The Nyquist plot in Fig. 5.21a shows that the interface of FTO, HTL, and perovskite had semicircle at the high-frequency region of the curve. The production of semicircle is as a result of capacitors being parallel with the resistor of the charge transfer. The resistance of the solution is represented by the area of high frequency at the beginning of the plots. The semicircle with high frequency at the layer of the hole transport originates from the charge transfer resistance (Rct) [124]. Furthermore, the inset in Fig.  5.21d was utilized in fitting the NiO-deposited films with using ZSimpWin software, where W, Cs, and Rs denote the Warburg impedance,

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Fig. 5.20  Energy bandgap spectra for the (a) NiO and (b) MALI spin-coated on NiO. (Adapted from ref. [117])

Fig. 5.21 (a) Nyquist plot with a circuit inset R(CR)(CR), (b–d) Bode plots of the NiO electrodes deposited at 10, 20, 30, 40, and 50 min, respectively, with an inset circuit in (d). (Adapted from ref. [117])

double-­layer capacitance, and ohmic resistance, in that order. From the graph, 0.20 and 0.13 kΩ were observed from the charge transfer and solution resistance, respectively. Three electrode systems were used to obtain the total ohmic resistance. The Warburg element gives the Rct in the samples and the rate of diffusion of the ions in the electrolytes. Efficient charge carrier transport that occurred at the interface of the NiO and the perovskite resulted to low resistances as obtained. The Bode-phase

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plots are shown in Fig.  5.21b–d, where the low- and high-frequency regions represent the carrier recombination and diffusion accordingly [125]. Due to better interfacial contact, carrier recombination is observed at the low region of frequency, while efficient carrier diffusion from the perovskite to the hole transport layer occurs at high-frequency region. This shows that incorporating NiO into perovskite material would decrease the rate of recombination. Figure 5.21d shows that as the frequency increases, constant-phase angle is observed. Thus, best electrochemical features and utmost phase angles were recorded by the NiO film and NiO/ MALI. Interfacial charge transfer and better electrochemical features were observed when NiO/MALI films were deposited at low deposition times. The solution resistance leads to approximate zero capacitance at high frequency, while high impedance was recorded at low frequency. PSCs fabricated using SnO2 as the electron transport layer (ETL) have yielded low performance due to many trap states when synthesized in low-temperature solution, and this has led to serious hysteresis. On this note, Xie et  al. [28] demonstrated a strategy to improve the electron transport system and the overall performance of the cells. They used novel method of incorporating little quantity of graphene quantum dots (GQDs) to enhance the electronic features of the SnO2. They spin-coated ETLs of SnO2:GQDs on ITO and annealed the samples at 180  °C.  The GQDs have zero- (0)dimensional confinement, and this makes it to have a tunable energy bandgap with distinctive properties compared to the conventional 2D graphene [126]. This shows that electrons generated can easily migrate to the SnO2 conduction band in presence of the GQDs. High-resolution TEM image of the 5–10  nm diameter GQDs is shown in Fig. 5.22a. The crystallized GQDs have 0.21 nm lattice spacing and that agrees with the (1100) graphite plane [127]. From Fig.  5.22b, an approximate 2.4  eV was obtained as the energy bandgap of the GQDs using UV–vis absorption spectrum, and this shows a very much smaller value than that of the SnO2. The dispersion of GQDs in water can easily be noticed in the inset in Fig.  5.22b, and this shows transparent yellow color. The photoluminescence spectra of the GQDs aqueous solution are shown in Fig.  5.22c. From the photoluminescence spectra, red-shift peaks were observed within 530–600 nm as the excitation wavelengths raises from 375 to 550  nm, and this confirms the dependency of photoluminescence on the

Fig. 5.22 (a) TEM image of GQDs, (b) energy bandgap of GQDs using UV-vis spectrum, and (c) photoluminescence spectra of GQDs in water [28]

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excitation wavelength. Thus, the GQDs exhibit these characteristics as a result of the inhomogeneous distribution of energy at the orbital [128]. From Fig. 5.23a, 400.1 eV peak was observed for the GQDs from the XPS spectrum, which corresponds to N 1 s as a result of nitrogenous functional groups. The pure SnO2 and the incorporated GQDs to the SnO2 revealed lower peaks. This is because the pure SnO2 films don’t have N ligands, while the GQDs and SnO2 mixture had little N ligand as a result of the low quantity of GQDs. The transmittance spectra of the SnO2 and SnO2:GQDs composite with same deposition thickness shown in Fig. 5.23b revealed that the transmission of the SnO2 is higher than that of SnO2:GQDs within 300–550 nm and also within 550–900 nm. The reason for this is because of the reduced bandgap energy of the GQDs that affected the SnO2 counterpart. ITO/SnO2(GQDs)/Ag design was used to evaluate the importance of incorporating GQDs into SnO2 during the conductivity check. At 0.1 V bias voltage in the dark, the J-V curve in Fig. 5.23c shows that the pure SnO2 had a slight lower conductivity when compared to the SnO2:GQD. On the other hand, the conductivity of the SnO2 in Fig. 5.23c is approximately 20 times lower than that of the SnO2:GQD film, which signifies magnificent improvement. The conductivity of SnO2:GQD alters with light off and on as shown in Fig. 5.23d. The saturation point took tens of seconds to be reached after illumination and returning to the dark took some minutes.

Fig. 5.23 (a) XPS spectra of N 1  s peaks of GQDs, SnO2, and SnO2:GQDs, (b) transmission spectra of SnO2 and SnO2:GQDs, (c) J–V curve of the device ITO/SnO2:GQDs/Ag under dark and AM 1.5  G simulated sunlight, and (d) device current responses for ITO/SnO2:GQDs/Ag under dark and AM 1.5 G simulated sunlight [28]

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The valence band (EV) and conduction band (EC) between the SnO2 and GQDs are shown in Fig. 5.24a. The lower EC of the SnO2 allows the movement of electrons from GQDs to the SnO2 without barrier in energy. Bandgap energy above that of GQDs is generated by the charge carriers upon illumination. Carrier lifetime and electric field influence the electron recombination and ejection or both [129]. The GQDs have long carrier lifetime than the SnO2 due to its high level of energy and discrete energy level, which might lead to high and fast injection of electron to the SnO2 to avoid recombination and subsequently averts the trapping of electrons [130, 131]. The work function of the SnO2:GQDs was carried out under illumination as revealed in Fig. 5.24b, and it was observed that after illumination, the Fermi level (F) of SnO2:GQDs reduced from 4.35 to 4.01 eV, which is closer to EC while in the same way, the EF also reduced from 4.37 to 4.29  eV for the SnO2. The carrier mobility and trap density were obtained using space-charge-limited current method, which helps to comprehend the conductivity mechanism [132]. Here, the electron trap-state density before and after illumination was studied as shown in Fig. 5.24c, d. The J-V curve showing ohmic response at low bias is indicated with green line. The square-law area ends in a suddenly increasing current emanating from the trap filling due to the increase in the voltage. The electron trap-state density of pure SnO2 had insignificant decrease, while for the SnO2:GQDs it is shown that it decreases from 4.30 × 1016 to 1.23 × 1016 cm−3 after illumination. The device’s dark current

Fig. 5.24 (a) Electron transfer mechanism from GQDs to SnO2, (b) change in work function for the SnO2 and SnO2:GQDs, electron traps and mobility of (c) SnO2 and (d) SnO2:GQDs [28]

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can be fitted very well using Mott−Gurney law when operating in the trap-free region as indicated by the blue line in Fig. 5.24c, d. Figure 5.25a shows the SEM micrograph of the perovskite device with SnO2 film having clear thin and uniform layer of about 40 nm between the perovskite layer and the ITO. The changes in SnO2:GQDs upon illumination can be illustrated with the energy bandgap in Fig.  5.25b. The Voc of the cell was found to improve because of the reduction in the EF for the SnO2:GQDs, and generally the SnO2:GQDs devices had better performance than the pure SnO2 for all the photoelectrical parameters. The PCE histogram over 40 samples for the SnO2 and SnO2:GQDs shown in Fig.  5.25c showed that the device fabricated with pure SnO2 had lower PCEs (16.6 ± 0.9%) than the ones with SnO2:GQDs (19.2 ± 1.0%). Using two different ETLs, the J-V features were shown in Fig. 5.25d, and this was evaluated using reverse and forward scan. The SnO2-fabricated device had 17.91% and 15.84% PCE from the reverse and forward scan respectively [133]. Respective 17.47 and 20.23% steady-state PCEs were obtained from the ETLs of the SnO2 and SnO2:GQDs as shown in Fig.  5.25e. The external quantum efficiency spectra of the device is shown in Fig. 5.25f revealing integrated photocurrent densities of 22.47 and 21.36 mA cm−2 for the ETLs of SnO2:GQDs and SnO2, respectively. This study has confirmed that efficiency in carrier transport of materials can be achieved by adding GQDs to them, which will later reduce electron traps.

Fig. 5.25 (a) Cross-sectional SEM micrograph of the device showing some structures, (b) device energy level diagram showing the result for before and after illumination, (c) PCEs histogram for SnO2 and SnO2:GQDs cells, (d) J–V curves of the best PSCs based on SnO2 and SnO2:GQDs, (e) time-dependent-stabilized PCEs, and (f) EQE spectra of the best-performing cells based on SnO2 [28]

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5.10  Conclusion This chapter discussed in detail the optically active metal oxides as it emphasized their importance in photovoltaic technology. Metal oxides can be synthesized using different methods such as hydrothermal, sol-gel, coprecipitation, etc. The PCE of conventional solar cells has been observed to be much higher than that obtained from DSSC, perovskite, and hybrid solar cells due to the recombination effect, electron trapping, and detrapping in the electronic band and instability of the cells, which leads to reduced electron mobility. When compared to inorganic metal oxides, the mobility of organic metal oxide is low, and this has been observed as a major disadvantage of this type of material. It is expected that hybrid materials incorporating both the organic and inorganic metal oxide would mitigate this problem. Low bandgap materials have been used to dope metal oxides so as to enhance the absorption efficiency of wide bandgap metal oxide materials so that a greater percentage of energy can be absorbed within the wide range of the electromagnetic spectrum. The active layers of metal oxide made very thin have been used to reduce the effect of recombination. Materials such as ZnO and TiO2, etc. that are processed in solution form have been used to improve the collection of electron effectively. On the other hand, an increase in the surface area and roughness of the electrode materials makes greater area available for efficient collection of charges. Growing different seed layers from solutions containing different chemicals and annealing the films at different temperatures alter the properties of the materials and enhance the performances of the metal oxide cells. Further research abilities are needed to improve the overall performance of metal oxide materials used for photovoltaic applications.

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Chapter 6

Structural and Electronic Properties of Metal Oxides and Their Applications in Solar Cells Agnes Chinecherem Nkele, Sabastine Ezugwu, Mutsumi Suguyima, and Fabian I. Ezema

6.1  General Introduction Oxides are the least free energy materials formed due to metallic corrosion and interactions [1]. Metal oxides are remarkably studied because of their high carrier mobility, relatively cheap, optical transparency, highly sensitive, and tolerance to mechanical stress [2]. Metal oxides are crystalline solids that comprise of a metal cation and an oxide anion held together by electrostatic force. They are classified into acidic (SO2, CO2), basic (Na2O, CaO, MgO), amphoteric (ZnO, Al2O3), and neutral oxides [3]. Basic oxides react with water to form bases and with acids to form salts. SnO2 and ZnO are usually n-type semiconducting materials that are electrically resistive due to defects [1, 4]. MgO and Al2O3 are insulating wide bandgap materials that are resistant to dopants. Metal oxides with high valence band

A. C. Nkele Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu, Nigeria S. Ezugwu Department of Physics and Astronomy, University of Western Ontario, London, ON, Canada M. Suguyima Department of Electrical Engineering, Faculty of Science and Technology, Tokyo University of Science, Yamazaki, Noda, Japan F. I. Ezema (*) Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu, Nigeria Nanosciences African Network (NANOAFNET), iThemba LABS-National Research Foundation, Somerset West, Western Cape Province, South Africa UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa (UNISA), Pretoria, South Africa e-mail: [email protected]; [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_6

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undergo decomposition, while those with semiconducting features are effectively utilized in catalytic systems [5–8]. Transition metal oxides (TMOs) can exist in the form of semiconductors, insulators, or superconductors. They have well-defined structures, great electronic features, high work function, excellent optical and magnetic properties, stable mechanical and thermal properties, and an electronic structure that is made up of two separate energy bands [9]. The high work function induces surface inversion in heterojunction cells. It has been shown that the work functions of metal oxides are affected by the material used at the interface and the film quality [10]. Their properties are also affected by defects, the repulsion between the electrons, disordered lattice, and elemental composition. They can manifest different physical features like resistive switching, high temperature, high superconduction, and magnetism. Some of the TMOs of interest include titanium dioxide (TiO2), nickel oxide (NiO), manganese oxide (MnO2), cerium oxide (CeO2), cobalt oxide (CoO), and molybdenum oxide (MoO3). These TMOs exhibit a wide bandgap energy, are mostly chemically unreactive, and have high physical and optical stability, which enhance energy harnessing [11, 12]. To properly synthesize TMOs, the stoichiometry of the oxygen and cation components should be carefully put in check. Several physical and chemical methods have been adopted in synthesizing metal oxides like hydrothermal, sol-gel, pulsed-laser deposition, spin coating, chemical precipitation, sputtering, chemical vapor deposition, and electrodeposition methods [2]. The addition of additives and diverse synthesis methods aid the modification of the structure, morphology, and size of metal oxides. For instance, single metal layers can be deposited on the surface of the oxide to improve its sensitive and conductive nature [13]. Engineering their bandgaps through composite formation or doping improves their overall properties and in particular, light absorption in the visible electromagnetic spectrum [14]. These properties can be understood from experimental investigations and theoretical methods. Such studies shows that the metal oxide surfaces and interfaces are affected by its complex chemical nature, thereby increasing its defect and making the oxide reproduction difficult [1]. The defects present at the surface of these materials are responsible for the varied oxidation states, chemical, and catalytic features exhibited by metal oxides [1]. The surface features of metal oxides can be tailored to improve its interaction with other molecules. Modified surfaces would be more thermodynamically stable with minimal defects, and better suited for application in devices such as solar cells [15]. To improve the electrical conductivity of metal oxides, the intrinsic features of metal oxides are regulated, and metal oxide composites with high electrochemical performance are fabricated. The overall electrochemical performance of metal oxides can be improved by modifying their chemical buildup, architecture, morphology, dopant addition, alloy formation, etc. [16]. The increasing energy demand due to the explosive population rate has made renewable energy sources to be fast-growing in the power sector. The materials required for the synthesis and fabrication processes are available and affordable, especially with the use of earth-abundant TMOs [17]. Silicon heterojunction solar cells are becoming an interesting area of research that substituted the p-n structure.

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The contacts at the heterojunctions serve as semipermeable membranes that selectively allow one carrier to permeate through while blocking the other carrier [18]. To fabricate a TMO/silicon heterojunction, the TMO is deposited on the silicon substrate through any choice synthesis method. The silicon architecture has a thin intrinsic layer that creates a carrier transport pathway. The quality of the heterojunction interface determines the rate and efficiency of photo-generated charges [17]. Heterojunction-silicon solar cells provide selective carrier paths for efficient charge migration. Properly aligning the energy bands of the TMOs and the silicon yields devices with higher fill factor, reduced energy bands, and increased device performance [19]. TMOs can be applied in catalytic devices due to their great redox abilities, making alcohols in chemical factories, surface passivation, conversion of carbon dioxide for energy production, splitting of water during photocatalysis, and gas-sensing devices. TMOs also serve as electrochromic materials and photoelectrodes in solar cells [20].

6.2  Structural Properties of Metal Oxides Studying metal oxide surfaces can be a complicated process due to its complex chemical composition, physical characteristics, various oxidation states, and crystal structure [1]. For instance, vanadium oxide has several phases, and molybdenum oxide exists in different structures. Metal oxides exhibit polymeric structures [3]. The structures of oxides reveal the coordination numbers existing between the metals and their ions. Defects encountered in bulk metal oxides may be point defects, planar defects, or dislocation/line defects. Point defects occur due to interstitial sites, dopant, and ion vacancies. Crystallographic properties of metal oxides can be improved by adding additives, carbon nanotubes, and dopants [2]. The structures of some metal oxides have been outlined in Table 6.1. Structural parameters of metal oxides can be evaluated using advanced computational methods like Hartree-Fock formalism or local density formulae [1]. Ionic models have been successful in qualitatively determining some structural parameters using interionic potential. Ionic models can also resolve defect problems by introducing ions as polarizable quantities. Shell models also handle defect issues by treating the ions singly as negative charges (shells) that are connected to a positive core. Mott-Littleton method aids defect treatment by treating the ions individually using interaction parameters. Browne, Sofer, and Pumera adopted 2D metal oxides for the conversion of electrochemical energy [21]. TMOs fall into this category because it has a flexible bonding structure, different oxidation states, and weak stacks of 2D sheets bonded by Van der Waals forces. The structural features of TMOs give room for exfoliation of these materials into the 2D structure synthesized through diverse methods like conversion process, self-assembly [21]. Huang et  al. synthesized metal oxides with hollow structures for use as catalysts [16]. Hollow structures produce metal oxides with reduced density, wide surface area, more active sites, and flexible chemical and

150 Table 6.1  Metal oxides and their crystal structures

A. C. Nkele et al. Metal oxides ZnO TiO2 Fe3O4 ReO3 TiO2 MoO3 Cu2O NiO MgO ZrO2 Al2O3 CoO CaO CeO2 BaO SnO2

Crystal structures Wurtzite Tetragonal Spinel Cubic Tetragonal Orthorhombic Cuprite Cubic Cubic Tetragonal Hexagonal Rock salt Cubic Fluorite Cubic Rutile

Fig. 6.1  Crystal structure of (a) spinel oxides (b) perovskites [24]

physical features. Spinel oxides like cobalt oxide, perovskites, etc. with hollow structures connect their metal cations with oxygen anions to form tetrahedrons as shown in Fig. 6.1 [16, 22, 23]. Controlling the thickness of the shell and volume of the hollows would increase their activities and performance. Metal oxides with octahedral configurations aid the acceleration of charges during redox reactions [16]. Varied synthesis methods adopted and the numbers of interlayer spacing have enormous influence on the stability and performance of metal oxide catalysts. Templating is a method used to prepare hollow spheres of metal oxides having

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Fig. 6.2  Step-by-step templating approach for hollow metal oxide production. Each templating step is accomplished at an increased temperature. (Adapted from [24])

different structures and buildups as shown in Fig. 6.2 [24]. Engineering the structure of metal oxides at the atomic level greatly affects their physical, capacitive, carrier storage, and chemical features [25]. The engineered structure allows the electronic features of metal oxides to transit from insulating to metallic characteristics.

6.3  Electronic Properties of Metal Oxides The performances of metal oxides are greatly influenced by the mobility of the charge carriers. Electronic features of metal oxides include wide bandgap nature, electronic structure, carrier mobility, effective masses, etc. Some metal oxides like ZrO2 and Cu2O behave differently when doped, unlike most metal oxides as it causes a reduction of the surface [1]. Instead, bombarding them with ions normalizes the stoichiometry at the surface. Surface areas can be activated by minimizing the grain size to nanoscale, such that its gas-sensing features can be enhanced. Band theories have been used to describe the transport mechanisms of metal oxides. Materials with energy bandgaps above 3 eV are optically transparent in the visible region. Low optical absorption, high carrier mobility, and large energy bandgaps produce band structures with mobile carriers between the metal and oxygen orbitals. The large energy bandgaps of metal oxides can be reduced by integrating with plasmonic materials like silver so that the bandgap can be minimized toward the visible region [26]. Silver is relatively cheap, stable, shifts optical absorption to the visible spectrum, and produces high performance. Karimi-Maleh et al. synthesized titanium dioxide, zinc oxide, and nickel oxide integrated with silver and obtained high photocatalytic activities and low bandgap energies [26]. Wetting occurring at the interface of metal oxides influences the rate of adhesion of the metal oxide. The endothermic chemical reactions between oxygen and metals

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develop bonds that form metal oxides with minimal entropy and surface energy. The pressure, surface energy, temperature, and partial pressure of oxygen influence the bond existing between the metal and oxide. The independent features of the metals and their oxides affect the bulk properties of the synthesized materials. The kind of bonds that spring adhesion could be dispersive, diffusive, electrostatic, or chemical [27]. The number of bonds existing at the interface between metal and oxide is influenced by the roughness at the surface. Metal oxide interfaces are heterogeneous, crystallographically oriented, suffer defects at the surface, and suffer impurity and dislocation effects. Electronic properties of metal oxides can be modified by creating oxygen vacancies at high-voltage areas or using different synthesis methods so that carrier movement and electrochemical activities can be improved [28]. Hydrogenation can also be employed in modulating the electronic characteristics such that the photocurrent, light absorption, and stability can be enhanced.

6.4  A  pplication of Some Transition Metal Oxides in Solar Cells Some of the TMOs that find useful application in solar cells include titanium dioxide (TiO2), nickel oxide (NiO), manganese oxide (MnO2), cerium oxide (CeO2), cobalt oxide (CoO), and molybdenum oxide (MoO3).

6.4.1  Titanium Dioxide, TiO2 Titanium dioxide is an n-type semiconductor that exhibits bandgap energy that is dependent on its crystalline structure. It can exist in either the anatase, rutile, or brookite phases [7]. Each of the phases exists at specific temperature ranges and crystallization structures. The anatase and rutile phases exist as tetragonal, while the brookite phase exists in orthorhombic structure [29]. The high bandgap nature of the anatase phase hinders optical absorption in a wide range of the electromagnetic spectrum and affects the crystalline structure of the titanium dioxide film [30]. This limitation can be resolved by modifying its features such that light absorption can occur over a wide range of the electromagnetic spectrum. TiO2/Si heterojunction enhances high electron mobility without hole interference, to increase the efficiency of the solar cell device [31]. TiO2 can be synthesized through the evaporated electron beam method, sputtering technique, pulsed-laser deposition, chemical vapor deposition, and sol-gel method [32]. Titanium dioxide finds useful application in photodiodes and solar energy devices [33–35]. As shown in Fig.  6.3, TiO2/Si heterojunctions synthesized via radio frequency sputtering exhibited pyramid-like morphology, amorphous nature when grown on a silicon substrate that transformed into a crystalline structure upon annealing at 673  K, reduced transmittance and

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Fig. 6.3  Exhibited (a) morphological, (b) structural, (c) transmittance, and (d) band gap energy features of the TiO2/Si heterojunction film [37]

energy bandgaps after annealing, and hole-blocking effects at the junctions [36]. This morphology minimized surface reflectance, reduced transmittance, and bandgap energy could be attributed to annealing effects and density of film, while the crystalline structure influenced electron transport and optical features of the heterojunction formed.

6.4.2  Nickel Oxide, NiO Nickel oxide is a p-type semiconductor, abundant, relatively cheap, highly conductive, easily processed, improved carrier transport mechanism, thermally, and chemically stable [37]. Several syntheses methods have been adopted in the formation of nickel oxide [6, 38, 39]. Nickel oxide is a wide bandgap material, electrochemically stable, durable, and exhibits a small electronic conductivity. The conductivity of NiO is affected by interstitial sites and oxygen vacancies [40]. It exists in different oxidation states and morphologies like nanosheets, nanoflakes, nanotubes, nanowires, etc. Figure 6.4 shows the flow processes involved in forming a microstructure

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of nickel oxide [41]. The porous morphology is due to the formation process employed. NiO films can be usefully applied in solar cells [40]. NiO/silicon heterojunction fabricated by oxidizing silicon-deposited nickel by ultraviolet light yields more conducting films [42]. Fabricating heterojunction of nickel oxide and silicon at an annealing temperature of 400 °C produced an energy bandgap of 3.69  eV and high optical transmittance [43]. At the silicon junction, nickel oxide serves as an efficient electron-blocking material. The photodetector heterojunction exhibited good photoresponse, excellent current-voltage features, and good photovoltaic activities under illuminated light. Pulse-deposited NiO/Si heterojunction cells showed increased current when the magnetic field was externally applied [44]. Yang et  al. fabricated NiO/Si solar cells and obtained an increased built-in voltage, fast hole carrier transport, and power conversion efficiency of 9.1% [45].

6.4.3  Manganese Oxide, MnO2 Manganese is a metal that is environmentally friendly, nonpoisonous, exists in various oxidation states (+2, +3, +4, +6, +7), and forms mixed oxides with metals like iron, niobium, tantalum, etc. [46]. Manganese oxide is relatively inexpensive, less poisonous, naturally abundant, and its structure is normally crystalline [41]. MnO2 films can be synthesized via several methods [47, 48]. Varying the preparation conditions of the manganese precursor-like ambient temperature, hydrogen, surrounding air, etc. yields various structural oxides. The morphology and structure of manganese oxide films can be modified during the synthesis on silicon substrates to enhance its electrocatalytic performance [49]. The report shows the obtained results of manganese oxide films synthesized on silicon substrates by thermal evaporation

Fig. 6.4  Mechanisms involved in forming nickel oxide microstructure [42]

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under controlled annealing situations, with the MnO2/Si films annealed in vacuum had higher relative permittivity than those annealed in the air, which yielded no silicates [50]. Manganese oxide films find application in perovskite solar cells.

6.4.4  Cerium Oxide, CeO2 Cerium oxide is a rare earth metal oxide with a fluorite structure and diverse synthesis methods [51, 52]. Increasing the concentration of oxygen defects increases the conductivity of the ions by improving its rate of diffusion. Its hydrophobic nature allows it to readily adsorb organic compounds for better catalytic performance [53] and is optically transparent to visible light [54]. Cerium oxide exhibits a similar lattice constant and crystal structure as that of silicon [55]. Figure 6.5a, b revealed the nanocrystalline structure of the films, while Fig. 6.4c shows electronic transition as evident from the prominent absorption peak [56]. Synthesizing cerium oxide on the silicon substrate enhances effective charge transfer [55]. Cerium oxide can be applied in optical and solar cell devices [56].

Fig. 6.5 (a, b) Electron diffraction patterns and (c) spectroscopic ellipsometry spectra of the CeO2 films [57]

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6.4.5  Cobalt Oxide, CoO The oxides of cobalt exist in three different forms: vis cobalt (I) oxide, cobalt (II) oxide, and cobalt (III) oxide [57, 58]. In appearance, cobalt oxide is a black solid and its bandgap energy is approximately two electron volts. It is known to be resistant to corrosion and exhibits high performance and efficiency in optoelectronic devices [41]. It is a p-type semiconducting material with diverse synthesis methods and highly efficient hole extraction abilities [59]. The hole extraction and transport mechanisms of cobalt oxide in perovskite solar cell upon solar irradiation have been shown in Fig. 6.6. CoO films can also be applied in solar cells [57].

6.4.6  Molybdenum Oxide, MoO3 Molybdenum oxide, MoO3, is an efficient hole extractor with high work function, wide energy bandgap, small coefficient of absorption, low synthesis temperatures, etc. Molybdenum oxide can be synthesized via several methods [60]. It is an active oxide for decomposing hydrocarbons [60]. Molybdenum oxide layers when fabricated on silicon heterojunction cells manifested high transmittance, increased power conversion efficiency, high fill factor, and open-circuit voltage values. From Fig. 6.7, the heterojunction cell exhibited a charge-permeating surface, crystalline structure, reduced bandgap energies, a current density of 37.18  mAcm−2, and improved efficiency of the MoO3 cell having a silicon emitter [60]. MoO3 films find applications in organic and inorganic solar cells.

Fig. 6.6  Diagram showing the efficient hole transport mechanism of cobalt oxide [60]

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Fig. 6.7  The (a) surface, (b) structural, (c) bandgap energy, and (d) J-V features of the molybdenum oxide film [61]

6.5  C  harge Transport Mechanism in Metal Oxide/Silicon Solar Cells Heterojunction solar cells use thin wafers due to low-temperature processing techniques. These cells have minimized recombination effects and increased efficiencies, unlike those of homojunction cells. To increase the efficiency of solar cell devices, the ideality factor and voltage at maximum power point should be reduced. Heterojunctions create an interface that selectively allows carriers to be transported to the absorber material. Thermionic emission triggers the motion of the charge carriers [10]. This mechanism of transporting charges allows one kind of charge to pass through the interface while the other charge is restricted. Electron collecting materials like TiO2 and LiF and hole transport materials like WOx and MoOx have been deposited on silicon substrates with a high power conversion efficiency of 14.7% [61]. The high efficiencies obtained from silicon heterojunction cells are due to the layered stacks of passivating contacts that induce a potential at the surface for carrier transport [10]. The induced surface potential allows the light-­generated positive carriers to migrate to the silicon front surface and recombine with electrons at the interface. Incorporation of wide bandgap metal oxides helps to regulate the device

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photocurrent [10]. The wide bandgap nature of metal oxides gives room for high optical transparency. Efficient hole transport mechanisms are affected by the absence of charge barriers at the hole layer and the surface inversion rate at the silicon wafer [10].

6.6  M  ethods of Improving the Efficacy of Transition Metal Oxides The efficient performance of TMOs is affected by charge carrier recombination and surface reflective losses [62, 63]. The recombination effects can be minimized by creating a heterojunction between the semiconductor and the silicon material, while surface reflective losses can be reduced by incorporating surfaces with high refractive indices [64].

6.6.1  Addition of Dopant Doping TMOs would enhance hole transport from the silicon-absorbing material to the conductive surface. The cations in the TMOs can also be doped to tailor oxygen defects. The electrical conductivity of NiO can be enhanced and its resistivity reduced by doping with compounds like graphene, carbon nanotubes, and metal oxides. Doping nickel oxide also enhances the hole transport properties and improved efficiency, especially in perovskite solar cells [22]. Doping cobalt oxide with manganese or lithium helps to improve its crystalline structure, reduce particle size, and reduce the energy bandgap [65, 66]. The increased sensitivity of cobalt oxide upon doping makes it applicable in solar cell devices. Doping cerium oxide leads to increased diffusion rate and electronic conductivity.

6.6.2  Formation of Composites TMOs can combine with other materials to form composites. NiO can form composites with graphene, carbon nanotubes, other TMOs, and activated carbon. Nickel oxide-based composites improve their conductivity, stability, surface area, and energy storage ability [40, 67]. Mixed composites of metallic oxides also yield better performing devices [68]. Manganese oxide has high operating voltage and can form composite compounds with lithium and be potentially applied as cathode materials [49]. Manganese oxide can form composites with activated carbon and carbon nanotubes to yield a better performing material [69].

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6.6.3  Heat/Plasma Treatment The oxygen defects in TMOs can be tailored by heat treatment, especially in gaseous atmospheres, and the solution process adopted. Plasma treatment is a complex chemical and physical technology that eliminates film defects and improves adsorption at the surface, as it involves excited carriers, energized electrons, ions, and free radicals [70]. Subjecting the TMO films to plasma treatment creates vacant oxygen sites at the film surfaces, enhances the surface morphologies, crystal structure, and performance of the films [70].

6.6.4  Electroplating Electroplating films enhance their charge-retention ability by sandwiching the metal electrodes on the silicon substrates [71]. Manganese oxide films can be loaded on silicon substrates to form porous and conductive materials for use in solar cell devices [72].

6.7  Conclusion Research interest in metal oxides is growing steadily because of their high carrier mobility, relatively cheap, optical transparency, highly sensitive, and tolerance to mechanical stress. This chapter has given a general knowledge of metal oxides and their individual cum bulk properties. The structural and electronic characteristics of metal oxides have also been discussed. On the other hand, the unique properties of transition metals like well-defined structures, great electronic features, high work function, excellent optical and magnetic properties, and stable mechanical and thermal features have prompted enormous interest in optoelectronic devices incorporating these materials. Heterojunction architectures are important in establishing high efficiencies in solar cell devices with TMOs. They serve as the future hope for window layers in solar cells. Silicon heterojunction solar cells are interestingly researched upon and are substituting the p-n structure. This chapter has discussed some TMOs like titanium dioxide, nickel oxide, manganese oxide, cerium oxide, cobalt oxide, and molybdenum oxide. The features, synthesis procedures, and obtained experimental results were also discussed. The heterojunction interface of these transition metals when deposited on silicon substrates creates a carrier transport channel that enhances charge transfer operations. The performance of TMOs can be enhanced by adding dopants, forming composites, applying heat, treating with plasma, and electroplating. The properties of the TMOs have revealed interesting features that make them applicable in solar cell and optical devices.

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Chapter 7

Optically Active Metal Oxides for Photovoltaic Applications A. C. Lokhande, V. C. Lokhande, D. S. Dhawale, I. A. Qattan, Shashikant Patole, and Chandrakant D. Lokhande

7.1  Introduction The need for global energy demand with clean energy sources for various applications has increased at an astonishing rate [1]. The primary energy sources such as fossil fuels, natural gas, and coal are extensively used for satisfying the existing domestic and commercial energy needs [2]. However, these energy sources are limited and will be depleted in the coming future. Also, the consumption of these energy sources leads to serious problems such as global warming and pollution. Hence, there is a great need to develop alternative energy sources that can overcome these existing problems. In 2015, 195 countries reached an agreement to control greenhouse gas emissions, confront the impacts of climate change, and develop alternative energy sources at the Paris climate conference. Since then, efforts have been made to develop sustainable energy sources for human development. In recent years, the photovoltaic industry received significant research attention because of its potential to satisfy the current energy needs. The photovoltaic industry is exclusively based on the use of solar energy that is abundant, clean, renewable, and inexhaustible. As per the report of the Joint Research Centre (Europe), at present, the total worldwide energy generation from the photovoltaic industry nearly accounts A. C. Lokhande (*) · I. A. Qattan · S. Patole Applied Quantum Materials Laboratory (AQML), Department of Physics, Khalifa University of Science and Technology, Abu Dhabi, United Arab Emirates V. C. Lokhande Department of Electronics and Computer Engineering, Chonnam National University, Gwangju, South Korea D. S. Dhawale · C. D. Lokhande Centre for Interdisciplinary Research, D. Y. Patil Education Society (Deemed to be University), Kolhapur, India e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_7

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Fig. 7.1  Global energy production in the last 3 years. (Source: International Energy Agency)

for 20% and is predicted to reach over 50% in the next 60 years [3]. In the year 2017, 98.9 GW of usable energy was produced from photovoltaic, which was further increased to 629 GW in 2019 (Fig. 7.1), and is expected to increase further in the coming years. Hence, it is quite clear that the photovoltaic industry is emerging at an alarming rate and is competent enough to solve energy-related problems for the development of human society. Generally, the solar cells form the basis of the photovoltaic system for energy generation. The solar cell consists of a p-n junction heterostructure, in which the p-type semiconductor acts as an absorber layer and the n-type semiconductor as a buffer layer that converts the incident light into electricity through the electron-hole transport mechanism. The solar cells are fabricated on a single layer (single junction) or multiple layers (multijunction) of light-absorbing materials to exploit the advantages of induced charge generation processes. With the consistent development in the research technology, the production cost of solar cells has reduced by 70% within a decade, thereby emerging as a reliable and cost-effective system for the energy industry [4]. The solar cells are classified into three categories, namely (I) the first-generation, (II) the second-generation, and (III) the third-generation cells (Fig. 7.2). The first-generation solar cells are based on semiconductor crystalline silicon (c-Si) that exists in a single crystal (monocrystalline) and multiple crystals (polycrystalline) forms. Currently, the solar industry market is based on first-generation silicon (Si) solar cells that have achieved a record high efficiency of 26.7% [5]. However, there are several issues with Si solar cells for commercialization such as the necessity of thick absorber layers (100–200 μm), high-cost complex processing methods, the existence of impurities, and poor performance under diffused light conditions [6]. On the other hand, the second-generation solar cells are based on

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Fig. 7.2  The classification of solar cells. (Redrawn from [4])

thin-film technology, where the solar cells have a lower absorber layer thickness (1–2 μm), operate efficiently in diffused light conditions, exhibit simple processing steps, and are relatively cheap as compared to Si solar cells [1]. Such unique advantages render the second-generation solar cells as the preferred choice for the photovoltaic industry. The second-generation solar cells based on the chalcogenides (sulfides, selenides, and tellurides) such as CIGS (CuInGaS) and CdTe absorber compounds have attained the highest efficiencies above 20% that is comparable to the efficiency of Si solar cells [7, 8]. However, the existence of toxic (Cd) and rare-­ earth elements (In, Ga, and Te) in these compounds produce costly and toxic solar cells, thereby restricting their scope for mass commercialization. Alternatively, the recently developed Cu2ZnSnS/Se4 (CZTSSe) absorber compound has gained attention for thin-film solar cell applications. CZTSSe is a kesterite compound composed of earth-abundant elements and exhibits similar optical and electrical properties to its counterpart (CIGS and CdTe) solar cells. The highest reported efficiency based on CZTSSe-based thin-film solar cell is 12.6%, processed using the hydrazine-­ solution approach [9]. However, this achieved efficiency is quite low as compared to its counterparts (CIGS and CdTe) and needs to be enhanced further. Even though CZTSSe is an emerging photovoltaic material, there are several challenges associated with it for mass commercialization. The existence of a toxic element (Se) and the involved hazardous chemical (hydrazine) processing imparts a serious threat to human health and the environment [10]. Also, the existence of various intrinsic defects and secondary phases significantly alter the physical, electrical, optical, and chemical properties, thereby restricting it from delivering efficient performance. The third-generation solar cells such as perovskites, dye-sensitized solar cells (DSSC), organic solar cells, and quantum dot solar cells have also been actively researched [4]. These solar cells are the extended version of second-generation solar cells as they employ thin-film optical layers with slight modifications in the device

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architecture and components. Though these solar cells have achieved reliable performance, several challenges such as high cost and device stability issues are still needed to be addressed for their mass commercialization. The primary goal of the photovoltaic industry is to produce thin-film solar cells that are cost-effective, nontoxic, exhibit simple processing steps, and deliver efficient performance [11]. Hence, it is vital to properly select the active materials that form the components of the solar cell. With this given background, metal oxides can be considered as emerging materials for thin-film solar cell applications due to their unique optical, electrical, structural, chemical, and physical properties. It has been reported that substituting a stacked layer with oxide results in improved solar cell efficiency (>1%) [12]. Moreover, the amalgamation of various properties such as cost-effectiveness, nontoxicity, high absorbance/transmittance, efficient charge generation/separation, high structural/chemical stability, and defect passivation enables metal oxides to be used as active material in various stacked layers of thin-­ film solar cells. Thus, it is clear that metal oxides have a significant role in the development of thin-film solar cells and are worthy of further research. Hence, the present chapter is aimed to provide a brief insight into the current status, ongoing developments, associated issues/solutions, and potential future of metal oxides in thin-film solar cells.

7.2  Structure of Thin-Film Solar Cells Before proceeding further on the use of metal oxides in photovoltaic, it’s important to understand the basic structure of thin-film solar cells. As seen in Fig. 7.3a, the typical thin-film solar cells are fabricated in a stacking structure, where a specific layer with a special function is deposited over other layers.

Fig. 7.3 (a) The typical structure of thin-film solar cells and (b) the schematic representation of the working mechanism of the solar cell

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• At the bottom, a conducting substrate is used (back contact). Its role is to provide support to the entire solar cell structure and to extract the photo-generated charge carriers to the external load. • The absorber layer absorbs the incident sunlight (photons) and converts into charge carriers. Generally, p-type absorber layers are used to generate holes. • The buffer layer also functions as the absorber layer. N-type buffer layers are used to generate electrons. Also, the buffer layer avoids shunting paths between the absorber layer and the transparent conducting oxide (TCO) layer. • The TCO layer promotes the passage of visible light to the p-n junction for electron-­hole pair generation. TCO also extracts the generated charge carriers to the external load through top metal contact. • In case if the absorber layer is too thin to absorb the incident light, back reflectors are used to avoid the wastage of incident light and divert it back to the absorber layer. • Antireflection coatings (ARC) are used to reduce the effects of backward reflection losses by enhancing the refractive index of incident light as it travels through different mediums (air to absorber layer). The basic working mechanism of the solar cell is based on the photovoltaic effect in which a potential difference is generated across the heterojunction in response to the incident light (electromagnetic radiation). As the light is incident at the p-n junction, the electron-hole pairs are generated. The generated electron-hole pairs flow across the junction resulting in net current flow (Fig. 7.3b). The generated charge carriers are then extracted by the metal contacts.

7.2.1  I deal Material Properties Requirement in Thin-Film Solar Cells To identify suitable materials for the solar cell application, it is important to understand the properties required to develop efficient solar cells. As most of the materials used in the solar cell are semiconductors, their optical, electrical, and structural properties should be considered. Also, as the solar cells are expected to work with full efficiency for 20  years, their mechanical and chemical stabilities (resistant against harsh environment) are important. The desired properties of the various components of thin-film solar cells are discussed below. 1. Back and Top Metal Contacts: The primary function of the metal contacts is to extract charge carriers generated in the semiconductors and supply it to the external load. To accomplish this task, the metal contact should have high electrical conductivity and high work function. As per the semiconductor theory, the significance of high work function is such that it enables the formation of ohmic contact with the semiconductor [11]. As the ohmic contact between the heterostructure is made (metal semiconductor), the Fermi level aligns perfectly with

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the band bending of valence and conduction band edges, thereby promoting efficient charge transfer paths [13]. Typically, metal contacts with high work functions such as Ni (5.12 eV), Mo (5 eV), and Al (4.28 eV) are widely used for charge extraction [14]. The exact work function values of the absorber material are unknown, but as per some simulation reports, its value is generalized to ~5  eV.  Hence, for efficient charge extraction, metal contacts exhibiting work function values above 5  eV should be used. The metal contacts must exhibit high-temperature stability and high chemical inertness as to not form any undesirable products or introduce secondary phases in TCO and absorber layers during the processing of solar cells. The coefficient of thermal expansion of the metal back contact must be similar to that of the absorber layer to maintain good surface adhesion and promote efficient charge transport at the interface. In addition, the metal contact should exhibit high reflectivity and smooth surface to reflect the incident light passing through the absorber layer for enhanced photo-­ carrier generation. 2. Absorber Layer: Absorber layer forms the heart of thin-film solar cells. As the incident photon energy is converted into charge carries by the absorber layer, its optical properties should be considered of prime importance. A typical absorber compound should have an optimal band (1–1.5 eV) with a high absorption coefficient (10−5 cm−1). In addition, the absorber layer thickness should be sufficient enough to completely absorb the incident photon energy to avoid optical losses. Literature reports have suggested the optimum absorber layer thickness between 1 and 2 μm [1, 15]. The internal defects in the absorber compound should be minimal as possible to reduce charge recombination losses. The microstructure of the absorber must be compact and smooth as it helps to form an intimate contact with the n-type buffer layer for efficient charge carrier transport across the junction (Fig. 7.4). If the microstructure is porous, it generates numerous rough surfaces, where the effect of charge-trapping becomes prominent. The structural and compositional properties of the absorber compound should also be

Fig. 7.4  The schematics of contact between P-N junction interfaces. (Reprinted with copyright permission from [1])

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c­ onsidered. The crystallinity of the absorber compound should be high as it improves optical absorption ability and enhances minority carrier lifetime. Furthermore, close control of the absorber elemental composition must be maintained as the compositional inhomogeneities generate numerous secondary phases and defects, thereby affecting optical and electrical properties. 3. Buffer Layers: Like absorbers, the buffer layer properties are equally important as they also accomplish the task of carrier generation. For efficient carrier generation, the bandgap of the buffer layer should be between 2 and 3 eV. The suggested buffer layer thickness should be between 30 and 70  nm to reduce the negative impacts of series resistance in solar cells. The proper selection of the buffer layer is very important as its band alignment with the absorber layer must be compatible. Usually, when a junction is formed (p-n), either a “cliff-like” or “spike-like” band structure is formed at the interface. The “cliff-like” band structure is undesirable as it reduces interface bandgap and lowers the activation energy than bandgap resulting in interface recombination at various trap states [16]. On the contrary, the “spike-like” band structure is desired as it equalizes the activation energy and bandgap and reduces interface recombination by promoting efficient electron-hole transfer by thermionic emission [1]. In the photoelectron yield spectroscopy (PYS) study conducted by Sato et al. [17], it was proved that the CdS buffer layer forms a “cliff-like,” while the oxide buffer layers such as IZO, ZTO, SnO2, ZnO form a “spike-like” band structure with chalcogenide Cu2SnS3 (CTS) absorber compound [17]. Hence, the proper selection of a buffer layer with the absorber layer is necessary for realizing an efficient solar cell performance. 4. Transparent Conducting Oxides (TCOs): The TCOs perform two major tasks in thin-film solar cells, namely (a) charge extraction from the semiconductor p-n junction and its transfer to metal contacts for external load and (b) protecting the underlying stacking layer from external effects (dust and humidity) and allowing only visible light to enter into the solar cell. TCOs should have a high electrical conductivity (>103 S/cm) and wider bandgap (>3.2 eV) with an optical transmissivity above 90%. Besides these properties, other important factors to be considered are as follows: 5. Material Availability and Processing: The material used for the photovoltaic technology must be abundant to exploit its usage for mass scale. Also, the processing of the material should be less energy-intensive, simple, and cost-­effective to make it compatible with the current technological demands. The terms “material availability and processing” are strongly interrelated and go hand in hand. For example, silicon is the second most abundant material with active photovoltaic properties. However, it cannot be directly used in solar cells as it requires various complex processing steps (extraction from ores and purification) to produce its pure form that is suitable for photovoltaic. In such a case, the processing may be costlier than the other available less abundant materials [12]. Hence, the solution would be to use minimal material as possible while maintaining its photovoltaic properties or develop novel cost-effective processing systems.

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6. Material Toxicity: Material toxicity is one of the important factors that need stringent regulations. The material and its associated processing steps should be environmentally friendly to ensure the safety of humans, flora, and fauna life. 7. Environmental and Chemical Stability: The commercial solar cells are expected to work efficiently for 20  years. During this longtime span, solar cell performance is partially or completely degraded due to environmental attacks such as temperature variation and humidity (moisture). The temperature variation alters the junction properties between various stacking layers. This results in the reduction of the voltage output across the junctions, and the net effect results in lower power output. The moisture attack is more prominent as it reacts with various components of solar cells and forms undesirable oxide compounds that significantly reduce the performance. Therefore, the material employed for the solar cell application must have a high-temperature coefficient and high oxidation resistance. In case of a material does not exhibit these desired properties, its performance can still be maintained by using encapsulators for moisture prevention and cooling systems for heat dissipation.

7.3  Metal Oxides in Solar Cells Metal oxides have an important role in the development of solar cells since the inception of photovoltaic technology. Metal oxides are oxygen-based compounds and are formed when a metal/metalloid donates two electrons to the 2P orbital of an oxygen atom that forms the valence band of the oxide. Oxygen is a naturally occurring element and exists in two forms, namely diatom (O2) or ozone (O3). Based on the oxidation states of oxygen (O−2, O2−2, and O2−1), it is further classified into oxides, peroxides, and superoxides [12]. The electrical and optical properties of the oxides are mainly dependent upon the type of bonding that exists in its atomic structure. Usually, oxides with ionic bonding have high dielectric constant value and are insulating in nature (wider optical bandgap), while the oxides with covalent bonding have mediocre electrical conductivity (narrow optical bandgap) and can be used in photovoltaic applications [18]. The compositional variation (under or over stoichiometry), oxidation states, and the coordination structure are also believed to influence the properties of metal oxides. The nontoxic nature, high natural abundance, and free availability of oxygen make oxides highly suitable in solar cell applications for long-term sustainable development. The oxides are classified based on the metals (with similar properties) with which they form compounds such as alkali metals, rare-earth metals, refractory metals, ferrous metals, precious metals, fusible metals, semimetals, and mixed metals. In the review article on the “applications of metal oxides in photovoltaic devices” by Sonya Calnan, a detailed classification on the types of oxides is highlighted [12]. Therefore, the readers are encouraged to peruse the article for further insights. Alkali metal oxides based on all group 2 elements are unstable and highly reactive with moisture. Furthermore, their processing and handling is difficult, and thus

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cannot be employed in solar cell applications. Rare-earth metal oxides exhibit unique optical properties and are emerging materials for solar cells. However, their high cost due to lower natural abundance and the involved complex processing steps restrict their application in the solar industry on a commercial basis. Refractory metal oxides are based on the use of various transition metals such as Cu, Ti, W, V, Hf, Nb, Mo, etc. [12, 19–23]. The variable oxidation states of these transition metals impart intriguing optical and electrical properties to their oxides. Metal oxides such as CuO and Cu2O exhibit p-type conductivity with tunable bandgaps in the range of 1.4–2.2  eV [24, 25]. These metal oxides have been actively used as an absorber layer in thin-film solar cells. Other metal oxides such as TiO2, MoO2/3, and WO3 have been widely used as a buffer and/or TCO layer. TiO2 exhibits high visible light transparency and exhibits n-type conductivity due to the existence of interstitial Ti atoms and oxygen vacancies. MoO2 exhibits metallic conductivity, while the MoO3 exhibits an insulating/semiconducting nature (wide bandgap) due to the existence of excess oxygen vacancies [20]. The stoichiometric WO3 is a wide bandgap stable semiconductor (transparent to visible light) with high conduction electron concentration and could be used as an active TCO layer [26]. Ferrous metal oxides are based on ferrous materials such as Fe, Co, Cr, and Ni. Nickel oxide exhibits p-type conductivity with wide bandgap (>3.5 eV), high work function (>5 eV), and high chemical stability, enabling it to be possibly used as a TCO layer. The oxides of Fe, Co, and Cr absorb visible light and are used as absorber layers [12]. Precious metal oxides based on Ag and Au are used in solar cells due to their high conductivity and high reflectivity [27, 28]. However, the higher costs of these metals restrict their application in industrial systems on a large scale. Fusible metal oxides are based on fusible metals that have a relatively low melting temperature. These metals include mercury, zinc, cadmium, indium, gallium, tin, antimony, bismuth, and lead. The fusible metal oxides are wide bandgap semiconductors (>3 eV) and are widely used as TCO in solar cells [12]. Oxides such as SnO, Bi2O3, and PbO exhibit p-type conductivity [18], while oxides such as ZnO, In2O3, Ga2O3, and CdO exhibit n-type conductivity [29]. Currently, the research is mostly focused on developing TCOs with high transparency and high conductivity. This objective is achieved with n-types semiconductors, while the research on developing p-type semiconductors is still in the primordial stage. The n-type ZnO semiconductor exhibits wide bandgap (>3 eV), and the intrinsic i-ZnO is obtained by Zn interstitials and/or oxygen vacancies. It has also been reported that intrinsic doping is attained due to various preparative experimental conditions and unintentional hydrogen doping. Extrinsic doping is attained in ZnO by Al [30], B [31], and Ga [32] doping at Zn sites. Efforts have been made to prepare p-type ZnO by nitrogen doping. However, the attained electrical conductivity of the fabricated p-ZnO is low and needs to be enhanced further for realizing its practical application [30]. SnO2 is an n-type semiconductor with a wide bandgap of 3.5 eV. When fluorine (F) is doped in it, the F− ion displaces O2− ion and a free electron is released in the conduction band, and the conduction band minimum (CBM) shifts close to the Fermi level, thereby reducing its bandgap [33]. The fluorine-doped SnO2, that is, FTO exhibits high electrical conductivity and high work function (4.5–5.5 eV) making

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them suitable in solar cells as back contacts. Similarly, tin-doped In2O3 (ITO) exhibits similar properties to FTO and is also widely used as back contacts in solar cells. Tin doping in other metal oxides also results in enhanced electrical properties. Tin doping in Ga2O3 results in improved electrical conductivity while maintaining its wide bandgap (>4.85  eV) [34] and in the case of CdO, its electrical mobility is enhanced making them suitable for TCO and buffer layer applications, respectively [35]. However, these oxides cannot be used in the industrial system due to their toxic nature. Semimetal oxides are based on metals such as Si, C, Al, Ge, and B [12]. These metal oxides exhibit ionic bonding, and thus are insulating in nature. These metal oxides are generally used as diffusion barrier layers and for mechanical/chemical protection. Especially, the Al2O3 is highly preferred as it exhibits wide bandgap, high chemical stability, high hardness, and high dielectric constant [36]. Mixed metal oxides are based on multiple elements in which either one is a transition metal. These metal oxides exhibit unique properties due to multiple elements and their associated oxidation states. Mixed metal oxides such as Al: ZnO [37], Al: Zn0.87Mg0.13O [38], B: Zn0.88Mg0.12O [39], In2ZO4 [40], Cd2SnO4 [41], and CdIn2O4 [42] exhibit wide bandgap. Especially, the Cd-based oxides exhibit high conductivity and could be used as an effective TCO material [12]. Other mixed metal oxides, namely LaTiO3 [43], YTiO3 [43], BiFeO3 [44], KNbO3 [45], and LaAlO3 [45], have been effectively used in perovskite solar cells. Thus, unlike metal chalcogenides, the metal oxides exhibit versatile functions (as back contact, the absorber layer, buffer layer, and TCO), and hence have a great potential in advancing the solar cell technology for a sustainable future.

7.4  Application of Metal Oxides in Thin-Film Solar Cells As discussed above, it is clear that metal oxides have versatile functions in solar cells (first-, second-, and third-generation solar cells). As per the scope of the present chapter, only those metal oxides that have been used in second-generation thin-­ film solar cells are considered for a brief discussion.

7.4.1  M  etal Oxides as Back Contact and Intermediate Barrier Layers in Thin-Film Solar Cells Generally, metal contacts are expected to have a high work function to form an ohmic contact with the absorber layer for efficient charge transport [46]. The well-­ reported metal oxides, namely ITO and FTO, exhibit the desired properties and have been used as back metal contacts in thin-film solar cells. Comparatively, ITO exhibits more electrical conductivity than FTO and has shown positive impacts in terms of enhanced solar cell performance. Hossain et al. [47] demonstrated the impact of

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Fig. 7.5 (a) The J − V characteristics of CdTe-based thin-film solar cells based on ITO and FTO substrates, (b) the cross-section FESEM image of CZTSe solar cell based on 80 nm MoSe2 intermediate barrier layer, and (c) the schematic representation of the flat band structure of CdTe cells with/without back contacts. (Reprinted with copyright permission from [47–49])

metal back contacts on the performance of CdTe-based thin-film solar cells. It was found that the ITO back contact-based solar cell exhibited slightly better performance than the FTO back contact solar cell due to lower series resistance (Rs) (Fig. 7.5a). Moreover, the ITO, as compared to the FTO, servers as an effective barrier for the undesired hole flow, thereby enhancing shunt resistance, fill factor, and the voltage output. Besides electrical conductivity, the thermal stability of the back contacts is also important. ITO back contacts are prone to “elemental diffusion” (In and Sn) in the adjacent absorber layer during the processing of solar cells through thermal treatments. This results in the formation of spurious phases that significantly reduce the solar cell performance. In such cases, an intermediate barrier layer can be used to restrict the diffusion and restore the performance. In the case of bifacial CZTSe solar cell using ITO as back contact, Mo intermediate barrier layer

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used by Kim et al. [48] enabled to attain a high efficiency of 5.21% by forming a thin ohmic MoSe2 layer (80 nm) as compared to the reference solar cell (without Mo layer) that attained a power conversion efficiency of 0.40% (Fig. 7.5b). It should be noted that the thickness of the MoSe2 layer is very important. The thin MoS/Se2 layer (100  nm) forms nonohmic contact (rectifying) that introduces serious electrical resistance and reduces the solar cell performance. In cases where the formation of thick MoS/Se2 layers is unavoidable, MoO2 can be used as an intermediate barrier layer, as it exhibits high work function, high conductivity, and nonreactivity with S or Se vapors. The MoO2 intermediate barrier layer used in CZTS solar cells significantly reduced the formation of thick MoS/Se2 layers and enhanced the performance by improving the minority carrier lifetime (from 2.9 ns to 7.9 ns). Furthermore, the MoO2 layer promoted Na diffusion that significantly enhanced the absorber layer qualities by forming compact and void-free microstructure with the uniform elemental distribution. Other metal oxides such as ZnO, WO3, SiOx, and Al2O3 have also been reported as intermediate barrier layers in thin-film solar cells. Two-dimensional (2D) material such as reduced graphene oxide (RGO) has also been used as a back contact in thin-film solar cells. RGO exhibits high mechanical/ chemical/temperature stability, high transparency, and better hole collection efficiency, ideally suitable for back contacts. The unique atomic structure arrangement (sp2 cluster form) in RGO promotes efficient charge injection and transport with the contact metal electrode. The oxygen functional groups in RGO tend to reduce its electrical conductivity, which can be increased through various elemental doping approaches. Considering these properties, Zhu et  al. [49] employed Sb-doped RGO-Au back contact in CdTe solar cells and achieved efficiency as high as 13.6% as compared to 8.8% efficiency obtained for the pristine Au back contact. The enhanced performance was attributed to the synergistic effects, in which the Sb enhanced the electrical conductivity of RGO, the intense chemical interactions between the delocalized C 3pz orbitals of RGO and the polarized 6s/6p valence orbitals of Te resulted in efficient carrier (hole) injection from CdTe absorber to Sb-RGO back contact, and the formation of ohmic contact between the CdTe absorber and the Sb-RGO-Au back contact as compared to the Schottky contact formed with the pristine Au contact (Fig. 7.5c). Hence, it is clear that metal oxides have shown reliable performance in thin-film solar cells as back contact and intermediate barrier layers.

7.4.2  M  etal Oxides as Absorber Layers in Thin-Film Solar Cells Among the reported metal oxides as absorbers in thin-film solar cells, copper oxide has attracted significant attention due to its suitable properties. It exists in two forms, namely CuO and Cu2O. CuO has p-type conductivity with high carrier con-

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centration (1017−1022 cm−3) and exhibits an indirect optimal bandgap in the range of 1.1–1.7  eV [50]. Despite such favorable properties, the efficiency of CuO-based thin-film solar cells is quite low (1.2%) and is mostly attributed to the presence of high-density recombination centers and poor charge collection efficiency [51]. Cu2O exhibits p-type conductivity with a direct optical bandgap in the range of 2–2.5 eV and is composed of earth-abundant cost-effective, nontoxic elements [52– 55]. The theoretical efficiency of Cu2O is about 20% and up to now, the highest efficiency of 8% is attained with the GeO2 buffer layer [56]. The major reason for the restricted lower performance is due to low hole concentration leading to higher electrical resistivity. Hence, there is a great scope to conduct further research and enhance the efficiency to its theoretical limit. Cu2O has been synthesized using various techniques such as electrodeposition (ED) [57], spray pyrolysis [58], sputtering [59], chemical vapor deposition (CVD) [60], and hydrothermal [61], etc. The ED of Cu2O is easy and favorable, as compact and uniform coatings are attained depending upon the applied deposition conditions. Ismail et al. [55] studied the impact of substrates on the microstructural and optical properties of electrodeposited Cu2O. Three substrates such as FTO, ITO, and Au were used, and it was found that the Au substrate produced the best results in terms of film morphology and crystallinity. The film deposited on the Au substrate exhibited uniform, compact, and smaller grain-sized morphology with superior crystallinity as compared to FTO and ITO substrates (Fig. 7.6). The attained superior crystallinity was recognized from the enhanced charge separation properties as determined from the ­photoluminescence spectroscopy study. These enhanced characteristics were attributed to the higher electrical conductivity of the Au substrate that promoted high-density cationic deposition with the least electrical resistance. The attained current density of the Cu2O film on the Au substrate (3.5 mA) determined from the dark current-­voltage analysis was significantly higher than the one attained with FTO (250 μA) and ITO (60 μA) substrate. Hence, the electrical conductivity of the substrate is important and should be considered while electrodepositing absorber films. The thickness

Fig. 7.6  The SEM image of surface morphologies of Cu2O thin film deposited on different substrates. (Reprinted with copyright permission from [55])

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control of the Cu2O is also important as it affects the solar cell performance. It has been reported that the electrodeposited Cu2O films exhibit low carrier concentration (1013-1014  cm−3) [62]. At such low concentrations, very limited charge carriers (holes) will be generated, and thus lower performance will be attained. Therefore, in the study conducted by Tran et al. [54] on the electrodeposited Cu2O films, it was suggested that the optimum thickness should be around 3–3.5 μm, as it will help restore its electrical properties, probably up to some extent. A high-­efficiency value of 0.52% was attained for the suggested absorber layer thickness. In the case of lower (3.5 μm) absorber layer thickness, lower performance (0.22 and 0.37%, respectively) was attained due to lower shunt and higher series resistance, respectively. Besides controlling the absorber layer thickness, the elemental doping approach should also be considered for improving the carrier concentration. Various elements such as Si [63], In [64], Co, Mn Fe, Ni [65], and Na [53] have been doped in the Cu2O absorber layer, and out of these, the Na doping has shown promising results. Ke et al. [53] reported that the Na doping in sputter-deposited Cu2O absorber film significantly improved the hole concentration up to 2.11  ×  1018  cm−3 from 9.31 × 1014 cm−3 for the pristine film and reduced the resistivity from 2.43 × 103 to 6.8 Ω cm. Furthermore, the morphology of the Cu2O film also improved in terms of enhanced crystallinity and compactness, thereby reducing charge recombination centers and improving minority carrier lifetime (Fig. 7.7). The highest efficiency of 1.68% was attained for the Na doped Cu2O solar cell. A similar effect of improved electrical property (reduced resistivity) was obtained by Minami et  al. [66] for Na-doped Cu2O film, where the highest power conversion efficiency of 6.1% was attained. Doping other alkali elements such as Li, K, Rb, and Cs should also be considered as they have shown promising performance in chalcogenide-based thin-­films solar cells. Especially, an increment in the carrier concentration can be expected as these alkali elements form cationic vacancies in the absorber crystal structure. Noble metal nanoparticles such as Ag, Au, and Pd can be introduced in the absorber layer to enhance the light trapping effect and enhance the performance [67]. In addition,

Fig. 7.7  The FESEM image of Cu2O thin film deposited with Na doping. (Reprinted with copyright permission from [53])

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hole transport layers (HTLs) and back surface field (BSF) layers can be introduced to enhance the charge separation capabilities and minimize recombination losses [68, 69]. Hence, the synergistic effects of noble metal nanoparticles, HTL, and BSF are expected to deliver efficient performance. Based on this, the mixed metal oxide based on Cu2O:N/CuO:N/Pd:CuO absorber layer exhibited a higher power conversion efficiency of 8.3% [70]. Another metal oxide such as CoFeO4 exhibits the desired optical properties (optimum bandgap (1.5 eV) and a high absorbance coefficient) and should be used in a thin-film solar cell application [71].

7.4.3  Metal Oxides as Buffer Layers in Thin-Film Solar Cells The buffer layer plays an important role in creating a depletion layer at the junction and avoiding the shunting paths between the TCO and the absorber layer. For efficient charge transport, it is expected that the CBM of the buffer layer should be higher than that of the absorber layer. The most commonly used metal oxide buffer layers include ZnO, Mg: ZnO, Zn(S, O), IZO, ZTO, SnO2, and In(O, S, H) [12, 17]. Cd compound-based oxides have also been used as buffer layers. CdO has a direct bandgap of 2.2–2.5 eV, exhibits high electrical conductivity, has high transparency in the visible light region, and possesses high thermal stability (melting point 1500 °C) [72–74]. Doping various elements (In, Cu, Ga, Sn, and F) has resulted in significant changes in the properties of the CdO buffer layer [29, 75–78]. A recent study based on p-Si/n-CdO heterojunction solar cells demonstrated the impact of Sn, Sb, and Se elements on the properties of the CdO buffer layer [79]. The morphology of the CdO buffer layer drastically improved as it became compact and smooth after elemental doping (0.5 wt%). This effect was translated by the reduction in roughness values from 11.3 nm (pristine CdO) to 6.93 nm (CdO:Sb), 6.47 nm (CdO:Sn), and 5.93 nm (CdO:Se) (Fig. 7.8a). The formation of a buffer layer with smooth and compact morphology is highly favorable as it forms intimate contact (absorber/TCO) for efficient charge migration and reduces recombination losses. Moreover, in a similar trend, the bandgap also reduced from 2.1 eV (CdO) to 1.7, 1.7, and 1.6  eV for Sb-, Sn-, and Se-doped CdO films, respectively, reflecting improved electrical properties. As the CdO:Se exhibited more superior properties than other dopants, it attained the highest efficiency of 3.5% as compared to CdO:Sb (2.24%), CdO:Sn (1.44%), and CdO (0.78%) (Fig.  7.8b). A similar effect of improved properties has been attained by Cu doping in the MoOx buffer layer [80]. The Cu doping resulted in strong hybridization among the Cu 3d, O 2p, Mo 4d states. This promoted the electrons to occupy the antibonding d-d* band, thereby shifting the Fermi level within the conduction band. As an outcome of this, a significant reduction in the bandgap values from 3.22 to 2.29  eV (improved electrical conductivity) and the reduction in Schottky barrier (associated with defects) was attained, eventually resulting in improved performance (14.63%) as compared to undoped MoOx buffer layer (9.1%) in CdTe solar cells. Hence, it’s clear that ele-

Fig. 7.8 (a) The AFM-3D images of undoped and (Sb-, Sn-, and Se-) doped CdO film with 0.5 wt% concentration and (b) the J − V patterns of the fabricated CdO/Si solar cell under dark and illumination for undoped and (Sb-, Sn-, and Se-) doped CdO solar cells. (Reprinted with copyright permission from [79])

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mental doping in metal oxide buffer layers has proved beneficial in enhancing the solar cell performance.

7.4.4  Metal Oxides as TCO Layers in Thin-Film Solar Cells The TCOs form the important component of thin-film solar cells. Generally, TCOs are expected to have high electrical conductivity and high transparency for delivering efficient performance. This is achieved by doping various elements such as Ge, Ga, In, Al, B, and Ce in the n-type ZnO layer [81–83]. Out of these dopant elements, Al doping has been found to impart profound effects in the ZnO layer to work ideally as a TCO layer. Al-doped ZnO TCO layers have been fabricated using both physical and chemical techniques. Physical techniques include RF sputtering, solid-­ state reaction, and E-beam evaporation, while chemical techniques include sol-gel, chemical bath deposition, nanoparticle dispersion inks, successive ionic layer adsorption, and reaction (SILAR), and spray pyrolysis. In the case of physical techniques, various preparative parameters influence the properties of the fabricated TCO layer. Li et al. [84] demonstrated the impact of varied oxygen (O2) content on the properties of Al:ZnO (AZO) TCO layers. As the O2 content was varied from 0 to 2%, the electrical properties varied accordingly. The resistivity increased from 2.2 × 10−3 to 5.9 × 10−1 Ωcm, while the mobility and carrier concentration reduced from 6.1 to 1.1 cm2/Vs and 4.8 × 1020 to 9.7 × 1017 cm−3, respectively (Fig. 7.9). The

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Fig. 7.9  The electrical properties of AZO films with different oxygen concentrations. (Reprinted with copyright permission from [84])

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O2 vacancies in the AZO compound are the source of charge carriers. Hence, the lower density of O2 vacancies (attributed to the higher oxygen content) naturally results in lower carrier concentration and increased resistivity. Furthermore, the reduced mobility was attributed to the existence of stupendous grain boundaries due to higher O2 content. Thus, it is clear that the AZO properties can be easily tuned by varying O2 content. The high conductivity of the AZO is attributed to the existence of Al3+ ions at the substitutional sites of Zn2+ ions and the existence of interstitial atoms. Hence, the Al content also significantly affects the electrical properties of the AZO layer. Higher Al content leads to lower carrier mobility, lower carrier ­concentration, and higher resistivity, and hence, optimum Al concentration should be maintained. Zhu et al. [85] suggested the optimum Al concentration to be around 2.5 wt% as the high efficiency of around 17% was obtained for CIGSe solar cells. It should be noted that vacuum-based thermal treatments (>300 °C) are required to enhance the optical and electrical properties of the TCOs. Furthermore, the thermal treatment enhances the uniformity and adhesion of TCOs on the substrates, while the vacuum treatment improves its purity. In the study conducted by Asemi et al. [86], the impact of vacuum-based thermal treatment on the electrical properties of AZO was studied. Keeping the temperature constant, the AZO films were fabricated at varied vacuum pressures (103, 10−2, and 10−5  bar). The outcome of the study revealed that as the pressure reduced from normal (103 bar) to low vacuum (10−2 bar), the carrier concentration increased (from 6.7 × 1019 to 8.4 × 1020) and the resistivity decreased (from 2.1 × 10−3 to 1.8 × 10−4). The enhancement in the electrical properties was attributed to the low vacuum pressure that knocked out oxygen atoms creating numerous oxygen vacancies. In AZO, a single oxygen vacancy donates two electrons to the conduction band of AZO, and thus high-density charge carriers are generated. With further reduction in vacuum pressure (10−5 bar), the electrical properties of the AZO deteriorated as it resulted in structural deterioration and interlayer diffusion, thereby introducing defect scattering centers in its crystal structure. Hence, the vacuum pressure control is an important parameter during the AZO synthesis and should be optimized carefully. As mentioned before, the AZO deposition is usually accomplished at high temperatures (>300 °C) to enhance its properties. If the deposition is carried at room temperature, it results in the formation of porous, low crystalline, and carbon-based impure films that significantly lowers its electrical properties. However, very recently Liu et al. [37] developed a two-step chemical-based technique for the synthesis of AZO at low temperatures (300  °C). In another study conducted by Hassani et  al. [87], the AZO film fabricated using glycol-­based colloidal nanoparticle ink also exhibited 75–90% optical transparency (Fig. 7.10c). Liu et al. [88] fabricated a solution-processed based composite TCO electrode based on Ag nanowires and AZO nanoparticles for CZTS solar cells. Initially, the Ag nanowires were spin-coated on a glass substrate, followed by air drying at 120 °C. Then, AZO nanoparticles were spin-coated and annealed at 150 °C to form the composite TCO electrode (Fig. 7.11a). The spin-coated composite TCO electrode exhibited superior morphological, optical, and electrical properties. The AZO nanoparticles effectively filled the interlayer gaps between the randomly oriented Ag nanowires for efficient electron charge transport between wires. The tightly formed composite structure produced uniform morphology (Fig. 7.11b, c), and the collective properties of AZO nanoparticles and Ag nanowires resulted in high opti-

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Fig. 7.11 (a) The schematic representation of the fabrication of Ag nanowire-based AZO nanoparticle composite film (AgNM/AZO-NPs), (b, c) the surface and cross-sectional FESEM image of the AgNM/AZO-NPs, respectively, (d, e) the schematic structure and the cross-sectional FESEM image of the solar cell based on AgNM/AZO-NPs TCO, respectively, and (f, g) the J − V curves and the corresponding EQE curves of the solar cell based on the on AgNM/AZO-NPs TCO, respectively. (Reprinted with copyright permission from [88])

cal transparency (>91%), low electrical resistivity (150 °C) which is requisite for the synthesis of other metal oxides (Fig. 8.15). Kim et al. [75] used an organic modifier layer composed of [6, 6]-phenyl C61 butyric acid methyl ester (PCBM) on ZnO layer further to boost the PSC’s performance. The device exhibited PCE of 12.2%. The Voc increased to 1.03 V for PCBM/ ZnO device from 0.83 V for just ZnO due to the modified electronic structure of the ZnO layer V as seen in Fig. 8.16. Wu et al. deposited Al-ZnO using sputter technique to fabricate PSC. According to the report, the Voc of the device increased by up to 1.07. One of the major hindrances in the development of ZnO ETL-based PSC is its chemical stability. ZnO is a basic oxide with an isoelectric point greater than 8.7. On account of its basic nature, ZnO causes rapid degradation in the perovskite absorber layer. The presence of the hydroxide group on the ZnO surface accelerates the degradation mechanism in the perovskite absorber layer. The organic halide layer quickly breaks down, forming zinc iodide and water.

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Fig. 8.15 (a) The device structure of the planar heterojunction perovskite solar cells and the molecular structures of the interfacial layers PCBM. A cross-sectional SEM image of the device is also shown. Energy-level diagram of the components of the perovskite solar cell. (Reprinted with copyright permission from [38])

Fig. 8.16  Current density–voltage characteristics of the planar heterojunction perovskite solar cells. (Reprinted with copyright permission from [38])

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Kelly et al. [36] systematically showed the deprotonation of methylammonium cation in the presence of hydroxide groups existing at the surface of ZnO nanoparticles. Consequently, it draws additional moisture in the cell, thereby causing a complete breakdown of the perovskite crystal structure. Even though ZnO has severe issues with stability, PCE over 20% has been recorded for ZnO-based PSCs [76, 77]. The key moving forward will be the use of different methods to suppress the degrading ability of ZnO. Doping, nanocomposites and use of protective layers can help in realizing ZnO as a potential alternative to TiO2.

8.3.6  Nb2O5 Nb2O5 has exciting properties for film fabrication, such as high refractive index, low optical absorption in the visible and near-infrared region, corrosion resistance and thermal and chemical stability. Some of its applications are sensors, electronic devices and optical interference filters. Nb2O5 is a polymorphic compound with more than 15 structural configurations. The different phases are temperature dependent. The optical properties also show deviation from average values based on the phase and temperature treatment it is subjected to. The bandgap of Nb2O5 lies in a range from 3.2 to 5 eV [78]. The conduction band position is a little lower (0.4 eV) than TiO2. Niobium oxide use in photovoltaics can be traced back to its application in DSSC. Although it enjoyed moderate success in DSSC but in PSC, it has mixed results. Wang et al. [79] deposited 30 nm of low temperature–processed Nb2O5 to achieve performance equivalent to TiO2. The Voc of Nb2O5 PSC surpassed the TiO2 PSC in terms of Voc. With PCE of 20.22%, this study demonstrates the applicability of Nb2O5 in high-performance PSC.  The study also revealed that Nb2O5-based PSC showed remarkable resilience to UV exposure and high Voc of 1.19 V. Figure 8.17 shows the degrading protection Nb2O5 offers over TiO2. Ling et al. [80] studied the effect of temperature on Nb2O5. The room temperature–processed amorphous a-Nb2O5 (PCE 17.1%) shows similar performance to that of high temperature–processed crystalline c-Nb2O5 (PCE 17.2%). From this study, it is evident that Nb2O5 can produce good results even when used in an amorphous state. The primary reason for this performance was the suitable band alignment with the perovskite absorber material, as seen in Fig. 8.18. Figure 8.19 shows the flexible PSC using amorphous Nb2O5 as ETL with 12.1% PCE. Yang et al. [81] demonstrated flexible PSC using e-beam evaporated Nb2O5 on the ITO-PET flexible substrate. They used dimethyl sulphide as an additive to the perovskite layer to improve the performance of flexible PSC. The stability of the device increased 1.72 times. PCE of 18.4% and Voc of 1.10 V for a flexible substrate cell are quite encouraging for the use of metal oxide in PSC. Figure 8.20 shows that Nb2O5 can be used in flexible PSCs and can withstand a large number of bending cycles indicating the excellent quality of the film.

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Fig. 8.17 (a) Normalized Jsc of un-encapsulated PSCs. Degradation diagrams of (b) Nb2O5 and (c) TiO2-based PSCs upon UV exposure. Corresponding optical photographs of the perovskite thin films on different substrates before and after UV irradiation. (Reprinted with copyright permission from [79])

8.3.7  HTL The function of HTL is fast extraction and transportation of hole carriers. The type of HTL, which can be used in a PSC, is dependent on the architecture of the cell. Synthesis temperature and processing techniques play an essential part. There are two types of HTL materials, organic and inorganic, which are used in PSC fabrication. The essential criteria include high hole mobility, easy of processing, low-­ ­ temperature synthesis, stability and valence band compatibility. The p-i-n architecture mainly uses the inorganic metal oxides as HTL since high temperature is tolerable during the first stage of fabrication. HTL layer is coated on the FTO substrate, followed by the perovskite layer. Perovskite layer is highly sensitive to moisture and high temperature. Thus, the use of metal oxide HTL on the perovskite layer is restricted due to these reasons. For n-i-p architecture, organic materials such as Spiro-OMeTAD or PTAA are used as HTL. Methods of synthesis of these materials have been developed using solvent engineering to make the process compatible with perovskite absorber material. However, with metal oxides, such process is quite complicated and pose a threat to the perovskite layer through degradation by corrosive precursors. Quite a few metal oxides are used as HTL in PSC. CoOx, CrOx, NiOx, MoOx and CuO have been reported as HTL in PSC.  Most of the devices reported with metal oxide HTL are based on p-i-n architecture. Very few reports involve n-i-p architectures. The performance of such devices is quite significant. With PCE rang-

8  Metal Oxides for Perovskite Solar Cells Fig. 8.18 Energy-level diagram of the PSCs with crystalline and amorphous Nb2O5. (Reprinted with copyright permission from [80])

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c-Nb2O5 –4.02 eV

–3.9 eV

–4.55 eV –4.31 eV FTO a-Nb2O5 –4.8 eV ITO

MAPbI3

–5.4 eV

SpiroOMeTAD –5.2 eV

–5.1 eV Au

Fig. 8.19 (a) J − V curves of the champion devices based on a-Nb2O5 and c-Nb2O5 under different scanning directions, (b) IPCE spectra of the champion devices, (c) the PCE distribution histogram of the PSCs using a-Nb2O5 and c-Nb2O5 as the ETLs and (d) photograph and J − V curve of flexible device. (Reprinted with copyright permission from [80])

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Fig. 8.20 (a) Illustration of the flexible (F-PSC) structure. (b) Cross-sectional SEM image of a completed flexible device with the MAPbI3–DS absorber. (c) J – V curves of F-PSCs under both reverse and forward scan directions using MAPbI3 and MAPbI3–DS as the absorber layer. (d) Stable current density and PCE of the F-PSCs based on MAPbI3 and MAPbI3–DS measured when the device was biased at 0.84 and 0.88 V, respectively. (e) The PCEs of flexible device based on MAPbI3–DS at different bending curvature radii after 5000 flexing cycles. (Reprinted with copyright permission from [81])

ing between 10 and 20%, such devices might open up new revolution in PSC. Stability has been a prominent issue with perovskite solar cells. Use of metal oxides as transport layers improves the lifetime of the device. Although the hole mobility in metal oxide is indisputably lower than organic material, lower cost of production, excellent stability and comparable performance make up for the shortcomings.

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8.3.8  NiOx NiO has a large bandgap of 3.5–4 eV with valence band edge energy at −5.3 eV. It can effectively transport holes and block electrons. NiO rose to prominence after a report by Irwin et al. in 2008 [82]. A thin coating of Ni (p-type semiconductor) was used as a substitute for PEDOT PSS in bulk heterojunction solar cell. Despite its low conductivity, NiO devices showed an increase in performance. Wang et al. [83] were the first to report the use of NiOx in PSC. They achieved a respectable 9.51% PCE with Voc of 1.04 V and Jsc of 13.24 mA cm−2. Yang et al. [84] demonstrated a p-i-n device comprising NiOx as HTL and ZnO as ETL. NiOx was synthesized from nitride salts. The ZnO nanoparticles were spin-coated on top of the perovskite absorber layer. They reported a maximum PCE value of 16.1% and 90% stability after 60 days of air storage. Nejand et al. [85] achieved an efficiency of 7.23% for n-i-p architecture device in which NiOx was deposited on the perovskite layer using sputter. The FTO/TiO2/PVK/NiO configuration helped in maintaining the stability of the device. NiO is synthesized using techniques like sputtering, ALD, sol–gel, electrodeposition and hydrothermal. Xie et al. [86] reported PCE over 20% for p-i-n PSC with NiOx as HTL and PCBM/TiO2 as ETL. The device exhibited Voc of 1.1 V and Jsc 23.09 mA cm−2. NiO layer was deposited by spin-coating NiO precursor solution (nickel acetylacetonate in acetonitrile and ethanol) and then heating at 500 °C. High temperature promoted crystallization in NiO. Comparative performance of NiO and Spiro-OMeTAD can be seen in Fig. 8.21.

8.3.9  CuOx Copper oxides are p-type semiconductor with a long history in photovoltaic application. Due to low cost, environmental friendliness, low toxicity and ease of synthesis, CuOx can be a potential alternative for organic HTL materials. Few reports on CuOx in PSC have been published. Based on the published reports, CuOx PSC devices have exhibited high Voc (0.9–1.1 V). Huang et al. [87] demonstrated that CuOx PSC can achieve high PCE of 17.1%. CuOx film was deposited by spin-coating a solution of copper acetylacetonate in dichlorobenzene. The coated substrate was then heated at 80 °C for 20 min. The p-i-n device was fabricated with C60 and BCP as the ETL layers. Yu et al. [88] also achieved similar PCE (17.43%) for solution-processed CuOx-based PSC. Nejand et al. [89] successfully deposited Cu2O using magnetron sputtering to fabricated n-i-p PSC.  The device achieved the highest PCE of 8.93%. Other metal oxides such as MoO3, CoOx and CrO are now garnering the attention of researchers. Due to their band position, Voc values have consistently crossed 1  V.  Although the significant hurdle is fill factor, the low hole mobility of metal oxide HTL is one of the factors.

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Fig. 8.21 (a and b) FE-SEM images and (c) J − V curves of FTO/TiO2/CH3NH3PbI3−xClx/NiOx/ Ni and representative FTO/TiO2/CH3NH3PbI3−xClx/Spiro-MeOTAD/Ag under light and dark with and without the perovskite layer. (Reprinted with copyright permission from [86])

8.3.10  Ternary Oxides Ternary oxides appear to be an attractive prospect when it comes to PSC application. Efficiently altering the properties of material depending upon the application by merely varying the composition seems to be a very convenient option. Due to the popularity of TiO2 in PSC, ternary Ti oxides such as Zn2Ti3O8, BaTiO3, SrTiO3 have been tried in PSC application. However, their performance was not as expected. The band mismatching and low charge-carrier mobility prompted researchers to abandon them in favour of zinc-tin oxides (ZSO) and barium-tin oxides (BSO).

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The problem, however, is the high synthesis temperature. It is difficult to achieve the required composition and purity of the crystal structure without using high temperature. The presence of two metal cations in metal oxide can produce intriguing properties (electrical and optical). However, ensuring the desired outcome at low temperature is a difficult feat to achieve. Nonetheless, efforts are being directed to lower the synthesis temperature of ternary oxides. ZSO has attracted attention due to its electrical and optical properties. The electron mobility is 30 cm2 V−1 s−1 while bandgap is 3.8 eV. The conduction band edge position is in an acceptable range. ZSO has excellent chemical stability, which makes them resistant to acids, solvents and other corrosive materials. The first report on ZSO-based PSC was published in 2014. In 2015, Shin et al. [90] demonstrated a low-temperature (below 100 °C) synthesis of ZSO particles. Flexible PSC using the ZSO ETL layer yielded PCE over 15%. The report explains the role of hydrazine and pH of the precursors in the formation of ZSO. Sadegh et  al. [91] achieved PCE over 21% for ZSO-based devices. ZSO was prepared using CBD method, which resulted in a compact and dense layer of ZSO. Uniform ZSO also leads to the formation of a perovskite film with better surface coverage, enlarged grains and reduced recombination losses. The cells also showed a notable increase in the Voc. BSO is another ternary metal oxide with high electron mobility (320 cm2 V−1 s−1) and acceptable optical properties (3.2 eV bandgap). Sun et al. [92] prepared BaSnO3 by the peroxide-precipitate route. BSO nanoparticles with varying size were prepared and deposited on the FTO substrate. At the optimum size, the FTO surface roughness reduced due to complete coverage by BSO particles. The particles accommodated the cracks and crevices of the FTO layer to smoothen it out. It has been reported that BSO is a cubic perovskite-type oxide and behaves as an n-type semiconductor. Excellent tetragonal perovskite layer formed because of the smooth underlying layer with cubic crystal structure. Depending upon the size of BSO, the Voc varied from 0.59 V to 1.02 V. The highest efficiency of 10.96% was observed with 0.98 V Voc. The relation between particle size and cell parameters is shown in Table 8.2.

Table 8.2  BSO thickness and its corresponding cell’s performance BSO films BSO-­ 30 nm BSO-­ 63 nm BSO-­ 79 nm BSO-­ 123 nm

Jsc (mA/ cm2) 3.99 ± 0.86 (4.38) 15.27 ± 1.56 (17.51) 15.68 ± 1.90 (17.45) 12.36 ± 0.02 (12.05)

Voc (V) 0.475 ± 0.176 (0.596) 0.793 ± 0.062 (0.746) 0.984 ± 0.005 (0.986) 0.991 ± 0.055 (1.020)

FF 0.248 ± 0.047 (0.364) 0.533 ± 0.063 (0.675) 0.628 ± 0.038 (0.637) 0.637 ± 0.062 (0.709)

PCE (%) 0.44 ± 0.21 (0.80) 6.45 ± 1.15 (8.82) 9.65 ± 1.06 (10.96) 7.77 ± 0.61 (8.72)

Rsh (Ω cm2) 73.5 ± 27.4

Rs (Ω cm2) 101.7 ± 51.1

533.7 ± 89.4

14.1 ± 1.59

618.9 ± 165.0 13.3 ± 2.1 360.0 ± 117.5 19.5 ± 2.3

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La doping in BSO (LBSO) demonstrated by Myung et al. [93] exhibited high mobility, photostability and structural stability. The LBSO has a unique crystal structure. The cubic perovskite structure without octahedral tilting gave rise to some interesting properties. UV photocatalytic property of the material was suppressed due to small dipole moment arising from its structure. This is especially important, as UV degradation of perovskite by ETL metal oxides is a serious issue.

8.3.11  Issues with Metal Oxides There are some issues involved in using metal oxide as ETL.  The main issue is perovskite degradation due to its photocatalytic activity of wide bandgap metal oxides. TiO2 and WO3 absorb UV radiation and oxidized water molecules forming hydroxyl group on the surface. This problem can also stem from the annealing of the oxide layer during synthesis. N2-annealed samples are known to suffer from these issues. The oxygen vacancies created can amplify the photocatalytic ability of metal oxides. Certain metal oxides are acidic or basic. TiO2, WO3, SnO2, Nb2O5, etc., are acidic while ZnO, NiO, CuO, etc., are basic in nature. Both types of metal oxide pose a danger to the perovskite layer by chemical degradation. This issue was prominent in ZnO PSC. Another problem is the migration of defects and oxygen vacancies under the electrical field towards the perovskite interface. The accumulation of these defects and vacancies at the transport layer–perovskite interface causes abnormal h­ ysteresis. The charge transfer between the interfaces is hindered. This can further result in recombination loss. The charge conductivity of metal oxides is two orders lower than the conducting polymers and is classified as inferior, excluding some exceptions. The charge conductivity is an important attribute for a transport material. Increase in conductivity will most certainly ensure better performance of PSC. To address these issues, researchers have come up with different strategies. Use of these strategies can solve one or many of the problems simultaneously. Doping, UV filters, buffer layers and surface passivation can help mitigate the issues with metal oxides. Doping metal oxide with another metal is the best way to fine-tune the electrical and optical properties of metal oxides. Doping also changes the band positions in the metal oxides. Depending upon the desired band position, a dopant can be introduced to alter the energy levels in the required direction. The nature of dopant is also crucial as it also alters the charge conductivity to certain extent and helps in the reduction of defects or deep traps in the metal oxide. Reports of ETL doping resulting in improved PCE and increase in Voc and stability are common. The suitable dopant can be found using computational methods based on density functional theory (DFT). DFT calculations can provide information related to the electronic struc-

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ture as well as material properties such as the density of states, carrier density, bandgap, band position, Fermi level, to name a few. UV filters can help in blocking the radiation reaching the PSC. Thus, an additional layer of filter on the PSC can prove to be a facile method to overcome the problem. This approach can be the PSC bulky and may not always be viable. Another strategy is down conversion by use of quantum dots layer. The quantum dot layer converts high-energy UV light to low-energy visible light. This approach is now becoming more popular since this can also be used to boost the conversion efficiency. Tavakoli et al. [94] used CdSe/CdS quantum dots to achieve this effect. Akin et al. [95] employed CsPbBr1.85I1.15 quantum dot interfacial layer to fabricate stable PSC with PCE over 21%. Use buffer layers to isolate the metal oxide layer and perovskite layer can prove an effective strategy in mitigation UV as well as chemical degradation. Buffer layers of insulation polymers or metal oxides such as Al2O3 and ZrO2 have helped in isolation of perovskite layer [96–99]. Increase in PCE and stability was observed in PSC employing buffer-insulating layers. Tan et al. [100] employed chlorine-capped TiO2 as ETL for planar PSC.  The surface passivation by chlorination helped in reducing traps and increasing the performance (PCE >20%) and stability of the cell.

8.4  Conclusions PSCs have displayed their potential to dominate the energy-harvesting sector. Apart from the stability issue, PSC has all the properties (high PCE, low cost, low-­ temperature synthesis, easy fabrication and scalability) to take over from Si technology and liberate the world from the clutches of polluting energy sources. Metal oxides are inherently stable and resistant to physical and chemical elements. One of the most significant advantages of metal oxides is the low cost and ease of synthesis. They can be synthesized in many shapes and sizes by different (scalable) methods. Currently, metal oxides are the preferred choice as ETL in PSC. High PCE has been achieved with metal oxide as ETL. TiO2 once held the position of an undisputed king as ETL in PSC. There are many emerging candidates. Exhibition of high-performance PSCs with use of metal oxides such SnO2, ZnO, Nb2O5, WO3 has shown that there is still scope for improvement. The research until now has highlighted some of the drawbacks of using metal oxides; these issues range from stability to mobility. With proper strategies, these problems can be overcome. HTL is another area where the use of metal oxide can significantly improve the stability of PSC. At present, polymers are considered a suitable option due to their carrier mobility and facile synthesis method. NiO, CuO, MoO3 and CoOx are good hole transport materials. PSCs with metal oxides as both transport layers have been fabricated. With their excellent stability performance and high PCE, metal oxides transport layers seem to be a good option moving forward.

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Acknowledgement  This work was supported by the Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education under Grant 2018R1D1A1B07048610.

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Chapter 9

Doped Metal Oxide Thin Films for Dye-­ Sensitized Solar Cell and Other Non-Dye-­ Loaded Photoelectrochemical (PEC) Solar Cell Applications M. D. Tyona

9.1  Introduction Metal oxides are fascinating group of materials that are widely studied nowadays due to their unique characteristics such as mechanical stress tolerance, high optical transparency, exceptional charge carrier mobilities [1]. Different kinds of metal oxides exist, a few include cuprous oxide (Cu2O), zinc oxide (ZnO), titanium oxide (TiO2), cadmium oxide (CdO), tin oxide (SnO2), and tungsten oxide (WO3). Prominent among these is ZnO due to its abundance in nature, nontoxicity, high environmental stability, and wide range applicability such as in photonic crystals, varistors, optical modulator waveguide, photodetectors, surface acoustic wave filters, [1–5], photodiodes, gas sensors, photovoltaic (PV) solar cells, and light emitting diode [6], among others and in addition its charge carrier mobility is high. ZnO can be identified as a group II–VI compound semiconductor with a wide and direct bandgap (about 3.20–3.70 eV, amorphous and crystalline) in the near-UV region of the electromagnetic spectrum [7]. It has a large free-exciton binding energy [1, 7– 10] which enables excitonic emission processes to persist at or even above room temperature [6, 11–15]. ZnO is mostly known as an n-type semiconductor due to the presence of native defects such as oxygen vacancy, zinc interstitials, and hydrogen interstitials in the ZnO lattice [14]; synthesis of the p-type is not generally easy [1, 16–20]. In its crystalline form, two major states can be identified namely: hexagonal wurtzite and cubic zinc blende [1]. The wurtzite structure as shown in Fig. 9.1 is the most stable state of ZnO, having lattice constants a = 0.3296 and c = 0.52065 nm, and hence it is most often used [1, 16]. The crystalline ZnO exhibit distinctive characteristics such as high surface-to-volume ratio, high optical transparency in the visible range, and quantum confinement effect [1], among others, which guarantee M. D. Tyona (*) Department of Physics, Benue State University, Makurdi, Benue State, Nigeria e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_9

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Fig. 9.1  A schematic diagram of the ZnO wurtzite crystal structure with lattice constants a and c [16]

the wide range of novel applications earlier mentioned. The structure of ZnO can best be described as having a number of alternating planes consisting of tetrahedrally coordinated O2− and Zn2+ ions, which are stacked alternately along the z-axis and have no central symmetry as shown in Fig. 9.1 [1, 16]. The crystal lattice of ZnO is in a close match with some other compound semiconductors such as GaN and as such, it can be used as a substrate for their growth process [16]. ZnO can also replace some semiconductors such as GaN in several applications like in optoelectronics when some characteristics of the material are enhanced [21, 22]. Another prime advantage of ZnO is that its epitaxial films are easy to grow. Both physical and chemical methods can be used to grow very good thin/thick films of ZnO [23, 24]. Most of the frequently explored methods to synthesize epitaxial films of zinc oxide are sol–gel process, electrodeposition, spray pyrolysis [25, 26], RF sputtering, successive ionic layer adsorption and reaction (SILAR), spin coating, chemical bath deposition (CBD), electron beam epitaxy, laser evaporation, and ion beam sputtering, among several others [27–29]. These methods are often used to synthesize alloys and compounds of zinc oxide [2, 30]. Choosing a particular method of synthesis would be guided by factors such as the advantages of the method, the kind of application intended with the synthesized material, and/or cost implication [3–5, 31]. Crystalline zinc oxide thin films are known among semiconductors to have the largest number of novel morphological structures, which include nanocages, nanohelixes/nanosprings, nanocombs, nanorods, nanorings, nanowires, nanobelts, nanotubes [26, 32, 33]. The efficient performance of thin film-based devices is leveraged strongly on the structural, electrical conductivity, and optical properties of the component films. An investigation of these characteristics and their reliance on the film properties is very crucial as it enables the optimization of film parameters for enhancing device applications. For photoelectrochemical (PEC) solar cell application of zinc oxide nanostructures, large internal surface area, highly porous surface, and high surface

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roughness are key to excellent penetration of electrolyte which is very useful for enhancing PEC performance of thin film electrodes [34, 35]. Chemical methods are facile, inexpensive, and very dependable for growing good-quality electrodes for PEC application [1, 23]. Herein, the chemical bath deposition (CBD) method is very appropriate for synthesizing large area films of zinc oxide with attractive characteristics for PEC solar cells [13, 17]. CBD method is well known for producing zinc oxide nanostructures on different substrates including stainless steel and microscope glass [20]. Regardless of the numerous advantages of ZnO, the electrical conductivity in ZnO is generally low as well as ZnO has very limited absorption spectra due to the wide bandgap, and consequently, its usage in optoelectronic devices is limited [1, 16]. Bandgap engineering is an effective method to control the electrical conductivity in zinc oxide and to improve its spectra absorption [23]. Introducing a little content of impurities (dopants) and native point defects (down to about 10–14 cm−3 or 0.01 ppm) can remarkably alter the electrical, morphological, optical, and structural properties of ZnO [1, 16]. Therefore, a knowledge of the effect of native point defects (i.e., vacancies, antisites, and interstitials) and doping with appropriate dopant is key to controlling the electrical conductivity in zinc oxide, which will effectively tailor the bandgap and enhance its performance [36, 37]. Growing alloys of ZnO with other metal oxides such as MgO or CdO is another surest means of achieving bandgap manipulation in zinc oxide [23, 37]. It is known that the addition of Mg into zinc oxide increases its bandgap, whereas significant bandgap shrinkage is realized with the addition of Cd [24, 28]. Although CdO and MgO crystallize in the rock salt structure [37], alloys of Cd1−xZnxO and Mg1−xZnxO with appropriate contents will retain the wurtzite structure of ZnO with remarkable variation in bandgap [1, 16, 37, 38]. The replacement of Zn2+ ions with group III ions (B3+, In3+, Ga3+, and Al3+) [36, 37] through doping gives rise to extra electrons and enhances electrical conductivity, optical, magnetic, and thermal properties of ZnO. In3+ and Al3+ have been among the frequently used dopants as a result of their small ionic radii and relatively low material cost. The replacement of Zn2+ ions with In3+ or Al3+ in ZnO lattice increases charge carriers concentration [36–38]. It is reported that the electron concentration can increase from 1016 to 1021/cm−3, thereby enhancing the electrical conductivity of the material [21, 22]. The charge carriers mobility is significantly influenced by scattering at the disorder locations created in the crystal structure as a result of doping. Hence, properly doped crystalline ZnO nanostructures, with pure phase, are very beneficial for achieving good electrical conductivity as well as structural and optical properties [39]. Another efficient method for enhancing spectra absorption in ZnO especially for PEC solar cell application is to adsorb an effective sensitizer dye onto the photoelectrode. Dyes are narrow bandgap materials with the capacity to absorb photons effectively in the visible and infrared portions of the solar radiation. A photosensitizer (“dye”) is a basic component of the DSSC. Its major function is to absorb the solar radiation and inject the excited electrons into the conduction band of the oxide substrate. The fundamental properties of a photosensitizer for efficient performance are

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the following: strong light absorption in the visible and near-IR region (for efficient light harvesting); high solubility in organic solvents (to improve the deposition from stock solutions); the availability of suitable anchoring ligands such as carboxylic or phosphonic acid groups (to promote the effective interaction with the oxide surface and thus the coupling of donor and acceptor levels); the lowest unoccupied molecular orbital (LUMO) of the dye must be sufficiently high in energy for efficient charge injection into the nanostructured photoelectrode, and the highest occupied molecular orbital (HOMO) must be sufficiently low in energy for efficient regeneration of the oxidized dye by the redox pair; and high thermal and chemical stability (to assure a high turnover number and a corresponding lifetime of the device). Moreover, electron transfer from the dye to the coated oxide layer (ZnO or TiO2) must also be rapid in comparison with decay to the ground state of the dye [40–44]. Therefore, this chapter is focused at discussing illustratively two techniques (i.e., doping and dye-sensitization of photoelectrodes) to improve key properties of ZnO (electrical conductivity, structural, and optical) that would enhance its PV performance. Consequently, a detailed review of the effects of doping Al and In in ZnO through a chemical method will be intensively carried out as well as the resulting impacts on PEC solar cells. Furthermore, the effect of a sensitizer dye (rhodamine 6G) on the PEC performance of AZO thin film electrodes will be clearly considered.

9.2  U  sing Doping as an Effective Method to Engineer Key Properties of ZnO for Enhanced Energy Harvesting Doping is the intentional introduction of impurity atoms (dopants) into the lattice structure of an intrinsic semiconductor with a focus on enhancing its electrical, optical, and structural as well as other key properties [1]. Impurity atoms can be referred to as donors or acceptors depending on the number of valence electrons they possess [6, 21]. Typical example of donor impurities includes arsenic, phosphorus, and antimony while indium, boron, and aluminum are acceptors. For typical semiconductors like germanium (Ge) and silicon (Si) (elemental), the commonly used dopants include indium, boron, and aluminum, (acceptors) and arsenic, phosphorus, and antimony (donors) [1]. Incorporating a little content of dopant in the lattice structure of a semiconductor brings about a remarkable transformation in their electrical characteristics [45], since the dopants generate free radicals (electrons or holes) in the material [10]. A doped semiconductor is referred to as an extrinsic semiconductor. The kind of extrinsic semiconductor that would be achieved in the process of doping is dependent on the number of valence electrons of the dopant [1, 45]. A Si semiconductor which is doped with an acceptor will produce a p-type, whereas doping with a donor will lead to an n-type [1] as shown in Fig. 9.2. In n-type semiconductors, the current is conducted by the electrons (majority charge carriers), while in p-type, holes are the current conductors [1]. In a properly doped Si crystal, ­electrical conductivity can be increased by a factor of 106 [1]. In an extrinsic semiconductor, the impurity atoms in the crystal lattice are the main providers of electric current conductors throughout the crystal [1, 16]. Extrinsic semiconductors have

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Fig. 9.2  A scheme of doped semiconductors: (a) n-type doping using antimony and (b) p-type doping using boron [16]

Fig. 9.3  A scheme of doped semiconductors indicating energy levels of: (a) n-type and (b) p-type dopings

proven to be electrically neutral. The terminologies of n- and p-type doping strictly refer to the major current conductors (majority charge carriers) [13]. Any positive or negative current conductors are associated with a fixed acceptor or donor (dopant) as shown in Fig. 9.3. Compound semiconductors can as well be doped; aluminum and indium or similar dopants can be used to dope TiO2 or ZnO.  Previous researches have demonstrated that structural, optical, and morphological characteristics of zinc oxide can be significantly modified by doping with Al, In, or Cu [13, 23, 39]. It has been established in several literatures that Cu, Al, and In dopants have tremendously lowered or widen the bandgap of zinc oxide [13, 23, 39].

9.3  Impacts of Al Impurities on Zinc Oxide Properties The wide bandgap in ZnO (3.3  eV) semiconductor is an advantage for materials applications like photoelectrochemical solar cells due to its environmental stability [13, 39]. Nevertheless, its light absorption capability is limited by the wide band-

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gap, seeing that only photons below a particular wavelength λg can be absorbed because the intensity of the solar spectrum is normally much intensive around 2.7 eV [1]. Hence, ZnO semiconductors can only absorb photons effectively within the ultraviolet spectrum [13, 39]. This drawback in the form of very limited light absorption in ZnO can be overcome through doping. Previous reports have shown that not only bandgap widening could be remedied but also other optical, as well as morphological and structural properties of zinc oxide, could be tailored through the use of suitable dopants [13, 23, 39], to meet pre-desired applications [39]. Al is prevalently studied as a suitable transition metal which could affect remarkable and attractive changes in the crystal lattice of ZnO which are beneficial in many applications, especially PV solar cells [21]. Al is a highly conducting metal and is known for green luminescence band enhancement by creating localized states within the bandgap of zinc oxide. It can also substitute Zn rapidly on zinc oxide crystal lattice because of its high ionization energy and less formation energy [33, 39]. Tyona et al. [10] studied the impact of Al dopant on optoelectronic characteristics of chemically grown zinc oxide thin films. The measured characteristics of the Al-doped zinc oxide (AZO) thin films were affected by different parameters which include, the content of Al in the chemical bath, growth conditions, and post-­ deposition annealing. Al concentrations of 0 at.%, 1 at.%, 2 at.%, and 5 at.% were considered for doping ZnO [10]. This quantity, although small (within a strict doping range of up to 10  at.%), could generate remarkable physical impacts in zinc oxide. It is known that concentrations of dopant above this range may be considered as alloys or composite growth [1, 10]. The AZO thin film electrodes were grown by a chemical route (CBD). Zn(NO3)2.6H2O (SD Fine Chemicals) and Al(NO3)3.9H2O (Loba Chemie, India) were used as sources of Zn2+ and A3+, respectively; the bath was complexed with NH3 solution (28%) (Thomas Baker). Detailed description of the growth process and other techniques for materials characterization have been given elsewhere.

9.3.1  Structural Studies Structural properties of the AZO electrodes in the work of Tyona et al. [10] were studied with a Philips PW1830 X-ray diffractometer, considered over a range of 20° to 80°. In this study, as shown in Fig. 9.4a, the XRD analysis of un-doped zinc oxide and AZO film electrodes showed that the peaks of diffraction exactly correspond with ZnO (JCPDS File No.036–1451). Those peaks correspond to lattice numbers (1 0 0), (0 0 2), and (1 0 1). The films showed hexagonal wurtzite structure and were polycrystalline in nature for all the samples. All the film samples were seen with high intensity peaks in the z-direction, which is the known densest plane in wurtzite ZnO; however, for films with Al:Zn of 1 at.% and 5 at.%, the peaks of diffraction decreased [10].

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Fig. 9.4 (a) XRD analysis of Al-doped zinc oxide thin films, (b) Spectra of absorption of Al-doped zinc oxide thin films. (Inset: Chemical structure of rhodamine 6 G dye), (c) Plot of Tauc’s relationship between the absorption coefficient, α, and the photon energy, hν, for AZO thin films with various Al contents. (Inset: The trend of bandgap with the variation of Al content), and (d) Spectrum of absorption for 1 mM rhodamine 6G dye in acetonitrile [10]

The shrinkage of the z-plane with the addition of Al into ZnO (i.e., Al:Zn of 1 at.%), as seen in Fig. 9.4a(ii), could be due to some influenced crystallographic defects that cause a diminution of the z-axis and thus lower the crystalline quality of the film [10, 46, 47]. Nevertheless, by raising Al content to 2 at.%, the mobility of charge carriers and concentration in the conduction band of AZO were increased [10, 46, 47]; hence, the crystallographic defects in the film were lessened. This result proposes an improvement in the crystalline quality of the AZO z-plane for 2 at.% Al as seen in Fig. 9.4a(iii) [10]. Increasing Al content to 5 at.% would lower the crystalline quality because of increased compressive strain in the films [17]. It could be reasoned that because of the high quantity of impurities being separated at the grain boundaries, contraction in grains size takes place [10, 36–38, 46, 47]. It, therefore, means that Al:Zn of 2 at.% would be the optimum content in zinc oxide that would yield the best Al-doped zinc oxide films. Additionally, there was a notable shift in diffraction peak of 0.6° in 2θ in the direction of higher angles for all Al-doped zinc oxide films as seen in Fig. 9.4a(ii–iv). This is an indication of effectual substitution of Zn by Al on ZnO crystal lattice [10, 46, 47]. By introducing Al impurities in ZnO, Al atoms substitute Zn atoms [10, 47].

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Thus, the substitution of zinc atom by aluminum atom yields one free electron since aluminum atom has one more valence electron than zinc [10, 47]. Consequently, substituting zinc ions with aluminum ions gives rise to n-type doping and hence, enhances the concentration of charge carriers, causing lowering of resistivity which indicates an enhancement in electrical conductivity [10, 46, 47]. There was a slight notable increase in the estimated crystallite sizes from 21 to 29 nm.

9.3.2  Optical Studies Figure 9.4b illustrates the spectra of optical absorption for Al-doped zinc oxide electrodes in the ongoing study of Tyona et  al. It is clearly seen that ultraviolet absorption is dominant in the absorbance spectrum of AZO with Al:Zn of 1 at.%. For Al-doped zinc oxide films with 2 at.% and 5 at.% concentrations, a red-shift in absorption band edges was noted up to 430 nm and 450 nm [10], respectively; these are attributed to increase in optical absorbance prompted by aluminum impurities, arising due to d–d transition [14, 23, 48]. The enhancement in absorption reported here is comparatively high. In their result, bandgap increased marginally from 3.09  eV to 3.12  eV (with respect to 0 at.%), showing high sensitivity to aluminum impurity concentration as shown in the inset of Fig. 9.4c. An uncommon and abrupt lessening and broadening of bandgap were seen in the course of doping which may consequently result from Burstein–Moss effect [46, 48]. Increase in bandgap with Al impurities also implies successful doping of aluminum into zinc oxide [10, 46]; notwithstanding, stress relaxation mechanism is a vital cause of the abrupt decrease in bandgap at lower impurity concentration [46]. Therefore, the bandgap of Al-doped zinc oxide has its maximum value at Al:Zn of 2 at.% which suggests that this would be the optimum Al impurity concentration for Al-doped zinc oxide. Dye molecules with the proper molecular structure are used to sensitize wide bandgap nanostructured photoelectrode. When a photon is absorbed by a dye molecule adsorbed to the surface of the nanostructured photoelectrode (ZnO or TiO2), the dye molecule is oxidized and the excited electron is injected into the nanostructured material for transportation through the external circuit to the counter electrode. In the study under consideration, rhodamine 6 G dye (1 mM in acetonitrile) was used as the sensitizer on AZO electrodes and its absorption spectrum as measured with a Shimadzu UV-1800 spectrophotometer is depicted in Fig. 9.4d, showing a flat band peak from 538 to 589 nm [10, 24].

9.3.3  Morphological Studies Typical SEM micrographs of the Al-doped zinc oxide electrodes in the study were acquired using a scanning electron microscope (SEM), JEOL JSM-6360, which are shown in Fig. 9.5. The micrographs depict high sensitivity of the morphology to

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Fig. 9.5  Scanning electron microscope images of Al-doped zinc oxide electrodes with various aluminum contents. (a) Pure ZnO (b) 1 at.%, (c) 2 at.% (d) 5 at.% [10]

aluminum concentration. The porosity and crystallite sizes depreciate with the initiation of aluminum impurities in the zinc oxide crystal structure, which is adduced to the contrast in the ionic radii of zinc and aluminum [10, 47]. At Al:Zn of 1 at.%, nano-dendritic shapes were seen (Fig. 9.5a), which appear more defined; nevertheless, a remarkable contraction in the average diameter of the dendritic rods arose (Fig.  9.5b). At Al:Zn of 2  at.% (Fig.  9.5c), the nano-dendrite structures became densest, more-defined with closely even rod size (diameter ~29 nm). At Al:Zn of 5 at.% (Fig. 9.5d), the morphology was modified to dense and randomly aligned nanorods of mean diameter ~17 nm [6, 15, 21].

9.4  T  he Impact of Al-Doped ZnO (AZO) Electrodes on Dye-­Sensitize Solar Cell (DSSC) Performance To investigate the DSSC activities of the AZO electrodes in the work of Tyona et al. [10], all experiments were executed in a single compartment cell in polyiodide solution as the electrolyte [24, 48]. Photoelectrode (AZO) area of 1 cm2 was illuminated

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Fig. 9.6  Experimental setup for the measurement of photoelectrochemical current of the DSSC: (a) photograph of the fabricated DSSC (b) a schematic diagram of the setup [10]

with a power input (Pin) of 80 mWcm−2 from a xenon arc lamp [10]. The PEC properties of Al-doped zinc oxide electrodes were measured by constructing model designed cells, n-AZO + rhodamine 6G dye/0.1  M polyiodide/platinum/SCE as shown in Fig. 9.6b [10]. The J − V characteristics of Al-doped zinc oxide DSSCs are illustrated in Fig. 9.7a. They explained that the dark currents measured at each electrode were practically nonzero because of partial illumination from the background light. Jsc and Voc measured at all the photoelectrodes under illumination were practicably high; with maximum values for the electrode with Al:Zn of 2  at.%. These Jsc and Voc were relatively higher than those measured for unsensitized Al-doped zinc oxide electrode (0.092 mWcm−2 and 372 mV) subjected to the same conditions. This proposes that Al-doped zinc oxide electrodes sensitized with rhodamine 6G dye would be more photoactive than the non-sensitized [10, 24]. It is clear that rhodamine dye has enhanced the photoresponse of the Al-doped zinc oxide electrodes into the visible part of the solar spectrum (Fig. 9.4d), which led to the absorption of more photons with appropriate energy [21]. In addition, the photoexcited dye molecules and the valence electrons of the Al-doped zinc electrode were simultaneously injected into the conduction band (CB) of the photoelectrode. These raise the toll of charge carrier transport into the CB, which leads to an increase in photocurrent [22, 24]. Jsc and Voc exhibit high sensitivity to aluminum content, with the maximum current and voltage measured for Al: Zn = 2 at.% [10]. The evaluated values of Jsc, Voc, power conversion efficiencies, and fill factors (FFs) for all the Al-doped zinc oxide electrodes are recorded in Table 9.1. Electrochemical impedance spectroscopy (EIS) of Al-doped zinc oxide electrodes was performed with a Frequency Response Analyzer connected to a Potentiostat in a two-electrode mode. Figure 9.7b–d shows the impedance spectra for Al-doped zinc oxide electrodes sensitized with rhodamine 6G dye, measured at bias potentials: 0.37, 0.51 and 0.45 V in the dark.

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Fig. 9.7 (a) J − V characteristics of Al-doped zinc oxide electrodes sensitized with rhodamine 6G dye with various aluminum contents and Nyquist plots of Al-doped zinc oxide DSSCs with various aluminum contents: (b) 1 at.% (c) 2 at.% (d) 5 at.% [10] Table 9.1  DSSC solar cell parameters of AZO electrodes sensitized with rhodamine 6G dye [10] AZO electrodes with % Al 0 1 2 5

Photocurrent density (Jsc) [mAcm−2] 0.09 0.32 2.46 1.55

Photovoltage (Voc) [mV] 306 374 515 454

Jmax (mA/ cm2) 0.07 0.22 1.79 0.75

Vmax (mV) 212 239 397 298

Efficiency η (%) 0.02 0.07 0.89 0.27

Fill factor (FF) 0.54 0.44 0.57 0.34

From their results, Fig. 9.7b, for Al:Zn of 1 at.%, the interfacial charge transfer resistance is shown by a semicircle in the high frequency part of the Nyquist plot. The impedance is fairly high, contributed largely by the photoelectrode and electrolyte [39]. The electron transport resistance Rt at the photoelectrode is low consequent to the low series resistance of the rhodamine dye-sensitized electrode [10, 12, 49]. The charge diffusion resistance at the platinum counter electrode/electrolyte/ photoelectrode interface in the mid and low frequency regions show low transfer

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activity as a result of fairly high charge transfer resistance of the electrolyte as seen in Fig. 9.7b and Table 9.2a. This explains the drop in photocurrent of the cell as previously seen in Fig. 9.7a. It was also noticed that the Al-doped zinc oxide electrode with Al:Zn of 2 at.% indicated the least charge transfer resistance, Rct (Fig.  9.7c), which inferred faster detachment of photogenerated charges. Elevation in Rct was seen in Fig.  9.7d for Al:Zn of 5 at.% as a result of rise in crystallographic defects in the film, which accordingly lessened charge separation and transfer in the cell and hence lowered PV activities. In Fig. 9.7c, d, the zero series resistance of the photoelectrodes as shown in the equivalent electrical analogs, Table 9.2b, c), also subscribes to the reduction in electron recombination activities in the photoelectrode and accelerates dye molecules injection and transfer; hence, improving PEC cell activities as noticed in Table 9.1.

9.5  Effects of Indium Dopant on ZnO Properties Indium is a post-­transition metal that forms about 0.21 parts per million of the Earth’s crust [50]. It has a melting point above sodium and gallium, but less than lithium and tin [50]. Chemically, indium is close to gallium and thallium, and it is mostly in-between the two in terms of its characteristics [4, 51]. Indium is often utilized in making alloys and is often considered as the “metal vitamin,” which means that a small quantity of indium can cause remarkable changes in an alloy [3, 51]. For instance, the addition of little quantities of indium to gold and platinum alloys causes them to be much stronger [3, 51]. Previous researches have shown that conductivity, length, diameter, and charge carrier density of ZnO nanorods can be increased by n-type doping with Ga, Al, or In, thereby improving its photovoltaic outputs [39]. It is also known that indium impurities produce more photoactivities than Al impurities in CuInSe2 solar cells [39]. Studies have shown that doping zinc oxide with indium decreases the room temperature resistivity of the films; nevertheless, in some cases, the resistivity tends to increase for heavily doped films [28, 29]. A similar rise and posterior lessening with indium impurities were revealed for optical properties [28, 29]. Tyona et al. [39] studied the impact of In impurities on the optoelectronic properties of chemically grown zinc oxide thin films. In their experimental procedures, In concentrations of 0–6 at.% from InCl3 were employed to dope 0.1 M Zn(NO3)2.6H2O to generate various specimens of In-doped ZnO (IZO). Aqueous NH3 (28%) was employed to tune the pH of the chemical bath to 10.5 during the experimental process with a bath temperature of 353 K. A pH of 10.5 was imperative in order to lower the precipitation of In(OH)3 since lesser or more pH would cause the growth of low adherent and low crystalline In-doped zinc oxide nanorods. Additional precipitation in the solution was eschewed by steering the bath all-round the duration of growth. Fine and well-adherent In-doped zinc oxide films were deposited on a microscope glass slide and stainless steel substrates after 5 h. Various characterizations were used to study the IZO thin film electrodes.

Rs

C1

R1

Rs = 31.60 Ω R1 = 1.12 × 104 Ω C1 = 6.96 × 10−06 Ω W = 8.16 × 10−05 Ω

(a)

w R2

R1

Q1

Q2

Qy1 = 2.07 × 10−10 Ω Qa1 = 1.45 Ω R1 = 1.42 × 103 Ω R2 = 46. 29 Ω Qy2 = 9.14 × 10−05 Ω Qa2 = 0.63 Ω

(b)

R2

R1 C2

C1 Q1 R3

C1 = 6.97 × 10−08 Ω R1 = 84.20 Ω R2 = 256.89 Ω C2 = 7.03 × 10−07 Ω Qy1 = 8.16 × 10−05 Ω Qa1 = 0.69 Ω R3 = 4.31 × 103 Ω

(c)

Table 9.2  Equivalent electrical analog adopted for modeling the electrochemical cells of AZO electrodes in a two-electrode arrangement with the attending variables realized for the dark condition at the potential bias of 0.374 V, 0.515 V, and 0.454 V (a) 1 at.% (b) 2 at.% (c) 5 at.% [10]

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9.5.1  Film Thickness Studies It is well known that many properties of thin films that govern their suitability in different applications are guided by film thickness; thus, the verification of film thickness is necessary [20, 39]. For PEC solar cell application, the model film thickness for superlative light absorption across the functional solar spectrum is in the range of 8–12 μm [20, 39, 52]. According to Tyona et al. [39], the variation in film thickness of In-doped zinc oxide with deposition time at a fixed bath temperature (353 K) is shown in Fig. 9.8. It was determined using a surface profilometer (AUB 105 Technology, Model XP-1). Their results showed a time-dependent film thickness that is also sensitive to the content of In in the bath. Increasing the temperature of the bath to 353 K, precipitation starts in the solution and metal ions were slowly freed to form a thin solid film of In-doped zinc oxide on the substrate [53]. Figure 9.8 also demonstrated the firm reliance of the film thickness growth on In impurity content in the bath [2–4]. It is explained that this could be attributed to the rise in the number of free metal ions in the bath due to addition of the In impurities which rises the density of nucleation sites, and thus supporting rapid growth of the thin films with porous and rough surfaces [50, 54, 55]. A scrutiny of the growth process of In-doped zinc oxide by chemical bath deposition method showed that heating the alkaline zinc nitrate bath to 353 K for about 1 h instigates precipitation, and growth of In-doped zinc oxide starts by heterogeneous reaction at the substrate surface, giving rise to film thickness increasing with time [39, 51, 56]. The growth period was varied from 2 to 5 h (Fig. 9.8). It is noticed that at the early stage of the growth procedure, the film growth rate was speedily; nevertheless, after 4 h, the rate slowed down until the final thickness was achieved at 5 h. This could be due to a decline in the number of free metal and chalcogenide ions in the bath [11, 39, 51, 57]. Records of the film thickness and surface roughness as measured are shown in Table 9.3.

Fig. 9.8 (a) Variation of film thickness with deposition time for different contents: 0 at.%, 2 at.%, 4 at.%, and 6 at.% (b) various film morphologies [39]

9  Doped Metal Oxide Thin Films for Dye-Sensitized Solar Cell and Other… Table 9.3  Records of measured film thickness and surface roughness of In-doped zinc oxide with various indium contents [39]

Thickness IZO electrodes with % In (μm) 0 5.70 2 6.71 4 9.46 6 11.56

249 Surface roughness (nm) 123.0 128.0 133.0 137.0

Fig. 9.9  XRD patterns of the In-doped zinc oxide synthesized with various In concentrations: 0 at.%, 2 at.%, 4 at.%, and 6 at.% [39]

9.5.2  Structural Studies Figure 9.9 represents their result for the XRD analysis of the In-doped zinc oxide nanostructures, synthesize with rising content of In (0  at.%, 2  at.%, 4  at.%, and 6 at.%). The XRD figures were match up to ZnO (JCPDS File No.036–1451), polycrystalline, with hexagonal wurtzite structures. High intensity peaks were observed along the c-axis (0 0 2) for all specimens, which confirm the anisotropic development of zinc oxide nanorods which SEM morphology would also establish (Fig. 9.11) [11]. Nevertheless, less intensive peaks corresponding to lattice numbers (100), (102), (102), and (103) were seen. The existence of a distinctive peak along the z-plane in all the film specimens indicates uniformity in the phase development of the In-doped zinc oxide. No peak of diffraction similar to In(OH)3 or In2O3 as suggested by some researchers was seen [11]. This shows that In replaces Zn atom without creating any impurity phase [58, 59]. It was also noticed that, by incorporating In into ZnO (i.e., 2 at.%), the prominent diffraction peak in the c-axis diminished apparently as shown in Fig. 9.9. This could be adduced to the introduction of stress in zinc oxide by In as a result of the contrasting sizes of In+3 (r = 0.080 nm) and Zn+2 (r = 0.074 nm) [11, 39]. Notwithstanding, increasing In concentration to

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4  at.%, the intenseness of the (002) peak was greatly increased, whereas others remain unaltered. This proposes that the rise in In concentration to 4 at.% causes stress relaxation in the film because of increase in charge carrier concentration in the conduction band of the semiconductor which enhances the crystallinity of the film [50, 55]. Additional rise in In content to 6 at.% caused remarkable shrinkage in the intensity of the (002) peak as shown in Fig. 9.9; hence, proposing that 4 at.% is the optimum In impurity content for achieving the best quality In-doped zinc oxide film. These results have definite consequences on enhancing the morphological structure, enlarging the thickness, specific surface area, and photoelectric activities of In-doped zinc oxide films [11, 39, 51]. The crystallite sizes (D) of the In-doped zinc oxide films were estimated in accordance with the full width at half maxima (FWHM) of peaks of diffraction using Scherrer’s equation, Eq. 9.1 [1, 3, 50], where β is the broadening of the diffraction lines measured at half maximum intensity (radians), λ = 1.5406 Å is the wavelength of the CuKα radiation. The crystallite sizes were found in the range of 16 to 29 nm for un-doped and In-doped zinc oxide along (002) plane.



D

0.9  cos 

(9.1)

9.5.3  Optical Studies Figure 9.10a shows the result for optical spectra absorption of In-doped zinc oxide films determined on a spectra range of 360–800 nm. As could be seen, the absorbance is dominated with ultraviolet absorption (360–400 nm) and very little absorption in the visible range for the specimens with In:Zn of 0 at.% and 6 at.%. They observed a red-shift in absorption band edges for In:Zn of 2 and 4 at.% films by order of 492 and 540  nm, respectively. This notable enhancement in absorbance initiated by indium impurity atoms could be assigned to the light-scattering impact of zinc oxide nanorods [39, 51, 58], which was improved by a rise in the rods diameter and length arising from In atoms as would also be established by the morphological structures (Fig. 9.11) [5, 19, 28, 29, 39]. This improvement in absorbance is ultimately higher than those currently existing in many reports for In doping. The inset of Fig. 9.10a shows the transmittance spectra, which is relatively high (42–74%) for all the In-doped specimens, with In:Zn of 4 at.% showing 72%. The high transmittance seen with the In-doped zinc oxide electrodes, especially for In:Zn of 4 at.%, is in accord with the improved absorbance earlier noticed. This justifies that the grown electrodes are appropriate for photovoltaic application [13, 23]. Figure 9.10b presents the plot of Tauc’s relationship between the absorption coefficient, α, and the photon energy, hν, as seen in Eq. 9.2, where αo is a constant. Eg is the bandgap and n is a constant that depends on the probability of transition (it takes values as 1/2, 3/2, 2, and 3 for direct allowed, direct forbidden, indirect

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Fig. 9.10 (a) Spectra absorbance of the In-doped zinc oxide films synthesized with various indium contents: 0 at.%, 2 at.%, 4 at.%, and 6 at.%. (Inset: transmittance spectra) and (b) Plot of Tauc’s relationship between the absorption coefficient, α, and the photon energy, hν, for In-doped zinc oxide films with different indium concentrations. (Inset: bandgap trend vs. In impurity concentration in zinc oxide) [39]

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Fig. 9.11  Scanning electron microscope images of the In-doped zinc oxide films with different indium concentrations on a magnification of X10000: (a) 0 at.% (b) 2 at.%, (c) 4 at.%, and (d) 6 at.% [39]

allowed, and indirect forbidden transition, respectively) [2, 4]. They showed that extending the linear portions of the curves in the Tauc’s plot to the photon energy axis (Fig. 9.10b) determines the bandgap energies (Eg) of the In-doped zinc oxide electrodes.





 o  h  Eg  h

n



(9.2)

The bandgap values revealed a lowering from 3.15 eV to 2.92 eV (with respect to 0 at.%), showing notable sensitivity to indium concentration as shown in the inset of Fig. 9.10b. The bandgap modification with indium contents, as illustrated in the inset, infers that doping zinc oxide with indium would regulate its bandgap appropriately with a focus to enhancing its absorption further into the visible spectrum as is shown with 4 at.%, which has the much suitable bandgap [53, 58]. The lowering

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of bandgap energy is possibly due to the replacement of Zn+2 by In+3 on the zinc oxide lattice, which leads to the development of new recombination centers with lesser energy state [31, 50]. The bandgap lessening as seen here can promote the excitation of more valence band electrons to the conduction band in the presence of visible light photons with lower energy state [31, 50]. This is advantageous for improving photovoltaic performance of a UV-active material [59].

9.5.4  Morphological Studies The scanning electron microscope images of the In-doped zinc oxide films in the work of Tyona et al. [39] are shown in Fig. 9.11. The images revealed randomly aligned nanorods morphology which corresponds to the zinc oxide hexagonal wurtzite structure as previously recognized by XRD. It is noticed from the side view of the In-doped zinc oxide film (inset, Fig. 9.11c) that all zinc oxide nanorods grew vertically and strongly fastened to the substrate [39]. The length, size, and density of the nanorods demonstrated high sensitivity to In concentration in the films. The un-doped specimen (0  at.%) was recognized with dense and randomly aligned nanorods of average diameter ~57 nm as shown in Fig. 9.11a. With the introduction of indium in the system, an evident rise in the average diameter of the nanorods to ~81 nm (Fig. 9.11b–d) was seen. The enlarged diameter of the nanorods at In:Zn of 4  at.% would significantly expand the specific surface area of the In-doped zinc oxide films, thereby enhancing its absorbance [27, 37, 60, 61]. The doped nanorods exhibited nonuniformity in rod sizes; this uneven development is not much familiar; nevertheless, Rathore and Sarkar proposed the consequence of the primary precursor, Zn(NO3)2.6H2O, on the growth development [53]. Additional scrutiny on the surface structures of the films revealed a little reduction in the density of nanorods at the incorporation of indium impurities (Fig. 9.11b). This could be consequent to the variance in the ionic radii of zinc and indium as previously proposed by XRD [2–4, 33, 54]. Rising the concentration of indium to 4 at.%, the nanorods showed the densest morphology (Fig. 9.11c); notwithstanding, at 6 at.%, there was little reduction in the size of the rods as is evident in Fig. 9.11d. There is a demonstration of enhancement in the crystalline quality of the zinc oxide nanorods doped with indium, wherein, 4 at.% content demonstrates the best crystal quality. This result favorably agrees with XRD investigation.

9.5.5  Surface Wettability Studies Surface wettability gives knowledge regarding the interactivity of liquids and solids in touch. It is designated by the value of contact angles, an infinitesimal variable [62]. It is a much essential variable in surface science that is easy and effectual for analyzing surface energies [31, 32, 62, 63]. Figure 9.12 illustrates the water contact

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Fig. 9.12  Contact angles variation with indium impurity contents of the In-doped zinc oxide films [39]

angles variation with indium impurity contents in In-doped zinc oxide films in the work of Tyona et al. [39]. The water contact angles exhibited remarkable responsiveness to indium content in the films. With the introduction of In impurity atoms in zinc oxide, the water contact angle reduced from 63.2° (un-doped Fig. 9.12) to 41.7°. Additional increase in the concentration of indium to 4  at.%, lessened the water contact angle to its base value of 23.9° which rose again to 33.8° as indium concentration rises to 6 at.%. This means that the porosity and surface roughness of the nanorods increase with indium impurities, which establish a rise in the average diameter of the rods, thus making the water to move into the pores and crevices and causing the contact angle much hydrophilic [31, 32, 63–65]. Furthermore, lesser values of contact angle are favourable for electrolyte percolation all-round the porous film, which is very beneficial for electrochemical solar cells for enhanced photovoltaic activities [39]. This result is in accord with preceding considerations with XRD and SEM.

9.6  T  he Impact of In-Doped ZnO (IZO) Electrodes on PEC Solar Cell Performance The performance of the electrochemical solar cell is based on the semiconductor– electrolyte junction. The role of electrolyte in this class of solar cells is of much importance because it functions as a medium for charge transfer connecting the photoelectrode and counter electrode [39, 64, 66]. The photoresponses of the

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In-doped zinc oxide electrodes in this work were considered by fabricating typical arrangement cells, n-IZO/0.1 M Na2SO4/platinum/SCE. The J − V characteristics of In-doped zinc oxide solar cells considered in the dark and under light (80 mW/cm2) are illustrated in Fig. 9.13. The dark currents measured against each electrode were almost insignificant. The photocurrent densities, Jsc, and open-circuit voltages, Voc, varied significantly with indium concentration in zinc oxide, with a maximum for 4  at.% electrode hence confirming the preceding properties of the In-doped zinc oxide. The Jsc and Voc of the In-doped zinc oxide electrodes were distinctively above those determined for un-doped zinc oxide electrode (0.53  mA/cm2 and 511  mV) subjected to the same conditions. This proposes that the elevated enhancement in optical properties of the zinc oxide electrodes induced by indium impurities caused the harvesting of many suitable photons, leading to improved photovoltaic activities. Power conversion efficiencies (η or PCE) and fill factors (FFs) were determined using Eqs. (9.3) and (9.4) [39, 67]. The Jsc, Voc, PCE, and FFs determined for all the In-doped zinc oxide electrodes are presented in Table 9.4.

 FF 

I max  A / cm 2 Vmax V  Pin W / cm 2 

(9.3)

I max  A / cm 2 Vmax V  I sc  A / cm 2 Voc V 

(9.4)

Fig. 9.13  J − V characteristics of In-doped zinc oxide electrodes with various concentration of indium: 0 at.%, 2 at.%, 4 at.%, and 6 at.% [39]

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Table 9.4  Measured and calculated photovoltaic solar cell variable for the In-doped zinc oxide electrodes [39] IZO electrodes with % In 0 2 4 6

Photocurrent density (Jsc) [mA/ cm2] 0.53 2.92 5.61 2.24

Photovoltage (Voc) [mV] 511 522 738 570

Jmax (mA/ cm2) 0.32 2.07 4.37 1.61

Vmax (mV) 352 292 522 254

Efficiency η (%) 0.14 0.76 2.85 0.51

Fill factor (FF) 0.42 0.40 0.55 0.32

9.7  Conclusions This chapter has carefully taken a review of doped metal oxide thin films with specific emphasis on ZnO. The chapter gave a general overview of ZnO, highlighting its advantages and drawbacks. Prominent among the drawbacks were low electrical conductivity and low spectra absorption due to wide bandgap of ZnO. Doping with Al and In and also dye-sensitization of the photoelectrode were proposed as key techniques to enhance the electrical conductivity and the spectra absorption of the ZnO thin films. The review considered chemical synthesis, characterization, and application of AZO and IZO in DSSC and non-dye–loaded PEC solar cells. Two works of Tyona et al. [10, 39] were used to illustrate the chemical synthesis, characterization, and application of AZO and IZO electrodes in DSSC and non-dye–loaded PEC solar cells. The study of Tyona et al. [10, 39] revealed that there was a notable enhancement in electrical conductivity of ZnO upon doping with Al and In, respectively, due to the increase in charge carrier concentration in the conduction band. Modification of bandgap of ZnO became very obvious upon doping, and absorption band edges were red-shifted into the visible spectrum, which confirmed improvement in spectra absorption. The study also revealed that with the sensitization of the AZO photoelectrodes with rhodamine 6G dye, the absorption band edges were further red-­ shifted into the visible region and that confirmed further improvement in spectra absorption of ZnO thin film electrodes. Structural, morphological, and other properties of ZnO such as film thickness and surface wettability were evidently seen to improve upon doping. These electrodes have proved very effective in DSSC and non-dye–loaded PEC solar cells applications. In conclusion, the significance of the enhancement in the key properties of ZnO with doping and dye-sensitization of the photoelectrodes is that new metal oxide semiconductor (ZnO) thin films has been achieved which are very effective in DSSC and non-dye–loaded PEC solar cells. These materials can also be suitable for other optoelectronic and electronic applications.

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References 1. Tyona MD (2019) Doped zinc oxide nanostructures for photovoltaic solar cells application. In: Zinc oxide based nano materials and devices. IntechOpen, Rijeka, Croatia. https://doi. org/10.5772/intechopen.78819 2. Wang H, Xie C (2006) Controlled fabrication of nanostructured ZnO particles and porous thin films via a modified chemical bath deposition method. J Cryst Growth 291:187–196 3. Kim MS, Yim KG, Kim S, Nam G, Lee DY, Kim JS, Kim JS, Leem JY (2012) Growth and characterization of indium-doped zinc oxide thin films prepared by sol–gel method. Acta Phys Pol A 121:217–220 4. Hu Z, Jiao B, Zhang J, Zhang X, Zhao Y (2011) Indium-doped zinc oxide thin films as effective anodes of organic photovoltaic devices. Int J Photoenergy 2011:1–5 5. Salh A, Hyeonwook P, Woo KK (2019) Optimization of intrinsic ZnO thickness in Cu(In,Ga) Se2-based thin film solar cells. Materials 12:1365–1381 6. Yao PC, Hang ST, Lin YS, Yen WT, Lin YC (2010) Optical and electrical characteristics of Al-doped ZnO thin films prepared by aqueous phase deposition. Appl Surf Sci 257:1441–1448 7. Drici A, Djeteli G, Tchangbedgi G, Deruiche H, Jondo K, Napo K et al (2004) Structured ZnO thin films grown by chemical bath deposition for photovoltaic applications. Phys Status Solidi (a) Banner 201:1528–1535 8. Li Y, Gong J, Deng Y (2010) Hierarchical structured ZnO nanorods on ZnO nanofibers and their photoresponse to UV and visible lights. Sens Actuators A Phys 158:176–187 9. Lao CS, Liu J, Gao P, Zhang L, Davidovic D, Tummala R et al (2006) ZnO nanobelt/nanowire Schottky diodes formed by dielectrophoresis alignment across au electrodes. Nano Lett 6:263–275 10. Tyona MD, Jambure SB, Lokhande CD, Banpurkar AG, Osuji RU, Ezema FI (2018) Dye-­ sensitized solar cells based on Al-doped ZnO photoelectrodes sensitized with rhodamine. Mater Lett 220:281–284 11. Hao Y, Yang M, Li W, Qiao X, Zhang L, Cai S (2000) A photoelectrochemical solar cell based on ZnO/dye/polypyrrole film electrode as photoanode. Sol Energy Mater Sol Cells 60:349–359 12. Lopes T, Andrade L, Ribeiro HA, Mendes A (2010) Characterization of photoelectrochemical cells for water splitting by electrochemical impedance spectroscopy. Int J Hydrogen Energy 35:11601–11608 13. Tyona MD, Osuji RU, Ezema FI, Jambure SB, Lokhande CD (2016) Enhanced photoelectrochemical solar cells based on natural dye-sensitized Al-doped zinc oxide electrodes. Adv Appl Sci Res 6:7–20 14. Alkahlout A, Al Dahoudi N, Grobelsek I, Jilavi M, de Oliveira PW (2014) Synthesis and characterization of aluminum doped zinc oxide nanostructures via hydrothermal route. J Mater 2014:235638–235646 15. Jambure SB, Patil SJ, Deshpande AR, Lokhande CD (2014) A comparative study of physico-­ chemical properties of CBD and SILAR grown ZnO thin films. Mater Res Bull 49:420–425 16. Janotti A, Van de Walle CG (2009) Fundamentals of zinc oxide as a semiconductor. Rep Prog Phys 72:126501 17. Snure M, Tiwari A (2008) Band-gap engineering of Zn1 420 nm) studies demonstrating high, stable on/off photocurrent activity. The schematic Fig. 12.15e exhibits PEC mechanism of ZnO/Au/Al2O3 nanostructures. Authors appealed that electrons from the defect states of ZnO transferred to Fermi level of Au and after visible light illumination, hot electrons from Au back to the conduction band of ZnO nanorods. Surface passivation of Al2 O3 caps photocorrosive defects present on ZnO leading to supression of charge recombination. Moreover, it provides minimum charge transfer resistance resulting in higher photocurrent for prolonged time and stability.

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Fig. 12.15 (a) LSV curves, (b) transient photocurrent density (at 0 V vs. Ag/AgCl), (c) chronoamperometry (at 0.5  V vs. Ag/AgCl with visible light with λ  >  420  nm), (d) photoconversion efficiency, and (e) schematic of charge transfer mechanism of ZnO/Au/Al2O3 heterostructure [88]

12.4  Decoration Vs. Doping As discussed in the earlier section, to improve the PEC performance by harvesting more and more visible part of electromagnetic radiation strategies like doping of metal or nonmetal ions in the matrix of the ZnO crystal or decoration of ZnO nanostructures by plasmonic metal was adopted by several researchers. As described in Fig. 12.16, doping of metal or nonmetal ion can be performed by three different ways [89]. First two cases are achieved by metal ion doping whereas third corresponds to nonmetal ion doping as shown in Fig. 12.16; (1) promotion of the valence

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Fig. 12.16  Modification in bandgap energies of semiconductor  due to  in (a) p-type and  (b) n-type doping, and (c) nonmetal ion doping [16]

band maximum to form acceptor level, (2) decrement in the conduction band minimum to form donor level, and (3) generation of the localized energy levels near to the valence band [16]. In the first two cases, metal ion tries to substitute Zn2+ ion that leads to the formation of intra bandgap levels states within the forbidden energy gap of ZnO leading to its narrowing. Doping gives rise to the creation of either an acceptor level beneath the original conduction band (Fig. 12.16a) or a donor level over the original valence band (Fig. 12.16b). Such shallow and deep level states in the energy gap of valence and conduction band of ZnO nanostructure cause absorption of visible light and increment in the charge carrier density leading to enhancement in the PEC performance. The position of metal dopants in the ZnO matrix whether it is substitutional or interstitial sites or segregated to grain boundaries is determined by the ionic radius size and concentration of the metal dopant. Though bandgap of ZnO comes under UV regime, the intra levels created within bandgap lead to absorption of waves of higher wavelength in the visible region. Looking at the literature, doping of the nonmetal ion in ZnO crystal structure is found to be more beneficial since it forms upper level shifting of the valence band. Moreover, bandgap reduction takes place without giving rise to the formation of intra band levels, therefore occurrence of charge recombination centers is highly impossible and enhanced PEC performance can be achieved [44, 90, 91]. Orbitals of nonmetal dopant make hybridization with oxygen 2p states of ZnO leading to the promotion of the valence band and thus bandgap narrows resulting in extended visible light absorption [44]. Several nonmetal ions with lower electronegativity than oxygen and similar size like ­nitrogen, carbon, chlorine, selenium are chosen by the researchers for doping them in ZnO crystal and utilized for PEC application [70]. We have enlisted few recent approaches by the researchers for the doping of metal or nonmetal ion in ZnO crystal as given in Table 12.2 where the morphology of ZnO nanostructure and its PEC performance values are also mentioned. Careful observation of Table 12.2 demonstrated that the photocurrent densities of many of metal ion doping can reach up to few microamperes. Only rare attempts were found where photocurrent densities can boost up to few milliamperes. Intra gap levels play a crucial role in determining final PEC performance (i.e., photocurrent density or efficiency) which may be the plau-

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Table 12.2  Metal and nonmetal ions doped ZnO nanostructures for PEC applications Sr. No. Doped ion 1. Nitrogen

2.

Nitrogen

3.

Silver

4.

Silver

5.

Copper

6. 7.

Aluminum Boron

8.

Carbon

9.

Chlorine

10.

Cobalt

11. 12. 13. 14. 15.

Cobalt Chromium Copper Copper Copper

16.

Copper

17.

Lead

18. 19.

Gallium Gallium

20.

Hafnium

21. 22.

Potassium Lithium

23.

Magnesium

ZnO morphology Nanowire

PEC performance Efficiency of 0.15% at an applied potential of +0.5 V (vs. Ag/AgCl) Tetrapods Photocurrent of 0.99 mA/cm2 at +0.31 V vs. Ag/AgCl Nanorod Efficiency of 1.1% and photocurrent of 0.96 mA/cm2 Oval aggregate Photocurrent of 65 μA/cm2 grains Spherical Efficiency of 0.1% at an grains applied potential of +0.5 V (vs. Ag/AgCl) – Photocurrent of 1.2 μA/cm2 Nanorod Efficiency of 2.6% and photocurrent of 2.054 mA/ cm2 vs. Ag/AgCl Hollow Photocurrent 1.32 mA/cm2 at microsphere 0.85 V (vs. RHE) Nanorod Photocurrent 2.16 mA/cm2 at 1.2 V vs. Ag/AgCl Nanorod Photocurrent 0.15 mA/cm2 at 1.33 V vs. RHE Nanograin Photocurrent 2.3 μA/cm2 Nanograin Photocurrent 55 μA/cm2 Nanograin Photocurrent 10.6 μA/cm2 Nanograin Efficiency of 2.16% Nanorod 0.068 mA/cm2 at 0.9 V vs. SCE Nanorod Efficiency of 0.35% and photocurrent of 0.92 mA/cm2 Circular grains Efficiency of 0.35% at 0.5 V vs. Ag/AgCl Nanorod Photocurrent 0.6 mA/cm2 Nanograins Photocurrent 0.14 mA/cm2 at 0.5 V vs. SCE Columnar Photocurrent 8 μA/cm2 at microstructure 0.6 V vs. NHE Nanorod Photocurrent 8 μA/cm2 Nanorod Efficiency of 0.24% at 0.7 V and photocurrent of 0.44 mA/ cm2 Nanorod Photocurrent density of 0.35 mA/cm2 at 1.5 V vs. Ag/ AgCl

Increment in the photocurrent Ref. ~23.5 times higher [90] than ZnO ~8.25 times higher [44] than ZnO ~3.4 times higher [91] than ZnO – [95] ~16.8 times higher [98] than ZnO – ~318 times than ZnO

[99] [100]

~10.5 times than [101] ZnO ~21 times than [102] ZnO ~2 times than ZnO [103] – ~20 times than ZnO ~12 times than ZnO – ~3.3 times than ZnO ~3 times than ZnO

[104] [105] [106] [107] [108] [70]

–

[109]

~3 times than ZnO [110] ~7 times than ZnO [93] ~2.5 times than ZnO – ~0.7 times higher than negative poled ZnO ~9 times than ZnO

[111] [112] [92]

[113]

(continued)

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319

Table 12.2 (continued) Sr. No. Doped ion 24. Nitrogen

ZnO morphology Branched nanorod

26.

Nanotube Nitrogendoped carbon with MOF Nickel Nanograins

27.

Selenium

Nanoplates

28.

Vanadium

Nanosheets

29.

Yttrium

Nanorod

30.

Yttrium

Nanorod

25.

PEC performance Photocurrent density of 1.7 mA/cm2 at 0.8 V vs. Ag/ AgCl Photocurrent density of 0.4 mA/cm2 at 0.4 V vs. Ag/ AgCl Efficiency of 4.2% and photocurrent of 4.6 mA/cm2 Photocurrent density of 0.2 mA/cm2 Photocurrent density ~35 mA/cm2 at 0.8 V vs. RHE Photocurrent density 1.59 mA/cm2 at 1.23 V vs. RHE Photocurrent density 0.84 mA/cm2 at 1.23 V vs. RHE and efficiency 0.4%

Increment in the photocurrent –

Ref. [114]

~13.3 times than ZnO

[115]

~1.78 times than ZnO ~13.3 times than ZnO ~1.6 times than positively poled ZnO 0.44 times than ZnO

[116]

0.47% times than ZnO

[120]

[117] [118]

[119]

sible reason behind discrepancies in the photocurrent densities as discussed above. Though, intragap level states produced due to metal doping give rise to higher light absorption of visible light range, the PEC performance may weaken because the new states can act as charge recombination centers [89]. On the other hand, nonmetal doping just shift valence band maximum upwards without forming deep intra levels resulting into low probability of recombination centers creation and thereby better photocurrent density can be achieved as compared to metal doping. ZnO nanostructures morphology plays a vital role in deciding final PEC performance; therefore, let us take a look at the difference in the synthesis process and morphology in doping and decoration processes. Generally, decoration of plasmonic metal nanoparticles over ZnO nanostructure is multistep, tedious, and complex procedure wherein doping of metal or nonmetal ion can be performed in one to two steps [58, 92]. For instance, [58] Zhang et al. decorated Au nanoparticles over branched ZnO nanorods (Fig. 12.17a) for PEC application and achieved high photocurrent of 1.45 mA/cm2 at 1 V vs. RHE. During the synthesis process, a seed layer was applied over the substrate for the uniform growth of ZnO rods. Thereafter, preferentially oriented nanorods were coated with ZnS and annealed at a higher temperature. Further smaller sized ZnO nanorods were grown leading to three-­ dimensional growth. Lastly, Au nanoparticle decoration over whole complex nanostructure and fine PEC performance was achieved. On the other hand, by using hydrothermal method, Lee et  al. [92] synthesized Li doped ZnO nanorods (Fig. 12.17b) by keeping the substrate upside down and after annealing high voltage polling of ZnO nanorods under application of the high electric field leading to the

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efficiency of 0.24% at ~0.7 V and photocurrent of 0.44 mA/cm2. As discussed in the earlier section, ZnO is usually found in the nanorod morphology due to its wurtzite crystal structure and it serves a critical role in charge transfer. However, most of the times doping makes changes in wurtzite crystal structure that leads to slight or major deviation in the nanorod morphology [93]. Dependent upon dopant concentration, doping introduces a certain degree of residual stress and reduction in the crystallinity. On contrary, highly and single-crystalline pristine ZnO crystal favors rapid delivery of photogenerated charge carriers to the conducting substrate [94]. Therefore, a higher concentration of dopant may produce variation in the pristine crystal structure of ZnO resulting in hampering of charge transport leading to the adverse effects on PEC performance. Additionally, in many cases, instead of nanorod morphology, nanograins nanostructures were also reported due to doping (increase in the concentration of doping) of metal ion [95]. Unlike to doping, ­decoration of plasmonic metal nanoparticle process is performed on the readymade ZnO nanostructures (mostly nanorod); thus, there is no question of a change in the crystal structure. In short, the decoration process retains the crystal structure of ZnO which helps in the extended light absorption and scattering as well as promotes charge transfer (as discussed above). Sticking to previous point of morphology of doped ZnO nanostructures, very few attempts managed retention of nanorod structures; however, further increase in the dopant concentration may lead to some critical issues like establishment of separate nanoparticle of dopant metal, decrement in the electron mobility, and increment in disorder [16, 96]. For instance, increment in the Ag precursor concentration may not give rise to uniform replacement of Ag ion in place of Zn ion. Also, ionic radius of Ag ion is greater than Zn ion; hence at certain dopant concentration, separate agglomerated particles of Ag could be formed in grown nanostructures which act as a flaw for PEC application [96, 97]. Therefore, in comparison of both strategies, decoration finds out to be effective than doping.

12.5  Outlook and Frontiers The upcoming era in PEC field will be controlled by scalable fabrication of multifunctional photoelectrodes of various designs. Therefore, an appropriate understanding of growth mechanism and simple methods of synthesis for complex architectures of photoelectrode is in progress. Owing to LSPR of plasmonic metal nanoparticles, their hybrids with ZnO can boost conductivity and visible light harvesting. Substantial efforts have been dedicated towards modification in relative energy positions and junction formation, light absorption range, short diffusion length, morphology, charge carrier conductivity, increase in photogenerated charge carrier separation, and inhibition to photocorrosion to develop powerful photoelectrode. However, few suggestions are presented to cultivate plasmonic metal nanoparticles decorated ZnO photoelectrode to deliver practical values of PEC performance by looking at eclectic literature.

12  Plasmonic Metal Nanoparticles Decorated ZnO Nanostructures… ZnS shell

321 ZnO

ZnO nanowire ZnO seed layer Step 1

Step 3

(a)

Ton exchange

FTO substrate

An nanoparticles

Branched ZnO nanowire

ZnO seed shell

Step 4

Step 5 Au deposition

ITO coated Glass substrate

ZnO

Annealing in air

Step 2

Hydrothermal

Hydrothermal

ZnO seed layer

(b)

_

OH

Zn2+

Li+

Fig. 12.17  Schematics of typical synthesis procedure of (a) Au metal decoration over three-­ dimensional branched ZnO nanostructures [58] and (b) Li ion doping in ZnO nanorods [92]

1. For decoration of plasmonic metal nanoparticle, three-dimensional morphological form of ZnO is advantageous since its enlarged surface area with cavities hold the ability to form light scattering centers and also integrated arrangement with recurring regularity provides single-crystalline pathways for charge transfer. Careful tuning of crystal planes, configuration, the aspect ratio of nanorods morphology, etc., will play a decisive role in determining PEC performance. 2. If researchers are interested to grow one- or two-dimensional structures then complexity in the nanostructure can be introduced by growing the photoelectrode over three-dimensional substrate. Additionally, it will facilitate all-­ embracing penetration of the electrolyte solution leading to thoroughgoing interaction between photoelectrode and electrolyte. In line with the comment, special attention must be provided concerning the use of different electrolytes for greater stability and efficient charge transport.

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3. To increase carrier concentration approaches towards the synergetic effect of doping and decoration on ZnO photoelectrode by retaining morphological advantages are expected. 4. Dedicated work is demanded towards mechanisms like FRET of LSPR. 5. Comprehensive strategies are required to develop thin and crystalline passivation layer so that light absorbance should not affect and smooth charge transfer takes place. 6.  New materials (semiconductors/nanocarbons) may be incorporated in metal– ZnO hybrids for the building of better interfaces (type II or Z-scheme) leading to rapid charge transfer and separation. Acknowledgments  Mangesh A. Desai is thankful to Council of Scientific and Industrial Research (CSIR), India for awarding senior research fellowship (SRF). Authors are thankful to University grants commission (UGC).

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Chapter 13

Oxygen-Deficient Metal Oxide Nanostructures for Photocatalytic Activities Rahul B. Pujari and Dong-Weon Lee

13.1  Introduction The United Nations Development Program in the 2000 World Energy Assessment found that the yearly capability of sun-powered energy was 1575–49,837 exajoules (EJ). This is a few times bigger than the complete total world energy utilization, which was 559.8 EJ in 2012. Photocatalysis reactions can utilize such an enormous amount of solar energy and make an available simpler and usable form of chemical energies such as methane (CH4) and hydrogen (H2). It is an accelerated photoreaction carried out in the presence of a catalyst. There are various types of photocatalysis reactions such as photoelectrochemical water splitting, which utilize solar energy for the conversion of H2O into simpler and usable chemical energies to mankind. Metal oxide semiconductors are most often used in photocatalysis reactions owing to their ease of preparation, high chemical stability, and natural abundance. Unlike the continuum electron (e−) states of metals, the semiconductors have conduction and valence bands (CB and VB) separated by a finite forbidden energy band gap, where energy levels are absent. When the photon with the energy equal to or greater than band gap of metal oxide is incident, the e− is excited to the conduction band by generating a hole (h+) in the valence band. The photogenerated e− and h+ pair is known to be exciton. Besides exciton formation in the conduction and valence bands, their separation and transportation through bulk material to the surface of semiconductor are two other important factors in the success of photocatalysis. Metal oxides such as TiO2 absorb ultraviolet (UV) region solar radiation and exclude visible and infrared (IR) owing to large band gap (~3.2 eV). The UV spectrum is nearly 4% in the entire electromagnetic spectrum of solar radiation [1]. R. B. Pujari · D.-W. Lee (*) MEMS and Nanotechnology Laboratory, School of Mechanical System Engineering, Chonnam National University, Gwangju, Republic of Korea e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_13

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Although low band gap (2.1 eV) Fe2O3 delivers high solar-to-hydrogen (STH) conversion efficiency, poor electronic conductivity, short h+ diffusion length, and charge carrier life time reduce photocatalytic performance [2]. To address these challenges of metal oxides, different nonstoichiometric preparation methods are encouraged for the improvement of electronic and optical properties. The metallic and nonmetallic element doping has been carried out for a long time for enhancing the electrical and optical properties of metal oxides. However, recently, a newly popular strategy has been utilized, i.e., preparation of intrinsic oxygen defects (vacancies) in the metal oxides for the modulation of electronic and optical properties.

13.2  M  ethods for Introducing Oxygen Vacancies in Metal Oxide Nanostructures The vacancies in the crystalline materials are the intrinsic defects caused by the absence of atoms/ions at the crystal lattice sites, which are different from impurities in the material. The anion vacancies produce localized defect energy state near to conduction band minimum in the forbidden band gap of metal oxide, which forms F center or color center and acts as a donor state for electrons. While the cation vacancy is formed as mentioned above, the valence band maximum acts as an acceptor site for holes. The oxygen vacancies are the most common point defects in the metal oxides and act as shallow donors. It increases charge carrier density in the metal oxide that helps to improve electronic conductivity and facilitates charge transfer between electrode/electrolyte and electrode/substrate interfaces. Different types of methods have been employed to obtain the oxygen vacancies in the metal oxides. The oxygen vacancies can be obtained during the synthesis or in the post-­ synthesis treatment of metal oxides.

13.2.1  Doping of Elements It is the oldest method for oxygen vacancy formation in the metal oxide nanostructures, and metals such as Au, Fe, Cu, and Ag non-metallic elements like C, I, F, N, P, and S, and metalloid like B are used for the formation of oxygen vacancies in the metal oxide. The overall charge neutrality principle is used to form an oxygen vacancy in the doping process of metal oxide. Thus, when low ionic state elements (in comparison with host ions) are doped in the metal oxide, oxygen vacancies are formed above the valence band maximum. For example, oxygen vacancies are produced in TiO2 nanotubes grown on Ti foil. The two-electrode cell with TiO2 cathode, Pt anode, and an aqueous solution of boric acid (H3BO3) (electrolyte) was used for B doping. An electrochemical treatment was executed on TiO2 nanotubes at 1.8 V for 30 min. The doped B ions at the interstitial sites of anatase TiO2 balance the

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residual charge in titania and lead to the alteration of the oxidation state of few titanium ions from Ti4+ into Ti3+ in order to maintain charge neutrality and related oxygen vacancy is formed [3]. It is discovered that the likelihood of a B3+ particle subbing grid oxygen is incredibly low [4]. Moreover, B3+ replacement makes the framework less steady; henceforth, the B at the substitutional positions moves to the interstitial and makes oxygen vacancy.

13.2.2  Chemical Reduction/Oxidation The different wet chemical methods or solid state redox reactions are used to create intrinsic oxygen vacancies in the metal oxides also known as self-doping of metal oxide, for example, self-doping of TiO2 by post-synthesis treatment with NaBH4 as reducing agents. In the chemical reduction process of TiO2 nanostructure, reductants convert Ti4+ into Ti+3 (oxygen vacancy), which is known as self-doping of TiO2, and it is different than extrinsic doping. Most importantly, self-doping process retains crystal structure or morphology of pristine metal oxide. The high self-doping of Ti+3 in TiO2 can form continuous energy band just below the conduction band (CB) minimum of TiO2. Kang et al. [5] reduced TiO2 nanotube arrays (NTAs) by one-step chemical method. The NaBH4 treatment of TiO2 introduced oxygen vacancies on the surface and bulk of TiO2 (Fig. 13.1). The high reducing ability of NaBH4 converted Ti+4 to Ti+3. The reaction for oxygen vacancy formation is shown by the following reactions:

NaBH 4  8OH   NaBO2  8e   6H 2 O

(13.1)



Ti 4   e   Ti3

(13.2)

The calculated donor densities (oxygen vacancies) in TiO2 before and after NaBH4 treatment are 8.93 × 1017 and 2.06 × 1018 cm−3, respectively.

Fig. 13.1  Oxygen vacancy representation in TiO2 semiconductor

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Similarly, hydrazine (N2H4) is used for self-doping of Ti3+ (oxygen vacancy) in TiO2 thin film using hydrothermal method [6]. The donor densities of TiO2 was increased from 6.9 × 1018 to 8.54 × 1019 cm−3. Rather than the reduction of metal oxides, an oxidation-based technique is additionally proposed for delivering ­high-­concentration doping of Ti3+ all through the mass and surface of metal oxides, for example, TiO2. Self-doped TiO2−x was synthesized by controlled oxidation of titanium hydride (TiH2) utilizing hydrogen peroxide (H2O2) as an oxidizing operator. TiH2 vanished step by step and was changed to TiO2−x with the prolongation of response time [7].

13.2.3  Electrochemical Reduction In the electrochemical reduction (ECR) process, metal ion accepts external electrons and reduces to low valence state. The charge balance is compensated by oxygen vacancy formation, and cation (C+) from electrolyte solution is intercalated into metal oxides. The mechanism can be summarized as:

xC  MO y  xe   C x MO y  z



(13.3)

where C+ can be Li+, Na+, or H+ and MO are metal oxides such as TiO2, ZnO, WO3, or BiVO4. Zhang et al. [8] employed ECR of TiO2 nanotubes (NTs). The negative potential of −0.4 V vs. RHE was applied to TiO2 electrode in 1 M Na2SO4 for 30 min. In the reduction process, Ti4+ ion captured e− and turned into Ti3+. The pristine gray TiO2 NTs were turned to dark blue after ECR treatment. The Ti3+ content after ECR of TiO2 NTs was increased by 6% examined by XPS. Meekins et al. [9] performed intercalation of H+ and Li+ into TiO2 by employing temporary electrical pulses at the potential of 1 ns, and H2-treated samples showed slower medium decay than pristine TiO2, reflecting the formation of long-lasting e−–h+ pairs. However, visible light pumping (450 nm) of H2-treated TiO2 showed fast initial decay contrary to UV-pumping, followed by almost relaxation to the baseline, which signifies the high-speed combination of e−–h+pair.

13.3.2  Time-Resolved Fluorescence Spectroscopy (TRFS) Fluorescence comprises light excitation of an e− to higher energy level followed by energy loss of excited e− by nonradiative ways, and finally emission of low energy photon. Thus, there are competitive processes of photon emission and non-radiative energy loss. Fluorescence spectroscopy comprises excitation of electrons in the molecules of a certain material with the beam of light and successive emission of light with the low energy photon. The emitted light is directed to a filter and extended to the detector. Based on the measurements of light in the detector, the corresponding change in the atom or molecule is obtained. Steady-state fluorescence spectrum is obtained from the excited molecules using the same source of light, and the graph is obtained with the intensity of emitted photon versus wavelength. A fluorescence emission spectrum is obtained by a single-­wavelength source of excitation, and a graph of emission as a function of wavelength is plotted. The TRFS is highly sensitive to TA spectroscopy, as it is measured in a zero background signal. Wheeler et al. [35] used TRFS technique for studying oxygen vacancies in the metal oxides. In the study, TiO2 NWs were treated with hydrogen, and triple-­ exponential fluorescent decays were obtained as shown in Fig. 13.10. The decay of pristine TiO2 was fitted to three different time constants of 35 ps, 120 ps, and >1 ns. The hydrogen-treated TiO2 showed a very fast decay in case of the fast component with enhancing annealing temperature. However, middle component decays with 200  ps lifetime, while long component showed more than 1  ns lifetime for all annealed TiO2 with 0.4 and 1.3 J cm−2 fluence runs. The relative intensity of fast component decays was increased from 68% to 77% with increasing annealed temperature ranges of 350–550 °C. A significant change in fast decay for hydrogen-­ treated TiO2 compared with bare sample reflects that hydrogen annealing changes quantity as well as the depth of trapping states inside band gap. Medium decay component behaves less fast in hydrogen-treated TiO2 (200 ps) compared with bare TiO2 (120 ps), attributing slow PL decay caused by much long-lasting trap state in the hydrogen treated, similar to deep trapped states [35]. Wheeler et  al. [35] used TRFS as well as TA techniques and found a model (Fig. 13.11) that gives energy states of both hydrogen-treated and bare TiO2. The slow decay seen in TA spectroscopy study of hydrogen-treated TiO2 when excited

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Fig. 13.10 (a) Intensity vs. time (ps) of H:TiO2 NWs along with fitting curves and (b) global fitting results of the PL lifetime of H:TiO2, annealed at 350, 400, 450, and 500  °C, showing the contribution from each lifetime component to the overall integrated spectrum at each wavelength. The H:TiO2 samples required triple-exponential functions to fit their decays. (Reproduced with permission from [35], The American Chemical Society 2011)

with UV radiation was compared with bare TiO2 attributes that hydrogen exposure increases the lifetime of charges, favorably in oxygen deficiency sites. Using the location of PL peak (∼460 nm, or 2.7 eV), single-fluorescent trap site was calculated as ∼0.3 eV just near to conduction band (CB), While oxygen vacancy site was calculated at ∼0.75 eV near to CB [2]. They roughly allocated 30 ps decay time, which was seen in TA spectroscopy study for e− relaxation from CB near to fluorescent confinement sites. After hydrogen exposure, the fast decay became slower to 60 ps, caused by oxygen vacancy sites or change in density of fluorescent confinement sites. The e− present in deeply existed trap sites might be longer lasting com-

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Fig. 13.11  Proposed model for energy states of TiO2 associated with optical properties and dynamics studies. (Reproduced with permission from [35], The American Chemical Society 2011)

pared with the e− present in shallow fluorescent confinement sites of 0.3 eV. When hydrogen-treated TiO2 is excited with visible radiation, the e− jumps from VB to the oxygen vacancy sites. The gross fast decay of charge carriers by visible light excitation is caused by a large multiple of states inside oxygen vacancy sites. No long-­ lived component was seen in the case of UV pumping. The absence of longer component was persistent with no photocurrent in the visible region because of the fast recombination of charge carriers [35].

13.3.3  Soft and Hard X-Ray Spectroscopy The oxygen vacancy-induced enhanced charge carrier dynamics in the metal oxides is easily identified by time-resolved laser spectroscopy techniques. Structural and related properties of the material are also equally important to understand the fundamental role of oxygen deficiencies [36]. Various X-ray probes (soft and hard) are used to obtain information about the surface and bulk of the material. For example, Xue et al. [37] used X-ray absorption near-edge structure (XANES) spectroscopy

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and studied O K-edge of hydrogen-treated TiO2 for the analysis of the electronic structure. They found the same feature in O K-edge spectra of treated and bare TiO2, except an apparent change in absorption intensity that reflected decreases in the transition of O1s in hybrid orbitals. The reason was deficiency of O in the hydrogen-­ treated TiO2 lattice. Also, intensity of t2g peak after hydrogen treatment was improved greater than 4 days that signified oxygen deficiencies was remained stable for consecutive 4 days and more. Similarly, XANES profile of Ti K-edge gave important structural information about the bulk of TiO2. The Ti K-edge profile of bare and hydrogen-treated TiO2 exhibited same absorption edge, with a decrease in absorption intensity for the first 4 days after hydrogenation, which is because of deficiency of O occurred in the TiO2 lattice. Moreover, electron paramagnetic resonance (EPR) study showed that g value did not change for the initial 4  days, as oxygen deficiencies have become stable. It was signified that oxygen deficiencies were formed on the TiO2 surface for the first 4 days [37], and afterwards, the positive shift was seen in absorption edges caused by oxygen vacancies formation in the bulk TiO2. Furthermore, radial structure functions (RSFs) of TiO2 before and after H2 treatment showed the same three coordination distances of 1.5, 2.5, and 3.4A° [37]. The peak shift was seen for hydrogenated TiO2 which is related to distorted octahedral symmetry of TiO2.

13.4  P  hotocatalytic Applications of Oxygen-Deficient Metal Oxide Thin Films Fujishima and Honda found that TiO2 can oxidize water (H2O) into pure oxygen and hydrogen using sunlight in 1972. Nowadays, intense research is focused on the purification of air and water using photocatalytic activities of metal oxide semiconductor materials.

13.4.1  Photocatalytic Water Splitting Schematic of producing hydrogen and oxygen by photocatalytic oxidation of water at the semiconductor electrode surface is shown in Fig. 13.12. Photocatalytic water splitting converts solar energy into more suitable chemical energies in the form of H2 and O2 (Fig. 13.12). Thermodynamics of overall water splitting plays important role in the water oxidation reaction. Theoretically, water splitting has Gibbs free energy of +237.13 kJ mol−1 to oxidize into hydrogen and oxygen molecules, which is not favorable thermodynamically, hence needs some energy to oxidize water, as shown in Eq. (13.6):



1 H 2 O  H 2  O2 DG 0  237.13 kJmol 1 2

(13.6)

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In photocatalytic water splitting, metal oxide (semiconductor) is excited with the radiations higher than band gap of metal oxide. As a result, e− jumps from VB to CB by creating h+ in VB. The e− reached the conduction band brought out the reduction of H+ while h+ oxidize H2O if the bottom of CB is positioned at the higher negative potential side compared with reduction potential needed for the conversion of H+ into H2, i.e., 0 V vs. NHE when pH has zero value. Also, the top of VB should be located at the greater positive potential than oxidation potential needed to converting H2O into oxygen (which is 1.23 V vs. NHE at pH value of zero). Also, reduction potential needed for the conversion of H+ to H2 and oxidation potential for converting H2O to O2 is based on the pH of the solution. Thus, metal oxide with higher band gap than 1.23 eV and band positions fulfilling the requirement of the redox potential of water is necessary for full water splitting into oxygen and hydrogen. Also, increment in STH efficiency is better with lesser band gap of metal oxide that helps to absorb larger visible radiations of the solar spectrum. But the reduction in band gap causes a decrease in an impulse, which is compensation for activation energy hurdle required for redox process. Thus, separation of charge carriers produced by photocatalytic reaction, their fast transfer to active surface sites of metal oxide electrode with less recombination, and photocatalytic reactions between adsorbed ions/molecules and photogenerated charges to give hydrogen and oxygen molecules are the required steps, which decide visible-light-stimulated photocatalytic water oxidation. There are some metal oxides, which are efficiently used for photocatalytic water oxidation like TiO2, WO3, and Fe2O3. However, their intrinsic properties do not support efficient STH conversion. For consideration, TiO2 has required band positions for water splitting, but its higher band gap (~3.2 eV) reduces visible light absorption and consequently lowers efficient STH conversion. On the contrary, a-Fe2O3 shows

Fig. 13.12  Schematic for water oxidation reactions in the photocatalysis process. CB and VB: represent the conduction and valence band [38], respectively. (Reproduced with permission from [38], The Royal Society of Chemistry 2012)

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low band gap (~2.1 eV); however, it has adverse intrinsic properties for carrying out catalytic water oxidation such as low charge carrier lifetime ( 400 °C), which they attributed to the larger particle size due to the Ostwald ripening effects [54]. With larger particle size, it becomes harder for the active materials to be completely utilized. During intercalation, the phase-­ transformation reaction will lead to a core-shell structure, with the formation of new phase taking place layer by layer at the shell. The newly formed shell/phase has low ion diffusion coefficient, shielding the crystal core from being utilize for storage reaction [55, 56]. To ensure full utilization of materials while achieving high crystallinity, nanostructure engineering is adopted. Nanostructured cathode materials will not be discussed in this chapter. Further readings on nanostructured cathode materials and techniques of synthesis can be found elsewhere [57, 58].

20.5  Composite Materials as Cathode for Li-Ion Batteries Few composite materials were extensively studied due to their excellent electrochemical performance, which will be discussed later. Undeniably, synthesizing cathode materials in nanostructures had significantly improved the capacity of the

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materials by improving the utilization rate of the active materials. Other than nanostructured cathode, doped as well as composite materials also demonstrated improved capacity, mostly via manipulating the electrical environment, improving the electrical conductivity of the materials. One of the best examples would be the LiNixCoyMn1−x−yO2 composite materials, which will be further elaborated later. Manipulating the ratio of Co2+ could significantly affect the conductivity of the material. Besides, with Mn4+ acting as the stabilizing agent, the structure is capable of accommodating more Li+ than their pristine counterpart without dissociating the structural integrity. In fact, synergistically mixing different cathode materials does not only affect its capacity but its electrochemical potential as well. The open-circuit voltage of a battery is determined by the chemical potential difference between the cathode (μC) as well as the anode (μA), expressed as [59]: Voc 



 A  C e

(20.8)

where e is the Coulombic electronic charge. Aside from the chemical potential of both cathode and anode, the potential window of the electrolyte (the gap between lowest unoccupied and highest occupied orbital) also played a pivotal role. It is crucial for the μA to lie below the lowest unoccupied molecular orbital (LUMO) and μC to lie above the higher occupied molecular orbital (HOMO) of the electrolyte. In the opposite scenario, a passivation layer will be formed at the electrode-electrolyte interface due to reduction (μA lies above LUMO) or oxidation (μC lies below the HOMO) of the electrolyte, as illustrated in Fig. 20.4. The passivation layer is known either as solid-electrolyte interface (SEI) or cathode-electrolyte interface (CEI), depending on the electrode interface where the passivation layer forms [60]. The passivation layer helps to improve the stability of the electrode materials, preventing

(a)

µA

Eg

HOMO Anode

(b) µA

LUMO

Eg

eVµ

µc

Electrolyte Cathode

(c) µA

LUMO

HOMO Anode

LUMO

Eg

HOMO

µc

Electrolyte Cathode

Anode

eVµ

µc

Electrolyte Cathode

Fig. 20.4 (a) The potential window of the battery extracted from the chemical potential of cathode and anode, respectively. (b, c) Oxidation and reduction of electrolyte when the chemical potential of cathode or anode does not lie within the energy gap of the electrolyte, forming passivating layers

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further loss of active materials during prolong cycling. However, as the thickness of the passivation layers increased, the internal shunt resistance of the battery increased. The formation of passivation layers also parasitically consumes the Li+, further reducing the number of ionic compounds and deteriorating the electrochemical performance of the battery. Other than that, the μC position is limited by the top of the anions p-orbital. Consider the case of transition metal with redox energy close to the top of anion p-orbital, such as the case in Li1−xCoO2 [61], the removal of Li+ during delithiation introduced holes into the Co4+/Co3+ couple, changing the state of the d-orbital symmetry from polaronic to itinerant when 0.5 400 mA h g−1 at 100 mA g−1. The redox mechanism of NiO is similar to that of other conversion-type reactions:

NiO  2 Na   2e   Na 2 O  Ni

(21.6)

The superior stability and low operating voltage of this hybrid nanocomposite delivered interest to put it in real applications.

21.2.8  Cupric Oxide (CuO) Cupric oxide is also an attractive source for battery applications due to high earth abundance, low cost, and high safety [65]. In contrast to the other TMOs, the CuO also suffers low conductivity, large capacity fading, and cyclic poor rate performance issues [65]. Y. Lu et al. synthesized micro-structured CuO/C sphere composite to address these issues [66]. They synthesized CuO/C nanocomposite via an aerosol spray pyrolysis followed by an oxidation process. As shown in Fig. 21.6a,

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Fig. 21.6 (a) Schematics of the assembly procedure of Ni-NiO/PCNs hybrid nanocomposites, (b) powder XRD patterns of as-synthesized Ni-NiO/PCNs hybrid nanocomposites, (c) discharge/ charge, and (d) retention profile of Ni-NiO/PCNs hybrid nanocomposites. (Reprinted from Ref. [64] with the permission from RSC publications)

first the resorcinol formaldehyde resin solution was prepared by polymerizing resorcinol (23.5 g) and formaldehyde (25 mL) at room temperature. The copper nitrate (2.55  g) and resorcinol formaldehyde resin were added in ethanol (275  mL) and then atomized in an atomizer at a 0.18 MPa pressure in an Ar. The drops of aerosol were then delivered in a quartz tube and calcined at 800  °C.  To obtain a CuO/C nanocomposite with a CuO particle dimension of around 10 and 40 nm, the acquired Cu-CuxO/C was then reheated at 260 °C for 3 h in the air. The obtained CuO, 10 nm CuO/C, and 40 nm CuO/C were confirmed by powder XRD pattern (Fig. 21.6b). Further, the obtained nanocomposite of CuO/C was assembled in a coin cell with CuO/C|1 M NaPF6/EC/DMC|Na cell configuration. The charging-discharging profile of these composites was recorded at 50 mA g−1 current densities. Figure 21.6c shows the CuO/C nanocomposite with 10 nm particle size delivered high and stable capacity in comparison with CuO/C nanocomposite with 40  nm particle size. However, the nanocomposite follows the following reaction:

CuO  xNa   xe   Cu II1 x Cu I x O1 x / 2  x / 2 Na 2 O



Cu II1 x Cu I x O1 x / 2   2  2 x  Na    2  2 x  e   Cu 2 O  1  x  Na 2 O



Cu 2 O  2 Na   2e   Na 2 O  2Cu

(21.7)

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A. B. Ikhe

21.2.9  Molybdenum Oxide (MoO3) Among the TMOs apart from the first transition metal series, the molybdenum oxide is the most studied TMOs for battery applications [67]. Among all the above discussed TMOs, molybdenum oxide (MoO3) has received particular attention due to low toxicity and its extraordinary theoretical capacity (1117 mA h g−1) [68, 69]. Conversely, due to the low conductivity and poor cyclic stability of MoO3, it needs to make composite materials with carbon sources. Researchers have addressed these issues by tailoring MoO3/graphene oxide or MoO3/reduced graphene oxide composites for SIBs, but the volumetric capacity loss due to high carbon contain was disappointing [70, 71]. Smartly, to avoid all these issues, Y. Jiang et al. have grown MoO3 nanotube arrays and dually functionalized MoO3 (P-MoO3−x) nanotube arrays on the Mo substrates [72]. As cartooned in Fig. 21.7d, they synthesized MoO3 nanotube arrays by anodization of Mo foil in the 0.4 M NH4F/glycerol: water (9:1) electrolyte at 5 V for 2 h. After anodization, the MoO3 film was washed with ethanol to remove electrolyte and then dried in a nitrogen flow. Moreover, 0.1 mg cm−2 nanotube arrays were grown on the Mo substrate after anodization. Further, to synthesize P-MoO3−x nanotube arrays, the as-synthesized MoO3 nanotube arrays were placed in the sodium hypophosphate (NaH2PO2 · H2O) containing quartz tube furnace, and then the furnace was heated to 300 °C for 60 min. The as-synthesized P-MoO3−x nanotube arrays were with nanotube diameter of 80–90 nm, wall thickness 10 nm, and height ~560 nm established by scanning electron microscopy (SEM) displayed in Fig. 21.7e. A piece of MoO3 or P-MoO3−x nanotube array on Mo foil was directly used as an anode in the coin cell. MoO3 or P-MoO3−x|1 M NaClO4/EC/DMC|Na cell configuration was used for electrochemical studies. Figure  21.7f compares the charging-discharging profiles of MoO3 and P-MoO3−x nanotube arrays delivering the revocable capacity of 434 mA h g−1 and 572 mA h g−1, respectively. The excellent stability of these materials was then confirmed by a retention profile near about 1500 cycles at 2000 mA g−1 in Fig. 21.7g.

21.3  T  ransition Metal Oxides for Non-aqueous Potassium/ Potassium-Ion Batteries KIBs have attracted much more interest as a substitute solution for LIBs due to its low reduction potential and high capacities. The use of potassium metal in the battery undergoes dendrite formation and other side reactions with cathode materials [73]. Therefore, researchers are developing alternative anodes with a high capacity for KIBs [73, 74]. Since the TMOs are already investigated in LIBs and SIBs, the TMOs have developed more interest due to low cost, low operating voltage, high capacities, and non-degradable nature [74, 75]. Only a few of the TMOs have been reported for KIB applications such as TiO2, Co3O4-Fe2O3 nanoparticles, CuO, MoO2/rGO hollow sphere composite, and zinc oxide or germanium dioxides

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Fig. 21.7 (a) Schematics of the aerosol spray pyrolysis gadget and the founding route of microand nanostructured CuO/C spheres, (b) powder XRD patterns of nanostructured CuO/C spheres, (c) charge-discharge profile of 10-nm CuO/C spheres. (d, e) Synthesis schematics and morphology of P-MoO3−x nanotube arrays grown on the Mo foil (inset is the cross-sectional images of nanotube array), (f) discharge-charge, and (g) retention profile of MoO3 and P-MoO3−x nanotube arrays. (Reprinted from Refs. [66] and [72] with the approval from RSC and ACS publications)

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adsorbed on silicon nanocage [76–80]. Apart from these anode materials for KIBs, V2O5 was examined too as a cathode material for Na+ storage [81]. However, only selected single-TMOs are discussed in here due to the chapter’s limitations.

21.3.1  Titanium Oxide (TiO2) For the first time, G.-W. Lee et al. reported Magneli phase Ti6O11 as an anode candidate for KIBs [76]. This Magneli phase was derived from anatase TiO2 through the electrochemical process. They incorporated multivalued carbon nanotubes (CNTs) to minimize low electronics conductivity issues in TMOs. The TiO2/CNT microspherical compound was synthesized by the spray drying method by incorporating 9.4 wt% CNTs. Figure 21.8a shows the high-resolution transmission electron microscopy image (HR-TEM) of the obtained TiO2/CNT composite with clear lattice fringes of ~0.35 nm. The lattice fringes were corresponding to the (101), suggesting the highly crystalline anatase TiO2. However, a coin cell with TiO2/ CNT|0.8 M KPF6/EC/DEC|K cell configurator was assembled for electrochemical and spectroscopic investigations. As shown in Fig.  21.8b, the cell delivered 460 mA h g−1 specific capacity upon the first discharge and then fall to ~150 mA h g−1 from the second cycle onward. The sudden drop in capacity was studied by ex situ XRD in Fig. 21.8c. The phase changes were noticed during the discharging-­charging process; Ti3O5 (18.8°, ICSD no. 010761066) and K3Ti8O17 (31°, ICSD no. 010721699) phases were observed at 0.9 V during the first discharge. The Magneli part Ti6O11 (26.3°, 28.9°, and 30.3°, ICSD no. 010851058) at 0.48 V and additional phases Ti0.5O5 (24.3° and 31.8°, ICSD no. 010716414) and K6Ti2O7 (31.3° and 34.1°, ICSD no. 010791757) were observed at fully discharged state. Upon charging, only the Magneli phase of titanium oxide was retained, indicating in situ conversion of the TiO2 phase to the Magneli Ti6O11 phase. However, there is a tremendous need to develop titanium-based materials to enhance capacities and high electrochemical stabilities in KIBs.

21.3.2  Cobalt Oxide and Iron Oxide (Co3O4-Fe2O3) As discussed earlier, the TMOs from the first transition metal series have high theoretical capacities, but the TMOs suffer low specific capacity and low electronic conductivity in KIBs. I. Sultana et al. prepared mixed phases of single-TMOs as an anode for KIBs [77]. They synthesized Co3O4-Fe2O3/C by the ball milling method to enhance the specific capacity. First, the CoCl2  ·  6H2O, FeCl2  ·  4H2O, LiNO3, LiOH · H2O, and H2O2 with a molar ratio of 0.01:0.01:0.1:0.02:0.05 was ball-milled and then heat-treated at 300 °C for 3 h in an air atmosphere. The solid mass was collected after natural cooling to RT and then washed away with deionized water multiple times to dissolve and separate unreacted parts in the solid. After overnight drying at 100 °C, the hybrid Co3O4-Fe2O3/C composite was prepared by mixing a super p-carbon black and Co3O4-Fe2O3 with the 1:2 ratios and then ball milling at

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Fig. 21.8 (a) HR-TEM images of TiO2/CNT microspherical composites, (b) discharging-­charging profile of TiO2/CNT composite at 50 mA g−1, (c) ex situ XRD schemes of TiO2/CNT composite during charging in the potential window 10 mV to 2.5 V vs. K/K+ at 50 mA g−1. (d) SEM images and (e) powder XRD patterns of Co3O4-Fe2O3 and hybrid Co3O4-Fe2O3/C, (f) discharge-charge, and (g) retention profile of Co3O4-Fe2O3/C composite at 50 mA g−1. (Reproduced from Refs. [76] and [77] with the permission from ACS and RSC publications)

rotation speed 75 rpm for 30 h. The Co3O4 and Fe2O3 phases in composites were confirmed by the SEM images and XRD patterns in Fig. 21.8d, e. As presented in Fig. 21.8d, the particle dimension of Fe2O3 was much bigger than the Co3O4 particles. However, the coin cell was assembled with the Co3O4-Fe2O3|0.7 M KPF6/EC/ DEC|K cell configuration and then cycled at 50 mA h g−1 to obtain a discharging-­ charging profile. As shown in Fig. 21.8f, the Co3O4-Fe2O3 delivered 780 mA h g−1 capacity upon the first discharge, but the capacity fading was observed with the following cycles. With the following possible potassiation/de-potassiation reactions, Ca. 200 mA h g−1 capacity was reserved after 50 cycles at a 50 mA g−1 current density (Fig. 21.8g):

Fe 2 O3  6K   6e   3K 2 O  2 Fe 



Co3 O 4  8K  8e  4K 2 O  3Co

(21.8) (21.9)

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21.3.3  Cupric Oxide (CuO) The conversion-type CuO anode with great theoretical capacity, little cost, and small volume expansion upon charging/discharging has been studied for SIBs and LIBs, but the electrochemical study for potassium storage remains mysterious [36, 65, 66]. Due to the advantage of low volume expansion upon charging-discharging process, K. Cao et al. reinvestigated the CuO as an anode material for KIBs [78]. They synthesized CuO nanoplates (~20 nm thick) by a hydrothermal method. The 70-mL deionized water containing 3  mmol of Cu(NO3)  ·  3H2O was stirred for 10 min at RT, and then 20 mmol of NaOH was added into that solution. Further, the mixture was transferred to an autoclave and heated at 80 °C for 12 h. Finally, the product was collected after multiple-time washing with water to remove NaOH from the precursor. The obtained pure crystalline CuO nanoplates were then confirmed with the powder XRD pattern by comparing with standard CuO (JCPDS no. 5661) in Fig. 21.9a. The thickness and shapes of CuO nanoplates were then measured by atomic force microscopy (AFM). Both spectroscopic investigations revealed the pure phase of CuO nanoplates with a thickness of 20  nm. Further, electrochemical studies were performed with the cell configuration of CuO|1  M KFSI/EC/DEC|K.  Figure  21.9c shows the charging-discharging profile of CuO nanoplates with delivering capacities of 351.9, 284.6, 227.1, 206.8, and 163 mA h g−1 at 200, 400, 800, 1000, and 2000  mA  g−1 with the following conversion-type reactions:

CuO  K  KCuO



KCuO  K  K 2 O  Cu



2Cu  K 2 O  Cu 2 O  2K

(21.10)

The cyclic stability was also confirmed by cycling up to 100 cycles at a current density of 1000 mA h g−1 (Fig. 21.9d). The obtained specific capacities and cyclic stability were higher than those reported in SIBs [65].

21.3.4  Molybdenum Oxide (MoO2) Excellent stability, high theoretical capacity, and cost-effective molybdenum oxides were also studied in LIBs and SIBs [68, 69]. For the very first time, C.-L. Liu et al. investigated MoO2 as an anode candidate for KIBs [79]. They synthesized MoO2/ rGO composite for K+ ion storing with high cyclic performance. First, the graphene oxide (GO) was synthesized by following the Hummers scheme from natural flakes of graphite oxide [82]. Further, as shown in Fig. 21.9e, the MoO2/rGO composite was prepared by ultrasonicating 80  mg of GO in 30  mL of deionized water for 60 min, and then 87.8 mg of ascorbic acid (C6H8O6) and 800 mg ammonium heptamolybdate tetrahydrate (AHM) were added to that exfoliated GO dispersion.

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Fig. 21.9 (a) Powder XRD pattern and (b) AFM image of as-prepared CuO nanoplates, (c) discharge-­charge profile and (d) retention profile of CuO nanoplates at various scan rates, (e) schematics of synthesis procedure of MoO2/rGO composite, (f) HR-TEM images, and (g) retention profile of MoO2/rGO composite at 50 and 500 mA g−1. (Reproduced from Refs. [78] and [79] with the approval from Wiley publications)

The solution was then heated at 180 °C for 15 h in a Teflon-lined stainless steel autoclave. The concrete precursor was collected after multiple-time washing with deionized water and overnight drying at 60 °C. Finally, the MoO2/rGO composite was obtained by calcination of the precursor at 400 °C for 3 h in an Ar. The HR-TEM image of MoO2/rGO in Fig.  21.9f shows the MoO2 particle incorporated in rGO hollow spheres. The electrochemical study of the as-prepared MoO2/rGO composite was performed with MoO2/rGO|0.8 M KPF6/EC/DEC |K coin cell. This composite delivered a specific capacity of more than 300 mA h g−1 at 50 mA h g−1 with excellent cyclic stability (Fig.  21.9f). However, further development in molybdenum oxide as an anode in KIBs is required to reach high specific capacity and cyclic stability before commercializing it.

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21.4  T  ransition Metal Oxides for Other Non-aqueous Multivalent Ion Batteries Secondary ion batteries with traditional monovalent batteries including LIBs, SIBs, and KIBs are not sufficient for the ever-increasing energy demands [20]. Due to low earth abundance of these monovalent elements, the monovalent batteries are costlier and not possible for large-scale productions. Consequently, the scientific interest has been diverted to the high-abundance multivalent metals such as magnesium (Mg), calcium (Ca), zinc (Zn), and aluminum (Al) [83]. Figure  21.10 shows the importance of multivalent batteries over the current commercial battery technologies. The Earth’s crust abundance of multivalent elements is much higher in comparison with that of monovalent metals. The conversion-type TMOs as an anode candidate in multivalent batteries are hardly possible due to the higher reduction potentials of multivalent elements than the monovalent elements (Fig.  21.10). However, several insertion-type TMOs have been investigated so far for these batteries such as TiO2, V2O5, MoO3, and MnO2 [87–89]. The selected TMOs for multivalent batteries are discussed one by one in here.

21.4.1  TMOs in Magnesium Metal Batteries Magnesium metal batteries are a current research trend due to higher volumetric capacities than that of the monovalent metals (Li, Na, K), high magnesium abundance, non-toxic nature, and highly negative standard potential (−2.375 V vs. SHE)

Fig. 21.10  Comparison between standard reduction potential, volumetric and gravimetric capacities, and Earth’s crust profusion of monovalent and multivalent metal anodes proposed for electrochemical storage systems. (Reprinted from Ref. [83] with the permission from Frontiers publications)

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[83–85]. Several cathode materials were studied for Mg2+ storage such as Mo6S8, TiO2, V2O5, MoS2, MoO3, TiS2, MnO2, and MgMn2O4 in magnesium metal batteries [84–89]. In here, commonly studied TiO2, V2O5, and MoO3 TMOs are reinvestigated for Mg2+ storage in Mg/TMOs cell. S. Su et al. studied commercial TiO2 for Mg2+ storage [87]. The primary spherical crystals with 25–30 nm in size (Fig. 21.11a) and pure crystalline phase (Fig. 21.11b) of TiO2 were confirmed with the HR-TEM images and powder XRD pattern. The commercial TiO2 was then assembled in coin cell with the cell configuration of TiO2|0.5  M Mg(BH4)2/LiBH4/Tetraglyme|Mg. Figure 21.11c displays the galvanostatic charging-discharging profile of TiO2/Mg cell in the voltage between 0.5 and 1.7 V at a current density of 33.6 mA g−1. The cell delivered a high specific capacity of 155.8 mA h g−1 with excellent cyclic stability and rate capability. However, thin-film cathodes are more profitable to reach high energy density batteries. Coupling a high capacity, high cyclic stability thin-­ film electrode with magnesium metal will be another milestone in battery technology. G.  Gershinsky et  al. synthesized extremely pure V2O5 and MoO3 thin-film electrodes in nanoscale fatness for Mg2+ storage [88]. First, V2O5 films were vacuum deposited on the Pt foil with the film thickness of 200 nm and then heated at 415 °C

Fig. 21.11 (a) HR-TEM images and (b) powder XRD patterns of commercial TiO2. (c) Galvanostatic discharge/charge shape of TiO2/Mg cell at 33.6 mA g−1. (d) HR-SEM image of the thin-film V2O5 electrode; (e) XRD patterns of pristine, magnesiated, and demagnesiated thin-film V2O5 electrodes; and (f) cyclic voltammogram of Mg2+ intercalation process into the thin-film V2O5 electrode at 0.1 mV s−1. (g) HR-SEM image of the thin-film MoO3 electrode; (h) XRD patterns of pristine, magnesiated, and demagnesiated thin-film MoO3 electrodes; and (i) cyclic voltammogram of Mg2+ intercalation process into the thin-film MoO3 electrode at 0.1 mV s−1. (Reprinted from Refs. [87] and [88] with the permission from RSC and ACS publications)

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for 4.5 h in an argon atmosphere. Figure 21.11c shows the morphological view of V2O5 thin-film deposited on the Pt foil. The 20–30-nm size particles were equally deposited on the Pt surface. The magnesiation and demagnesiation in V2O5 thin film were confirmed by the potentiodynamic techniques with the V2O5|0.1 M Mg(TFSI)2/ ACN|Mg cell configuration, and it was further confirmed with the powder XRD pattern (Fig. 21.11d, e). The cell provides a specific capacity of 150 mA h g−1 with excellent stability up to 50 cycles. Further, MoO3 films were electrodeposited on the Pt foil throughout the 50 mM molybdic acid aqueous solution, and then the thin films with 100 nm thickness were obtained by heating it at 350 °C for 3 h. The uniform deposition of MoO3 thin films on the Pt surface was confirmed by HR-SEM images in Fig. 21.11g. The MoO3 thin films with the film thickness ~100 nm were then assembled in cell configuration of MoO3|0.1  M Mg(TFSI)2/ACN|Mg. The magnesiation/demagnesiation in MoO3 was confirmed by the cyclic voltammetry technique and powder XRD pattern in Fig. 21.11h, i. MoO3 thin films delivered a specific capacity of 220 mA h g−1 with superior cyclic stability. The use of V2O5 and MoO3 thin-film electrodes in the magnesium battery gives more attention to commercializing it in the future.

21.4.2  TMOs in Calcium Metal Batteries As mentioned earlier, the research trend diverted to multivalent batteries due to the aforementioned advantages. The calcium batteries could be the alternative solution for recent LIBs due to its great earth abundance and relatively low reduction potential (−2.9 V), but this technology is still in the early stage to conclude [90–92]. The TMOs as an anode for CIBs have not been reported so far due to the early research stage, but some TMOs have been studied for Ca2+ ion storage for cathodes such as α-MoO3 and α-V2O5 [93–95]. M.  Cabello et  al. did electrochemical and spectrochemical investigations on calcium intercalation and deintercalation in the layeredtype α-MoO3 material [93]. Figure  21.12a displays the galvanostatic charging-­discharging profile of α-MoO3|0.5 M Ca(TFSI)2/1,2-dimethoxyethane|Ca cell at a current density of 10 mA g−1. Ca. 0.25 Ca2+ was reversibly intercalated in the layered α-MoO3. The HR-X-ray photoelectron spectroscopy (HR-XPS) in Fig.  21.12b revealed the step-by-step calcination in the layered MoO3 structure. However, with the detailed investigations of Fig. 21.12a–c, the redox reaction of α-MoO3 was revealed. The α-MoO3 follows the Mo6+/Mo4+ redox reaction during the charging and discharging process in the α-MoO3/Ca cell.

MoO3  Ca 2   2e   CaMoO3

(21.11)

The commonly known layered V2O5 was also studied with the detailed investigation of Ca2+ intercalation/deintercalation [95]. R. Verrelli et al. did electrochemical investigations on calcium storage in the orthorhombic α-V2O5. The Ca/α-V2O5 cells were assembled in the several calcium-containing electrolytes such as 0.3  M

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Fig. 21.12 (a) Galvanostatic discharging/charging profile of Ca/MoO3 at 10 mA g−1, (b) detailed view of the XPS in the Mo3d at different discharged and charged states, and (c) coordination of calcium in the layered CaMoO3 structure. (d) Galvanostatic discharge profile of Ca/V2O5 cells in 0.3 M Ca(TFSI)2/EC/PC (red line), 0.3 M Ca(ClO4)2/EC/PC (purple line), 1 M Ca(ClO4)2/ACN (green line), and 0.45 M Ca(BF4)2/EC/PC (yellow line) at 100 °C and (e) corresponding ex situ XRD patterns at room temperature. (f) Coordination of calcium in the layered CaV2O5 structure. (Reprinted from Refs. [93] and [95] with permission from ACS and Elsevier publications)

Ca(TFSI)2/EC/PC, 0.3  M Ca(ClO4)2/EC/PC 1  M Ca(ClO4)2/ACN, and 0.45  M Ca(BF4)2/EC/PC.  The Ca/α-V2O5 cells were failed at room temperature and 50 °C. So, the galvanostatic discharge profile of Ca/α-V2O5 cells was obtained in the aforementioned electrolytes that were cycled at 100 °C (Fig. 21.12d). Corresponding ex situ XRD pattern of the α-V2O5 at an altered depth of discharge (DOD) (at the point of a–i mentioned in Fig. 21.12d) recorded in Fig. 21.12e revealed that more calcium in α-V2O5 was intercalated throughout the 0.45  M Ca(BF4)2/EC/PC electrolytes. The coordination of calcium in layered α-V2O5 was then summarized in Fig. 21.12f. However, the use of TMOs in calcium batteries is still in an early stage to conclude it for commercializing standards. There is a tremendous need to develop electrolytes and compatible TMOs for calcium batteries.

21.4.3  TMOs in Zinc Metal Batteries The non-aqueous Zn battery chemistry was overlooked so far due to its lower theoretical energy densities, low potential window, and high reduction potential (−0.76 V vs. SHE and 2.3 V vs. Li/Li+) [96, 97]. Recently, S.-D. Han et al. explored the non-aqueous electrolytes for Zn ion plating/stripping on Zn metal to provide a

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Fig. 21.13 (a) Pulse electrodeposition of BL-V2O5 on the CFS and (b) corresponding SEM image (enlarged in the inset), (c) galvanostatic discharge/charge profile of electrodeposited BL-V2O5 at a current density of 14.4 mA g−1, (d) SEM images of floret-like δ-MnO2 (yellow arrow) carbon black particles (red arrow), (e) HAADF-STEM images of δ-MnO2, and (f) Discharge/charge profile of δ-MnO2/Zn cell at 12.3 mA g−1. (Reprinted from Refs. [99] and [100] with approval from Wiley and ACS publications)

wider electrochemical frame up to 3.8 V vs. Zn/Zn2+ [98]. This new approach of these electrolytes opened up doors for high-voltage multivalent rechargeable batteries. Inspiring with these electrolytes, P. Senguttuvan et al. proposed bilayered (BL) V2O5 for high non-aqueous Zn batteries [99]. The BL-V2O5 cathode material for the Zn battery was prepared via electrodeposition on the carbon foam substrate (CFS). First, the pulse electrodeposition on CFS was carried out in a three-electrode open-­ cell shown in Fig. 21.13a with Ag/AgCl as a reference and Pt mesh as a counter electrode in 0.2 M VOSO4/H2O solution at a 50 °C. The BL-V2O5 cathodes were then obtained by vacuum annealing at 120 °C for 20 h. The pulse electrodeposited BL-V2O5 on CFS was then confirmed by SEM images in Fig. 21.13b. Finally, the obtained electrode was directly flocked in a coin cell with the cell configuration of BL-V2O5/CFS|0.5  M Zn(TFSI)2/ACN|Zn. Figure  21.13c shows the galvanostatic charging-discharging profile of BL-V2O5/Zn cell with delivering the specific capacity 170 mA h g−1 after 100 cycles at a scan rate of 14.4 mA g−1. For the first time, such an excellent electrostability and high specific capacity for TMO in a non-­ aqueous zinc battery were reported. The Zn-ion intercalation chemistry which commonly used MnO2 was also studied by S.-D.  Han et  al. [100]. They synthesized nanostructured δ-MnO2 powder by dissolving 2.5 mmol potassium permanganate and 1.0 mmol manganese sulfate monohydrate into the 30 mL of milli-Q-water and refluxing for 2 h in air atmosphere. The obtained compound was washed with milli-­ Q-­water to eliminate side products mentioned in Eq. (21.12) and then dried at RT:

2KMnO 4  2 MnSO 4  H 2 O  4 MnO x  K 2 SO 4  H 2 SO 4

(21.12)

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The floret-like layered δ-MnO2 microstructure was confirmed by SEM images in Fig. 21.13d. The width of floret-like δ-MnO2 units was in the range of 0.2–0.5 μm. High-angle annular dark-field (HAADF) STEM images in Fig. 21.13e show the thickness of δ-MnO2 flakes was in the range of 5–15 (~7 Å each interlayer) stacked (001) planes. Finally, the galvanostatic charge-discharge of δ-MnO2 has been performed in the δ-MnO2|0.5  M Zn(TFSI)2/ACN|Zn cell configurations. The floret like δ-MnO2 delivered a specific capacity of 123 mA h g−1 corresponds to the 0.2 mol of Zn2+ ions insertion per mol of MnO2 (Fig. 21.13f). However, the capacity dying was detected after 30 cycles in Fig. 21.13f; this capacity dying was mainly attributed to the electrolyte decomposition or structural changes by K and Mn dissolution or loss of the electronics contact of δ-MnO2. Therefore, the pulse electrodeposited BL-V2O5 was the only TMO candidate so far to achieve the basic requirements in electrochemistry of zinc battery such as high specific capacity, high operating voltage, and high cyclic stability.

21.4.4  TMOs in Aluminum Metal Batteries Aluminum, the most abundant element on the Earth with a low reduction potential, is also attracted for future battery technologies to inexpensive large-scale productions [83]. So far, major electrochemical studies on the aluminum batteries were in the aqueous electrolytes, but it limits the potential window for making high-voltage battery [101]. Therefore, many groups studied and invented non-aqueous electrolytes, which opens the doors for developing high voltage aluminum batteries [102]. However, TMOs also play an important role in developing high-voltage aluminum batteries. Therefore, frequently used TMOs for multivalent ion insertions (Mg2+, Ca2+, and Zn2+) such as V2O5 and MoO2 were reinvestigated for Al3+ storage [103– 105]. S. Gu et al. synthesized V2O5 nanowires and reinvestigated Al3+ insertion/deinsertion by following the previous report [103, 105]. The V2O5 nanowires were prepared by hydrothermal method; 0.364 g of V2O5 and 5 mL of H2O2 were added to 30 mL of deionized water and then stirred for 30 min. The as-prepared mixture was transferred into a stainless steel reactor, and then the reactor was heated at 205 °C for 4  days. Further, the precipitate was collected, washed with deionized water, dehydrated at 60 °C under vacuum, and then reheated at 500 °C for 4 h in air atmosphere. Figure 21.14a shows the XRD pattern and FE-SEM images of as-­prepared V2O5 compounds; it was found the pure crystalline and nanowires of V2O5. The V2O5 nanowires were assembled in a coin cell with the cell configuration of V2O5|AlCl3/ [BMIM]Cl (1.1:1)|Al. In Fig. 21.14b, the galvanostatic charge-­discharge profile was recorded in between 0 and 2.5 V potential window delivering the specific capacity 55 mA h g−1. The structural changes during phase changes upon aluminum insertion/deinsertion were observed (Fig. 21.14c). However, the structural changes are not a problem here, because it was reversible upon charging and discharging. The specific capacity of V2O5 in the aluminum battery was much lower than the other aforementioned multivalent ion batteries. So, the structural modification of V2O5 is mandatory to further enhance the specific capacity. J.  Wei et  al. also studied

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Fig. 21.14 (a) Powder XRD pattern and SEM image (in inset) of as-prepared V2O5 nanowires, (b) galvanostatic discharge-charge profile of V2O5 nanowires, and (c) ex situ XRD patterns of V2O5 nanowires at different discharged and charged conditions during the first cycle, (d) XRD pattern, and (e) EDX spectrum of MoO2 sputtered and heat-treated nickel foam, (f) galvanostatic discharge-­ charge profile of MoO2@Ni, at a current density of 100 mA g−1. (Reprinted from Refs. [104] and [105] with approval from Elsevier publications)

aluminum chemistry in MoO2 [105]. They deposited MoO2 (ultrasonically cleaned in acetone, ethanol, and water) on the nickel foam to overcome a capacity fading problem due to the loss of particle contact. The deposition was implemented by radio frequency magnetron sputtering system at 3.0 W cm−2 power for 60 min, and then the MoO2 deposited foam was calcinated at 400 °C in a tubular furnace. The pure crystalline MoO2 deposits (~1.91 mg cm−2) were confirmed by XRD study and elemental mapping (Fig. 21.14e, f). Finally, the galvanostatic charge-discharge profile was recoded using a three-electrode Swagelok cell with the configuration of MoO2@Ni|AlCl3/EMImCl|Al (MO) (Fig. 21.14f). The specific capacity at the discharge state was about 90 mA h g−1 at a current density of 100 mA g−1 with the discharge plateau at ~1.9 V vs. Mo. MoO2 was the first TMO which achieved high voltage and high specific capacity in a non-aqueous aluminum battery. The high-­ voltage plateau and high energy density TMOs as a cathode material along with specially designed non-aqueous electrolyte develop an interest for commercializing point of view. However, there is countless room for the optimization of TMOs for aluminum batteries, making them an outstanding energy storage device.

21.5  Summary and Perspectives Remarkable achievements and progress have been made on TMOs in the post-LIBs including monovalent and multivalent batteries. The TMOs are considered important in batteries due to their safety, earth abundance, and inexpensiveness and potentially can be used in the large-scale productions. TMOs, mostly from the first

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transition metal series, used in post-LIBs are reviewed in this chapter. Most of them were synthesized to obtain a thin-film electrode for battery applications. All the metal oxides were characterized by spectroscopic methods and then elevated for battery performance. Some metal oxides showed extremely auspicious results in the KIBs and SIBs, but the same was not gained in multivalent batteries. Even though the TMOs electrodes in multivalent batteries are not so promising in comparison with multivalent ion batteries, but unquestionably the multivalent batteries will lead in the coming decade with these TMOs. Further, in-depth investigations of the electrochemical mechanisms and structural properties are mandatory to enhance the voltage and specific capacities in the future battery technologies. When the TMOs derive to practical applications in the batteries, few factors need to be considered such as power density, power density, energy efficiency, cycle stability, charging rate, safety, and cost. However, it is believed that the TMOs from the first transition metal series will play an important role in future battery technologies. Acknowledgment  The author gratefully acknowledges the financial supports and research facilities obtained from the Department of Printed Electronics Engineering under the Suncheon National University, Republic of Korea.

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Chapter 22

Nanostructured Metal Oxide-Based Electrode Materials for Ultracapacitors Chukwujekwu Augustine Okaro, Onyeka Stanislaus Okwundu, Philips Chidubem Tagbo, Cyril Oluchukwu Ugwuoke, Sabastine Ezugwu, and Fabian I. Ezema

22.1  Introduction The increasing growth of the world’s population and urbanization and the unfavorable impacts of the burning of fossil fuels have triggered the need for the exploration of sustainable, accessible, and environmentally friendly energy sources. The energy sources of free gifts of nature such as the sun and wind have drawn the interest of researchers as possible alternatives that are clean/green (pollution-free) and capable of meeting the ever-growing global energy demand. The dependence of these energy sources on time and weather constrains their full utilization and hence the need for the development of associated energy conversion and storage devices

C. A. Okaro · O. S. Okwundu · P. C. Tagbo · C. O. Ugwuoke Science and Engineering Unit, Nigerian Young Researchers Academy (NYRA), Onitsha, Anambra, Nigeria Department of Physics and Astronomy, University of Nigeria, Nsukka, Enugu, Nigeria S. Ezugwu Department of Physics and Astronomy, University of Western Ontario, London, Canada F. I. Ezema (*) Science and Engineering Unit, Nigerian Young Researchers Academy (NYRA), Onitsha, Anambra, Nigeria Nanosciences African Network (NANOAFNET), iThemba LABS, National Research Foundation, Somerset West, Western Cape, South Africa UNESCO-UNISA Africa Chair in Nanosciences/Nanotechnology, College of Graduate Studies, University of South Africa (UNISA), Pretoria, South Africa e-mail: [email protected]; [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_22

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that can store reasonable amounts of energy during peak periods and effectively deliver such as during low or off-periods [1–3]. Electrochemical energy storage and conversion devices such as batteries, ultracapacitors, and fuel cells have shown promising traits as sustainable, accessible, pollution-free, and highly efficient energy backup systems [4–6]. Of all, ultracapacitors, a.k.a. supercapacitor (SCs), have attracted more attention due to their high power density, long life cycles, fast charge/discharge rate, wide operating temperature ranges, and portability [7, 8]. As a result, the research community has and continues to witness progressive growth in SC-based researches, as evidenced by the increasing number of publications on SC, as demonstrated by Okwundu et al. [6]. Its popularity in the research world is followed by its vast application in many sectors of today’s technology, such as automobile, consumer electronics, power grids, and military [9–12]. The origin of today’s SC can be traced back to the first ever-known capacitor and then referred to as “the Leyden jar,”, which was developed in the year 1745 by Pieter van Musschenbroek [6]. The content of the Leyden jar is comparable to that of the modern capacitor, for instance, the Leyden jar contains a glass with its interior and exterior coated with metal foils which act as conductive plates and the glass jar that serves as the dielectric [7]. Since then, capacitor-based research has witnessed the increasing growth. The first SC used a carbon-based electrode and was reported in 1957 by Howard Becker in his patent granted by general electric [7]. Not until when the Standard Oil of Ohio (SOHIO) patented an electric double layer capacitor and the licensing by Nippon Electric Company (NEC) [7], SC was not commercialized in the market. Afterward, many research works have been focused on developing SCs with higher efficiency and new areas of application. However, the applications of the traditional SC are limited by the low energy density it offers, which is mainly as a result of the use of carbon-based electrode materials. In a quest to improve the energy density of SCs, the Conway’s group established the idea of pseudocapacitance in 1991 by using a metal oxide (RuO2) as the electrode material, which showed to have higher energy density and lower equivalent series resistance (ESR) [13]. This concept of pseudocapacitance is different from the carbon-based capacitor (EDLCs) due to the charge storage mechanism. The charge storage mechanism of EDLCs is more of charge accumulation on the electrode-electrolyte interface (non-­ faradaic), while pseudocapacitors combine the non-faradaic mechanism of EDLCs and a fast redox reaction [14]. The high capacitive performance of pseudocapacitors is attributed to the charge storage mechanism. Redox-active materials such as metal oxides (MOx), metal nitrides, and conducting polymers support pseudocapacitive storage mechanism, but the most widely used is MOx [15]. The present success of SCs in vast areas of applications is largely due to the utilization of MOx as the electrode material in one or both electrodes of SCs [16]. The suitability of MOx as electrode materials in pseudocapacitors is due to their ability to exist in multiple oxidation states [15]. Other desirable properties of MOx as pseudocapacitive electrode materials are low-cost, non-toxicity, availability, chemical stability, and workable morphology [17, 18]. The capacitive performance of pseudocapacitors is relatively dependent on the specific surface area (SSA) of the electrode material due to the storage mechanism of pseudocapacitors (surface redox

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reaction). By simply manipulating the structure and morphology of MOx, the capacitive performance of MOx has been shown to improve [19]. Furthermore, the utilization of nanoporous MOx as pseudocapacitive electrode materials have shown desirable improvement in the capacitive performance of SCs [20]; this is because the porous structure creates higher SSA, stimulates better infiltration of the electrode, and enhances more utilization of active sites by the ions [18]. Although SCs are endowed with promising properties such as high power density, long cycle life, fast charge/discharge rate, and ease in maintenance, their moderate energy density limits their prevalent utilization in everyday storage devices, using renewable and pollution-free energy sources [15]. To encourage the vast application of SCs, their energy density/charge storage capacity must be significantly improved. Finding precise strategies for improving the charge storage capacity of SCs without altering their high power density and long life cycle has been a trend in the research world [5]. One of the best strategies is the development of pseudocapacitive MOx electrodes for SC [15]. These MOx electrodes are cheap and have improved charge storage capacity, and hence commercialization can easily be achieved. In this chapter, the utilization of nanoporous MOx as pseudocapacitive electrode materials was discussed. The arrangement of the chapter is as follows: in Sect. 22.3, the components of SC is discussed, and the desirable properties of the electrode, electrolyte, separator, and sealant materials and their functions are also covered; Sect. 22.4 gives an overview of the fundamentals of pseudocapacitance; in Sect. 22.5, the performances of some electrode materials were summarized; Sect. 22.6 gives the applications of SC and lastly; in Sect. 22.7, the outlook and summary of nanoporous MOx supercapacitors were discussed.

22.2  Components of Supercapacitor A fully industrial packaged SC is made up of two electrodes, an electrolyte, two current collectors, and a sealant. Materials used for such components must possess certain properties before considered suitable. In this section, some properties of the electrode (with/without binding material), electrolytes, and current collector materials are considered.

22.2.1  Electrode The SC is composed of two electrodes: one is the positive electrode (cathode), and the other the negative electrode (anode). These two electrodes can be of the same material (symmetric SC) or different materials (asymmetric SC). However, certain properties are desirable of electrode materials for the maximum functioning of SC; this is because the electrode material used affects the overall performance of SCs.

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These properties are high conductivity, chemical stability, thermal stability, high specific surface area (SSA), and non-toxicity. Good conductivity is a must-have for any electrode material because the electrodes act as the pathway to the delivery of electrons during charge/discharge processes. A highly conductive electrode leads to higher power delivery, while a poor conductive electrode negatively affects SC performances [21]. Considering the effects of the second property on the performance of SCs, the life cycle of SCs is greatly influenced by the chemical and thermal stability of the electrode materials. This instability of the electrode material is primarily caused by dissolution, phase change, and the side reaction of the electrode during cycling [22]. The use of chemically inert material and compatible electrolyte gives room for a much more charge and discharge cycles of the SCs. The SSA of an electrode is basically the capacitive block of the SC since the charge storage of SCs mostly happens on the surface. Increasing the SSA leads to a corresponding increase in the capacitive performance of SCs [19]. The porosity of the electrode’s surface contributes to the capacitive performance of SCs, as a porous electrode leads to more utilization of the electrolyte ions [13]. Figure 22.1a shows the various components of SC with a porous electrode and SEM images for a clear view of the porosity. Carbonaceous materials, conducting polymers MOx, are examples of electrode materials for SC applications. Carbon-based materials exhibit double-layer capacitance (which will be discussed in Sect. 22.4), while conducting polymers and MOx are used in pseudocapacitors. The suitability of carbon-based materials for double-­ layer application is due to their inert nature, large surface area, high porosity, low cost, high conductivity, thermal stability, and environmental friendliness [24], but they offer lower energy density compared to the redox-active materials (conducting polymers and MOx) used in pseudocapacitance. Conducting polymers (CPs) show electrode properties comparable to that of MOx, but they are easily subjected to degradation and swelling during charging/discharging which in turn affects their

Fig. 22.1  Electrochemical supercapacitor: a schematic diagram showing the major components with illustrative SEM images, demonstrating the porous nature of SC electrodes—adaptation for high charge storage (a, b) an image of an industrially finished SC with cylindrical design [23]

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life cycle [21]. This undesired phase change and the short life cycle of CP act as the limiting factor to the commercialization of CP-based SCs [5]. MOx-based electrode materials exhibit higher energy density than carbon electrodes and longer lifecycle than CP-based electrodes. Aside from that, MOx electrode materials are readily available, chemically and thermally stable, inexpensive, and non-toxic [24]. The storage mechanism of MOx (as would be discussed in Sect. 22.4) is a surface phenomenon, and hence increasing the SSA leads to higher capacitive performance. The use of nanoporous MOx is a popular strategy for tackling the issue of the low specific surface area [25]. Most times, electrodes are deposited on the current collectors, but in places where deposition is not feasible, a binder is used to fasten the electrodes on the current collectors. Though binders help in maintaining the structural integrity and in the achievement of good adhesion between the electrode and current collectors, the effects of the content and property of binding materials have been observed to affect the performance of SC devices [2, 5, 26]. Effects such as lowering conductivity, reduced active surface, and wettability of the electrode have led researchers into developing SC devices that are binder-free [26]. Most commonly used binder materials are polyvinylidenediflouride (PVDF), Nafion and polytetraflouroethylene (PTFE) [27]. (a) A Brief on Electric Conduction in MOx. Generally, MOx have wide and varied electrical properties – some are insulative, metallic, and superconducting, while others are semiconducting. By manipulating the activity of oxygen (its partial pressure) in the gas phase during MOx synthesis, various non-stoichiometric MOx can be produced [28]. Electrical conductivities of MOx are strongly dependent on the non-stoichiometry in their compositions [29]. Upon generalizing that non-stoichiometric MOx usually have oxygen deficit with the corresponding reduction in the oxidation state of the predominant metal, Fog and Buck [30] noted that MOx stoichiometry can be altered slightly by heating to generate a lower valence oxide, or even the metal itself, in the absence of lower valence oxide. Doping with foreign reactive metal is effective as well. Although MOx materials could be electrically conductive, their conductivity in SC application cannot be compared to that of carbon [17, 18].

22.2.2  Electrolyte Another component that is considered as vital as the electrode is the electrolyte. The overall performance of SC is also reliant on the electrolyte used. The term electrolyte refers to the electrolyte salt and the solvent. The role of the electrolyte in SC performance ranges from provision to balancing of charges on the SC electrodes. Therefore the electrolyte should exhibit some desirable features to be considered suitable for SC application. These features are high conductivity, electrochemical and thermal stability, wide voltage window, high ionic concentration, low flammability, low tox-

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icity, low volatility, and corrosion resistance. A very important feature of a suitable electrolyte for SC application is the voltage window of the electrolyte, as the energy density of SC is directly proportional to the square of working potential. Hence SC made of electrolyte with a larger potential window and having other suitable operating conditions will exhibit higher energy density. Therefore, it is desirable to utilize electrolytes of higher potential window. Aside from the effect of low potential window of electrolytes on the energy density of SC, exceeding the potential window of electrolytes subjects the electrolytes to decomposition and instability which can in turn affect the overall performance of the SC [24]. For SC, the principal resistance is ESR (equivalent series resistance), which is composed of the electrolyte resistance and contact resistance between current collectors and electrodes [7]. Higher ESR value negatively affects the power delivery of SC devices. The electrolyte component of ESR is as a result of ionic conductivity (presence of free charge carriers), thermal instability, ion mobility, and viscosity of the electrolyte. A highly conductive electrolyte should exhibit high ionic mobility and low viscosity and hence a reduced ESR value of the SC [4]. The stability of electrolytes (both thermal and electrochemical) plays a key role in the performance of SC devices. The thermal stability of the electrolyte contributes to the ESR value of the SC, long life cycle, and the safe operation of ES. Electrochemical stability is related to the life cycle, larger potential window, and safe operations [4, 21]. SC electrolytes are classified as solid and liquid electrolytes. Solid-state electrolytes serve as an electrolyte and a separator material, while liquid electrolytes require a separator. Solid-state electrolytes are of two types: the polymer-based and ceramic-based electrolytes. The liquid-state electrolytes comprise of aqueous, organic electrolytes, and ionic liquids.

22.2.3  Current Collectors Though considered as a passive component of ES, the importance of current collectors in SCs cannot be neglected. The major role of the current collector is the transport of the stored charges in the electrode material to an external load. Hence, electric conductivity is a vital property of the current collectors for power delivery. Other properties such as electrochemical stability, compatibility, and thermal stability of current collectors affect SC performance characteristics (such as operating cell voltage, reliability, and lifespan) [26]. As mentioned earlier, the contact resistance between the current collectors and the electrodes contributes to the ESR value of the SC; therefore incompatibility between the electrode and current collector materials increases the ESR of the SC. Normally found within the cell, the compatibility of the current collector with the electrolyte is also a factor to consider when choosing collector materials. The mostly used collector materials are metals and metal alloys depending on the abovementioned factors (compatibility with both electrode and electrolytes). As a way of reducing the thickness of the electrode-to-­ collector interface, electrode materials or pastes have been deposited or grown onto current collectors [21].

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22.2.4  Separator The separator is a non-conductive membrane that comes in between the positive and negative electrodes, thereby preventing contact and electron transfer and allowing the movement of ions. For the excellent function of the separator, separator materials must be of good mechanical strength, ionic conductor, highly electrochemically stable, highly porous, small thickness, and sufficient surface wettability [2, 24]. The mechanical strength of the separator improves the durability of SC, while high porosity increases ionic conductivity, thereby reducing the internal resistance of the separator. The most commonly used separator materials are glass, mica, polyolefins, ceramic, cellulose paper, and so on, though their application is reliant on the choice of electrolyte and the temperature of operation [31].

22.2.5  Sealant Electrolyte degradation and surface oxidation of electrodes can easily occur in SC when exposed to contaminants such as air and water, thereby reducing the life cycle of the ES and affecting the smooth performance of the SC. A sealant helps in the prevention of the contamination of the SC by these impurities. Apart from that, proper sealing of the ES inhibits shunt resistance between nearby electrodes and cells especially in commercial applications where multiple SCs are linked together. The effect of shunt resistance SC performance is the creation of alternative current passage [32]. Polymeric materials are mostly utilized as sealant materials due to their flexibility, moisture resistance, mechanical stability, and electrical resistivity [5, 7]. The overall performance of SC is greatly influenced by the materials used in the buildup. Selecting the suitable electrode, electrolyte, separator, current collectors, and the sealant materials results in an SC of climax performance. From above, nanoporous MOx-based electrode materials have been deemed the best due to their unique properties such as higher SSA, workable morphology, and short path distance for the diffusing ions and electrons [24]. Furthermore, the compatibility of the electrolyte with the electrode and other components of the SCs is also a vital factor to consider in material selection, as the life cycle and the smooth performance of the SC relatively depend on that. Figure 22.1b shows the cylindrical form of an industrially finished SC, having all the components intact. SCs having miniaturized sizes and of different designs exist already in the market.

22.3  Fundamentals of Supercapacitance SC is an electrochemical energy storage device that is known for its fast-charging rate and long life cycles. Charging basically means the addition of charges, while the opposite is discharging. To charge an SC, a voltage source is connected to the

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positive and negative terminals of the SC (Fig. 22.1a has its cathode deposited on the yellow current collector and anode on the green): the positive terminal of the voltage source to the positive electrode (cathode), while the negative terminal to the negative electrode (anode). By this connection, electrons are drawn from the cathode to the anode, thereby making the cathode more electron deficit and the anode more electron-rich. This causes the drift/migration of the electrolyte ions; anions (negatively charged ions) will move toward the cathode and the cations toward the anode. So long as the electrolyte has abundant ions, the movement of charges would continue until the potential between the two electrodes equals that of the voltage source. When this happens, removing the voltage source leaves the system at that state. To discharge the SC, a load is connected to the terminals (in lieu of the voltage source), and the reverse process takes place. That is, electrons get drawn from the anode to the cathode through the load, and the electrolyte returns to its equilibrium state. Now, what happens when the electrolyte ions get to their respective electrode destinations (during the charging process)? In other words, how do electrochemical SCs store charges? Well, in this section, the storage mechanisms of SCs are discussed. The mechanisms are broadly classified into two: the electric double-layer capacitance (ELDC) and pseudocapacitance. It is based on these mechanisms that SCs are generally classified as electrostatic SCs or electric double-layer capacitors (EDLCs) and pseudocapacitors. While we consider EDLC in Sect. 22.4.1, we discuss pseudocapacitance among other redox processes in Sect. 22.4.2, in order to shun ambiguity.

22.3.1  Electric Double-Layer Capacitors (EDLCs) Electric double-layer capacitors are capacitors that store charges through the formation of double layers on the electrode-electrolyte interface. Double-layer formation is a phenomenon that occurs when an electrically active material (charged material) comes in contact with an electrolyte of a certain concentration of ions. During the charging process of SC, the cathode gains more positive charges, thereby creating an electric field which then attracts the same number of negative charges around the cathode. The structure formed on the cathode-electrolyte interface due to this charge balance is the electric double layer [21]. In a quest to retain the neutral state of the system, an equal number of positive charges of the electrolyte accumulate on the anode, hence the formation of another double layer. The existence of these two double layers (one on the cathode-electrolyte interface and the other on the anode-­ electrolyte interface) is the core of the capacitive performance of the EDLCs [21]. In order to explain the structure and what happens during the formation of a double layer, some theories were put forward. Hermann von Helmholtz was the first person to propose a theory that explains the double-layer formation in 1853. In his model, the charges of the electrode are counterbalanced by the formation of a rigid structure of opposite charges at a distance d away from the electrode. According to him, electrolytes of higher concentrations of ions have their excess charges stacked

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within the region of this rigid structure [33]. This layer is referred to as the Helmholtz layer. This layer is comparable to the traditional parallel plate capacitor where the solvent (water) molecule serves as the dielectric material [34], and the interlayer distance as the diameter of the molecule as shown in Fig.  22.3a. To explain the inadequacies of the Helmholtz model, the Gouy-Chapman model was put forward by Gorges Gouy (1910) and David L.  Chapman (1913) [33]. According to the Gouy-Chapman model, the ions compensating the charges of the electrode do not form a rigid structure as proposed by Helmholtz; rather they tend to diffuse into the electrolyte phase. This diffusion will continue until the opposing potential setup due to their exit constrains the possibility of exit of any further ion [35]. The exit of the ions toward the electrolyte is as a result of the electrostatic force and the thermal motion [33]. This layer (including the electrode surface as shown in Fig.  22.1b) formed due to the diffusion of the ions is called the diffuse layer. The Helmholtz and the Gouy-Chapman models were combined by Otto Stern in 1924. He encouraged the coexistence of both the Helmholtz and the diffuse layers [36]. In this model, the size of the ion was accounted to influence their approach toward the electrode ­surface, which was not considered in the Gouy-Chapman model [35]. According to Stern, these two layers are comparable to two capacitors in a series arrangement. Grahame modified the Stern model by considering the possibility of specific adsorption of ions occurring on the electrode-electrolyte interface. Grahame proposed that the Helmholtz layer is divided into the inner Helmholtz plane (IHP) and the outer Helmholtz plane (OHP). The IHP is the closest approach of the ions to the electrode

Fig. 22.3  Mechanisms of various electrochemical processes: (a) electrostatic charge storage for EDLC, (b) redox pseudocapacitance at the electrode surface or at a thin (nano) region into the bulk of electrode, (c) intercalation pseudocapacitance within the bulk of electrode (usually in-between crystal layers), and (d) underpotential deposition of metallic atoms (electrosorption)

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Fig. 22.2  Electric double-layer models: (a) Helmholtz model, (b) Gouy-Chapman model, and (c) Stern-Grahame model

surface and is made of solvent molecules and/or adsorbed ions at a distance dIHP away from the surface, while the OHP is the region immediately after the IHP at a distance, dOHP. In this model the capacitance of the Helmholtz layer is given as the series capacitance of the inner and outer Helmholtz planes [21]. Figure  22.2 offers the visualization of the above-explained models. From the models above, a satisfactory understanding of the double-layer formation and the capacitive behavior of EDLCs can be extracted. A look at Fig. 22.3a gives a practical view of the double-layer formation on the electrode-electrolyte interface. Altogether, the electrostatic or non-faradaic charge storage mechanism of SCs (EDLC) is not different from the charge accumulation on either plate of a conventional dielectric capacitor. The “super” nature of EDLC’s capacitance comes from the relatively large surface area of the electrodes, usually in the order of 103 m2/g [6], courtesy of advances in materials synthesis, and, secondly, the extremely small interlayer distance between the double layers (appreciable in the quantum realm).

22.3.2  Redox Processes Electrochemistry as a field of science aims at studying the structure and whatever goes on at the interface between an electrode (conductor of electrons) and an electrolyte (conductor of ions) or at an electrolyte-electrolyte interface [37, 38]. A host of such study has shown that redox-inert materials such as conducting carbon serve as pure EDLC electrodes, delivering non-faradaic current. It is worth noting that all

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SC electrodes either partially or fully deploy the EDLC mechanism [39]. Exempting the redox-inert materials where EDLC fully governs charge storage, there is always charge transfer across the electrode-electrolyte interface (across the double layer), leading to reduction and oxidation (redox) processes. SC electrodes that also deploy faradaic charge transfer/redox processes are said to be pseudocapacitive in nature [6]. The prefix “pseudo”, could either mean fake/false or approaching/almost; hence, the nomenclature, “pseudocapacitors” suggest that they are almost like capacitors; the resemblance would be appreciated as the section unfolds. With MOx materials, capacitive faradaic processes are encountered in both surface redox pseudocapacitance and intercalation pseudocapacitance. In addition, we would further consider underpotential deposition and the battery-type redox process, in order to clarify the standpoint of pseudocapacitance with MOx materials. (a) Surface Redox Pseudocapacitance In addition to mere surface adsorption/desorption of ions (EDLC), MOx electrode materials like RuO2 and MnO2 encourage charge transfer across the double layer with concomitant fast and reversible redox reaction at the electrode surface or at a thin (nano) region into the electrode as demonstrated in Fig. 22.3b. Such kinetically and thermodynamically favored faradaic process is limited to the surface and is never diffusion-controlled, hence the name surface-redox pseudocapacitance. Ions from the electrolyte get adsorbed on the electrode surface (crossing the double layer), and electron from the collecting plates flows through the bulk of the electrode to the surface region where it is consumed in the course of redox reaction or vice versa (the reaction could lead to the generation of electrons, which follows a reverse pathway), as illustrated in Fig. 22.3b. Equation (22.1) gives the general form of the redox reaction [39]. The specific cases of MnO2and RuO2 are given in Eqs. (22.2) and (22.3) [15]:

MOx  yA   ye   A y MOx



MnO 2  H   e   MnOOH



RuO 2  yH   ye   RuO 2  y   OH  y

(22.1)



(22.2)

(22.3)

Following from the general form Eq. (22.1), A+ can as well be a group 1 metal ion such as Li+, Na+, K+,etc. Since their redox processes occur at/near the surface and happens reversibly very fast, without alterations to crystal structure, redox pseudocapacitive charging/discharging occurs very fast with cycle life and power density akin to capacitors, but with greater energy density, due to the deployment of both faradaic and non-­ faradaic means. (b) Intercalation Pseudocapacitance Beyond the surface redox capacitance where electrolyte ions are limited to the surface, by employing electrode materials with inherent crystal-layered structure,

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the ions can migrate through the interlayer channels into the bulk of the electrode where they embark on similar fast and reversible reactions as described in the previous section. This mechanism is illustrated in Fig. 22.3c. Two points must be noted here: first, for the bulk faradaic process to be classified as pseudocapacitance, the diffusion process must be very fast such that the overall mechanism is not diffusion-­ controlled (or limited by diffusion), and, second, the intercalation (ion insertion) and de-intercalation must not result in crystal phase alteration [39, 40]. Because ions get inserted into the bulk of the electrode via tunnels or interlayer channels of the crystalline material, this mechanism is termed intercalation pseudocapacitance and is quite different from the battery-type mechanism [39–41]. V2O3 is a good for this mechanism [42]. (c) Underpotential Deposition While redox pseudocapacitance and intercalation pseudocapacitance are the widely recognized faradaic processes accountable for the charge storage in SCs [39, 43], the first of the three mechanisms is identified by Conway in the 1970s underpotential deposition [40]. This is similar to the adsorption of solvated ions encountered in EDLC, but for underpotential deposition, atoms (rather than ions) are deposited on electrode substrate, forming an atomic monolayer of a material different from the substrate material when a high potential is applied. The “underpotential” nomenclature is a full misnomer, because the electrosorption occurs at relatively high potentials than the Nernst potential [38]. Equation 22.4 describes the redox reaction that occurs during the electrosorption process, while Fig.  22.3d illustrates the mechanism:

S  An   ne   SA

(22.4)

where A could be a proton or metallic ion and the substrate (S) could be any conducting material (commonly metals) [21, 44, 45]. Underpotential deposition studies with MOx substrates [44, 45] show that deposition is not achievable, suggesting that the electrosorption process could not account for the pseudocapacitive behavior of MOx materials. Just recently, Ragoisha et al. [46] report the electrodeposition of metal adlayers on other metallic chalcogenides to be electrochemically irreversible—incapable of explaining their pseudocapacitance. As a result, underpotential deposition is not a plausible mechanism for pseudocapacitive MOx materials. (d) Battery-Type Redox Process When there is crystallographic phase altering intercalation of ions, followed by bulk faradaic reaction, the mechanism is no longer capacitive. Such redox processes occur in batteries and are diffusion-controlled. As a result, their charge/discharge rate is sluggish (relative to SCs), and even though they can store a relatively high amount of energy per kg, they a slow to deliver or acquire it; hence, their low power density. Batteries also have a relatively shorter life cycle than SCs due to the crystallographic phase wear and tear resulting from repeated intercalation and de-­ intercalation upon several charging-discharging cycles [6].

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22.3.3  A  ssessing the Electrochemical Mechanism of a Working (Active) Electrode Having described the capacitive (charge storage mechanisms), alongside the battery-­ type process, let us consider the practical means of electrochemically identifying various electrode types, based on their charge storage mechanism(s). Among the identification, methods are cyclic voltammetry. It provides data (both qualitatively and quantitatively) on the electrochemical activities taking place on the active electrode materials. It works by applying a known potential to the working electrode, relative to the reference one, with linear voltage sweeps within the specified potential range (determined by considering the electrolyte’s operating window). The scan provides data on electric current as a function of time [7]. A plot of the current versus potential is called cyclic voltammogram, and that serves as a tool for diagnosing electrochemical mechanisms. Figure 22.4a shows the characteristic cyclic voltammograms for electrode materials classified based on the various charge storage mechanisms. Rectangular voltammogram (I) theoretically represents an ideal capacitor, and the appearance of (redox) peaks is an indication that the mechanism involves faradaic means. From the figure, the resemblance of EDLC with conventional dielectric/parallel plate (ideal) capacitor is eminent. Deviations of the voltammograms from the ideal capacitive rectangular shape offer (to a reasonable extent) a means of qualitatively assessing their non-capacitive behavior. Surface redox pseudocapacitance (III) results in negligible peaks, while broad peaks come with the intercalation type (IV). Obviously, batteries (V) are far from capacitive behavior, while the pseudocapacitors fall between the SCs and batteries [7, 47]. The voltammogram of battery-type electrodes always reveals at least one anodic peak and one cathodic peak, well separated from each other along the potential axis—an indication of the irreversible nature of the battery-type redox process [48].

Fig. 22.4  Electrochemical identities: cyclic voltammograms (a) and galvanostatic chargedischarge plots (b) for the various energy storage mechanisms; where: I- parallel plate capacitor (an ideal electrostatic mechanism), II- EDLC, III- redox pseudocapacitance, IV- intercalation pseudocapacitance, V- battery-type bulk faradaic mechanism

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Another tool used for assessing the electrochemical fingerprint of electrode materials is the galvanostatic charge-discharge (GCD) technique. This works in a manner reverse to that of cyclic voltammetry. In the GCD technique, a fixed current density is applied to the active (working) electrode in the test cell, while the time-­ dependent potential response data is collected. The GCD plot is simply a plot of the response potential versus time [7]:



V dV I or   constant t dt C

(22.5)

Furthermore, the general shape of GCD plots for electrode materials based on some considered charge storage mechanisms is shown in Fig. 22.4b. Noting that capacitance (C), based on the manner of operation, is a measure of the quantity of charge (Q) stored per unit potential (V) across a pair of charge-stacked layers, C = Q/V and that electric current (I) is the rate of flow of Q per unit time (t), Q  =  It; hence, C = It/V. A capacitive material is therefore expected to conform to Eq. (22.5), for the constant current GCD analysis. In other words, the GCD plot of a supercapacitive material is expected to follow a sloping (ideally straight, or almost straight) line. From Fig.  22.4b, the charging and discharging processes with EDLC electrodes yield linear GCD plots (GCD profile II), while those of pseudocapacitors are significantly curved (GCD profiles III and IV) but still sloping, as a result of the faradaic reactions involved in their storage mechanisms. However, due to the diffusion-­ controlled redox processes encountered with battery-electrode materials, their charge-discharge profiles (GCD profile V) are far from the (straight and curvy) sloping profiles obtainable with EDLC and pseudocapacitive materials. Battery-type processes are easy to identify, for they portray apparent plateaus (constant voltage) in their galvanostatic charge-discharge profiles [7, 39, 47]. Nowadays, with advancements in nanotechnology, the bulk faradaic process in batteries, which is kinetically limited by ion-diffusion, is more often subdued by synthesizing the same battery-type electrode materials in such a manner that they are nanostructured (as nanoparticles), thereby reducing the bulk diffusion path for electrolyte ions. By this, several battery-type electrode materials have been tuned to exhibit pseudocapacitance. Such pseudocapacitive behavior which relies on ­material’s size is referred to as extrinsic pseudocapacitance [39, 40, 43]. A host of articles have stressed on the inappropriateness of using pseudocapacitance for battery-type materials [39, 47]. By cyclic voltammetry, Wang et al. [49] quantitatively established a relationship between pseudocapacitance and particle size, in the range of 5 to 10  nm, pseudocapacitance, as well as the total amount of charge, stored increased, appreciated with size reduction. In addition, a means of quantitatively estimating the contributions of both surface capacitive and diffusion-controlled faradaic processes (from the data obtained from cyclic voltammetry) is known [49]. But first, surface capacitive processes have b-values of 1 and the diffusion-controlled processes, values of 0.5, as described in Eq. 22.6. The equation expresses the response current (i, from cyclic voltammetry), in terms of voltage sweep rate (ν), with a and b, as constants. By combining the

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capacitive ν and diffusion-controlled √ν, Eq. (22.7) emanated with constants k1 and k2, indicating the respective contributions [39, 47, 49]:

i  a b i V   k1  k2 

(22.6)

(22.7)

22.4  Electrode Preparation Techniques The overall performance of the electrochemical supercapacitor is a function of the design and integration of functional electrode materials. Consequently, the electrode material’s selection and fabrication will be an important consideration in designing energy storage devices. The synthetic reproducibility and excellent conductivity in an electrolyte have marked nanostructured MOx a unique architecture of materials for supercapacitor applications [50, 51]. Hence, it would be of great necessity to converse on the trending inexpensive fabrication techniques that will ease the production route of the unified structure of nanomaterials with a high degree of morphological control for an SC application. Figure 22.5 shows the categories of deposition (or coating) techniques that have been used for the successful synthesis of MOx nanostructures. In this section, we briefly summarize the non-­ templated solution-phase growth approaches for the production of various nanostructures and further discuss on MOF-templated approach for the controlled preparation of nanoporous MOx architectures.

22.4.1  P  reparation of MOx Nanostructures Using Liquid-­Based Techniques Many techniques for the successful synthesis of MOx nanostructures have been introduced in the past years. These techniques can be classified as physical approaches and chemical approaches based on their various experimental constituents and setups. The liquid-phase deposition method also known as wet chemistry route has gained significant attention due to their various advantages over other methods which include its applicability in low temperature and pressure conditions, minimized energy consumption, inexpensive production cost, productiveness in MOx production, simple experimental setups, etc. [52–55]. Unlike the liquid-phase approach, the physical approach requires expensive vacuum equipment, higher temperature, longer reaction time, and often non-environmental-friendly constituents. We categorize the liquid-phase approach considered in this context into hydrothermal and electrochemical deposition, and aqueous solution deposition techniques. Figure 22.6 is the diagrammatic representation of various low-temperature solution-­

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Fig. 22.5  Categories of various deposition techniques for the production of nanostructured MOx

based techniques frequently used in the production of MOx nanostructures. Among the above mentioned techniques, the aqueous solution growth is widely employed for the synthesis of MOx nanostructures by various routes such as chemical bath deposition, dip-coating, spin-coating, successive ionic layer adsorption and reaction, liquid atomic layer deposition, etc. [57–60]. Hence, we briefly consider these methods based on their various procedures, factors that tailor the morphological features of the as-synthesized products, and their top advantages. 22.4.1.1  Hydrothermal This is a widely used liquid-phase approach employed in the growth and ­crystallization of nanoparticles from a heterogeneous solution under very high pressure for a very long time in an enclosed autoclave [61]. This technique makes use of a universal solvent as the medium of operation in the deposition process. A strong metallic reactor of a high thickness generally known as an autoclave is filled with the solution maintained at a temperature difference between the pressurized ends of the particle’s growth chamber. One end within the chamber is characterized by an elevated temperature to enable the disintegration of the solvent, while the opposing end that is relatively cooler is for the growth of the nanoparticles. The autoclave reactor is designed such that the interior walls are highly resistant to the corrosion that may be induced by some corrosive chemicals used in the process, thus capable of accommodating the high heat and extreme pressure of the system. The protective wall which is usually gold, silver, or platinum coatings is usually made after proper consideration of the type of solution and the amount of heat to which the system

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Fig. 22.6  Schematics of various liquid-phase techniques commonly used for deposition of MOx nanostructures: (a) Langmuir-Blodgett films, (b) aqueous chemical grow, (c) dip-coating, (d) spin-­ coatings, (e) spray pyrolysis, (f) inkjet printing, (g) electrodeposition, (h) liquid atomic layer deposition (Reproduced from [56] with permission from Wiley and Sons. Copyright 2015 Wiley-VCH Verlag GmbH & Co. KGaA)

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could be subjected to [62]. This process partially depends on the solubility of the solution under high vapor-pressured solvent at elevated temperature levels. Other factors to be considered while optimizing the morphology and crystallography of any functional materials under the hydrothermal process are the reaction duration, pH of the solution, type of solution (precursor), surfactants/catalyst, and solvent used. 22.4.1.2  Advantages of Hydrothermal Synthesis Over Other Methods The critical study of this method reveals its potentialities in comparison to other techniques and is summarized as follows: • Ability to control the content and composition of the materials through multiphase reactions during synthesis. • Ability to deposit particles with high vapor pressure during heat treatment. • The dynamic nature of the method makes it possible to incorporate other physicochemical processes during product optimization. • Highly controllable particle sizes and a high degree of crystallization are easily achievable. • Ability to optimize the disparity of the particles without necessarily undergoing the calcination process. • Highly crystalline-nanostructured materials that are not stable at high temperatures could easily be produced using this method. • Its adoption by industries for large-scale production gives it an edge over other alike. 22.4.1.3  Electrochemical Deposition This is a unique deposition process, whereby a substrate is necessarily a conductive material to allow the simple electrolysis of the cell solution to be accomplished either by oxidation or reduction process and bring about coating onto the substrate by application of a mild voltage between the electrodes [63, 64]. Here, the applied electric field is the agent that derives the ionic particles from the counter electrode (anodic terminal) through an electrolyte solution to the working electrode (cathodic terminal or substrate) where the deposition takes place. The term electrochemical deposition is often used interchangeably with electrophoretic deposition because both approaches usually take place within the electrochemical cell; however, swift chemical bonding process occurs in the former approach unlike the latter approach [65]. This is an ideal but one of the oldest methods employed for a smooth fabrication of nanostructured MOx materials. With this technique, the particle sizes and crystalline structures could be perfectly controlled at the electrode/substrate’s surface without the use of surfactants rather by simply altering the conditions of the electrochemical process [66]. Factors that could influence the morphological features of the synthetic

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product are the deposition time, the potential difference between the cell’s electrodes, the pH of the electrolyte, the molar concentration of the solution, and the constituents of the electrolyte. Hence, this process is categorized as the class of low-temperature deposition method and is an isothermal technique with lots of advantages [67]. 22.4.1.4  Advantages of Electrochemical Deposition The top advantages of electrochemical deposition are as follows: • Its ability to deposit uniformly aligned nanoparticles on substrates with various orientations. • It can easily be utilized in the laboratory for small-scale materials fabrication because the experimental equipment is not scarce and is less expensive. • The easy-to-control electrical parameters make it possible to control the film thickness, structures, stoichiometry, doping, etc. of the deposited particles by controlling the thermochemical parameters. • It is highly essential for the synthesis of heterojunctions by varying the cell electrolyte. • The deposition in this approach occurs almost at equilibrium in contrast to other methods at elevated temperatures, and this method can also be used for the deposition of compositionally nonequilibrium materials. • It establishes a wide range of industrial applications ranging from electroplating through materials purification to exploitation of various materials properties. • It does not require chemicals that are harmful or lethal to human health, thus making it one of the most secure approaches for the fabrication of MOx nanostructures. • For the fact that electrochemical activities happen only at the cathodic terminal (i.e., at the substrate site), cell solution is efficiently utilized, thus reducing the risk of chemical wastage.

22.4.1.5  Aqueous Solution Deposition This technique can also be referred to as aqueous chemical growth and has largely been applied in the growth process of MOx-nanostructured materials through various routes such as dip-coating, spin-coating, conventional drying/calcination, chemical bath methods, etc. The forerunner of this approach is the sol-gel route that involves the use of various molecular precursors, normally metallic salt and alkoxides. Precursor in this context is simply prepared by the chemical combination of metallic reagents, a suitable solvent, and any other chemical that may act as a catalyst to bring about as desired nanostructured architecture. In this approach, the precursor solution undergoes smooth and controlled hydrolysis and poly-condensation reactions creating an interconnected polymeric arrangement with metal ion forming its center and oxygen group forming the side linkages all inside the sample solution

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Fig. 22.7  Figurative representation of steps involved in aqueous solution deposition: (a–d) schematic illustration of the four-step processes and (e–f) emblematic representation of the process (a–d) at the molecular level (Adapted from [53]. Pubs. RSC. under the CC BY 3.0 license)

at mild temperatures [53]. After the breakdown of the organic network after deposition, well-oriented nanocrystalline MOx is obtained by post-deposition annealing at relatively low temperature. Figure 22.7 shows the models of step-by-step processes involved in this synthetic approach with the corresponding development in molecular structures of the material at every stage. After obtaining a precursor solution at the first step, the deposition of the as-prepared solution on a suitable substrate will form a wet layer on the substrate. In the second step, the concretion of the gelatinous wet layer will begin to occur due to the fast evaporation of the solvent in the wet layer, and this compaction is caused by the viscoelastic flow of the deposited substance at room temperature. The breakdown of metal precursors and the subsequent evaporation of the remaining solvent and other unstable substances still within the system require some heat treatment at mild temperatures which forms the basis of the third step. The low-temperature annealing process is very essential in that it will favor the condensation process and expedite the evolution of metal-oxygen networks in an amorphous state. At the fourth step, there is a need for further mild pyrolysis or calcination of the amorphous state product to shorten the diffusion distance between molecules, thereby alleviating an arrangement of the metal-­ oxygen networks to form lattice-like structures of MOx. Various MOx nanostructures such as CuO, Co3O4, ZnO, etc. have successfully been synthesized using this approach [68]. The uniqueness of this one-step route in the preparation of MOx nanostructures lies in exploiting the solvent that can readily

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provide oxygen or hydroxy groups to bring about gradual condensation reaction which will in return support the controlled crystallization of MOx nanostructures. These agents could be alkoxy groups, alkanols, ethers, carboxylic acids, and ketones, other than water. Some other factors that must not be neglected while employing this method in the preparation of nanostructures include substrate treatment, proper stirring during precursor preparation, and proper heat treatment throughout the synthetic process. 22.4.1.6  Advantages of Aqueous Solution-Based Deposition • It is comparatively cheaper among other various low-temperature deposition techniques. • The architecture of the final product is easily engineered in situ of the precursor solution, thus as a result of the chemical composition. • Requires less complicated equipment or manufacturing system for small- and large-scale production. • The stability and morphological features of the final product can easily be engineered by altering the stoichiometry of the solution when the precursor is being synthesized. • Large- and small-area depositions are easily accessible using various substrates like metal planes, glass, semiconductor materials, plastics, bare papers, grapheme, metal-organic frameworks, cotton textiles, etc. • The desired film thickness is easily achieved by cycling the deposition process with subsequent pyrolysis. • It is very easy, cheap, clean, and reliable to carry out the deposition process both on a commercial scale and on a laboratory scale.

22.4.2  N  anoporous MOx from Metal-Organic Frameworks (MOFs) The discovery of metal-organic frameworks (MOFs) in the late 1990s by Yaghi and Li has brought about the latest advancement in the fabrication of highly porous MOFbased nanostructures such as MOx, metal chalcogenides, bimetallic composites, etc. [69]. MOFs are multi-template crystalline materials having a porous structure with a tunable surface area. MOFs have widely been employed as a versatile precursor in the manufacture of various nanoporous materials for many applications which include energy storage and conversion devices, sensors, etc. [70–73]. The use of nanoporous MOx materials has been at the cutting edge of energy material technology for its application as electrode materials in electrochemical energy devices due to their huge potentials to accommodate steady electrochemical processes (e.g., faradaic redox reactions) [72]. Their use in supercapacitor applications has effectively increased the device capacitance due to the increased surface area induced by porosity.

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Functional nanoporous materials for energy storage devices have been successfully prepared using template-free techniques of controlled precipitation approach [74]; hydrothermal [75], solvothermal [76, 77], electrochemical depositions [78– 85]; etc., and a series of MOx such as cobalt oxide (Co3O4) with mesoporous morphology have also been synthesized through “facile binary-solution route following thermal decomposition at normal atmospheric pressure” [86]. However, the stable structures and optimized performances of the final products remain challenging. It was shown experimentally during performance evaluation and other electrochemical analyses that the capacitance of these architectures is relatively limited when compared to MOF-derived counterpart for supercapacitor applications [87]. Nanoporous aluminum oxide (Al2O3) has been synthesized through spray pyrolysis of its precursor in a salt solution; however, solid product was only observed at the temperature greater than 900 K which by consequence imposes more challenges in the material syntheses. The use of a self-templated approach that uses physicochemical processes to convert a temperate if its material to various porous architectures has been used in the past to synthesize an active hollow structure of MOx for energy storage applications. However, the practical performance and the structural stability of the produced material are still not optimized, and this may lead to the breakdown of the structures or phase changes after many cycles. Thus, there is a need for an enhanced method of fabricating highly stable nanoporous structures of MOx with controllable phase, surface area, and dimensions which will consequently influence its performance in supercapacitor application. MOF answers the question “Is the stable structure of nanoporous MOx for supercapacitor application achievable?” Method  MOFs are made by a strong chemical combination of metal ion (or inorganic cluster) and the organic linkers to become a rigid structure. MOFs can be used to prepare nanostructures in two distinct routes: either by direct pyrolysis under an inert condition of the atmosphere at higher temperature or by two-step pyrolysis which involves first treating in an inert gas (argon or nitrogen) and lastly treating in the air at a relatively lower temperature [73]. Figure 22.8 is a clear pictorial explanation

Fig. 22.8  Conventional representation of processes involved in the fabrication of highly porous and large surface area MOx nanostructures from MOFs composing of metal ions and organic ligands as the precursors. (Reprinted with permission from [72]. Coprright©2017, American Chemical Society)

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of how a highly porous MOx nanostructure is evolved from a sacrificial template (MOFs) under controlled pyrolysis in nitrogen gas and air, respectively. The chemistry of this process is not complicated. As the MOFs undergo controlled thermal annealing under an oxygen-rich atmosphere, in the air; for instance, the total breakdown of the rigid organic linkers will occur so that the metal ions remain with the pores. These remaining metal nodes are further oxidized to form a highly porous MOx structure. Thus, the resulting material is termed MOF-derived porous MOx. MOF-derived MOx nanostructures normally replicate the structure and porosity and establish the same surface area as that of the parent frameworks under controlled heat treatment conditions. Factors such as annealing temperature levels and average time of the treatment can be used to alter the structural composition, pore sizes, surface area, and the molecular orientation of the final structures. Figure 22.9

Metal oxides 600 Co-BTC MOF-71

MIL-125 (Ti)

Co-MOF

400

[Co3L2(TPT)2×G]n ZIF 67 ZIF 67 FE-MIL-88B

ZnO

ZIF-8 MOF-5

TiO2

Mn2O3

PB

Ni-BTC Ni-MOF

MgO Mg-aph

PB

Fe2O3

CeO2 Ce-BTC

Co-V-MOF

Co3V2O8

PBA

Co3O4

0

MOF-199

100

ZIF 67

200

Mn-MIL-100 Mn-BTC

NiO

300

CuO

Surface area (m2 g–1)

500

Fig. 22.9  Bar chart illustrating the recent progress in the preparation of MOF-derived MOx porous structures from the in situ precursors with their corresponding surface area of the obtain product; each bar showing various precursors used in the fabrication of the corresponding MOx at the base of the row. (Reprinted with permission from [72]. Copyright©2017, American Chemical Society)

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shows various MOF templates used in the production of porous MOx with their corresponding surface area. In the same way, the MOF templates possess highly tunable metal clusters and organic linkers, as well as highly controllable porosity and surface area, so they can be applied as effective templates or precursors for the construction of nanoporous MOx with highly tailorable compositions, structures, and properties to suit industrial needs [73]. That is to say that the major functional properties of as-prepared MOF-based nanostructures could be traced back to the original MOF precursor. Thus, MOFs precursor could be referred to as a sacrificial template or material for the fabrication of the highly tunable nanoporous structure.

22.5  Performances of Metal Oxide Supercapacitor Electrode The use of nanoporous MOx such as TiO2, WO3, NiO, ZnO, RuO2, ZnCo2O4, NiMoO4, MnFe2O4, etc. as electrode material for SCs is favorable in achieving excellent electrochemical performance and as well in providing an accessible route for insertion of electrolyte ions [88]. RuO2 has unique features that made it to be considered as the best among other electrode materials used in SCs. These features include long cycle stability, high capacitance, good reversibility, a wide potential window, and excellent proton conductivity. Despite these unique features, its commercialization is limited by high cost and toxicity. Other single MOx such as MnO2, TiO2, WO3, NiO, ZnO, etc. serve as alternative electrode materials due to low cost, eco-friendly, variable oxidation states, and chemical stability. Although they show good features, they are observed to have low electrical conductivity resulting in low specific capacitance. The low specific capacitance is curbed by tailoring them into various morphologies (such as nanospheres, nanorods, microspheres, nanofibers, core-shell structure, etc.), doping with various atoms, and incorporating them with other transition MOx to achieve a composite [89]. Table 22.1 shows the synthesis and performances of various electrode materials (single MOx, binary MOx, the composite, and hybrid) for SCs. It is obvious that the ZnO electrode has a very low specific capacitance of 5.87 F/g and as well MnO2 has low specific capacitance of 311.52 F/g, but if both are incorporated together (i.e., ZnO@MnO2), a high capacitance up to 600F/g or above is achieved. Apart from single MOx electrode, binary MOx such as NiCo2O4, MnCo2O4, ZnCo2O4, NiMoO4, MnFe2O4, etc. are used as electrode material for SC. Among them, NiCo2O4 and MnFe2O4 are the most studied binary MOx electrode for SCs. They offer higher specific capacitance compared to the single MOx; this is due to their high electrical conductivity [96]. Binary MOx SC electrodes offer a high specific capacitance up to 1000 F/g, and when a composite with single MOx is formed a higher capacitance is achieved; for example, ZnO@CoFe2O4 nanocomposite proves to be one of the best electrodes showing a whooping capacitance of 4050.4 F/g but offers low energy density. The capacitance and energy density of MOx or composite electrodes can as well be improved by incorporating it with carbon material or carbon derivate to form a hybrid electrode. The inclusion of carbon material or its derivate (e.g., graphene) in

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MOx enhances the ease of movement of electrolyte into the electrode due to the porous nature of the graphene sheet [112]. In return, MOx are stocked in between the layers of graphene preventing it from stacking back. From the table, Co3O4/ MnO2@GO hybrid electrode offers an improved capacitance and energy density as 1718 F/g and 420.91 Whkg−1, respectively. As described in Sect. 22.2, an SC is composed of a separator, two electrodes, and an electrolyte. Both the electrode and electrolyte affect the performance of an SC. Studies have shown that the major drawback of SC is low energy density. The energy density of an SC can be improved close to that of a battery without interfering with power density by fabricating an SC having MOx or its composite as a positive electrode and carbon as the negative electrode [113]. This configuration of SC is known as an asymmetric supercapacitor (ASC). The high energy density is produced by the faradaic process of positive MOx electrode, whereas the carbon electrode maintains the power density. An ASC with SiC@NiCo2O4/Ni(OH)2//SiC@ Fe2O3 electrodes offers an improved energy density of approximately 103 Whkg−1 and maintains the power density of 3.5 KWkg−1 when charged for just 6.5 s. On the other hand, the choice of the concentration of electrolyte also affects the performance of an SC. Various electrolytes (such as ionic organic, aqueous, and liquids) have been employed in the fabrication of SCs. Among them, an aqueous electrolyte is mostly used due to high capacitance and high electrical conductivity. Ivanova et al. [114] investigated the performance of ASC in KOH, K2SO4, and KOH + LiOH aqueous electrolytes and found that an ASC with mixed electrolyte shows higher discharge capacitance compared to those with KOH and K2SO4 due to the high electrical conductivity of KOH and the addition of Li+ facilitate the redox reaction. Generally, while selecting electrode materials for SC, some factors are to be put into consideration. They include the conductivity, porous nature, surface area of the electrode, and the wettability, resistance, and concentration of electrolytes.

22.6  Applications of Supercapacitor Most electrical and electronic devices (e.g., electric vehicle, electric rail transit system, mobile device, micro-grid, chemi-resistive senor, and many more) have relied on electrochemical energy storage (EES) system such as batteries and SCs as energy buffer. SC with exceptional features has stood out as the best among others. When incorporated into those devices, it shows a high-efficient energy buffer.

22.6.1  Electric Vehicle (EV) The interest of people has been increased toward the use of EVs due to their reduction in the emission of greenhouse gasses. The most common EVs are pure electric vehicles (PEVs) and hybrid electric vehicles (HEVs). PEVs require an electric motor for its movement, whereas HEVs require both internal combustion (IC)

Annealing in air and Core/shell H2 atmosphere nanoheterostructures Electrochemical Core-shell nanocables depositions

H–Ni/NiO

ZnO@MnO2

SILAR

Spherical middle-like particles Nano-grains

–

–

–

–

116.4

1500

–

5000

5000

4000

Cycle life 1000

0.5 M Na2SO4

1 M Na2SO4 1 M Na2SO4 1 M KOH

–

–

–

–

5000

1200

1000

5000

40.6

–

1000

170.6

3 M KOH. 37.061 3000

1 M Na2SO4 1 M Na2SO4 6 M KOH

2 M KOH

3 M KOH

SA/ Electrolyte m2g−1 3 M KOH –

Mesoporous 2 M KOH network-like Octahedra1 nanosheets 6 M KOH

Porous Cluster Nanospheres

Mn3O4

Ni(OH)2/Cu2O/ HT-assisted redox CuO reaction Ultrasonic spray NiCo2O4 pyrolysis Solvothermal and NiCo2O4 annealing Facile solvothermal Cu2O@ Ni(OH)2 reaction Co3O4-ZnO HT and annealing

Flower-like

Nanoparticles/ nanoflakes Nanoflakes

Electrodeposition

HT

Morphology Nanocomposite/ microspheres Nanoneedles

Synthesis HT

Co3O4/MnO2@ HT GO CBD Co3O4

Electrode material ZnO@ CoFe2O4 NiCo2O4@ MnO2 Co(OH)2/RuO2

Table 22.1  The performances of various nanoporous MOx electrode materials

613.5

717

786.2

833

892

931

1007.4

1474

1567

1718

2168

2353

S.Cap /F g−1 4050.4

– 70.82 ~57

245.98#/0.028ѣ 23#/1.5ѣ – – – 204.09#/0.0068ѣ 669.8#/8.119ѣ 49.35#/7.9ѣ –

0.3ʎ 15ʎ 5Ⱡ 3Ⱡ 1Ⱡ 0.8ʎ

2ᴧ 1Ⱡ

5ᴧ

420.91#/0.02755ѣ 82.23

1Ⱡ

89.8

95

85.2

78.7

82

–

86

–

1Ⱡ

87.4

Capacity retention/% 90.9

–

ED /PD 77.01#/0.56ѣ

1Ⱡ

10ʎ

Current density/ scan rate

[102]

[101]

[100]

[99]

[98]

[97]

[96]

[95]

[94]

[93]

[92]

[91]

Refs. [90]

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Nanocrystals Nanorods

Nanoneedle arrays/ nanosheet arrays

Nanoflower

Nanosheets Hexagonal rod like

CBD HT

Sol-gel method Facile HT

HT and electrodeposition

SILAR

Facile HT HT

CuO WO3

TiO2 α-MnO2

SiC@NiCo2O4/ Ni(OH)2// SiC@Fe2O3 Cu(OH)2-// rGO NiO ZnO 3 M KOH 6 M KOH

1 M KOH

1 M KOH 1 M Na2SO4 0.5 M KCl 1 M Na2SO4 2 M KOH

116.9 –

68.5

–

82.11 114

39.58 –

SA/ Electrolyte m2g−1 1 M LiNO3 –

242

312 311.52

576 570

S.Cap /F g−1 587

3000 200

81.67 5.87

10,000 165

5000

500 2000

– 5000

Cycle life 5000 – 93 ~100 80 86.6

80

19.65#/0.92ѣ 99#/0.45ѣ 882#/1260ѣ _ 103#/3.5ѣ

66.7#/5.698ѣ – –

5ᴧ 0.5Ⱡ 10Ⱡ 0.3Ⱡ 4Ⱡ

2Ⱡ 0.5Ⱡ 5ᴧ

78.5 –

Capacity retention/% 92

ED /PD 100.8#/0.399ѣ

Current density/ scan rate 0.5Ⱡ

[110] [111]

[109]

[108]

[106] [107]

[104] [105]

Refs. [103]

HT hydrothermal, CT calcination, CBD chemical bath deposition, SILAR successive ionic layer deposition, S.Cap specific capacitance, ED energy density, PD power density, SA surface area ᴧ = mV s−1, Ⱡ = Ag−1, ʎ = mA cm−2# = Whkg−1, ѣ = KWkg−1

Morphology Nanobelt/nanoflake arrays Woollen-like Nanofibers

Synthesis Two-step HT

Electrode material V2O5/TiO2

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engine and electric motor [115]. The power needed for EVs to accelerate is much higher than the one needed to keep it on a uniform speed, and the reduction in electrical power supply results in poor movement. This has been a major threat to the battery. The battery has shown promising features but has several drawbacks such as long charging time, low power density, and large internal resistance [2, 3]. The unique features of SCs such as high power density, fast charge, long lifetime, wide temperature range, and excellent electrical properties have made EVs the most fascinating technology which contributes tremendously to the reduction of fuel consumption to zero or near zero.

22.6.2  Electric Rail Transit System This system makes use of an electric power supply. The electric power is produced by generating stations, conveyed to the railway system, and then distributed to the moving train through a conducting cable that runs alongside the rail track. The conducting cable is either a third rail (situated at the rail track) or an overhead line fixed in between the two poles. The power source is majorly a DC voltage source. The two major problems that affect the DC voltage source in the electric rail transit system include under-voltage conditions and regenerating power braking systems [116]. When a train accelerates, it absorbs power in a large quantity which produces the required torque/force to keep the train in motion. The excess power absorption of the train might give rise to an under-voltage condition. When the DC source constantly falls below the required voltage for several periods, the train tends to disconnect from the power source and only reconnect back when the required voltage is supplied. This protective mechanism might reduce the efficiency or lead to damage to the train if the arising problem persists. On the other hand, another issue arising in the electric rail system is the regenerating braking system, and this occurs during the deceleration of trains. The excess energy absorbed during the acceleration of the trains is regenerated while decelerating. The regenerated energy is induced in the conducting cable, thereby increasing the voltage. The induced voltage is absorbed by the accelerating neighboring train. If the voltage is high enough, it can be released in the form of heat within the surrounding. These drawbacks are curbed by the adoption of energy buffers such as flywheels, SC, and battery. Among them, SC is the most promising storage device due to its unique features. It has a record of low energy density and a high power density up to 6 KW/Kg or above [117]. It serves as a voltage enhancement if the train experiences an under-voltage condition. During deceleration, the scientist has developed an electric motor that powers the train to reduce the regenerating braking system [118]. SC serves as the voltage source for the electric motor and as well absorbs the excess voltage dissipating as heat.

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22.6.3  Mobile Device The mobile device has been one of the easiest means of delivering massages, but the function of the mobile device is not hinged on that. It can as well be used in playing music and taking photos. Consumers always wish to have a mobile device with a quality camera comparable to a digital camera. Most mobile device cameras make use of LED flashlight, but it gives low-quality photos in dim light due to the low generation of flashlight. To achieve a better flashlight, high power is needed and the battery cannot provide such [119]. An alternative technology has emerged, thanks to the xenon flash tube. Xenon flash tube has attracted the use of SC in the mobile device, and this has reduced the high demand for energy in mobile devices. However, SC energy buffer is used to power mobile devices and as well used to enhance the performance of its camera. Also, it has been proven that SC is suitable for wireless charging [120].

22.6.4  Memory Device SC has shown an excellent technology in the area of memory devices. Traditionally, most data storage devices have solely depended on UPS for power back up. UPS is made of battery and it doesn’t sustain for a very long time. An alternative device such as redundant arrays of independent disk (RAID) and dynamic random access memory (DRAM) [121] which make use of SC has been developed. They show promising features far better than UPS and are designed especially for data preservation. In the case of power failure, data stored temporarily in volatile memory are saved in nonvolatile memory. This whole process is achieved within some seconds. RAID and DRAM use battery; this has led to their drawbacks due to the short lifespan of the battery, and SC serves as the replacement for the power backup. The long cycle of SC helps RAID and DRAM to preserve data for many years.

22.6.5  Wearable Electronic Device Wearable electronic device is one of the most trending technologies in the world. An electronic device such as electronic sensors, health monitors, mobile phones, smart cameras, and swatches is said to be wearable if it gives comfortability when worn like clothes or cap. Hence, it must be flexible. SC which serves as an energy buffer to these electronic devices has the ability to be stretched without deformation [122]. A team of researchers from Duke University and Michigan State University [123] was able to fabricate an SC with good mechanical strength that has the tendency to be stretched eight times its original size and yet offer good electrochemical performance. Therefore, SC is easily incorporated with wearable electronics. As shown in Fig. 22.10b, someone can comfortably wear an SC fabric cap, harvest energy, and play music while taking a stroll without running short of energy.

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Fig. 22.10 (a) SC powered EV [124] and (b) wearable wireless

22.6.6  Micro-Grid Micro-grids are small-scale electrical networks that operate independently or along with a power generation unit. The power generation unit includes the electrical grid and renewable energy system (e.g., solar cells, wind turbines, etc.). When interconnected with a renewable energy system, it offers a unique voltage supply, but its efficiency is limited due to the intermittent nature of renewable energy sources. Therefore, an energy buffer is required. SC is used in the stabilization of micro-grid voltage (when there is a voltage drop as well as when excess energy is harvested) [125]. Similarly, a micro-grid can as well be connected along with the electrical grid; the system is referred to as smart gird. In a smart grid, the micro-grid maintains the frequency of the voltage supply and integration of energy buffer, e.g., an SC in micro-grid helps keep the power supply on if there is a sudden loss of power in the electrical grid, thereby making the micro-grid act as a short-term power supply.

22.6.7  Chemi-Resistive pH Sensing The chemi-resistive sensor is a special type of electrochemical pH sensor that measures the pH of a solution by generating a change in conductivity or resistivity on chemi-resistive material (sensitive electrode). The change in conductivity or resistivity brought about the rise or fall in the measured voltage. The chemi-resistive material is not self-induced. Therefore, for current to flow, an external energy buffer is required [126]. Since the sensor is usually miniaturized, the bulky nature of the battery makes it unattractive. Hence, SC serves as an energy buffer.

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22.7  O  utlook and Summary of Nanoporous Metal Oxide-­Based Supercapacitors Green energy is globally desired. Considering the erratic nature of most green energy sources, the rising demands for consumer electronics, and the desired portability of electric-powered gadgets (including automobiles), energy storage is paramount. SCs have attracted a significant attention as charge storage devices because of their salient features: fast charging, high power density, long cycle life, etc. However, the low energy density of SCs has limited its applicability. Without many compromises, faradaic SCs (pseudocapacitors) combine the desirable attributes of both EDLCs and batteries. MOx materials have proven to be the best candidates for high-energy SCs, due to their abundance, low cost, eco-friendly nature, ease of fabrication, design flexibility, chemical and thermal stability, and most importantly their redox-active nature. However, they suffer from low electric conductivity and cannot attain surface areas as high as those of carbon [17, 18]. In this chapter, the basic components of SCs and the mechanisms of charge storage were discussed. The possibility of achieving energy densities comparable to those batteries, by triggering extrinsic pseudocapacitance, using nanostructured MOx, made it pertinent to consider various methods of fabricating nanoporous MOx-based electrodes, including those derived from metal-organic frameworks (MOFs). Also, the performances of some nano-MOxbased electrodes were reviewed, followed by some applications of SCs. The overview of MOx electrode materials performance showed that binary MOx electrode portrays greater specific capacitance when compared with single MOx. Also, composite and hybrid electrodes seemed to be promising. The utilization of two MOx electrodes in SCs has attracted little attention [40]. This is a result of the difficulty in obtaining suitable MOx material pairs, capable of giving high storage capacity and cycling stability. Few works are reported on aqueous electrolytes, but none for nonaqueous electrolytes [40]. This gap calls for research attention. It is worth noting that attractive nanoporous structures with the versatile design are achievable with MOx derived from MOFs [72]. However low electrical conductivity and lack of control over the pore size were reported [127]. Optimizing existing fabricating methods and developing new ones could be helpful. Meanwhile, an in-depth understanding of the charge storage mechanism of the MOx-based electrode is required, for that would serve as an eye-opener toward the invention of novel MOx materials of higher energy density for SC applications. So far, performances of SCs are chiefly hinged on investigations on the electrode material (composition, structure, and size); there is a need to equally divert attention to the electrolyte. Investigations toward achieving nano-pore wettability, high mobility of ions (down to the wetted regions), wide voltage window, electrolyte stability, etc., including safety and economics, are crucial and should be treated seriously. The operating potential windows of electrolytes have a great effect on the energy storage capacity of SCs. Selection and optimization of electrolyte materials should not be skipped when going into SC-based researches. Selection in the sense that the electrolyte chosen should be compatible with the electrode materials while optimizing

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the operating potential window of already-existing electrolyte materials or the development of novel materials should be one of the focal points of future researches. Moreover, computational simulation of SCs, focusing on the electrode-­electrolyte interface, as well as considering the full cell, is highly recommendable. Of course, computational tools reduce the burden of performing a series of experiments. Computer simulation could be indispensable in predicting the low limit of pore size needed to achieve good electrolyte wettability for a given electrode-electrolyte pair, and such information would be indispensable in the fabrication of electrode material. Simulation of the electrochemical process at the electrode-electrolyte interface would aid in choosing rightful electrode-electrolyte pairs for achieving a desired capacitance, power, or energy density. The same goes for the simulation of the full SC cell, whereby one could rightfully decide on the choice of symmetric, asymmetric, or hybrid SC and the rightful electrodes and electrolyte. Just as the use of computational fluid dynamics has encouraged and promoted research in transport phenomena, adopting computational tools into electrochemistry would push electrochemical energy conversion and storage to another level, making it feasible for SCs to gain wider applicability. Considering automobile application, Formula One (F1), the most famous scientifically advanced sport in the world [128], governed by the Fédération Internationale de l’Automobile (FIA), still deploys internal combustion (IC) engine-powered cars [129]. While the fastest road car motors can barely speed beyond 6000  rpm, the motor of a typical F1 car can attain 15,000 rpm [130, 131]. Out of a desire for high-­ speeding cars, Renault introduced a V6 turbocharged engine as early as May 1977. Thereafter, came the rumors of “one-lap wonder qualifying engines,” also known as “grenades”—the most powerful F1 car engines ever [132]. The turbo-engine technology got banned in 1989 (for safety reasons), sailed through a tough gestation, and reincarnated in 2014 [132, 133]. Aside from the turbo-engine, nowadays, the powertrain of F1 cars also comprises energy store (ES, a collection of connected lithiumion batteries, which usually weighs between 20 and 25  kg), an energy recovery system (ERS, which feeds energy to the ES and comprises of motor generator units for reclamation of waste kinetic and thermal energies emitted upon braking the car and from the turbo-engine(s), respectively [134]), and then control electronics. In this season, F1 cars are constrained to have at most three (1.6 L V6 IC) engines, three turbochargers, three each of the thermal and kinetic ERS components, and two control electronics [135, 136]. SCs are yet to grace the sport. However, in envisioning electric F1 sports cars (akin to the EVs), bearing in mind the huge power demands (especially during acceleration phases), SCs are obviously inevitable. Someday, SCs would be celebrated in F1 sport, just like the turbo-era. Until their overall benefit practically outweighs that of batteries, that day would never come. Lastly, the trending deliberation on whether a typical battery-type electrode material which exhibits fast surface redox behavior when nano-tuned should be qualified as pseudocapacitive or not is a mere trivia. Knowledge of the mechanism of “extrinsic pseudocapacitance” is necessary for the advancement of charge storage with nanomaterials, but not the nomenclature. Scholars should rather focus on achieving even higher energy storage capabilities with the proven promising nanoporous MOx materials.

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Chapter 23

Nanoporous Metal Oxides for Supercapacitor Applications Ved Prakash Joshi, Nitish Kumar, and Rahul R. Salunkhe

23.1  Introduction to Nanoporous Metal Oxides Nanoporous metal oxides (NMOs) having high surface area and controllable porous structure are useful to obtain high energy density (ED) along with long-term stability. Over the last few decades, many porous materials like zeolites [1], porous silica [2] or alumina [3], ceramic-based materials [4], and metal oxides (MOs) (RuO2 [5], MnO2 [6], Co3O4 [7], and ZnO [8]) have been reported in the literature. The NMOs produced via templated methods have decent porosity and high surface area values. However, their complicated synthesis process, limited availability, and high cost of natural templates, and the fact that they cannot be used for the large-scale production, restricted their practical applicability [9]. Hence, the suitable synthesis method for the development of NMOs is an intricate problem. The nanomaterials are the materials having a particle size of range 1–100 nm. They can be crystalline/amorphous, organic, or inorganic materials. These materials show unique properties (physiochemical, magnetic, and electrochemical) owing to small particle size and large surface-to-volume (A/V) ratio. Depending upon the particle shape, size, and formation, the nanomaterials can be further divided into different subclasses. The nanoporous materials are the one subclass of nanomaterials. The nanoporous materials are solid materials that have voids throughout the material. These voids improve the surface area of the material and make them suitable for utilization as a supercapacitor electrode. These porous materials have high structural stability, vast surface area, and highly conductive behavior. The vast surface area of the electrode can provide good stability, while the large pores can provide better mobility for ions [10]. V. P. Joshi · N. Kumar · R. R. Salunkhe (*) Materials Research Laboratory, Department of Physics, Indian Institute of Technology Jammu, Jammu, Jammu and Kashmir, India e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_23

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The chemical synthesis methods are the simple and low-cost methods applied for mass scale production of MOs. Over the last few decades, low-cost chemical processes, such as chemical bath deposition (CBD), electrodeposition, hydrothermal method, sol gel method, etc., have been studied to develop MOs. These routes provide superior control over the nucleation and growth of the NMOs. Also, these methods can operate in low-temperature conditions (CBD, electrochemical deposition) and offer large-scale and cost-effective production of various MOs. Before discussing different synthesis methods in detail, we will discuss the basics and working of supercapacitors so that readers can relate these concepts when we will be using them later in this chapter. A supercapacitor (SC) also known as an electrochemical capacitor or ultracapacitor is a clean and green energy storage device. SCs have many applications, including the automobile sector, mobile, lifts, cranes, and laptops. The SCs can provide high ED (5–10 Wh·kg−1) compared to conventional capacitors and excellent power density (PD) (5–50 kW·kg−1) compared to batteries [11–13]. The basic construction of SCs has two parallel plate electrodes separated by a separator which conducts electrochemically but not electrically (ceramic, paper) and an electrolyte (aqueous, organic, and ionic liquid) that ionically connects these electrodes. Depending on the charge-storing principle, SCs are divided into two categories: electric double-layer capacitor (EDLC) and pseudocapacitors. In an EDLC, the charge gets store via separation of charges in the Helmholtz double layer at the electrode-electrolyte interface. Whereas if a faradic charge transfer happens during the charge storage, including a redox reaction, then it is known as a pseudocapacitor. Carbon-based materials mainly show an EDLC behavior, while metal oxides/ sulfides/nitrides/phosphides and conducting polymers mostly show a pseudocapacitor behavior. The pseudocapacitors can be classified into two types: redox pseudocapacitor and intercalation pseudocapacitor. The materials that offer redox pseudocapacitance mainly include MOs (RuO2, MnO2, CuO, Fe2O3), binary MOs [14] of the form XY2O4 (X = Ni, Co, Mn, Cu, or Zn; Y = Mn, Co, or Fe), CoMoO4 [15], and NiMnO3 [16]. One the other side, metal nitrides such as Ni3N, Ta3N5, and Co3Mo3N [17]; metal phosphides including amorphous Ni5P4 and Co2P [18]; metal sulfides such as NiS2 and CoS2 [19]; and conducting polymers [20] are also redox pseudocapacitive in nature [11]. However, the perovskite MOs, T-Nb2O5, and TiO2-B are the typical examples of intercalation pseudocapacitor materials [21]. The faradaic charge transfer process at the electrode-electrolyte interface for the redox and intercalation-type pseudocapacitors is illustrated in Scheme 23.1a, b, respectively. In redox pseudocapacitors, ions adsorbed onto the surface of electrode material store energy through faradic charge transfer, whereas intercalation occurs when counterions are reversibly inserted into and extracted from the layered host material. Now let us understand the energy-storing process in detail. During the faradaic process, the oxidation and reduction reaction happens at the electrode material. Depending on the voltage applied to the electrode material, the oxidation and reduction current flows in the sample. Suppose if we provide the negative potential to the electrode material, then its electron energy increases. When it

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Scheme 23.1  Pseudocapacitor types and charge transfer process during oxidation and reduction reactions. Basic types of pseudocapacitors: (a) redox pseudocapacitor and (b) intercalation pseudocapacitor. The scheme showing (c) charge transfer from electrode representing a reduction current (d) represents an oxidation current flow. Here MO is the molecular orbital; the scheme is showing the highest occupied and completely filled molecular orbitals in the solution

reaches a particular level, the transfer of electrons from the material into the vacant electronic state of the electrolyte species occurs (Scheme 23.1c). In this situation, the flow of electron is from the electrode to the electrolyte solution. Hence the reduction current flows in the system. Similarly, if a positive potential is applied to electrode material, then its electron energy decreases, and now the direction of electron flow is from the electrolyte solution to the electrode material. This current is known as an oxidation current shown in Scheme. 23.1d. The MOs have been explored widely for SC application since they exhibit a good ED and low-cost and readily available materials. The performance of the MOs can be further boosted via developing a well-controlled pore structure in the material. In recent years, the NMOs have gained enormous research attention due to their unique

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physicochemical (high thermal stability, chemical stability, and flexibility) [22] and electrochemical properties (high specific capacitance) [23]. NMOs found a variety of applications like supercapacitors [23–25], batteries [26], solar cells [27], catalysis [28], separation technologies [29], and hydrogen storage [30]. These materials have an extensive range depending upon their composition with oxygen. When the metal molecules are connected with the oxygen molecules to develop MOs, their lattice structure changes, and accordingly they behave as a conductor, semiconductor, or insulator. MOs suffer from poor electrical conductivity, limiting their ED and PD [31]. The pore structure and surface area are important factors for improving the ED of the MOs. According to IUPAC classification, the nanoporous materials are divided into three groups: microporous materials (50 nm) [32]. Figure 23.1 shows the electrolyte counterions residing between pore walls of different sized pores. Figure 23.1a shows the electrolyte ions in a macropore cavity. Due to their large pore size, the pore walls behave like a planar surface, leading to the formation of EDLC. On the other hand, the mesopores do not have sufficient space for the formation of EDLC,

Fig. 23.1  IUPAC classification of nanoporous materials. As macroporous (pore size >50  nm), mesoporous (pore size is between 2 and 50 nm), and microporous (pore size should be 100 nm·min−1 [91]. CBD technique is a low-cost technique with full control over film formation, and surfactants can fully control the surface area of nanoporous MO [92]. This method has the potential for replacing traditional methods owing to reasonable control over the developed material properties, low cost, and highly reproducible process. Finally, the sol gel synthesis method is another eminent chemical technique having advantages like the development of impurity-free product [93]. It is a versatile route that has been used significantly for producing NMOs. Using the sol gel process, one can develop the desired shape of nanoporous MO structures like nanospheres, nanorods, nanofibers, and nanotube [94]. Stable surface and high surface area are the most promising benefits of the sol gel technique [95]. Apart, these methods can able to produce thin films as well as powder samples of materials, which gives an advantage of the commercialization of these materials.

5

–

Hydrothermal

Electrochemical plating

–

NPG substrate

1500

1016

360

810

550

316

476

62.2

101

750 207

263 401

Specific capacitance (F·g−1)

10

1

–

–

–

2

–

–

2

– 5

– –

–

–

0.82

10

1

–

0.5

–

1

10 –

1 0.5 mA·cm−2

Discharge Scan rate current (mV·s−1) (A·g−1)

0.5 M H2SO4

0.5 M Na2SO4 6 M KOH

2 M KOH

0.5 M K2SO4 5 M LiCl

2 M KOH

0.5 M K2SO4 1 M KOH

2 M KOH 1 M KOH

6 M KOH 2 M KOH

50.3

–

48.8

21

–

43.8

42.3

–

10.3

– 27.8

– –

Energy density Electrolyte (W·h·kg−1)

It is evident from the tabulated data that after creating the pores in the material its supercapacitive performance increases

13

12

Mesoporous NiO/rGO RuO2/NPG

–

10

11

HI-ULM nano-paper Glassy carbon Stainless steel Ni foam

1.8

Carbon / Solvothermal Ni(OH)2/NiO Sol gel MnO2

Graphite plates Ni foil

0.9

–

Ni foam

GNS/ZnO

6

–

10

V2O5

5

0.1 –

Crystallization technique Nanoporous metal oxides 7 Nanoporous Solvothermal Co3O4 Precipitation 8 V2O5 technique Hydrothermal 9 MoO3

RuO2 TiO2

Ni foam Graphene sheet Ni foam Titanium foil –

Hydrothermal Microwave-­ assisted method Biosynthesis Anodization method Coprecipitation

~3 0.8

Substrate

3 4

No. Sample Metal oxides 1 Co3O4 2 NiO

Mass loading Synthesis method (mg)

Table 23.2  Comparative performance of traditional metal oxides and NMOs

[81]

1.56 kWkg−1

[86] [87]

84 kWkg−1

[85]

[84]

[83]

–

1161

7200

–

[82]

[80]

8.1 kWkg−1

–

[79]

[77] [78]

[75] [76]

168.8

– 1.42 kWkg−1

– –

Power density (W·kg−1) Refs.

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23.3.2  T  oward the Commercialization of Nanoporous Metal Oxides The nanoporous MO is not readily available in the market due to the restricted choice of commercially existing soft templates, including surfactants; most of the research work is applied to the industry level yet to commercialize. Only a few of them are commercially available, but that is too expansive. Some commercially available strong hydrophilic/hydrophobic contrasts are Pluronic block copolymer, PMMA. However, commercially accessible block copolymers generally fail to produce NMOs due to the small difference in hydrophilicity/hydrophobicity difference and low glass transition temperature, which cannot easily replicate homogenous porous structure [45]. For the commercialization of NMOs, research needs such general synthesis that will apply to the industry level. In this field, all research is mostly focused on the property and application part of NMOs. Here we have discussed some most used synthesis methods for developing NMOs. The traditional techniques like template methods are not suitable for large-­ scale synthesis and don’t meet the industrial requirements. However, the chemical synthesis methods like CBD, ED, hydrothermal, and sol gel methods are low-cost methods that can be effortlessly scaled up to the commercial level. Although these methods have a good potential, they are not utilized broadly for the synthesis of NMOs. These methods can develop porous materials with the desired properties (large surface area, well-controlled pore structure, and good conductivity) for SCs. If these techniques can be utilized wisely to synthesize good-quality NMOs, and the developed material is used for SCs, breakthrough results can be achieved.

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Chapter 24

Nanoporous Transition Metal Oxide-Based Electrodes for Supercapacitor Application U. M. Patil, V. V. Patil, A. S. Patil, S. J. Marje, J. L. Gunjakar, and C. D. Lokhande

Abbreviations 3DG Three-dimensional hierarchical graphitic carbon AC Activated carbon ACEP Activated carbon prepared from enteromorpha prolifera BET Brunauer-Emmett-teller BJH Barret-Joyner-Helena CBD Chemical bath deposition Ccp Close-pack array CFs Carbon fibers CMK-3 Mesoporous carbon CNF Carbon nanofibers CNT Carbon nanotubes C-TEX27 Activated carbon cloth CV Cyclic voltammetry CVD Chemical vapor deposition EDLCs Electrochemical double-layer capacitors EG Expanded graphite FLG Few-layers graphene FSS Flexible stainless steel GC Graphite carbon GF Graphene foam GNS Graphene nanostructure GO Graphene oxide GPC CeO2 on 3D graphene IL-G Ionic liquid-graphene

U. M. Patil (*) · V. V. Patil · A. S. Patil · S. J. Marje · J. L. Gunjakar · C. D. Lokhande Centre for Interdisciplinary Research, D.Y. Patil Education Society (Deemed to be University), Kolhapur, Maharashtra, India © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 F. I. Ezema et al. (eds.), Chemically Deposited Nanocrystalline Metal Oxide Thin Films, https://doi.org/10.1007/978-3-030-68462-4_24

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IO IONCC

Iron oxide Iron oxide nanoparticle on nitrogen-doped activated carbon composite IUPAC International Union of Pure and Applied Chemistry MNNGP Manganese oxide nanofibers/Ni nanotube/graphite/paper MO Metal oxide MOF Metal-organic framework MSN Mesoporous silica nanoparticles MWCNT Multiwalled carbon nanotubes NCO-NW Nickel cobalt oxide nanowires NiO/CS-PNT NiO decorated on polypyrrole nanotube through chitosan NPs Nanoparticles PANI Polyaniline PC Porous carbon PCs Pseudocapacitors PPDA p-Phenylenediamine PSAC Pongam seed shell-derived activated carbon RCGA g-C3N4@reduced graphene oxide aerogel composite REMOs Rare-earth metal oxides rGO Reduced graphene oxide SCs Supercapacitors SILAR Successive ionic layer adsorption and reaction SS Stainless steel TMOs Transition metal oxides

24.1  Introduction The utilization of highly efficient, clean, and sustainable energy conversion technologies and renewable energy resources is receiving much attention due to increasing energy consumption at high levels in technologically passionate world. The era of modern world belongs to the hybrid electric vehicles and wearable and handy electronic appliances (smart watches, cell phones, sensors, implantable medical equipments) which transformed our lifestyles [1]. For long duration use, the efficient energy storage systems are prime requirement for these smart electronic devices. Besides existing energy storage devices such as capacitors and batteries, supercapacitors have gained a tremendous focus as an alternative energy storage device because of their greater specific energy compared to traditional capacitors and greater specific power with longer cycle life compared to rechargeable batteries [2]. In general, SCs are categorized into two main groups: electrochemical double-­ layer capacitors (EDLCs) and pseudocapacitors (PCs). These categories are generally based on their charge storage mechanism. In EDLCs the electrodes are carbon-based materials, which are electrochemically inactive; that is, while

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charging the electrode materials, there are no electrochemical reactions that can occur, but at the electrode/electrolyte interfaces, only the pure physical charges build up. The EDLC-based supercapacitors especially provide great power densities (~15  kW  kg−1) with low energy densities (~5–10  Wh  kg−1), and that limits their applications including electric vehicles, portable electronics, stop-and-go systems, power tools, etc. [3–6]. Meanwhile, a typical pseudocapacitive material contains conductive polymers as well as transition metal oxides (TMOs)/hydroxides/sulfides which show greater specific capacitance than EDLC-based electrode materials [7]. In the process of achieving high energy along with high power of devices, at the same time pseudocapacitive materials arisen into profile with the completely diverged electrochemical features from EDLC and Faradaic reaction (like batteries). The morphology of MOs with well-defined porous, nanocrystalline structure influences the electrochemical behavior and consequently the performance. In the supercapacitor application, a variety of organic and inorganic porous materials such as carbons, MOs, organic-inorganic hybrid materials, and polymers were used as electrode materials. According to the International Union of Pure and Applied Chemistry (IUPAC), porous materials could be distributed into three types by the pore sizes: microporous (~2 nm), mesoporous (~2–50 nm), and macroporous (